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lllilllli'l PUBLICATIONS

NIST Special Publication 847

Machining ofAdvanced Materials

Proceedings of the International Conference

on Machining of Advanced Materials

July 20-22, 1993

Gaithersburg, Maryland

Said Jahanmir, Editor
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'At Boulder, CO 80303. ^Some elements at Boulder, CO 80303.
NIST Special Publication 847

Machining ofAdvanced Materials

Proceedings of the International Conference

on Machining of Advanced Materials

July 20-22, 1993

Gaithersburg, Maryland

Said Jahanmir, Editor

Ceramics Division

Materials Science and Engineering Laboratory

National Institute of Standards and Technology

Gaithersburg, MD 20899

Sponsored by

National Institute of Standards and Technology,

National Science Foundation, and U.S. Navy

Issued June 1993

U.S. Department of Commerce

Ronald H. Brown, Secretary

National Institute of Standards and Technology

Arati Prabhakar, Director

For sale by the U.S. Government Printing Office Superintendent of Documents, Mail Stop: SSOP, Washington, DC 20402-9328
National Institute of Standards U.S. Government Printing Office

and Technology Washington: 1993

Special Publication 847

Natl. Inst. Stand. Technol.

Spec. Publ. 847

532 pages (June 1993)

CODEN: NSPUE2

For sale by the Superintendent

of Documents

U.S. Government Printing Office

Washington, DC 20402
PREFACE

Advanced materials, such as ceramics and composites, are being increasingly used in various

applications. The high cost associated with machining and damage generated during machining are major

impedements to the implementation of these materials. In some cases, current machining methods cannot be

used and innovative techniques or modifications of existing methods are needed.

The present volume contains papers presented at the International Conference on Machining of

Advanced Materials, held at the National Institute of Standards and Technology, Gaithersburg, Maryland,

July 20-22, 1993. The goal of this conference was to strengthen communication and technology transfer among

researchers and engineers involved in various aspects of machining as related to ceramics and composites. The

following topics are included in these proceedings:

• traditional and nontraditional machining and finishing techniques,

• mechanisms of material removal,

• nondestructive evaluation and characterization of machining damage,

• effect of machining damage on performance and properties,

• sensors for in-process measurement of surface quality,

• special cutting tools and cutting/grinding fluids for advanced materials,

• new machine tool designs,

• novel concepts for machining of advanced materials, and

• machining of advanced materials for specialized applications, including biological implants.

Planning for the conference began in January 1992. Following the preparation of a preliminary

program, several acknowledged leaders involved in research and technology development in machining of

advanced materials were consulted to determine if such a conference was warranted. Since the response was

positive, two committees were organized to assist in the planning and execution of the conference. The call-for- paper was mailed to more than 4000 engineers and research scientists worldwide with interest in the conference

topic, and was published in Ceramic Industry, Mechanical Engineering, and Ceramic Bulletin. A total of 115

abstracts were received. Following review, 91 abstracts were accepted based on their technical content and

compatibility with the goal of the conference. A total of 58 manuscripts were submitted to NIST. These papers were subjected to a thorough and critical review, and a total of 48 manuscripts were accepted. The revised

manuscripts sent to NIST in a camera-ready format are included in this volume.

To accommodate the large number of papers without having parallel sessions, the papers were divided

into oral presentations and posters. In addition to the published papers in these proceedings, 15 other

presentations were included without publication.

This conference was co-sponsored by the National Institute of Standards and Technology - Ceramics

Division, the National Science Foundation - Materials Processing and Manufacturing Program, and the U.S.

Navy - Manufacturing Technology Program. The support and encouragement received from Steve Freiman,

Bruce Kramer, and Phil Nanzetta are gratefully appreciated.

iii
The members of the Organizing Committee and the Advisory Committee, listed on the following page,

provided much needed expertise and assistance in formulating the plans and bringing the conference together.

This conference was endorsed by Ceramic Industry, the American Society of Mechanical Engineers,

and the American Ceramic Society. These organizations assisted in publicizing the conference. I am grateful

to Pat Janeway, Mildred DeVita, Shiv Kapoor, Paul Holbrook, and Ron Barks for their efforts and endorsement

of the conference.

I am indebted to Kathy Kilmer and Lori Phillips of the NIST Conference Office, who made all the

detailed arrangements for the conference. Many thanks to Vanessa Williams for her assistance in preparing the

conference program and assembling the proceedings.

Said Jahanmir

National Institute of Standards

and Technology

iv
Conference Chairman

S. JAHANMIR

National Institute of Standards and Technology

Organizing Committee

R. L. ALLOR

Ford Motor Company

C. J. EVANS

National Institute of Standards and Technology

K. V. KUMAR

GE Superabrasives

J. MEYER

National Institute of Standards and Technology

V. P. THOMPSON

University of Maryland at Baltimore

Sponsoring Organizations

National Institute of Standards and Technology

National Science Foundation

U. S. Navy

Participating Organizations

American Ceramics Society

American Society of Mechanical Engineers

Ceramic Industry

Advisory Committee

D. K. ANAND

University of Maryland at College Park

K. J. ANUSAVICE

University of Florida

T. HOWES

University of Connecticut

R. KOMANDURI

Oklahoma State University

W. KONIG

Fraunhofer Institute of Production Technology

J. E. MAYER, JR. Texas A&M University

P. A. MCKEOWN

Cranfield Precision Engineering, Ltd.

J. MOSKOWITZ

Ceradyne, Inc.

H. - J. RENKER

Studer Company

S. SALMON

Advanced Manufacturing Science and Technology

K. SUBRAMANIAN

Norton Company

N. P. SUH

Massachusetts Institute of Technology

T. NAKAGAWA

University of Tokyo

B. F. VON TURKOVICH

University of Vermont

T. V. VORBURGER

National Institute of Standards and Technology

V
COVER ILLUSTRATION

The cover illustrations were taken from the papers published in these proceedings. These illustrations indicate the broad coverage of topics presented in the conference.

DISCLAIMER

Information on product names, manufacturers, or suppliers are included in this report for clarification. This does not imply endorsement of the products or services by the National Institute of Standards and

Technology.

vi
TABLE OF CONTENTS

GRINDING OF TECHNICAL CERAMICS

Machining of Ceramic Components: Process-Technological Potentials 3 W. Konig and A. Wagemann, Fraunhofer Institute of Production Technology,

Aachen, Germany

Highly Efficient Finishing of Ceramics by Helical Scan Grinding 17 K. Suzuki, Nippon Institute of Technology, Saitama, Japan; T. Uematsu, Toyama Prefectural

University, Toyama, Japan; and T. Nakagawa, University of Tokyo, Tokyo, Japan

Grinding of Ceramic Materials: A Model for Energy Consumption and Force

Transformation 21

S. X. Liang and O. F. Devereux, University of Connecticut, Storrs, CT

On Performance of Metal-Bonded Single-Layer Diamond Abrasive Tool 33

A. K. Chattopadhyay and H. E. Hintermann, CSEM, Neuchatel, Switzerland

A Systems Approach for Grinding of Ceramics 43

K. Subramanian, P. D. Redington, and S. Ramanath, Norton Company, Worcester, MA

The Role of Material and Grinding Parameters in Grinding of Alumina Fiber-Reinforced

Aluminum Alloys 55

H. Chandrasekaran and J. 0. Johansson, Swedish Institute for Metals Research,

Stockholm, Sweden

Elastic Stress Fields Caused by Sliding Microindentation of Brittle Materials 71

Y. Ahn, T. N. Farris, and S. Chandrasekar, Purdue University, West Lafayette, IN

On-Machine Electro-Discharge Trueing for Metal Bond Diamond Grinding Wheels for

Ceramics 83

K. Suzuki, Nippon Institute of Technology, Saitama, Japan; T. Uematsu and T Yanase,

Toyama Prefectural University, Toyama, Japan; and T. Nakagawa. University of Tokyo,

Tokyo, Japan

Ceramic Grinding with Self-Sharpening Diamond Wheels 89

R. S. Hahn, Hahn Engineering, Auburn, MA

Chemomechanica! Effects of Boric Acid in Core-Drilling of Alumina 101

5". Jahanmir and H. Liang, National Institute of Standards and Technology, Gaithersburg, MD

CHARACTERIZATION OF MACHINED SURFACES

Ultrasonic Measurement of Surface and Subsurface Structure in Ceramics 117

J. A. Slotwinski, N. N. Hsu, and G. V. Blessing, National Institute of Standards

and Technology, Gaithersburg, MD

vii
Surface Characterization of Ultrasonic Machined Ceramics with Diamond Impregnated

Sonotrode 125

H. Dam, S. Jensen, and P. Quist, Danish Technological Institute, Taastrup, Denmark

Measurement of Residual Stresses in Machined Ceramics Using the Indentation Technique .... 135

Y. Ahn, S. Chandrasekar, and T. N. Farris, Purdue University, West Lafayette, IN

Detection of Subsurface Defects in Machined Silicon Nitride Ceramics by Optical

Scattering Methods 147

W. A. Ellingson, D. M. Ayaz, M. P. Brada, and W. O'Connell, Argonne National Laboratory

Argonne, IL

Chatter and Surface Pattern Detection for Cylindrical Grinding Using a Fluid Coupled Acoustic

Emission Sensor 159

Y. P. Chang and D. A. Dornfeld, University of California, Berkeley, CA

EFFECT OF GRINDING ON STRENGTH

High Efficiency Grinding of Structural Ceramics 171

H. G. Wobker and H. K. Tonshoff, University of Hannover, Hannover, Germany

Effects of Ceramic Microstructurcd Character on Machining Direction - Strength Anisotropy 185

R. W. Rice. W. R. Grace & Co. , Columbia, MD

Diamond Grinding of Silicon Nitride Ceramic 205

J. E. Mayer, Jr. and G. P. Fang, Texas A&M University, College Station, TX

Machining of Silicon Nitride: Experimental Determination of Process/Property Relationships . . . 223

R. L. Allor, T. J. Whalen, and J. R. Baer, Ford Motor Company, Dearborn, MI;

and K. V. Kumar, GE Superabrasives, Worthington, OH

Damage Determination and Strength Prediction of Machined Ceramics by X-Ray Diffraction

Techniques 235

W. Pfeiffer and T. Hollstein, Fraunhofer Institute fur Werkstoffmechanik, Freiburg, Germany

Influence of Microstructure on Grindability of Alumina Ceramics 247

P. Roth and H. K. Tonshoff, University of Hannover, Hannover, Germany

Effect of Grinding on Strength and Surface Integrity of Silicon Nitride: Part I 263

S. Jahanmir, T. J. Strakna, G. D. Quinn, and H. Liang, National Institute of

Standards and Technology, Gaithersburg, MD; R. L. Allor, Ford Motor Co.,

Dearborn, MI; and R. D. West, Therm Advanced Ceramics, Ithaca, NY

Effect of Grinding on Strength and Surface Integrity of Silicon Nitride: Part II 279

5". Jahanmir, T. J. Strakna, and G. D. Quinn, National Institute of Standards

and Technology, Gaithersburg, MD; R. N. Kopp, Norton Co., Worcester, MA;

S. C. Yoon, Cincinnati Milacron, Cincinnan, OH; and K. V. Kumar, GE Superabrasive

Worthington, OH

viii
PRECISION MACHINING

Tool Flank Wear in Single Point Diamond Turning 295

J. S. Strenkowski and W. C. Larson, North Carolina State University, Raleigh, NC

Depth Profiling of Residual Stress Along Interrupted Test Cuts in Machined Germanium

Crystals 303

R. G. Sparks, Alcoa Technical Center, Alcoa Center, PA; and M. A. Paesler,

North Carolina State University, Raleigh, NC

Design of Ultraprecision Grinding Machines for Ductile Mode Grinding of Brittle Materials ... 317

H. Shimura, Tokyo Metropolitan College ofAeronautical Engineering, Tokyo Japan;

F. Hashimoto, The Timken Co., Canton, OH; and J. Yoshioka, NUTR K. K., Chiba, Japan

Precision Form Truing and Dressing for Aspheric Ceramics Mirror Grinding 325

T Kuriyagawa, K. Syoji, and L. Thou, Tohoku University, Sendai, Japan

Electrochemical Dressing of Bronze Bonded Diamond Grinding Wheels 333

E. Welch, Y. Yi, and T. Bifano, Boston University, Boston, MA

Effect of Ductile-Regime Grinding on the Strength of Hot-Isostatically-Pressed

Silicon Nitride 341

L. K. Ives, C. J. Evans, S. Jahanmir, R. S. Polvani, T. J. Strakna, and M. L. McGlauflin,

National Institute of Standards and Technology, Gaithersburg, MD

Influence of Ductile Mode Grinding on the Strength of Silicon Based Ceramics 353

J. M. Boettger, United Technologies Research Center, East Hartford, CT; M. K. Ker,

Cranfield Institute of Technology, Bedford, England; P. Shore, Cranfield Precision

Engineering, Ltd., Bedford, England; and D. J. Stephenson, Cranfield Institute of Technology,

Bedford, England

Efficient and Precision Grinding Technique for Ceramics with Electrolytic In-Process

Dressing (ELID) 359

H. Ohmori, The Institute of Physical and Chemical Research, Tokyo, Japan

FREE ABRASIVE MACHINING

Ultrasonic Shaping: An Economical Procedure to Machine Complex Ceramic Components .... 385

R. Hahn and C. Lost, University of Tubingen, Tubingen, Germany; and P. Schulze,

Erosonic AG, Watrwil, Switzerland

A Model for Hydro and Hydro-Abrasive Jet Machining of Carbon Fiber Reinforced Plastic

Composites 397

R. Groppetti and A. Cattaneo, Universita degli Studi di Perugia, Perugia, Italy

Efficient Mechanochemical Polishing for Silicon Nitride Ceramics 409

T. Uematsu, Toyama Prefectural University, Toyama, Japan; K. Suzuki and M. H. Wu,

Nippon Institute of Technology, Saitama, Japan; K. Suzuki, Toyama Prefectural University,

Toyoma, Japan; and O.Imanaka, Gunma Polytechnic College, Gunma, Japan

ix
Magnetic Abrasive Finishing of WC-Co Curved Surfaces 415

M. Anzai and E. Kawashima, University of Tokyo, Tokyo, Japan; H. Otaki,

Nihon University, Narashino, Japan; and T. Nakagawa, University of Tokyo, Tokyo, Japan

TURNING AND MILLING

Machining Ceramic Materials for Dental Restorations 425

E. D. Rekow, University of Maryland at Baltimore, Baltimore, MD; G. M. Zhang, University of

Maryland at College Park, College Park, MD; and V. P. Thompson, University of Maryland

at Baltimore, Baltimore, MD

Turning of Fiber and Particle Reinforced Aluminum 437

K. Weinert and D. Biermann, University of Dortmund, Dortmund, Germany

Turning of Carbon and Glass Fiber Hybrid Cloth Composite Material 447

A^. L. Cupini, UNICAMP/FEM/DEF, Campinas, Brazil; and J. R. Ferreira, EFEI/IEM/DPR,

Itajuba, Brazil

Laser-Assisted Hot Machining of Ceramics and Composite Materials 455

W. Konig and A. K. Zaboklicki, Fraunhofer Institute of Production Technology,

Aachen, Germany

Study of the Formation of Macro- and Micro-Cracks During Machining of Ceramics 465

G. M. Zhang, D. K. Anand, S. Ghosh, and W. F. Ko, University of Maryland

at College Park, College Park, MD

The Machinability of Different Dental Materials for CAD/CAM Systems 479

K. H. Kunzelmann and R. Hickel, University of Munich, Munich, Germany

Mechanics of Tool Wear During Machining of Advanced Fibrous Composites 489

G. Santhanakrishnan, R. Krishnamurthy, and S. K. Malhotra, Indian Institute of Technology,

Madras, India

The Metallurgy of CBN and Its Wear in High Speed Machining of Ferrous Materials 501

G. Rai, Smith Megadiamond, Provo, UT

LASER AND ELECTRICAL DISCHARGE MACHINING

Crack-Free Processing of Hot-Pressed Silicon Nitride Ceramics Using Pulsed YAG Laser 517

A^. Morita, Chiba University, Chiba, Japan

Electrical Discharge Machining of Advanced Ceramics 525

N. M. Faulk, Tennessee Technological University, Cookeville, TN

Effect of Grain Size on Wire Electrical Discharge Machining of Polycrystalline Diamond 535

5". Z. Wang, K. P. Rajurkar, and J. Kozak, University of Nebraska-Lincoln, Lincoln, NE

Laser Machining of Silicon Nitride Ceramics 543

A. G. Solomah, SAC International Ceramics Ltd. , Mississauga, Ontario, Canada

X
Technology in Pulsed Nd:YAG Laser Drilling of an Al-Li Based/SiC Metal Matrix Composite

T. M. Yue and W. S. Lau, Hong Kong Polytechnic, Hong Kong; and C. Y. Jiang, Nanjing

Aeronautical Institute, Nanjing, China

xi

GRINDING OF TECHNICAL

CERAMICS

1

MACHINING OF CERAMIC COMPONENTS:

PROCESS-TECHNOLOGICAL POTENTIALS

W. KONIG and A. WAGEMANN

Fraunhofer Institute of Production Technology, Aachen, Germany

In the production of ceramic components machining in the green state bears many advantages according to flexibility and economy. This holds true especially for small lot sizes

when compared to forming processes like moulding and pressing. But unfortunately in

the ceramic industry, in most cases, the process design in green machining is based only on experience. Therefore, approaches will be presented for an economic and quality

consistent green machining process by applying the potentials of modern machine tools

and cutting tools like PCD.

The situation is different in the area of fme machining of ceramics where numerous approaches have been presented in recent years towards optimizing the machining

processes, which are already applied in industry such as grinding and lapping. From these

investigations, limits can be derived with respect to process performance and quality.

Improvement of flexibility, economy and quality in fine machining is connected with the

development of new ceramic machining technologies. Potentials will be illustrated for

grinding with increased cutting speed and for electro discharge machining of conductive

ceramics.

Introduction

A close-up look at the production of ceramic

parts reveals that manufacturing cost and

flexibility are determined less by the material

used, or by processing, but rather by shaping.

This process comprises not only the actual initial shaping (injecting, casting, pressing) but also

green and white machining and finish-machining

after sintering as shown in Fig. 1.

The technique of machining ceramic components

before sintering offers a great potential to make

shaping more flexible and less cost-intensive (1,

2). Green machining, in the case of small to medium quantities, permits the shaping process to be both flexible and economical in comparison

with other shaping processes as the unit tool costs are lower. The deployment of geometrically

simpler semi-finished products permits good

levels of homogenity and distribution of density

to be achieved in the green part, and thus high

part quality after sintering. It also allows high

levels of accuracy to be obtained in the sintered

parts. This is of vital importance as after

sintering, only a small amount of material must

be removed in order to obtain the final geometry

required; thus, shortening effectively the very

time consuming and costly finish-machining step.

Finish-machining of the functional surfaces of a part is, however, unavoidable in most cases. The

difficulties involved in machining sintered

ceramics can result in up to 90 % of the part

costs being incurred during this stage. A number

of investigations have, therefore, already been

undertaken into the common finishing methods

such as grinding and lapping with a view to making finish machining more economical and,

3
powder

• extruding • Injection

moulding • casting • pressing

• milling • turning • cutting • grinding •

• grinding • lapping • polishing •

processing

forming

green machining

sintering

finish machining

V

workplece

• plastlf icatlon • mixing • granulation •

reaction sintering • pressure

sintering •

Fig. 1 : Production steps of ceramic components

at the same time, more reliable in terms of

quality.

The main technological potential offered by

further improvements lies in the deployment of

innovative machining methods. High speed

grinding, electro-discharge machining and

laser-assisted turning will be presented here as examples.

Green Machining

Green machining of small numbers of units in

order to produce more flexibly at lower cost is an approach already adopted by the ceramics

industry. Research shows, however, that there is

little, if any, systematic understanding of the

process design necessary for green machining.

Company know-how is based mainly on

experience; the technological reserves have not

been tapped to a sufficient degree.

Technological reserves (Fig. 2) are not limited to the area of process control, i.e. the selection of

machining parameters and the partition of

machining steps. The use of suitably designed

tools and tool geometries specially adapted to the

material is important, particularly for drilling or

milling. Another factor of fundamental

significance is the selection of the cutting tool

material, or rather, of its exact specification. The

example in Fig. 2 illustrates clearly the degree to which tool life and tool costs and, therefore,

manufacturing costs can be influenced by the

selection of cutting material.

A major problem which arises in green

machining is the low level of strength of the

material and the resultant problems in clamping

and handling the parts (3). Clamping devices

commonly used in metal working are frequently

used to hold the parts. These devices, however,

have major disadvantages with regards to the

transfer of force and sensitivity. The use of

vacuum clamping or flexible clamping elements

offers options for improvement.

The situation concerning the machine tools

available is similar to that of clamping devices.

Here too, machine design and components are not

adapted to meet the special requirements of green

A
• process strategy • tools • new cutting materials • ceramic adapted systems

for clamDlng and handling • adapted machine tool

and oeriDherle N

turning of Dfe-sintered SK

1

1 A HI /.^ .. ^ - i //

A f it

II

• mm-— i

C » — ij 1 II

' 10 20 30 ^00 200 m m

cutting length

V - 50 V,. = 1 00 f

m mn] [fn/nitn][ mm/r]

K c r

mm

A 0,1 HM K10 6° Q' Q" 90' 50'' 0,8

0,1 PCD 6'' 0' 0° 72,5'- 55" 0,8

Fig. 2: Technological potentials

machining. The high abrasion of the particles

removed, calls for corresponding designs to ensure the protection of the bearing and guide

parts of the machine tool. Additionally, the use of

modern cutting tool materials such as polycrystalline diamond (PCD) makes very high

cutting rates and spindle speeds necessary.

Extraction and, in certain cases, separation and

re-use of the materials extracted (i.e. chips)

represents another important area in connection

with green machining (4). At first sight, this

appears trivial, but even extraction of the

particles removed, direcdy at the point of cut

presents a problem when modem machine tools

are in use due to a lack of space.

Some of the potentials indicated can be

considered independently of the material to be

machined (machine design and periphery,

clamping tools), whereas the other areas are determined largely by the composition or the

properties of the green material (in this context

the term "green material" also shall include

materials, which are pre-sintered to increase their

strength). Consequently, knowledge of

material- specific behaviour is vital to the

determination of the machining parameters. In order to be able to define a machining strategy,

it is, therefore advisable to have a method of

characterisation of the machinability. It would

then be possible to draw up definitions of ranges

in which machining can be carried out profitably.

Two examples will now be presented of an approach which permits appropriate ranges for

feed rate and infeed to be defined effectively in

the context of turning. Figure 3 defines the

variables, used in the following to describe the

turning process.

The test, referred to in the following as cone

turning test involves turning off cylindrically a pre-tumed cone made of the material in question

so that the infeed changes in the course of the

cut. The relationship between infeed and surface

quality is now instantly identifiable for each feed

rate. In the case of pre-sintered SiC (Fig. 4), it emerges, for example, that surface quality

deteriorates with diminishing infeed. In practice

this means that the concluding finish cut, which

is appropriate in some other cases, actually

5
results here in lower surface quality. When this interrelationship is considered for

section A-B

= K (1+ ) ^ f tool cutting edge inclination

Fig. 3: Characteristics of turning process and -tool

raw geometry

(cone)

finish geometry (cylinder)

material : SIC pre-slntered

tool : PCD |aoj>o |>^s |<r |«:r \ h

\ lOi 0° 1 0°l 90°l 85io,4

feed f : 0.2 mm/r cutting sDeed Vc : 500 m/mln

= 0

t

4

"

7 m

ap= 0,7mm Rj= 22,8 ]m ap = 1 ,8 mm Rtm= 17,7 ;jin ap = 2,6 mm = 15,5 \m

Fig. 4: Cone turning test at pre-sintered SiC

6
various feed rates (Fig. 5), it becomes clear that

the infeed required, increases with feed and that

there is a feed range in which the infeed is not

critical (in Fig. 5 for f = 0.1 mm/rev).

machining process. In contrast, when AI2O3 is machined, short chips are formed due to the

approximately 8 % proportion of binder in the

green part. Material removal is characterised here

>

Io*->

ICO

(Ua.

o>O)

CO

i_>

(0

f = 0 3 mm/r

f = 0 2 rm/r • f = 0 1 um/r

material : SIC pre- sintered

tool PCD 1 1 : coarse 10° 0° 1 0° 1 9or|7or

cutting speed : 500 tn/mln

0 0,5 1 1,5

Infeed a^

Fig. 5: At pre-sintered SiC the roughness decreases with increasing infeed

Monomer-cast AI2O3 behaves quite differenly.

This material has very favourable properties as regards green machining since it combines high

strength with almost plastic chip formation as is shown in the detailed photo in Fig. 6, on the left. Here too, there is a clear interdependency

between infeed and surface quality which is, however, exactly the reverse to that in the case of

the SiC shown. The almost ductile chip formation

is conducive to the achievement of a surface

without any spalling, the roughness of which is characterised only by feed grooves.

It seems likely that the completely different

dependencies between infeed and surface quality

obtainable are attributable to the fundamentally

different chip fomiation mechanisms. Pre-sintered

SiC is susceptible to brittle fracture and

unconnected particles are removed in the

less by brittle fracture than by ductile chip

formation.

The two examples described, indicate that the cone turning test is a simple means of defining

favourable ranges for the parameters feed and

infeed and can thus supply standards for process

design. Besides feed and infeed, cutting speed is an important influencing variable in the

machining process.

In contrast to feed and infeed, cutting speed

exerts little influence on surface quality. By

increasing the cutting speed, the proportion of

cutting time of the machining process can be

reduced without impairing the quality.

The reduction in cutting time is not the only

benefit to be gained from increased cutting speed.

Higher cutting speed also results in reduced tool

7
• = 60 m/mln

V(. = 100 m/inln \ = 200 m/mln

material SIC pre-slntered

tool PCD 1 Co 1 1 J £r

1 coarsel 4° | 0" 1 0° llOm 7ffl 0 feed f : 0.2 mm/r

Infeed a^ : 2 mm

1000 2000 3000 MOOO 5000 6000 7000 m 8000

cutting length Ic Fig. 7: The tool wear decreases with higher cutting speed

wear. It should be noted at this point, however,

that the level to which cutting speed can be

raised is limited by the machinery currently

available and by the strength of the material

8
(stress caused by centrifugal forces).

The interrelationships so far described relate only

to the influenece of process parameters. Other

means of influencing the process behaviour are

tool design and the selection of cutting tool

materials. The comparison of different cutting

materials in Fig. 2 shows tool life of

approximately 30 m for pre-sintered SiC. In the case of PCD, as shown in Fig. 7, tool life is prolonged by a factor of 100 while cutting edge wear remains unchanged. Additionally, as previously described, substantially higher cutting

speeds can be reached thus further cutting costs. Despite its higher purchase price, PCD is likely

to prove economically viable as a cutting material

for this application.

type of bonding phase offers an additional means

of influencing wear via the cutting tool material.

Another way of exerting influence on tool life is by adapting the tool geometry, in particular that

of the tool corner radius since this is the area most affected by wear. Evidence of this is provided by a comparison of two tools made of

identical material but with different cutting edge

geometries. A sharp cutting edge corner is subjected to particularly high stress and,

therefore, demonstrates considerably more

pronounced levels of wear than the rest of the

cutting edge. As shown in Fig. 9, the greatest

level of wear appears at the corner of the cutting

edge. In the case of a tool with a corner radius,

the stress exerted on the corner is distributed

along the circumference of the corner radius. The

tool

PCD

4" I 0"l 0'

coarse grain • ' fine grain

lOdI 7Cfl 0

material SIC Dre-slntered

feed f 0.2 mm/r Infeed 3p 2 mm cutting speed : 500 m/mln

0,09

mm

0 06

0.03

1000 2000 3000 4000 5000 6000

cutting length Ic

Fig. 8: 'Tool wear is influenced by PCD-grain size

7000 m 8000

The decision as to the cutting material to be used

involves the selection of a group of cutting tool

materials as well as its exact specification. The

wear curves plotted in Fig. 8 show that in the case of PCD, tool life can be influenced by the

grain size used. Finer diamond grain result in

better tool life behaviour. In the case of PCD, the

cutting edge and corner, therefore, demonstrate

uniformly low signs of wear. This relationship is of major importance since a

tool which has a corner radius does not suffer

much from shifting of the cutting edge; the

dimensional accuracy of the workpiece is consequendy high.

9
These interdependencies show that when the

technique of green machining is applied, analysis

of the cutting process in conjunction with the use

of modern tools and machine tools adapted to the

requirements of green machining offers

considerable potential for more flexible and more

economical production of ceramic components.

While these mechanisms can be set by the creep

feed grinding process in surface grinding

operations, this variant cannot be applied to the

ubiquitous peripheral grinding operations. The

cutting forces which occur as a result of the

increased contact length in peripheral creep feedgrinding are so high that a workpiece either can

tool

1 ^0

1

1 10° 0' 90° 85'

• edge radius =0,4 mm B edge radius rj = 0 mm

material ; SIC ore-slntered

feed f : 0,2 mm/r

Infeed ap : 2 mm

cutting speed : 200 m/mln

S 0 10

0 1000 2000 3000 4000 5000 6000 7000 m 8000

cutting length Ij

Fig. 9: The flank wear is substantially determined by the corner radius

High Speed Grinding

Comprehensive investigations into the grinding of

ceramic have shown that by ensuring favourable

cutting conditions at the individual cutting edge

when grinding advanced ceramic, it is possible to effect a process with minimum surface damage.

Depending on material and the process

parameters, removal can be mainly ductile or

brittle. While brittle material removal is more

favourable in terms of energy, as grinding can be

carried out with low cutting forces,

predominantly ductile material removal is currently believed to be necessary in order to obtain high part quality. This is achieved by

ensuring small chip thickness at the individual

cutting edge and a relatively large rounded tip

radius of the cutting edge.

no longer be held between centers or is destroyed. In general, very strong workpiece

deformation occurs due to the high cutting forces.

This results in time-consuming finishing

operations in order to obtain the level of

dimensional and form accuracy required.

Chip thickness can, therefore, be reduced only

via the following process parameters:

- higher diamond concentration

- smaller diamond grain size

- increased cutting speed

- reduced workpiece rotational speed

in the case of peripheral grinding at a constant

material removal rate.

With regards to tool specification and workpiece

rotational speed, changes can be made only

10
within certain ranges. Tools with fine grains (D < 91 pm) cannot be used for roughing as the

diamond grains would become overstressed and

would break off. It is very difficult to dress tools

which are too highly concentrated ( > C 100) as

the increased number of diamond grains offers

considerable resistance to the dressing tool.

Reduced workpiece rotational speed results in

high infeed per revolution in plunge grinding

operations. The effect of this is increased contact

distance, and thus, a higher number of cutting

edges engaging at one time. At a constant

removal rate, this transition to creep feed

grinding results in forces which are too high and,

therefore, in substantial form deviation.

Our own investigations into high speed grinding

of ceramic support the findings related to enhanced surface quality initially derived

theoretically, as Fig. 10 shows using traverse grinding as an example.

This result is independent of tool type. While use

of tool-type A (resin-bonded) is approved up to a cutting speed of only 63 m/s, when higher

cutting speeds are required, the best surface

qualities can be obtained using tool type-D

(vitrified bonded) at v^ = 100 m/s. These surface

qualities correspond approximately to those of

dressed functional surfaces in spite of the

relatively high removal rate. Tool wear is not

increased in comparison with conventional

grinding conditions.

6

workpiece material coolant Infeed/stroke

feed rate workpiece peripheral speed material removal

HIPSSN

Oil ag : 20 Mm

Vfg : 300 mm/mln

: 0,5 m/s, up grinding

8000 mm-

|_J

L V/

L

1

\

•^tm Ra

D151 Type D CI 00 A D181 Type A CI 00

1 1

1.5

1

0.5

20 40 60 80 100 120

cutting speed v^-

Fig. 10: Higher cutting speeds permit improved surfaces

In contrast, theoretically at least, higher cutting

speed permits decreased chip thickness and, at the same time, lower cutting force, at a constant

material removal rate. Higher cutting speed

results in reduced chip thickness at a constant

material removal rate and thus in lower forces on

the individual grain.

Increased cutting speed from the currently

standard range of 20 - 50 m/s to up to 150 m/s,

however, requires grinding wheels specially

designed for such speeds and adapted grinding

machines. In terms of the grinding machine, this

means the use of spindle and balancing systems

for higher rotational speeds and stronger machine

protection. This does not pose any major

11
problems due to the steadily expanding

application of CBN (Cubic Boron Nitride) high

speed grinding, but does, in general, mean higher

machine costs.

The tool itself, however, is a critical factor.

Initially the problem lies in the higher thermal

stress to which the diamond is subjected by the

high cutting speeds. Additionally, the centrifugal

force exerted by high rotational speed, subjects

the diamond layer to high stress levels. This

poses problems particularly for the multi-layer

diamond tools.

Knowledge of the radial expansion occurring

when CNC-machines are used is, however,

extremely important. Due to safety regulations

geometrical reference values are frequently

determined when the machine is at a standstill.

At a cutting speed of 120 m/s an absolute change

in geometrical dimensions of 30 pm is caused.

High speed grinding is beneficial especially when

CNC machines are used. The high removal rates obtainable using this method reduce the

machining time. This is of major significance

where high machine hourly rates are concerned.

material

outside wheel diameter

inside wheel diameter

wheel width

AlMg15MuF27

500 mm

203,2 mm

10 mm

c4-1 b

< /

>

< V

4 /

/

20 40 60 80 100

30

20

10

cutting speed v.

1

1

1/

1

1 -A1/

1

1 1-^

1 /1

20 40 60 80 100

cutting speed

Fig. 1 1 : Higher cutting speed causes increased radial tool strain

An appraisal of the radial strain and of the

equivalent stress on the outer diameter of a conventional grinding wheel to be used for

machining ceramic illustrates the problem (Fig.

11). A mathematical model was used for this

purpose (5). The equations on which the

calculations rest, are based on the

two-dimensional equilibrium analysis of the

forces occurring in a co-rotational frame.

In this example, the model indicates that the

stress to which the outer diameter of the tool is subjected to centrifugal force is negligible.

Electro Discharge Machining

Complex moulds or holes either cannot be

manufactured at all using industrially established

methods such as grinding or lapping, or can be

manufactured only in time-consuming, expensive

operations. The deployment of electro discharge

machining, which has already proved its worth in the manufacture of tools in the metal working

industry, is a viable option in such cases. Although this method requires the material to have some degree of conductivity (6), it is independent of the material's mechanical

12
characteristics (6).

The thermal material removal process effected by

EDM is characterised by the removal of material

in a solid, liquid or gaseous state. The energy

required for the cutting operation is supplied in

the form of temporary electrical discharges (arc).

The material removal process occurs in a

dielectric fluid (Fig. 12) which serves on one hand as an msulator between workpiece and tool

and on the other as a means of removing the

particles caused by the removal operation.

surface cut by electro discharge machining (7).

The free silicon between the SiC grains is removed and whole grains break out as a result. The reason for this is the high conductivity of the

free silicon. This encourages the base of the

plasma channel to form on the Si-bonding phase

and to result in material removal there.

The roughness achievable can be influenced

mainly via the impulse frequency of the

discharges. Due to a rise in the contamination of

the work gap, the higher impulse frequency

^ED-cuttlng;^

Wire

tool

Fig. 12: Principle of Electro Discharge Machining

In general, a distinction can be drawn between

electro discharge-sinking and electro dischargecutting (Fig. 12). Electro discharge cutting has

already been mastered to a high level and is, therefore, the most suitable for industrial

application. This method permits prismas and

holes of any contour with flank inclination angles

of up to 30° to be manufactured.

Silicon-infiltrated silicon carbide, is by virtue of

its 10 - 15 % ratio of free silicon particularly

suitable for electro discharge machining. Figure

13 shows the polished cross-section of a SiSiC

results in enhanced re-ignition conditions for the arc and thus in more pronounced removal of the

Si phase, or rather, in increased roughness. The

rate of discharge and impulse duration are of less

importance to surface roughness. However the

risk of thermally induced crack formation causing

damage to the surface rises with increasing

impulse duration.

Electro-discharge machining also exerts influence on part strength. Machining tests show that

bending strength changes with roughness.

Investigations into other machining methods

13
wire

Parameters :

wire run off speed 35 mm/s discharge current le 100 A open circuit voltage 300 V puis frequency 7 kHz discharge duration ti 1 .5 JJS workplece SiSlC

Fig. 13: Relief of ED-cut SiSiC surface

machining method :

- ED - sinking

- ED - cutting

workplece material

- EDH-SIAION

surface quality :

- Ra = 0,8 /Lim

Fig. 14: ED-cut gear pump

show, however, that it may not be concluded that

any correlation exists between roughness and

strength. Strength and Wei bull modulus levels are

low in comparison with^reep feed grinding or

lapping. The Weibull modulus can, however, be

increased to a value of 40 with a uniform degree

14
of strength by trimming the contours at low

impulse frequency (8).

The gear pump in Fig. 14 is an example of the

excellence this method has to offer. However it has not yet established itself industrially for the

machining of ceramics. In view of the fact that

this method perinits the economically viable

manufacture of parts and batch sizes which

cannot otherwise be manufactured economically,

any deterioration in workpiece properties which

may arise from the adjustment of the electric

conductivity is acceptable. . „ .

Laser Assisted Machining

Provided due account is taken of the strength

behaviour of ceramic which is dependent on

temperature, it is possible to machine sintered

parts with a defined cutting edge. In the case of

silicon nitride, there are amorphous areas at the

grain boundary of the largely crystalline structure which result in a softening of the material at a temperature of approximately 1100°C. This

permits the heated material to be machined in a turning operation, for example.

suitable for use as a source of energy to heat the

material. The laser beam is connected via an

optic on the second carriage of a lathe. The beam

is thereby directed onto the shearing zone via the

tool cutting edge. A CBN turning tool carries out

the cutting operation.

When suitable machining parameters are selected,

surface qualities are achieved comparable to those obtained in grinding operations. Even

cracks, which would be expected to form as a consequence of the thermal stress on the part

during the process have not, to date, been in evidence.

Laser-assisted tuming offers advantages over the

established grinding operations - particularly in the manufacture of parts with complex

geometries since it is much easier to use suitable

forming tools and to change tools. The

interrelationships explained, indicate that the

technology of laser-assisted hot machining can be

successfully deployed even in hard machining

operations involving ceramic components.

7777777777777777777777. ^^experimental set-up

///////////////////////

suction pipe cutting tool focusing optic

n : number of revolutions

V,: feed rate

Fig. 15: Principle of Laser Assisted Turning

Figure 15 shows a design of a machine suitable

for such applications (9). The very high energy

density, good dosability and the narrow

heat-affected zone of a laser make it eminently

Summary

The examples of different methods given, show

that in the field of shaping, there is considerable

room for improvement in terms of both quality

I 15
I

and economic viability in green machining and

finish machining operations after sintering. As

regards green machining, the lack of a systematic

understanding of the process is the greatest

drawback. Additionally, the tools and machine

tools used are insufficiently adapted to the

requirements of green machining.

The area of finish machining has already been

the subject of a number of research studies and

has advanced continuously. It should be noted

that processes such as high speed grinding,

electro discharge machining and laser-assisted

turning, for example hold out prospects of the

deployment of innovative machining technologies

as well as of continued optimization of

industrially estabhshed methods.

Literature

1. W. Konig, E. Verlemann, A. Wagemann,

"ZeitgemaBe Konzepte fiir die Keramikfertigung,

" Fortschrittsberichte der

DKG, Vol. 7, No. 4, pp. 27-39,

1993

2. D.L. Halcomb, M. C. Rey, "Ceramic

Cutting Tools for Machining Unsintered

Compacts of Oxide Ceramics," Ceramic

Bulletin , Vol. 61, No. 12, 1982

3. W. Konig, A. Zaboklicki, "Laserunterstutztes Drehen von Keramik," vorgesehen zur Veroffentlichung in VDI-Z, No.

7, 1993

4. G. Spur, "Keramikbearbeitung,"

C. Hanser Verlag, Munchen, Wien 1989

6. W. Konig, D.F. Dauw, 0. Levy, U.

Panten, "EDM, "A Future Step Towards

the Machining of Ceramics," CIRP

Annals, Vol.37, No. 2, 1988

7. W. Konig et al., "Quahtatssichere

Endbearbeitung in der Serienfertigung

keramischer Bauteile," AbschluBbericht ' zum BMFT Projekt Nr. 03M2011B 0, Aachen 1989 I

8. U. Panten, "Funkenerosive Bearbeitung von elektrisch leitfahigen Keramiken,"

' Dissertation RWTH Aachen 1990

9. W. Spang, "Mechanische Bearbeitung \cn

Hochleistungskeramik," Internationale

Zeitschrift fiir Technik und Wirtschaft, Vol. 43, No. 2, 1989

I

5. W. Konig, F. Ferlemann, "CBNSchleifscheiben fiir 500 m/s Schnittgeschwindigkeit," IDR, No. 4, pp. 222-

251, 1990

J
HIGHLY EFFICIENT FINISHING OF CERAMICS BY

HELICAL SCAN GRINDING

K. SUZUKI

Nippon Institute of Technology, Saitama, Japan

T. UEMATSU

Toyama Prefectural University, Toyatna, Japan

T. NAKAGAWA

University of Tokyo, Tokyo, Japan

A highly efficient finishing method in grinding of ceramics has been developed.

In this method, named "Helical Scan Grinding", a straight wheel or a mounted

wheel is fed at an angle from the direction of grinding velocity. Such motion can

be realized easily on an NC grinding machine. This method has a feature that a wheel with coarse diamond grains can be used in finish grinding, leading to high

efficiency finishing. Experiments showed that surface roughness was improved

with the angle. When a cermet workpiece was ground with a mounted diamond

wheel of mesh size 100/120, while the surface finish was Rmax=4.2|im(maximum

peak to valley height) in the normal grinding or at an angle of 0 degree, the

surface finish in the helical scan grinding at an angle of 45 degree was

Rmax=0.6!Am, about one seventh of that in the normal grinding. The helical scan

grinding produced similar effects on silicon nitride, silicon carbide, and hard

metals. This method is applicable to both flat and curved surfaces, and also to polishing of stainless steel with a wheel type abrasive brush.

1. INTRODUCTION

Ceramic materials have many superior

properties; it is expected that they will play an

important roll in the field of mechanical parts

in near future. In fact, some kinds of

engineering ceramics have already been used

successfully in, for example, automobiles and

machine tools. It is well known that

machining of ceramic materials is very

difficult. Usually, they are machined on a grinding machine with diamond grinding

wheels; but the efficiency is extremely low.

The efficiency has been improved by adoption

of rapid feed deep grinding and by the

development of a grinding center(l).

Recent studies on grinding of ceramics are directed to obtain very smooth surfaces.

Conventional methods of obtaining smooth

surfaces are to use grinding wheels with small

abrasive grains. But, the efficiency is not

sufficiently high enough to meet the demands.

17

In this paper, a highly efficient finish grinding

method named "helical scan grinding method",

which utilizes an inclined spindle or special

movement on an NC grinding machine, is described.

2. HELICAL SCAN GRINDING METHOD

Finish grinding has conventionally been

carried out in the following manner: (Dby

adopting a grinding wheel with fine abrasive

grains, (Dby adopting slow feed speed, and

©by adopting small depth of cut. But, stock

removal rate by this approach is inevitably

very low. The proposed "helical scan

grinding" can be implemented to obtain better

surface finish at high rates of stock removal.

This method is realized by 'inclining a wheel

axis relative to feed direction, as shown in

Figure 1(a), which results in reduction of

actual distance between adjacent abrasive

grains. Accordingly, this method has a feature
that a wheel with coarse grains can be used in

finish grinding, leading to high efficiency

finishing.

3. EXPERIMENTAL SETUP AND

CONDITIONS

For realizing the proposed finish grinding

method, a specially designed machine tool with an inclinable spindle is necessary. But

currently, such a grinding machine is not

available. In the present experiments a

horizontal machining center was used to yield

the same effect as an inclined spindle. A

straight wheel or a mounted wheel was fed, as shown in Figure 1(b), so as to form an angle

between the feed direction and the direction of

grinding velocity by utilizing two-dimensional

movement of an NC machine tool. Metal bonded diamond grinding wheels were used because they have a sufficient

strength to endure high axial grinding forces.

Workpiece materials were cermet, silicon

nitride, silicon carbide, tungsten carbide and

aluminium oxide. Grinding fluid was chemical

soluble type grinding fluid(dilution:l/50).

4. RESULTS AND DISCUSSION

4.1 Influence of Helix Angle

Figure 2 shows the effect of the helical scan grinding method in grinding cermet

material with a rather coarse wheel(SDlOO)

having D=30mm in diameter and 20mm in

length. Figure 3 illustrates photographs and

profile curves of surfaces ground at different

helix angles. Judging from the photographs,

the surface finish seems to be rather worse. But the roughness proved that the helical scan

grinding improved surface finish significantly.

Grinding by a normal method, with helix angle

a=0° , resulted in a surface roughness of

Rmax=4.2!im(maximum peak to valley height).

However, grinding by the helical scan grinding

method with helix angle of a=45° , improved

the surface roughness to Rmax=0.6|im. This

indicates that the helical scan grinding method

could yield 7 times superior surface finish.

Though much improvement in surface

roughness has been achieved in the proposed

method, micro waviness is observed on the

surface in certain conditions. This

Grinding wheel

Work

Feed direction

(a) Inclined spindle method

(b)Hellcal scan grinding on NC machine

Fig.1 Schematic principle of helical scan

grinding method.

Wheel{D=30mm,width=20mm,SD100,bronze)

Vs=565m/min,Vw=100mm/min,a=0.1mm,Cermet

0 20 40 60

Helix angle a(deg)

Fig. 2 Improvement of surface roughness in helical scan grinding.

phenomenon may be due to the lack of

straightness of the wheel along its axis.

However, waviness is better than conventional

grinding.

18
Figure 4 shows a relation between feed

speed Vw and surface roughness ground by

helical scan grinding at a helix angle of

a=45° and depth of cut a=0.1mm. Figure 5

shows a relation between depth of cut and

surface roughness at a=45° and

Vw=100mm/min. The smaller the depth of cut or/and feed speed, the smaller the surface

roughness and waviness. It is worth noting

that such a smooth surface with Rmax=0.3|j,m

was obtained inspite of the coarse grain wheel

with 100 mesh size.

10° 20° 30°

0.5mm

4.2 Helical Scan Traverse Grinding Method

In grinding of a wide surface, traverse grinding is usually used. Here the helical scan

grinding method was adopted in traverse grinding. Figure 6 shows the surface profile curves obtained by grinding in three different

grinding methods, i.e., conventional grinding,

helical scan grinding(a=45° ) and traverse

helical scan grinding(a=45° , traverse pitch

p=0.2mm) for the cermet material with 100

mesh diamond wheel. The waviness was

successfully reduced in traverse grinding,

though surface roughness in both helical scan

grinding methods improved to ahnost the same

extent. Figure 7 shows the result of the three

grinding methods for cermet with a straight

grinding wheel(D=125mm, b=13.5mm, SD200).

In this case, the surface roughness by helical scan method is Rmax=0.3|xm, whereas it is higher at Rmax=1.8|im in conventional

grinding. Both helical scan grinding methods

show little waviness with almost the same

surface roughness values, probably because

straightness of the wheel was good by an

appropriate trueing by "on-machine

electrodischarge trueing method"(2).

4.3 Influence of Workpiece Material

As a next step, the helical scan grinding

method was applied to sintered tungsten

carbide(WC), silicon nitride, silicon carbide,

and aluminium oxide. For the former three

materials, the helical scan grinding method

yielded remarkable effect in improving surface

roughness. But in case of aluminium oxide,

the improvement of surface roughness was

=0

6)L/m

Fig. 3 Surface appearance and profile curves in helical scan grinding.

E

5.E _

05 3" w m

<u to c <u

2 ^

<D "a o c

CO ra t:

CO

3.6

2.4

1.2

Wheel(D=30mm,width=20mm,SD100,bronze)

Vs=565m/min,a=0.1 mm,a=45deg, Cermet

Waviness

50 100 150 200

Feed speed Vw(mm/min)

Fig. 4 Relation between roughness, waviness

and feed speed.

19
only by a factor of two, which was rather a small improvement compared with the other

materials. This result indicates that the effect

of the helical scan grinding method in

improving surface roughness may depend on

the workpiece material.

4.4 Application to Finishing of Curved Surface

Automobile parts made of ceramics are increasing. They often require rather smooth

surface finish. Here, the helical scan grinding

method was applied to finish grinding of the

curved surface of rocker arm pads made of hot

pressed silicon nitride.

A straight wheel(D=150mm,b=15mm,SD400)

was fed on a machining center along an NC-programmed 3D curved path to make a curved surface with a certain helix angle.

Under conditions of helix angle of 45° and

feed speed of 50mm/min, a satisfactory curved

surface with roughness under Rmax=0.6jim was

obtained.

5. CONCLUSIONS

The proposed helical scan grinding method

proved to be effective for realizing high

efficiency finish grinding of ceramics and

cermets with a diamond wheel of coarse grain

size. Surface roughness was improved by

increasing the helix angle; at a helix angle of

45 degree, it was about seven times better than

in ordinary grinding. It was also found that

the proposed method can be successfully

applied to grinding of curved surfaces. It is expected that various other applications of the

helical scan grinding method will be found for

attaining high efficiency finishing.

[ACKNOWLEDGEMENTS]

The authors would like to express their

heartfelt thanks to Niigata Iron Works Co.Ltd.

They also thank to Messrs. T.Shimoda,

T.Yanase, and K.Ohseki for their experimental

assistances.

[REFERENCES]

(l)T.Nakagawa, K.Suzuki, and T.Ueraatsu,

"Highly Efficient Grinding of Ceramics and

Hard Metals on Grinding Center," Annals of

the CIRP 35/1, pp.205 - 210, 1986!

E

E _

E

10 -3» w in <D w £ C

2 $ o "a o c

CO 03

13 CO

2.4 -

1.2

Wheel(D=30mm,width=20mm,SD100,bronze)

Vs=565m/min,Vw=100mm/min,a=45cleg,Cermet

Waviness

Depth of cut a(mm)

Fig. 5 Relation between roughness, waviness

and depth of cut.

(a)Conventional grinding

(b)Single pass helical scan grinding ^

(a=45deg)

(c)Traverse helical scan grinding

(a=45deg,traverse pitch=0.2mm)

Fig. 6 Decrease of waviness in traverse helical scan grinding. (Mounted wheel, D=30mm,

b=20mm,SD1 00,cermet)

-H K—

(a)Conventional grinding

(b)Single pass helical scan grinding(a=45deg)

(c)Traverse helical scan grinding

(a=45deg,traverse pitch=0.2mm)

Fig. 7 Surface profile curves in helical scan

grinding with a straight wheel.

(D=125mm,b=13.5mm,SD200,cernnet)

(2)K.Suzuki, T.Uematsu and T.Nakagawa,

"On-machine Trueing/Dressing of Metal Bond

Grinding Wheels with Electro-Discharge

Machining," Annals of the CIRP 36/1, pp.115 - 118, 1987.

20
GRINDING OF CERAMIC MATERIALS:

A MODEL FOR ENERGY CONSUMPTION AND

FORCE TRANSFORMATION

S. X. LIANG and O. F. DEVEREUX

University of Connecticut, Storrs, CT

A simple model is introduced for relating grinding forces to operational equipment parameters: The normal force on the wheel is assumed proportional to the normal component of the

velocity of the workpiece relative to the wheel, and the tangential force on the wheel is assumed

essentially frictional and, thus, proportional to the normal force. Both force components are shown to be proportional to the product of the workpiece velocity and the depth of cut. Comparably simple expressions are derived for average normal and tangential stress at the

grinding interface and for specific grinding energy. The normal and tangential forces are related

to measured horizontal and vertical force components via the effective locus of application of

these forces on the wheel, i.e., by an angle <{) from the vertical which is shown to be closely

represented by 2fi)/3, where fi is the angle subtended by the grinding interface. Experiments on

silicon carbide using an instrumented surface grinder with a resin-bonded diamond wheel and a synthetic coolant showed excellent correspondence to this model over a wide range of table feed

rates and grinding depths.

Introduction

The economic use of ceramic materials in applications traditionally reserved for metals and alloys is often restricted by their shaping and finishing, processes normally performed by grinding. Methods for

study of ceramic grinding include a "machining" ap- proach, entailing analysis of grinding forces, grinding

energy and chip formation (1), and an indentation

mechanics approach, in which grinding is modelled as a multiplicity of individual indentations or individual

abrasive grain scratches. (1, 2, 3) In a mechanistic

study of grinding the independent variables necessarily

relate to force components at the grinding interface,

while the parameters amenable to measurement are

different, e.g., force components experienced by the

table, and those subject to control are still different,

e.,g., depth of cut and table feed rate in the case of a simple surface grinding operation. The development

of a simple model relating these sets ofparameters is an

important aspect of this study in that it permits an

association of material failure conditions with simple

machine parameters.

Analysis of Grinding Forces

Grinding behavior may be correlated with the

normal and tangential force components at the arc of

contact between the wheel and workpiece, as distin- guished from the horizontal and vertical force components that are amenable to measurement but not di- rectly relevant to the process. The normal force component may be associated with the penetration load

applied to the abrasive grains; the tangential component with the friction of grain/workpiece contact and a pseudo-viscous effect of the chip-laden coolant. In this model the normal force is assumed proportional to the

normal component of the velocity of the workpiece

relative to the wheel, V^, and the frictional component

of the tangential force to be proportional to the normal

21
'w

Figure 1: Definition of parameters at the wfieel/

workpiece interface.

force component. Thus, for the increment of contact

arc defined by the angles G and 9 + d9 , shown in Figure 1,

Fn =Cn Vw • r • sin 9 - de

= Cn • VwT» (l-COS P) =Cn*Vw'a (g)

and

Ft = Ct •

j Vw • r • sin 9 • d0

Jo

= Ct* VwT* (l-cos p) = Ct»Vw»a (7)

We may define corresponding values of the average

grinding stress, and o^, as the normal and tangential

components of the grinding force per unit width di- vided by the length of grinding arc:

dFn = On • Vw • sin9 • r • 49

and

dFt = q • Vw • sin9 • r •

(1)

(2)

where F^^ and F^ are the normal and tangential force

components, r is the wheel radius and c^^ and c^ are

constants. (The corresponding force components per

unit width are given by F^^' = ¥Jb and F^' = F^/b where

b is the wheel width or specimen width, whichever is smaller.) The wheel contacts the workpiece over the

range 0< 9 < P where, if a is the depth of cut, p is defined

by

cosB =

r

and, therefore,

(3)

sinp = 1-1 2a . Zi]

r 71

1/2

or

sinp = p = (2aj

1/2

(4)

(5)

for small a. Integrating equations (1) and (2) over the

range 0 < 0 < p.

Path traced by

next grain

(Figure 2a

(Figure 2b

Figure 2: Idealized chip geometry.

22
On = Fn _ Cn ' Vw ' a

r • p r • b • 3

and

at = =

' ' ^

r • 3 r • b • 3

(8)

(9)

where B is expressed in radians. Substituting from

equation (5),

Cn«V, 1/2 (10)

We might anticipate that the specific energy of the

grinding process, U, defined by

y ^ Energy consumption rate _ ^

Material removal rate Z*

(16)

would increase with decreasing chip thickness. As a

first approximation in dealing with brittle materials we

may neglect plastic deformation energy and treat the

energy consumption, E, as the sum of the fracture

energy.

and

(11)

where = 2»r is the wheel diameter.

The grinding energy is related to the geometry of

the chips formed during grinding. These may be

idealized by the shape depicted in Figure 2, with a maximum thickness given by the grain depth of cut, t, which in turn may be inferred from Figure 2. The

horizontal velocity of the woiicpiece into the wheel,

V^, may be resolved into normal and tangential components, V^'sinP and V^»cos(3, respectively. The

thickness of a given chip is defined by the time interval

tVVg, where is the peripheral velocity of the wheel

and h is the circumferential distance between adjacent

abrasive grains, during which the relative movement of

the wheel into the workpiece, i.e., in the normal direction, is given by

t= Jl. Vw-sin 3

Vs

= h Vw.(2ar

(12)

(13)

The length of the chip side, s, is given by

s = r. 3 (14)

Substituting the material removal rate, = a • b,

t may be expressed

t- 2h

V»'b (D,.a} 1/2 (15)

Efrac = n • Ai • Y

and the frictional energy

Efric=Fn • |i* Vs

or

E = n • Ai • Y + Fn • • Vs

and

U

_ n Ai ' Y + Fn ' ti* Vs Zw

(17)

(18)

(19)

(20)

where n is the number of chips created in unit time; A^

,

the surface area of one chip; y- the surface tension; and

|i, the relevant coefficient of friction between the

workpiece and an abrasive grain.

The rate of chip formation may be approximated by the area ofthe wheel which passes through the

workpiece per unit time, • b, times the density of

particles on the wheel surface, h-^, or

n = Vs'b

(21)

The surface area of the chip, as depicted in Figure 2, may be approximated by that ofthe lateral surfaces, i.e., the contribution of the edges may be neglected. Fur- thermore, we may approximate the width of the chip by

the abrasive particle size, m. Thus, the area Aj is given

by

Ai = 2 • r . 3 • m (22)

Substituting equations (6), (21) and (22) along with 3

23
Vs

Ft Fn

Vw

a. Down Feed Grinding

Fh

Fv

Vw

b. Up Feed Grinding

Figure 3: Force relationships in up feed and

down feed grinding.

= sin(3 = 2(a/Dy^ into equation (20), we have

1 Ds ) Zw

_ 4 » Vs* r« b « m« Y

/ ^ u/2 ^Cn'ia - Vs

h^'Zw IdJ b

(23)

U = C"Vw + C

At constant table feed rate:

U = C"'a^/2 + C

At constant material removal rate:

u = C\'^ + c

(25)

(26)

(27)

Transformation of Force Coordinates

Grinding force is usually measured as the vertical

and horizontal components of the force, using a three- component dynamometerplaced beneath the workpiece.

We have defined the normal and tangential grinding

stress values, and a^, which may be integrated along

the grinding arc, yielding effective net force components, ¥^ and F^. In conventional grinding, which is characterized by a small depth of cut, little error is entailed in equating these components with the vertical

and horizontal force components, respectively, but

appreciable error may be seen when the depth of cut is large as in creep feed grinding. The relationships

between the measured vertical and horizontal components and the defined normal and tangential components for up feed and down feed grinding are apparent

in Figure 3. Qearly

Ft = Fh • cos 0 + Fv • sin <}>

and

Fn = Fv • cos ({) - Fh • sin <()

for down feed grinding, and

Ft = Fh • cos ({) - Fv • sin (()

and

(28)

(29)

(30)

or, in terms of table feed rate,

^ ^ 4 Vs » r » m » Y ^ Cn » H « Vs ^24)

• Vw '{Ds • a}^'^ b

Equations (23) and (24) allow us to express several

basic dependencies. At constant depth of cut:

Fn = Fv • cos (j) + Fh • sin (j) (31)

for up feed grinding, where ^ defines the effective

locus of application of and F^.

The value of (|) is a function of the distribution of

normal and tangential stresses at the grinding interface. In this work we have assumed the normal force to be

24
proportional to the normal component ofthe velocity of

the woiicpiece relative to the wheel, and the tangential

force to be proportional to the normal force. These

relationships are given by equations (1) and (2):

dFn = Cn • Vw • sinG • r • d6

and

dR = Ct • Vw • sinG • r • dG

(1)

(2)

Thus, with reference to Figure 1 , the horizontal (x) and

vertical (y) components of these force increments are given by

dFnx = Cn • Vw • r • sin^G • dG

dFny = Cn • Vw • r • sin 6«cos G • dG

(32)

(33)

and

dFtx = Ct • Vw • r • sin G • cos G • dG

-2/-

(34)

dFiy = Ct • Vw • r • sin^G • dG (35)

The respective force components are, thus.

= Cn

2 Vw • r • sin G • dG

On • Vw T

p -

sin (2p) (36)

Fny = Cn •

j Vw * f • Sin G • COS G • dG

^LlX^sin^P

(37)

Fu = Ci Vw • r • sin G • cos G • dG

:^-^--^sin^P (38)

_2/

Fty = Ct • Vw • r • sin G • d9

_ Ct ' Vw ' r sin(2p) (39)

Clearly

P-lsin(2p)

sin^p

(40)

Fu Fny

Expanding the trigonometric functions, this reduces to

(41)

As the next term in the expansion is -2pVl5, this approximation is of the order of 10% in error at p = 0.7

radians = 40°.

Empirical Conflrmation of Model

Experimental System

Grinding tests were performed on sintered silicon

carbide specimens, 0.0254m long x 0.00635m wide x 0.0127m high (1 in x 0.2 in x 0.5 in), with a density of

3.1 g/crn^, a Knoop hardness of 2800 kg/mm^, an

elastic modulus of 4 1 0 GPa and a fracture toughness of

4.6 MPa-m^^. Tests were performed with a surface

grinder with a static stiffhess of 1 1 • 1 0^ N/m. To permit

a range of table feed rates the original hydraulic feed

table was replaced with roller bearing slide-way with a

travel length of 0. 1 m (4 in.), a load capacity of 1607 N

(360 lb.), and a straight line accuracy of 2.8x10'^ m

(0.00011 in.). The table was driven by a stepping motor

with a resolution of 12,800 steps/revolution, a static torque of 0.55 Nm, a velocity range of 0.01 - 50 rev/s in either direction and an acceleration range of 0.01 - 1000 rev/s^. The wheel was a resin bonded diamond

wheel with a diameter, Dg, of 0. 1778 m (7 in.); a width,

b, of 0.00635 m (0.25 in.); and an abrasive concentration of 25 volume percent of 170/200 mesh, medium

grade diamond. The wheel operated at a peripheral

speed, Vg, of 24.2 m/s; it was trued with a single point

diamond using a 2.54x10'^ m (0.1 mil) depth of cut, a transverse feed rate of 4.32x10"^ m/s (0.17 in/s) and a runout of 5.08x10'^ m (0.2 mil), and was dressed with

25
C

T5

(5

50

40

30 .

20 .

10

A

Fn' (N/mm)

F,' (N/mm)

1 2

Depth of Cut a (mm)

Figure 4: Grinding force per unit widtti vs. depth

of cut, a, at constant table feed rate,

Vw = 3.38 mm/s (8 in/min).

CcO

20 -r

Q Fp' (N/mm)

Ft (N/mm)

T

,

1 "

I

1.5, 2.0 2.5 Depth of Cut a (mm)

3.0

1 I

1 n 1

1 64.5 32.2 21.5 16.1 12.9 10.7 Table Feed Rate Vw (cm/min)

Figure 6: Grinding force per unit width vs. depth of

cut, a, and table feed rate, V^, at constant material removal rate.

-w- Vw*a'b = 34.2 mm^/s (0. 125 in^/min).

35

30

25

20

15

10

5

Q Fn (N/mm)

Ff (N/mm)

—'—r—r—T—1—I—I—1—I—I—1- 0 50 100 150 200 250 300

Table Feed Rate Vw (cm/min) '

Figure 5: Grinding force per unit width vs. table

feed rate, V^, at constant depth of cut.

a = 0.5 mm (20 mil).

2.5

2.0 .

55 1.5 -

^ 1.0 C

X)

•i 0.5

0.0

A On ( N/mm2 )

( N/mm2 )

0.0 0.5 1.0 Va (VfTTfTl)

1.5 2.0

Figure 7: Grinding stress vs. depth of cut, a'^,

at constant table feed rate, Vw = 3.38 mm/s (8 in/min)

26
Tabe Feed Rate Vw (cm/mi n) (min/cm)

Figure 8: Grinding stress vs. table feed rate. Vw , at pigyre 10: Specific grinding energy, U, vs. ^w-i at constant deptti of cut constant depth of cut,

a = 0.5 mm (20 mil). - a = 0.5 mm (20 mil).

0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5 0.0 0.5 1.0 1.5 2.0 2.5 3.0

1/Va (l/Vmrfi) 1/Va (lWmm)

I Figure 9: Grinding stress vs. deptti of cut, a" at Figure 1 1: Specific grinding energy, U, vs. a '^^ at constant material removal rate, constant table feed rate,

~Zn = 34.2 mm^/s (0. 125 in^/min). ^ 3 35 ^nm/s (8 in/min).

27
COE

E

a>c

LUD)C

TJC

o

it= o

<Da.

(O

0.4 0.8 1.2

Vi" (Vrrim)

2.0

Figure 12: Specific grinding energy, U, vs. a^^^at

constant material removal rate, Zw = 34.2 mm^/s (0. 125 in^/min).

an alundum rod. A synthetic coolant was applied at a

rate of 0.0631 L/s (1 gallon/min) at 6.895x10^ Pa (100

psi) and 24°C. Three components of grinding force were measured with a dynamometer employing a piezoelectric transducer. Specific energy, in N^m/m^, was determined as

a • b

(42)

Slide-way control and data acquisition were provided

by an on-line computer.

Observations

The experimental woric on silicon carbide described

in the foregoing section provides empirical support for

the simple force model presented in this manuscript.

For example, Figure 4 shows a correlation of the basic

assumptions of the model as described in equations (6)

and (7) with empirical observation, i.e., a linear relation

between both normal and tangential force, F^^' and F^' and depth of cut, a, at constant table feed rate, V^.
Figure 5 shows the linear relation, also assumed in equations (6) and (7), between the forces, F^^' and F^' and at constant a. The correlation of grinding force

with depth of cut and table feed rate is shown in Figure

6 at constant material removal rate, = V^«a»b. The

corresponding correlation of normal and tangential

grinding stress, o„ and a^, as described in equations

(10) and (1 1), is given in Figures 7, 8 and 9. Similarly,

equation (25) predicts a linear relation between specific (grinding) energy, U, and V^'^ at constant depth

of cut, shown experimentally in Figure 10; a linear

relation between U and a '^'^at constant V^, shown in Figure 11; and a linear relation between U and a.^^ at constant Z^, shown in Figure 12.

Discussion

The foregoing model represents two simple as- sumptions that appearto be appropriate for the grinding

ofbrittle materials, i.e., that the normal stress is proportional to the normal component of the velocity of the

woricpiece relative to the wheel, and that the tangential

stress is proportional to the normal stress and thus

arises from frictional causes. The proportionality in- herent in this model also causes the resultants of the

normal and tangential forces to be applied at the same

locus.

dPnx = Cn • Vw • r • sin^ 0 • d0 (32)

dPny = On • Vw T • sin 0 • COS 0 • d0 (33)

(41)

Although the model is, to a considerable extent, justified by the correlations shown in Figures 4 through 12,

it clearly represents an extreme in grinding behavior.

Another extreme might be represented by "plowing,"

in which the tangential force may be modelled as that due to a Newtonian viscosity. In this case the previous

expression for the normal force may be retained.

and

dFu = Ci • r • cos 0 • d0

dFiy = Ct • r • sin 0 • d0

(44)

(45)

The respective force components for this model are, thus.

'nx =Cn Vw T • sin^0 • d9

, Cn • Vw T sin (2(3) (36)

Fny = Cn • Vw • T Sin 0 • COS 0 • d0

(37)

Ftx =Ci r • COS 0 • d 0

= Ct T • sin P (46)

= CtT(l - cos P) (47)

dPn = Cn • Vw • sin0 • r • d 0

but now

dR = Ct • r • d0

(1)

(43)

i.e., the tangential force component is independent of

position along the grinding arc. The horizontal (x) and

vertical (y) components of these force increments are given by

Cleariy the locus of application of the normal force is unchanged, that is,

p- lsin(2p)

{an(\,n =Pnz_= '.—2_^ (40')

Fny sin P and

(41)

shown as point a in Figure 13. In accord with its lack

29
of dependence on the angle 9, however, the effective

site of application of the tangential force is the average

value of 0 along the contact arc,

.

, Fly 1 - cosB ,„ B

ian(j)t = —= —= lan—

F« sinB 2

or

(48)

(49)

shown as point b in Figure 13. Being applied at different loci, the net normal and tangential forces are not orthogonal. The effective locus ofapplication, (]), of

the resultant force is shown as point c in Figure 13,

where Fj^ may be resolved into orthogonal components

and F^ Clearly

(50)

2 3 with the actual value being determined by the material

and machine parameters c^^, c^ and V^.

As with any body, the forces acting upon the grinding wheel may be represented as a iranslational

force, F^ and a torque, T. Since the normal force,

expressed either as F^^ or Fj^, is by definition radial and

passes through the axis of rotation, the torque applied

to the wheel is clearly given by

T = Ft • r = Ft • r

i. e., from equation (43)

Ft = Ft = Ct • r »

(3

(51)

(52)

TTie normal force, F^, at the locus of application, <j), is determined by

Fi = F^ + F^= (Fnx + Fu'f + (Fty -Fny'f (53)

Resolution of equation (53) is facilitated by expanding

the trigonometric functions of equations (36), (37),

(46) and (47) and keeping only those terms of order 3 or less:

Fnx =

Cn • Vw • r (54)

Fu = CtT. jp- 1]

(55)

(56)

(57)

Substituting and discarding terms of order p^,

Fn= LL^.{9c„W + 6cn-Ct. vw-3ct')l (58)

6

The angle (j) may be determined through balancing

force components at the effective locus of application:

Fn • sin <t) + Ft • cos ((> = Fnx + Ftx

Ft • sin ({) - Fn • cos (j) = Fty - Fny

(59)

(60)

Multiplying equation (59) by F^ and equation (60) by

Fj and adding:

.

, Fn • (Fnx + Fu) + Ft • (Fly - Fny) . ^^^x

sin()) = i i-=(t) (61)

Fn + Ft

where F^, F„,, F„y, F^-, F^ and F^ are given by equations (58), (54), (55), (52), (56) and (57), respectively.

Equation (61) is seen to approach the limits expressed

by equation (50); that is

lim

ci ^0 4) = 2 . p

3

Cn

(62)

(63)

Brach et al. (5) studied the relationship between F^ and Fj, in cases in which the ratio of depth of cut to wheel diameter was large (or the normal force on the

wheel was large compared to the tangential component). Using the trigonometric relationships between

Fh and F^, and Fj and F„, with data obtained in the sawing of granite they empirically found (f) = 0.69P in up feed grinding and (|) = 0.54^ in down feed grinding.

Ohishi (6) used ({) = 0.5P as the effective point of

application of the resultant force and and found this to be consistent with observations on the grinding of a

30
bearing steel, while Shafto (7) found empirically that P/2 < 4) < 2p/3 for a nickel alloy and <)) = 0.7p for tool

steels.In practice the value of 4) may depend on such

parameters as the sharpness of the grinding wheel, the

coolant used and the grinding mode (up feed or down

feed). Also, while for brittle materials the value should

be close to 2|3/3, for ductile materials where the grinding may be characterized by plowing the value of(() may

approach 0.5(3. However, uncertainty in the point of

resultant force application does not yield comparable

uncertainty in the values of Fj and F,,. From equations

(28) and (29)

dPi = {-Fh • sin(j) + Fv • cos4>) • dtj) (64)

and

dFn = (-Fv • sin(j) - Fn • cos())) • d<t) (65)

or, if (j) is quite small,

dFt = Fv • d(t) (66)

and

dFn = -Fh • d(l) (67)

In our experiments the depth of cut, a, was of the order

of 2.5x10"^ ~ 2.5xl0'3 m (0.01" ~ 0.1"). Using the

wheel radius, r = 0.089 m (3.5"), and equation (3) the

angle p varied from 0.076 radians to 0.24 radians.

Taking the latter case and assuming a maximum value

for ((),(}> = 2p/3 = 0.16 radians, a 10% uncertainty in <^ would yield an uncertainty in the force components of

i AFj = 0.01 6Fy and AF^^ = -0.016Fjj.

i

! Summary

The importance of using F^ instead of Fj^ in the

calculation of spindle power consumed is well known,

I especially for creep feed grinding of brittle materials.

* However, we typically are able to measure Fj^ and F^, not Fj and F^; values of the latter parameters can be

determined only if the effective locus of application, c, of the resultant force can be estimated. In our model we

assumed that F^ behaved as a frictional force, i.e., F^ a

Fj^ and determined that (J) = 2^/3, where p is the angle

subtended by the wheel/workpiece interface. This

assumption enables a simple calculation of F^ and Fj^

via equations (28) and (29) or (30) and (31), and with

the additional assumption that the normal force, F^^, is proportional to the normal velocity of the workpiece

relative to the wheel the specific grinding energy is expressed by equation (23) or (24). Certain dependences of force and energy on process parameters

such as depth of cut and table feed rate are inherent in

this model. Experiments with surface grinding of

polycrystalline silicon carbide showed a general confirmation of these relationships, thus providing ex- perimental support to the model. If proportionality

between Fj and F^^ cannot be assumed as in the case of

ductile materials, the net tangential force and net normal force no longer have the same locus of application. In this case the angle (j) is still defined and may

be determined, but its dependence on process parameters is considerably more complex.

References

1. S. Malkin, and J.E. Ritter, "Grinding Mechanisms

and Strength Degradation for Ceramics," Winter

Annual Meeting of ASME. Chicago, pp. 57-72,

1988.

2. H.K. Tonshoff, and H. Trumpold, "Evaluation of

Surface Layers of Machined Ceramics," CfRP Reports and News. Vol. 38, No. 2, pp. 699-708, 1989.

3. H.K. Tonshoff, and R. Telle, "Chip Formation and

Material Removal in Grinding of Ceramics," 4th

International Grinding Conference. SME, MR 90-

539, 1990.

4. Alfred Lammer, "Surface Topography and Strength

of Ceramic Materials after Grinding," 4th International Grinding Conference . SME, MR 90-541,

1990.

5. K. Brach, D.M. Pai, E. Ratterman, and M.C. Shaw,

"Grinding Force and Energy," J. of Eng. for Indusla, Vol. 110, pp. 25-31, 1988.

6. Susumu Ohishi, "Study of Creep Feed Grinding,"

PhD Thesis . Tokyo Metropolitan University, Japan,

1984.

7. G.R. Shafto, "Creep Feed Grinding," PhD Thesis. University of Bristol, England, 1974.

31
Nomenclature

a depth of cut A, surface area of one chip

b grinding interface width

c^, Cj characteristic parameters of wheel and

workpiece

C\C" constants C", C" constants D wheel diameter

> E energy consumption rate normal force on grinding interface

tangential force on grinding interface

F

' normal grinding force per unit width

Fj' tangent grinding force per unit width

F, horizontal force on table

h F vertical force on table

V Fjj normal resultant force on wheel

F^ tangent resultant force on wheel

F^ resultant force on wheel

h average distance between two adjacent

abrasive grains

m average abrasive particle size n rate of chip production

r radius

s length of chip side

t maximum grain depth of cut U specific grinding energy

|j. coefficient of friction between the wheel and the workpiece

V table feed rate w peripheral speed of the wheel

Z material removal rate w p

angle subtended by the grinding

interface

9 angle between 0 and P

(J) effective locus of application of F^

0^ effective locus of application of F^

<j)^ effective locus of appHcation of F^

normal stress applied on the grinding

interface

shear stress applied on the grinding

interface

Y surface tension

32
ON PERFORMANCE OF METAL-BONDED SINGLE-LAYER

DIAMOND ABRASIVE TOOL

A. K. CHATTOPADHYAY and H. E. HINTERMANN

CSEM, Neuchatel, Switzerland

The dependance of performance of metal-bonded single-layer diamond abrasive tools on the topography of working surface is reported in this paper. The behaviour of such monolayer

diamond abrasive tools continued to change irreversibly during grinding of aluminium oxide

primarily because of grit wear. The rate of increase in normal force and power was higher for

tools with a smaller grit size in comparison to that for tools with a larger grit size. On the other

hand, single-layer diamond abrasive tool having low and uniform bond level demonstrated

distinct advantage in offering steady behaviour while grinding GG40 cast iron at a high depth

of cut. The normal force and the spindle power remained virtually constant for tools with a larger grit size over a long grinding span. However, these two parameters recorded higher

values and a slow increase with cumulative infccd for tools with a smaller grit size. But, a material like cast iron could cause premature failure of such diamond tools, even having a wide and uniform grit spacing, through loading when bond level was high.

1. Introduction

The remarkable properties of diamond such as high wear resistance and excellent cutting edge

stability are being used creatively to fabricate tools having different geometries with just a single

-layer of diamond grit bonded to a metal substrate

by a galvanic or a brazed metal layer. The essential

feature of such single-layer tool is large diamond

exposure compared to coventional composite type

metal-bond wheel, which results in free-cutting

action.

However, it is contended that grit protrusion of

galvanically bonded tool may not always be

sufficient to avoid intensive bond-chip or bondwork interactions causing force and heat build-up

and loading. This problem can be further

aggravated by the presence of "nodules" which

refers to galvanic nickel build-up in the space

between the grits (1). In contrast, in brazed tools

the very nature of bonding is supposed to keep a low bond level in the space between the grit thus,

allowing large chip clearance space, efficient

coolant flow and cool cutting action (1). A

recently developed vacuum sintering process also

permits the use of high protrusion of diamond grit in the single-layer tool unlike the galvanically

bonded counterpart (2). It has been further

revealed that commercially fabricated brazed tools

do not necessarily provide high diamond

exposure as claimed by the manufacturer because

of accumulation of brazed material in the space

between the grit during brazing (3). On the other

hand, control of bond level and maintenance of its uniformity which determines the initial

topography of the working surface depends on

various constructional as well as process

parameters of fabrication (3, 4, 5). The

performance of a single-layer tool depends

largely on the initial topography of the working

surface but may be significantly influenced by the

work material characteristics and cutting

conditions.

The principal objective of the the present

investigation was to evaluate the dependence of

performance of brazed metal-bonded monolayer

diamond abrasive tools on the topography of its working surface in grinding of an oxide ceramic

and cast iron.

2. Test Procedure

Brazed bonded monolayer diamond grinding

wheels with a cup shape were used on a tool and

cutter grinder to conduct experiment in face

grinding mode. Diamond grit were of two sizes

viz. D251 (250/212 ^m mesh) and D 126 (125/106

|j,m mesh). The diameter and width of working

surface of the wheel were 75 mm and 5 mm,

respectively.

33
Table 1 Different types of diamond abrasive tools

Tool Grit size Grit protrusion, |Jin Qiaracteristic feature

Type A D251 185-205 Metal-bonded single-layer

Type B D126 80-95 Metal-bonded single-layer

TypeC D126 50-60 Metal-bonded single-layer

TypeD D126 30-40 Metal-bonded composite type

200 jim

Fig. 1 Surface topography of diamond abrasive

tool of Type A.

Fig. 3 Surafce topography of diamond abrasive

tool of Type C.

si

Fig. 2 Surface topography of diamond abrasive

tool of Type B.

Fig. 4 Surface topography of diamond abrasive

tool of Type D.

34
40

oT 30

ui_o

t 20

COE

I

1 0

0

I 100

o

Q.

O

o 50

'a. V)

1 0 20 30

Infeed, |im

40 50

Fig. 5 Normal force and spindle power versus infeed for abrasive tools of Type A, B and C in dry grinding

of oxide ceramic.

The normal force and the power drawn by the

spindle during grinding were recorded, using a piezoelectric dynamometer and a wattmeter

connected to a two channel recorder. SEM

observations were made on the wheel surface

before and after grinding tests. Representative chip

samples of the work materials were collected

during grinding and also observed under SEM.

The work materials were as follows : sintered

aluminium oxide (hardness : 1200 HV| q, density

: 3.925 gm/cm^) and GG40 cast iron (hardness : 184 HV| q)

The workpieces were round bars of 12 mm

diameter, the end face of which was ground. The

grinding conditions were as follows : wheel speed,

30 m/s; table speed, 2 m/min; infeed, 10, 20, 30, 40

|im; environment, dry and wet (5% soluble oil).

3. Results

Designation of the diamond wheels used in the

present investigation along with their grit size and

grit protrusion are indicated in Table 1. Surface

topography of a brazed monolayer diamond tool

(designated as Type A) with grit size D251 is shown in Fig. 1. The grit can be seen to be

effectively wetted by the bonding matrix. The

relative height between the tip of grains and mean

bond base was found to be in the order of 185-

205 iim. Figure 2 visualizes a brazed single-layer

diamond tool (designated as Type B) with grit size

D126. In this case the grit protrusion was found to be in the order of 80-95 |im. Figure 3 shows the

surface of another brazed monolayer diamond

wheel (designated as Type C) which was very

similar to the wheel of Type B in terms of grit size,

distribution pattern and density but differing in bond level. It can be easily observed that bond

level was higher than that used in the wheel of

Type B. This tool showed grit protrusion of the

order of 50-60 |im.

To understand clearly the role bond level or

crystal exposure, a composite type metal bonded

diamond wheel (designated as Type D) with grit

size D126 was also included. Figure 4 reveals the

surface of the same obtained after truing and

dressing. This was done to expose the crystals

which formed a composite structure with the metal

matrix. Numerous cutting points were generated on the crystals after dressing. Truing and dressing

of the cup metal bonded diamond wheel were

performed according to the recommendation of

the wheel manufacturer (6).

35
0 10 20 30 40 50

Infeed, |im

Fig. 6 Normal force and spindle power versus infeed for abrasive tools of Type A, B and C in dry grinding

of cast iron.

Figure 5 shows that the normal force and spindle

power were always smaller for wheel of Type A in comparison to wheels of Type B and C. Diamond

wheel of Type B with a larger grit protrusion,

however, showed lower magnitude of force and

power in comparison to similar tool (Type C) but

with smaller grit protrusion. This difference was

more pronounced at higher infeed values.

Figure 6 shows the variation of normal force and

spindle power with infeed during dry grinding of

cast iron by diamond tools of Type A, B and C. The tool of Type A again developed lower

normal force and required lower spindle power in

comparison to tools of type B and C. The fact which draws attention was disproportionate

increase of normal force and spindle power at relatively high infeed values for the Type C tool. Comparing Figures 5 and 6 one can arrive at the

conclusion that spindle absorbed more power in

grinding cast iron than in aluminium oxide

ceramic.

The wheel of Type D was ineffective in dry

grinding of cast iron. This is quite evident from

Fig. 7. With infeed of 10 |im, the normal force

and spindle power was as high as 23N and 138 W,

respectively. At 20 |im infeed, these two

parameters escalated to 95N and 540W just after

34 grinding passes. Figures 6 and 7 together

indicate that during wet grinding of cast iron, tool of Type D exhibited higher force and power at aU

infeeds in comparison that exhibited by the

wheel of Type C even under dry grinding. Again,

this difference was more evident at higher infeed

values. Figure 7 further reveals that the force and

power absorbed during dry grinding of ceramic

by tool of Type D were far below that absorbed

during dry or wet grinding of cast iron by the same tool. Figure 7 also shows that the normal

force and spindle power could be effectively

reduced for tool of Type D in grinding ceramic

through application of cutting fluid. Such

reduction of force and power was of particular

relevance when grinding is conducted at high

infeed.

To understand the difference in performance

between wheels of Type A and B in grinding

oxide ceramic and cast iron, tests were conducted

for a long duration with 40 |im infeed.

Figure 8 shows that the force and spindle power

increased with cumulative infeed and rate of

increase was higher for tool of Type B (grit size

D126) in comparison to the tool of Type A (grit

size D251) and gap between the force and power

curves continued to widen as is evident from the

36
0

Ceramic (Dry)

Ceramic (Wet)

Cast Iron (Dry)

Cast Iron (Wet)

1 0 20 30 40

Infeed, |iim

Fig. 7 Normal force and spindle power versus infeed for abrasive tool of Type D.

50

same figure. Figure 9 on the other hand indicates

that normal force and power drawn by the spindle

during grinding of cast iron by tool of Type A

remained virtually constant with cumulative

infeed. The same figure further reveals that these

two parameters recorded higher values and slow

increase with cumulative infeed for tool of Type

B.

Figures 10 (a) and (b) show the types of chips

produced during grinding of oxide ceramic and

cast iron. Ceramic was ground mainly through

formation of fine powder-like chips as a result of

crushing. Cast iron, though formed fragmental

chips, underwent some degree of plastic

deformation before fracture as can be seen from

Fig. 10 (b). The size of cast iron chips was much

higher than those produced in grinding of

ceramic.

Figure 1 1 shows topography of wheel of Type A

after removing 76 mm from the ceramic bar. The

grit appear to be mostly worn flat. Grit fracture or

grit pull out was also not evident. The powder-hke

chips did not cause any wheel loading. Figure

12 reveals a different topography of the same

wheel of Type A after removing 76 mm from the cast iron bar. The active grit appear to have a fractured surface in total contrast to what can be seen in Fig. 11. However, there was no sign of

macrofracture of grit which takes place at the

bond level. Similarly grit pull out or wheel

loading was also not evident. Wheel loading did not occur in wheel of Type B while grinding cast iron but Fig. 13 clearly establishes the limitation

of Type C wheel while grinding cast iron only

because of loading. Micrographs of figure 13

clearly reveals heavily deformed chips of cast iron

adhered to the surface of Type C wheel during

grinding at 40 |j.m infeed.

4. Discussions

The difference in requirments of normal force

and spindle power for monolayer diamond tools

with different grit sizes may be related to increased number of cutting points of the tools

with smaller grit sizes. But variation in bond level in tools with two different grit sizes could also

influence performance. Single-layer diamond

37
80

60

40

o

o

u.

E

o 20

0

200

Tool : Type A ^ Tool : Type B

10 20 30 40 50

Cumulative Infeed, mm

60 70 80

Fig. 8 Normal force and spindle power versus cumulative infeed for abrasive tools of Type A and B m dry

grinding of oxide ceramic.

50

2 40

<3i O

0 30

1 20

o^ 1 0

0

^- 200

0 150

1 100

c

^ 5 0

Tool : Type A

Tool : Type B

10 20 30 40 50

Cumulative Infeed, mm

60 70 80

Fig. 9 Normal force and spindle power versus cumulative infeed for abrasive tools of Type A and B in dry

grinding of cast iron.

38
(b)

Fig. 10 Types of chip produced during grinding

of (a) oxide ceramic and (b) cast iron.

tools with same grit size and similar density of

distribution performed differently because of a wide variation in bond-chip and bond-work

interactions. A high bond level could have an

overidding influence.

Performance of an abrasive wheel also depends

on the workmaterial characteristics. The spindle

power which directly reflects tangential force was

rather high for cast iron and more than for

ceramic. As mentioned earlier ceramic was

ground mainly through crushing where as cast

iron underwent some degree of plastic

| .

" ' 200 ^im

Fig. 1 1 Surface topography of abrasive tool of

Type A after dry grinding of oxide ceramic.

Fig. 12 surface topography of abrasive tool of

Type A after dry grinding of cast iron.

deformation before formation of fragmented

chips. The plastic deformation mainly accounted

for the differences in spindle power and

tangential force between ceramic and cast iron

grinding. Apart from this severity of bond-chip

interactions, bond-work interactions could

influence the response of the wheel during

grinding of ceramic and cast iron. Chip

accommodation was easier with powder-like chips

of ceramic than with relatively "large" fragmented

39
13 SEM micrographs showing loading of abrasive tool of Type C in dry grinding of

40
chips of cast iron under identical conditions of

grinding. Similar observations (7,8) were also

made during grinding of various materials, which

varied widely in terms of ductility.

In ceramic grinding attritious wear was the major

grit wear caused by progressive grit abrasion by

Sie oxide ceramic resulting in rounding-off of the

cutting points and development of wear flats on a large proportion of grit. Such change in grit

geometry led to an increase in normal force

developed during penetration of the work

material and the energy required to remove

material during grinding. Flatening of the cutting

points could also increase friction at the gritworkpiece interface. In cast iron grinding, grit

fracture was the major cause of wear. Such

controlled brake down of the grit kept the wheel

rather sharp and normal force and spindle power

were virtually constant for tools with grit size

D251 and increased very slowly for tool with grit

size D126. Diamond abrasive tools with grit size

D126 and high bond level experienced a premature failure in grinding cast iron at high

infeed because of chip accommodation problems.

Because of high bond level, cast iron chips were

squeezed in the space between the bond and the

workpiece and adhered to the tool surface. Such

unfavourable change in wheel topography caused

escalation of force and power to very high values.

5. Conclusions

1. Single-layer diamond abrasive tool did not

suffer from grit pull out during grinding of oxide

ceramic. Also, there was no evidence of chip

accommodation problem with fine powder-like

chips of oxide ceramic even at high infeed.

However, the behaviour of such tool continued to change with passage of grinding time. This

occured because of attritious wear of the diamond

grit, caused by aluminium oxide resulting in

rounding off of sharp points and generation of wear flats.

6. References

1. L. M. Peterman, "Diamond Tooling in NonMetallics", Proc. Superabrasive 85 Conference , Chicago, IlHnois, pp. 1-12, 1985.

2. E. D. Kizikov, "Vaccum Technology for

Diamond Toolmaking", Industrial Diamond

Review . Vol. 1, pp. 20-23, 1991.

3. A. K. Chattopadhyay, L. Chollet, H. E. Hintermann, "On Performance of Brazed Bonded

Monolayer Diamond Grinding Wheel", Annals

CIRP. Vol. 40/1, pp. 347-350, 1991.

4. A. K. Chattopadhyay, L. Chollet H. E. Hintermann, "Experimental Investigation on

Induction Brazing Of Diamond with Ni-Cr

Hardfacing Alloy under Argon Atmosphere",

Journal of Materials Science , Vol. 26 pp. 5093-

5100, 1991.

5. H. E. Hintermann, A. K. Chattopadhyay, "New

Generation Superabrasive Tool with Monolayer

Configuration", Diamond and Related Materials,

Vol. 1, pp. 1131-1143, 1992.

6. Tyrolit (Austria), Technical Information on Diamond and CBN.

7. D. M. Pai, E. Ratterman, M. C. Shaw, "Grinding

Swarf', Wear, Vol. 131, pp. 329-339, 1989.

8. A. K. Chattopadhyay, L. Chollet, H. E. Hintermann, "On Performance of Chemically

Bonded Single Layer CBN Grinding Wheel",

Annals CIRP, Vol. 39/1, pp. 309-312, 1990.

2. The rate of increase of normal force and

spindle power was higher for tool with smaller grit

size (D126) than that for tool with larger grit size

(D251).

3. The bond level or crystal exposure has a predominant influence on the performance of an

abrasive tool in grinding those materials which

tend to stick to the wheel surface.

4. Comparative evaluation of single-layer brazed

diamond wheel with coventional metal bond

wheel may be made in grinding those materials

which may not create a critical chip storage

problem but, may cause irreversible change in the

wheel behaviour through rapid wear of the singlelayer diamond crystals.

41

A SYSTEMS APPROACH FOR GRINDING OF CERAMICS

K. SUBRAMANIAN, P. D. REDINGTON, and S. RAMANATH

Norton Company, Worcester, MA

Every manufacturing process involves four key elements: machine

tool, work material, processing tool, and operational factors. These

four input elements are related to output of the manufacturing process

through process interactions. This systems description is also

applicable for the grinding process. Through careful understanding of

the input elements, their role in process interactions and the

consequent output behavior, significant progress has been made in the

grinding of ceramics. Several of the features of input elements, i.e.,

machine tool, diaimond wheels, work material, and operational factors

influence the process interactions, which result in plastic deformation

and/or brittle fracture. The grinding system for cersumics is optimized

when brittle fracture is minimized and plastic deformation is m^aximized

by simultaneously controlling several of the input elements.

Successful use of this systems approach for grinding of ceramics has

resulted in production viable grinding of ceramics. A description of

this systems approach is followed by several examples. These technical

results will be useful if they are accompanied by recognition of key

cost elements and their control. A total cost approach to this end is

also described with examples.

Introduction ; Every manufacturing process involves four key elements: machine tool,

work material, processing tool and

operational factors. These four input

elements are related to output of the

manufacturing process through process

interactions. This systems description

of the grinding process is shown in

Figure 1. Each of the input elements

are composed of large number of variables. It is nearly impossible to

characterize each of these variables to

the last detail. Further, such approach

is expensive and time consuming. However, it is critical that the role of each

input variable is assessed in terms of

the process interactions: cutting,

ploughing, rubbing/sliding, at least

qualitatively. Quantitative assessment

of such microscopic interactions are difficult, but should be pursued whenever possible, using macroscopic or measurable process interactions such as forces and energy. These process interactions in turn influence the output of

the grinding system, which may be described in technical terms or as system

output, when economic factors are also

included. Thus every grinding operation

in a systems approach becomes a process

of simultaneously manipulating selected

few input variables in order to maximize

the cutting component of the process

interactions and minimize the tribological components: ploughing and rubbing/-

sliding in order to achieve the desired

output results. Through careful understanding of the input elements, their

role in process interactions and the

consequent output behavior, significant

progress has been made in the grinding of

ceramics. After careful understanding of

the technical elements of the grinding

process, it is also necessary to apply

them systematically through attention to

a few strategic elements of the total

cost. This paper attempts to highlight

some of the results achieved in these

areas.

43
3

Q.

i

</i UiO

o

3

a.3

o

MACHINE TOOL

PROCESSING TOOL

(GRINDING WHEEL)

WORK MATERIAL

OPERATIONAL FACTORS

MACROSCOPIC

PROCESS

VARIABLES

(FORCES,E?vJERGY)

PROCESS

OUTPUT

<WHEEL WEAR.

CHIP PRODUCED,

COOLANT EFFECTS,

RETAINED STR.ETC.)

MICROSCOPIC

PROCESS

INTERACTIONS

(CUTTING.PLOWING,

RUBBING/SLIDING)

SYSTEM

OUTPUT

<PART QUALITY,

TOTAL COST,

PRODUCmON

ECONOMICS,ETC.)

Grindability of Cereunio V8. Metala ; Ceramics include a wide range of

materials with their thermophysical

properties dependent on the type, com- position, microstructure, and the processing methods. In this regard, ce- ramics are a family of materials, just

as we recognize metals as a family of

materials. Once we recognize this

commonality, many of the principles

applicable to abrasive finishing of

metals—an industry with over 100 years

of history—become applicable to the

abrasive machining of ceramics. However, for the saume properties considered

(thermal conductivity. Young's modulus,

high temperature resistance, fracture

toughness, etc.), metals and ceramics

fall into distinctly different ranges.

If we recognize these differences, then

the principles of abrasive machining of

metals can be modified for successful

abrasive machining of ceramics.

The strength of cereimic material

varies widely, depending on the material

chosen. Even for a given material, like

silicon nitride, for exsimple, the

strength depends on the sintering aids

used and the sintering methods (pressureless sintering, hot pressing, hot

isostatic pressing (HIP), etc.) applied.

In this regard, the ceramic materials

are analogous to metals in which the

composition, microstructure, and

strength influence their grindability.

Table 1 compares typical properties of

selected ceramic materials with those of

metals and a typical polymer.

In general, ceramic materials have

higher stiffnesses (Young's modulus, E)

than metals. This would imply that

elastic deformation of ceramic materials

during grinding would be less than

metals for the same applied grinding

forces. This is indeed true, and as a result, ceramic materials can be machined to closer tolerances, more precise geometry, superior finish, and

better flatnesses and parallelism than

is possible for metals. This principle

is the basis for choosing ceramic materials over metals in precision instrument parts, machine tool beds or ways,

as well as for gage blocks.

Ceramics are generally more chemically stable than metals. Hence the

burn sometimes observed on metals during

grinding is rarely observed on ceramic

materials. Furthermore, the hot hardness and recovery hardness of metals are

FIGURE 1

A SYSTEMS APPROACH FOR

PRODUCTION GRINDING

much lower than those of ceramips. In

other words, any large scale thermal

softening that aids in the grinding of

metals may be rarely counted on in the

grinding of ceramics.

Thermal conductivity of ceramic

materials varies widely. This property

determines the ability to conduct heat

away from the grinding zone. Poor thermal conductivity in metals such as Titanixim and Inconel can lead to high temperatures and great difficulty in grinding.

Similarly, it is more difficult to grind

poorly conducting ceramics and the problem becomes worse when they are also poor

in thermal shock resistance.

One characteristic that significantly distinguishes ceramics from metals

is their low fracture toughness. With

all the above similarities between metals

and cereimics in mind, it would appear

that it is possible to achieve successful

grinding of ceramics if the generation

and propagation of cracks during the

grinding process can be minimized.

44
Mechanisms of Material Removal in the

Grinding of Ceramics ; Recognizing the low fracture

toughness of ceramics, models have been

proposed which describe initiation and

propagation of microcracks during the

grinding of ceramics. These models are generally described as "Indentation

Fracture" models. ^'''^ It has also been

proposed that ductile deformation of

ceramics takes place under extremely

small depth of cut. This model described as "ductile regime grinding" as- sumes occurrence of plastic deformation

only and seeks to achieve this condition

using extremely low material removal

rates. Under practical production

grinding conditions, it would appear

that both these models operate, dependent on the force per grain, volume

removed per abrasive particle, size of

the abrasive particle, etc.'^ The description of the occurrence of both

brittle fracture and plastic deformation

and their optimization to achieve desired production rate and part quality

is shown in Figure 2 . Several of the

features of input elements, i.e., machine tool, diamond wheels, work material, and operational factors influence

these process interactions, i.e., brittle fracture vs. plastic deformation.

The ceramic grinding system description

and the input/output relations are shown

in Figure 3. Referring to Fig. 2, it

would appear that the mechanisms of

material removal in the grinding of

ceramic material are associated with

both plastic deformation and brittle

fracture. It may be nearly impossible

to isolate one mechanism or the other

unless extreme grinding conditions are

chosen. These are identified as "coarse

grinding" (where brittle fracture mechanism dominates) and "fine or ductile

regime grinding" (where plastic deformation mechanism dominates). The conse- quence of such choices are low strength

and poor reliability in the case of

"coarse grinding" and extremely small

material removal rate, high forces and

specific energy in the case of "fine or ductile regime grinding" . Referring to

the process interactions described in

Figure 1, it may be envisioned that

plastic deformation is desirable to in- crease the cutting component. However,

excessive deformation may result in

inefficient cutting process accompanied

by excessive ploughing and rubbing/sliding components. While brittle

fracture may loe a means of generating

surfaces at the lowest forces and enerTABLE 1

PHYSICAL AND MECHANICAL PROPERTIES

Coefficient

of

Thermal

MATERIAL Expansion

gp-^/k)

CERAMICS

Aluminum

Tensile

strength

(MPa^

OF SELECTED CERAMICS, METALS AND

A TYPICAL POLYMER

Melting/

Decomposition Thermal

Tempe- ConducMATERIAL Density rature tivity

(a/cm^) ( °C) W/m.K

CERAMICS

Aluminum

oxide 3.4-4.0 2050 27

Silicon

Carbide 3.0-3.2 2500 63-155

Silicon

nitride 3.2-3.5 1900 9-30

Zirconium

oxide 5.8 2500 2 Ferrite 5.0-6.0 2000 8

Fused

Quartz 2.2 1670 1.4

METALS

Inconel

718 8.1 1400 12.4

Tool steel 7.84 1500 52

52100

bearing 7.85 1500 30-40

steel

Aluminum 2.2 660 204

POLYMER

Polyethylene 1.1 120 0.32

Modulus

of

Elasticity

oxide 8 205-550 385

Silicon

Carbide 4-5 415-550 420

Silicon

nitride 3 620-1100 305

Zirconium

oxide 9-10 965-1380 140

Ferrite 7.5 140-170 205

Fused

Quartz 0.55 110 77

METALS

Inconel

718 7 1240-1405 215

Tool steel 7 1405-2000 210

52100

bearing

steel 11 415-550 210

Aluminum 13 275-550 70

POLYMER

Polyethylene 167 7-20 3.5

45
gy, it is certainly not desirable where

highest strength and reliability are required. Thus optimizing the ceramic

grinding system becomes a process of

maximizing the cutting component through

plastic deformation, while minimizing

the grinding forces and energy input.

This can be accomplished through a careful and simultaneous selection of

input parameters identified in Figures

2 and 3. While this approach is common

with metal grinding, the additional

factor in cersutnics grinding will be the

control of force per grain. Under such

process interactions, selection of

suitable grinding cycle including rough

and finish grinding steps is a matter of

trade off between process economics and

surface characteristics desired. There

are several examples of the successful

use of this systems approach for grinding of ceramics. Few of these results

are described here.

Applications of the Systems Approach for

Grinding of Ceramics : Production Viable Grinding ; Table 2 describes the results towards production

viable grinding of ceramics. Material

removal rates (MRR) of l-30mm^/mm/sec

(.1-3 in^/min/in) are now routinely

achieved; the higher MRR are used for

rough grinding processes while the lower

values are used for finish grinding.®

Such MRR, if they are achieved using

coarse abrasive particles ( 100/j or larger), predominantly result in brittle

fracture with poor surface c[uality (low

strength and poor surface finish)

Finer abrasive grains can be successfully used to achieve the high MRR' mentioned earlier, if the machine tool has

rigidity and power to transmit the

forces and energy required. This may, on occasion, also require unique wheel

design to overcome the limitations of

part geometry and fixturing.

Mirror Finish Grinding ; Plastic deformation dominated grinding of ceramics

using fine abrasive particles often

permit controlled finishing of ceramic

surfaces. As a result, extremely smooth

surfaces have been produced in various

grinding modes (e.g.) surface grinding,

cylindrical grinding, creep feed grinding. For example, mirror finish grinding of ceramics in external cylindrical

grinding has resulted in finish of the

order of 7nm (.3l/ inches) in hot pressed

Silicon Nitride*^'. Such superior finish

is often dependent on the work material

grain size, uniformity and strength.

"7

FIINISH GRINDING (PROPOSED)

ROUGH

GRINDING (PROPOSED)

COARSE

GRINDING

(PRESENT)

INPUT M/C TOOL

LOW VIBRATION

PROCESSING TOOL

SMALL GRAIN SIZE SMALL GRAIN TIP RADIUS

WORK MATERIAL SMALL GRAIN SIZE HIGH FRACTURE

TOUGHNESS

OPERATIONAL FACTORS

LARGE CUTTING VELOCITY

SMALL DEPTH OF CUT

SMALL CHIP THICKNESS

LARGE

LARGE

LARGE

LARGE

LOW

SMALL

LARGE

LARGE

PROCESS VARIABLES

I SMALL FORCE PER GRAIN LARGE

I LARGE SPECIFIC ENERGY SMALL

f LARGE GRINDING FORCE SMALL

OUT PUT

HIGH

FINE GOOD

LOW

PART STRENGTH LOW

SURFACE FINISH POOR

EDGE CHIPPAGE POOR

MATERIAL REMOVAL HIGH RATE

FIGURE 2 GRINDING OF CERAMICS

(GOVERNING MECHANISMS

AND CONTROL VARIABLES)

Pull out or localized removal of work

material while mirror finish grinding is an evidence of inhomogeneity in grinding

forces caused by limitations in one of

the four input elements.

46
Grinding from Solid ; When the grinding

system is integrated as described earlier, with attention to process interactions that minimize the force per grain,

high MRR grinding conditions can be used

to grind components of complex geometry

from simple solid shapes. Figure 4

shows several components that were

ground from simple shaped raw material

blanks.

Figure 4a shows four fluted end mill

ground from a solid cylindrical rod of

16mm diameter. The entire end mill was finished in 35 minutes, which compares

well with 5 to 15 minutes required for

grinding steel end mills from cylindrical blanks. The details of grinding

conditions are shown in Figure 4a.

Figure 4b shows a spring produced by

helical slot grinding on a hollow tube.

Figure 4c shows a rod plunge ground from

a nominal size of 10 to 15mm to less

than 1mm diameter. This was carried out

through a plunge cylindrical grinding

process using a vitrified bonded diamond

wheel. The details of grinding conditions are shown in Figure 4c. Such thin

sections are difficult to produce in the

grinding of metals. However, ceramic

materials due to their higher stiffness

do not deflect as much and hence, can be

machined to achieve extremely thin sec- tions or small cross-sections (if brittle fracture during grinding is minimized) through the systems approach.

Figure 4d shows a range of parts, most

of which are produced from simple shaped

unfinished blanks.

Microgrindinq of Ceramics ; The principles of controlled chip formation in

finishing of ceramics can be extended to

emulate precision finishing processes

used in metal components. Hyperlap is

a fixed abrasive finishing process in

which grinding wheels operating at low

velocities replace loose aibrasive lapping. This results in higher MRR achievable in the grinding processes while

maintaining flatness and parallelism

achievable in lapping processes. This

technology has been extended in the

finishing of ceramics using a process

called Microgrinding. Sometimes this

application is also described as Flat

Honing^'°'. Table 3 shows some of the

results achieved in microgrinding of

ceramics on a variety of work materials.

In addition to the obvious productivity

and quality advantages, this process is

environmentally benign as simple coolC/D CO

LU8

DC

Z)

D.

[-Z)O

MACHINE TOOL

MULTI AXIS CNC GRINDER WITH

HIGH RIGIDITY AND HP

PROCESSING TOOL

FINE GRIT DIAMOND

,.,^r,,^ ..*-r.-r„», GRINDING WHEEL WORK MATERIAL SIMPLE SHAPED CERAMIC BLANK

OPERATIONAL FACTORS

GRINDING CYCLE DESIGN

PART HOLDING/FIXTURING

CREEP FEED GRINDING

ON MACHINE TRUING

AND DRESSING

COOLANT SYSTEM/FILTRATION

CHIP THICKNESS CONTROL

DECREASE

FORCE PER GRAIN; POWER

TOTAL FORCE

AND

^

FORCE VARIATIONS

{MACROSCOPIC)

MAXIMISE THE

"CUTTING" PROCESS

(THROUGH PLASTIC

DEFORMATION VS

.BRITTLE FRACTURE)

(MICROSCOPIC)

HIGH MRR

COMPLEX

PARTGEOMETRY

MINIMUM DAMAGE

GOOD SURFACE

FINISH MAXIMUM

STRENGTH

(PROCESS OUTPUT)

RELIABILITY LOW TOTAL

^ COST/PART

HIGH YIELD COST EFFECTIVENESS

COMMERCIAL

VIABILITY

(SYSTEM OUTPUT)

FIGURE 3

SYSTEMS DESCRIPTION FOR THE

PRECISION GRINDING OF CERAMICS

TABLE 2 PRODUCTION VIABLE CERAMICS GRINDING

Past

Material 7

Removal

Rate

(mm^/sec/mm) .1-1.0

(cu. in/min. /in. ) 0.01-0.1

Surface Finish

(/jm) .750-1.27 . (microinch) 30-50

Now

1-30

0.1-3

0075-.125

.3-5

Cycle Time:

6 Hours —> 1 Hour

Hrs —> 35 min.

48 Min —> less than 10 min.

47
ants or water can be used without the

need for lapping fluids. Also oil bearing abrasive slurry need not be disposed

of in microgrinding processes as none is

required

TABLE 3 MICROGRINDING/FLAT HONING CASE STUDIES

FLAT- THICKSURFACE NESS NESS

CASE FINISH TOLE- TOLENO. MATERIAL WHEEL MRR Ra RANCES RANCES

(mm'/sec) (jim) (/xm)

1 SAPPHIRE 320 GRIT

38Xia' .50 5-10 2.5

VITRIFIED BOND

Result: 5 TO 10 TIMES FASTER COMPARED TO

LAPPING WITH CONVENTIONAL ABRASIVES

SILICON 320 55 .20 2.5 2.5 NITRIDE GRTT

VITRIFIED BOND

Result: 5 TIMES FASTER COMPARED TO LAPPING WITH

CONVENTIONAL ABRASIVES

3 SILICON 320 27 .25 2.5 2.5 CARBIDE GRIT

METAL BOND

Result : 3 TIMES FASTER COMPARED TO LAPPING WITH

CONVENTIONAL ABRASIVES

4 ALUMINA/ 600 2.2 Tie GRTT

VITRIFIED BOND

.125 NO CHIPPING

AT 200X

FERRTTE 320 GRTT

METAL BOND

55 .40 2.5 2.5

Every ceramic component starts with

the ceramic raw material, which in most

cases, will be in powder form. The

amount of ceramic material used to produce one component is termed as the material cost. This powder is often processed in one or more ways to achieve the

unfinished blank. These processes may be

hot pressing, hot isostatic pressing,

sintering, etc. We shall call these

"Ceramic Processing" and the cost associated with this as the cereunic processing

cost. The unfinished blank may then be

finished often using grinding or other

abrasive processes such as ultrasonic

machining. The cost associated with

this, we will call as the "Finishing

Cost". Thus the total cost of a ceramic

component (Fig. 5) will' be:

Total Cost = Finishing Cost -I- Ceramic Processing Cost + Material Cost (1)

In reality, these elements of costs

are not independent. "Yield" is defined

as the fraction of the total number of

parts produced to arrive at one good or usable component. Yield in every processing step influences the total cost.

The yield becomes increasingly critical

as we move from raw material to finished

part. When a part is rejected at the

finishing stage, the total cost of a good

part produced increases not only due to

the loss in the finishing, but also due

to all other costs incurred up to this

time. This compounding effect is shown

in the equation below:

Finishing Cost Ceramic Processing Cost

Total Cost = +

(Y.Fin.) (Y.Fin.) X (Y.Proc.)

Note: Work Material : See Table Machine : Double Sided Lapping/Grinding Machine Wheel : SOOmra Dia; See Table for grit size and bond type OPERATIONAL PARAMFTERS: Wheel Speed : 1 to 5m/sec at O.D. Work Speed : 0.1 to 5m/sec

Coolant: Mineral Oil (or) Water Based Coolants Total Cost Approach for Grinding of

Ceramics

Having briefly looked at the technical results of the systems approach

for ceramic grinding, let us now briefly

review some economic aspects. There are several recent publications which address this subject Our objective

here is to highlight the key cost con- trolling elements and how to attack them

strategically.

Material Cost +

(Y.Fin.) X (Y.Proc.) (2)

where,

Y.Fin. = Yield in the finishing pro

cess

Y.Proc. = Yield in the ceramic pro

cessing operation.

Finishing Cost consists of the cost of

machining, inspection, set up or fixturing. Thus:

Finishing Cost = Machining Cost -I- Inspection Cost -t- Set Up or Fixturing Cost (3)

as illustrated in Fig. 6.

48
FIGURE 4a

CERAMIC END MILLS FINISHED BY GRINDING

OF SOLID CYLINDRICAL BLANKS

Work Material ; Hipped Silicon Nitride

Machine ; 5 axis CNC Grinder

Wheel ; SD320R100 BXL6145

OPERATIONAL PARAMETERS ; Wheel Speed ; 35 m/sec

Work Speed ; 50 mm/min

Depth of Cut ; Roughing; 2.5mm

Finishing; 0.25mm

Coolant ; 100% oil applied at 2001/min;

.828 MPa

FIGURE 4b

CERT^IC SPRINGS GROUND FROM HOLLOW

TUBES USING SUITABLE GRINDING METHODS

Work Material ; Zirconia - Machine ; CNC Grinder

Wheel ; SD180R100 BXL6145

OPERATIONAL PARAMETERS ; Wheel Speed ; 35 m/sec

Work Speed ; .6 mm/min

Depth of Cut ; .37mm

Coolant ; 100% oil applied at 2001/min;

.828 MPa

FIGURE 4c

THIN DIAMETER CERAMIC RODS GENERATED

FROM STANDARD CYLINDRICAL RODS BY

PLUNGE GRINDING

Work Material ; See Figure

Machine; Okuma CNC Grinder

Wheel ; SD3201150 VX223C

OPERATIONAL PARAMETERS ; Wheel Speed ; 45 m/sec

Work Speed ; 5.0 m/min at O.D.

Depth of Cut ; .05 mm/rev

Coolant ; Water Soluble Oil

FIGURE 4d

A SELECTION OF COMPLEX GEOMETRY

CERAMIC PARTS WITH SURFACES GENERATED

BY DIAMOND GRINDING PROCESSES

49
Let us now focus on machining

costs, and it consists of components due

to machine tool, labor, abrasive tools

and overhead costs. Thus:

Machining Cost = Machine Tool Co§t + Labor Cost + Overhead Cost + Tool Cost (4) The Machine Tool, Labor and Overhead

Costs are directly linked to the time it

takes to produce one part, called the

cycle time. Thus:

Machine Tool Cost (M.T.C.)=Machine Rate ($/hr) x Cycle Time (5)

Labor Cost (L.C.) = Labor Rate ($/hr) x Cycle Time (6)

Overhead Cost (OHQ = Burden Rate ($/hr) x Cycle Tuor (7)

Tool Cost = Abrasive Wheel Cost/No. of Parts Machined per Wheel (8) Since cycle time has a profound

impact on the finishing cost, let us look into its components briefly. It is

defined as the time consumed in finishing per part. Thus:

Cycle Time =Total Time Consumed (T)/No. of Parte machined

in Time T. (N) (9)

Total Time (T) = Set up time+Inq>ection Time +Grinding Time (10)

Grinding Time = Wheel/Work Contact Time + Wheel

Preparation Time

Wheel preparation time is required for

truing and dressing. Part load and

unload time is included in the set-up

time.

Grinding time directly influences

the number of parts machined per wheel

mentioned in equation 8. Thus cycle

time indirectly influences the tool cost

as well.

The above equations describe the

elements of costs and their inter-relationship in the manufacture of ceramic

components. Attempts to minimize each

cost element individually is tedious and

often futile. It is imperative to

recognize the key cost elements and

attack them strategically. This approach cannot be solely based on economics, but a rather careful combination

of economic and technical factors will

be required. In the following, few of our experiences are described:

yield in the finishing process,

described in equation 2 would appear to

be a key cost controlling element.

CERAMIC

POWDER

ffiATERlAL^

CERAMIC

FABRICATING

PROCESS

CERAMIC

COMPONENT

TOTAL COST PER PART=

MATERIAL COST+

CERAMIC PROCESS COST-^

FINISHING COST

FIGURE 5 COMPONENTS OF COST FOR A CERAMIC PART

TABLE 4 INFLITENCE OF CERAMIC HtOCESS AND

FINISHING PROCESS YIELD ON THE TOTAL COST

Fini- Ceramic shing FmiProcess Process Material Process shing Total CASE Yield Yield Cost Cost Cost Cost L%1 £%I

Hypo- 100 100 5 5 10 20

thetical

Actual

(a) 20 100 25 25 10 60

(b) 100 20 25 25 50 100

(c) 20 20 125 125 50 300

Step 1 100 100 6 6 12 24

Step 2 100 100 7 4 4 15

Note : All coste are in relative unite. As an example in the hypothetical

case, the totel material cost (20 unite) consiste of 25 % in material cost

(5 unite), 25% in process cost (5 unite) and 50% in finishing cost (10

unite). Subsequent examples are illustrated as modifications using the hypothetical case as the reference or base. The above table is based on

our experience in a wide range of advanced ceramic componente

grinding.

Step 1 : Yield in the finishing process is optimized through systems approach to grinding which includes modifications to the geometry of

the unfinished blank and ceramic processing steps.

Step 2 : Step 1 is followed by changes in component design (without

altering the fimction) that lower the ceramic process cost and accelerate the finishing process.

50
Telble 4 shows a hypothetical case and

three situations which represent the

range of realistic conditions commonly

observed. While the exact values are

not critical, the trends are important.

The total cost noted in our exeunples

(a), (b) and (c) represent the range of

costs quoted for cereunic components,

which are 3 to 15 times that of ideal or anticipated future costs for ceramic

components. However, it is possible to

systematically attack these total costs

using step 1 and step 2 also demonstrated in Table 4. Figure 4 shows actual cases where step 1 has been demonstrated

successfully.

TABLE 5 TWO STEP APPROACH TO

REDUCE FINISHING COST

CASE: PRESENT STEP1 STEP2

FINISHING COST = MACHSNING COST

+ INSPECTION COST + SET UP COST

FIGURE 6

TOTAL COST APPROACH TO

REDUCE THE FINISHING COST

Through this and other exeimples, it has

been learned that the yield can be maximized by careful understanding of the

grinding interactions and optimizing them

using the systems approach described in

Figures 2 and 3. Finishing cost and

cycle time are two other elements which

require close scrutiny in the total cost

approach. Finishing cost can often be

minimized by controlling the inspection

and set-up costs. This is pursued simultaneously with cycle time reduction using

the understanding of grinding process

interactions described earlier. These

approaches for total cost reduction are shown in Figure 6 and Table 5. While

some of these steps may appear to be obvious, it is critical that these steps

PRESENT STEP 1 STEP 2 QPT IIP

• ^VERAL TO SINGLE OR FEW • SAME AS

MANY SET UPS. SET UPS STEP1

* POOR FIXTURING GOOD FIXTURE • PART DESIGN

- LACK OF MULTl- DE9QN PERMITS RAPID

-AXtS MACHINtNG MULTI AXIS CNC SETUP.

CAPABILITY MACHINES USED

* LIMfTED PROCESS CAREFULLY

PLAN^BNQ OR DESlGfCD PROCESS WHEEL DESIGN. PLAN AND WHEEL

INSPECTION

SINGLE OR FEW * SAME AS STEP

• MULTIPLE SET UPS SET UPS REDUCE STEP 1. REQUlfe REPEATED CUMULATIVE • PART DESIGN

INSPECTION FOR ERRORS AND PERMITS RAPID GEOMETRIC ACCURACY INSPECTION TIME. INSPECTION AND CUMULATIVE MACHINE TOOL AND REDUCED

ERRORS. ACCURACY AND INSPECTION.

• SMALL STOCK WITH PRECISION USED NEAR NET SHAPE AND TO BUILD CLOSE

ABSENCE OF LOCATING TOLERANCE IN TVE SWFAtXS REQUIRE PARTfCDUCINQ

FreOUENT INSPECT INSPECTION TIME.

-ION.

MACHINING

* PART DESIGN LESS * SAME AS PRESENT * PART DESIGN TOLERANT FOR MODIFIED TO

VARIATIONS IN THE PROVIDE

MACHINING PROCESS LOCATING

* PART DESIGN NOT SURFACES FOR

MACHINING CONOUaVEFOR

Hhl-tHENCE. RAPID MACHINING.

• SIMPLE SOLID • PART DESIGN NOT

SUITABLE FOR SWGtE SHAPE WITH CONSISTANCY STEP OR MULTI AXIS

• RAPID MACHI- MACHINING

NtNG METHODS

51
25

n

100 100

20

CYCLE FINISHING YIELD FINISHING

TIME(HRS.) RATE($/HR.) (%) COST($/PART)

RELATIVE RELATIVE

UNITS UNfTS

PAST

AFTER SYSYETMS APPROACH

INCLUDING TOTAL COST CONSIDERATIONS

FIGURE 7 COST BENEFITS BASED ON

THE SYSTEMS APPROACH

are pursued rigorously in order to find

early commercialization of the advanced

and emerging materials. Figure 7 illustrates an actual case where the cycle

time is dramatically reduced and the

yield is improved, thus decreasing the

total cost by a factor of 25. In order

to achieve these results, state-of-theart grinding system is recpiired which

increases the finishing rate (machine

rate + labor rate + burden rate) by a factor of 4. However, such increases are totally justified with high return on investment or short payback period,

because of increased throughput (cycle

time reduction) and finishing cost reduction. In addition, the component also

exhibited higher strength and reliability

during performance tests.

Surprisingly, methods to lower

finishing costs are also associated with

sound finishing practices necessary to

achieve reliable grinding of high

strength parts! Thus cost effective

finishing of ceramics to achieve low

total cost per part would appear to be

concurrent with reliable finishing methods for ceramics.

Summary ;

It has been described earlier in

this paper that machine tools, diamond

wheels, work materials and operational

factors play a key role in the grinding

of ceramics. Integrating these four key

elements into a grinding system has been

shown to be required both for technical

reasons as well as economic or total cost

considerations. The biggest challenge in

the introduction of advanced materials

like ceramics is the difficulty in bringing about such system integration based

on sound technical reasoning and economic

considerations. The risks, uncertainties

and cost avoidance considerations, always

seem to dominate over reasonable technical experiences whenever new manufacturing systems are to be implemented.

Grinding, like any manufacturing

process, is an interdisciplinary effort,

often requiring expertise in each of the

four input areas, i.e., machine tool,

wheel, work material and operational

factors. Integrating the grinding system

is far more effective when expertise in

these four elements are brought together.

Such interdisciplinary efforts facilitate

systems integration far more than individual and interconnected efforts.

Indeed such interdisciplinary and cooperative effort decreases the risks and

52
uncertainties. Several redundant costs

are avoided or eliminated in this process. Such effort makes the total cost

approach — where changes in materials

and ceramic processes are acceptable to

increase finishing yield and decrease

total cost — practical to implement.

There are attempts to bring sUaout interindustry colleUsoration of machine tool

builders, ceramic manufacturers or users

and wheel manufacturers to achieve the

necessary integrations of grinding systems for cereunic components''^'^. Widespread efforts to achieve such system

integration through sound technical and

economic reasoning will be necessary to

achieve early commercialization of

ceramic components.

List of References

1. K. Subramanian and S. Ramanath,

"Principles of Abrasive Machining

Processes", Section 5, Vol. 4,

Ceramics and Glasses . ASM Hand

Book, 1991.

2. B.R. Lawn, A. A. Evans, and D.B.

Marshall, Elastic/Plastic Indentation Damage in Ceramics: The

Median/Radial Crack System, J. Am.

Ceram. Soc, Vol. 63 (No. 9-10),

1980, p. 574-581.

3. I. Inasaki, Grinding of Hard and

Brittle Materials, Ann. CIRP, Vol.

36 (No. 2), 1987, p. 463-471.

4. R.W. Rice and J.J. Mecholsky, The

Nature of Strength Controlling

Machining Flaws in Ceramics, The

Science of Ceramic Machining and

Surface Finishing II, Special

Publication 562, National Bureau

of Standards, 1976, p. 351-378.

.5. M. Miyashita, Ultraprecision Centerless Grinding of Brittle Materials, First Annual Precision Engineering Conference, North Carolina

State University, Raleigh, 1985.

6. T.A. Bifano, T.A. Dow, and R.O.

Scattergood, "Ductile Regime Grinding: A New Technology for Machining Brittle Materials", Proceedings of the Intersociety Symposium

on Machining of Advanced Ceramic

Materials and Components, ASME,

11/12/1988, pp. 113-120.

7. K. Subramanian and S. Ramanath,

"Mechanisms of Material Removal in

the Grinding of Ceramics", Proceedings of the Symposium on Precision Machining: Technology and

Machine Development and Improvement, ASME Publication, PED-Vol.

58, 1992.

8. R.N. Kopp, et.al, "Progress in the

Finishing of Technical/Advanced/Fine/Electronic Ceramics", SME Paper No. M91-178,

1991.

9. Unpublished internal research

results, Norton Company.

10. J.H. Indge, "Double Sided Precision Abrasive Machining of Cereimic

Materials", SME Paper No. EM90-

139, 1989.

11. S. Jahanmir, et. al, "Cereimic

Machining: Assessment of Current

Practice and Research Needs in the

United States", NIST Special Publication 834, June 1992.

12. S. Das and T.R. Curlee, "The Cost

of Silicon Nitride Powder and the

Economic Viability of Advanced

Ceramics", American Ceramic Society Bulletin, Vol. 71, No. '7,

7/1992, pp. 1103-1111.

13. K. Subramanian and P.D. Redington,

"Total Cost Approach for Ceramic

Component Development", Ceramic

Engineering and Science Proceedings, American ceramic Society,

Vol. 14, #1-2, 1993, pp. 309-320.

14. K. Subramanian and P.D. Redington,

"Ceramic Machining Technology

Center -An Innovative Approach for

Precision Production Grinding",

SME Paper No. EM91-256.

15. "The Challenges of Ceramic Machining Continue", Edited by Laurel M.

Sheppard, American Ceramic Society

Bulletin, Vol. 17, No. 11, November 1992, pp. 1590-1603.

Acknowledgement : The authors would like to thank Mrs.

Chris Merithew for her assistance in the

preparation of this manuscript.

53

THE ROLE OF MATERIAL AND GRINDING PARAMETERS

IN GRINDING OF ALUMINA FIBER-REINFORCED

ALUMINUM ALLOYS

H. CHANDRASEKARAN and J. O. JOHANSSON

Swedish Institute for Metals Research, Stockholm, Sweden

Results from a grinding study of AI2O3 short fiber reinforced Al-alloy MMC

is presented in this paper. Variants of MMC material were tested for a range of

grinding conditions and abrasive wheel types. Grinding forces and surface

finish were measured and qualitative study of the ground surfaces using optical

and scanning electron microscopy was also carried out. While the grinding

conditions appear to play only a limited role upon the measured surface finish,

the abrasive type on the other hand displayed great influence. These indicate

that the measured surface finish parameters do not reflect the true slate of the

surface. Force considerations and qualitative surface studies revealed the overall

superiority of diamond and cubic boron nitride wheels over the conventional

abrasives. The two important features, namely disturbance of the near surface

fibers and the development of micro-cavities along the contour of the fibers in

the grinding plane observed by us could have critical implication upon the

surface integrity of these MMC materials.

1. INTRODUCTION

In the course of the past few years the class of

materials identified as metal matrix composites

(MMC) consisting of high-performance reinforcements in a metallic matrix have increased in their range. The main attraction of these materials is the possibility of tailoring the mechanical and

physical properties between those of the rein- forcement and the metallic matrix. The cost of the

materials are dictated by the matrix, reinforcement and the manufacturing process used. In this context alumna/silicon carbide fibre/particle

based aluminium composites using some type of

pressure infiltration or melt mixing appear to be

cost competitive (1). Accordingly these may be

expected to find specific applications in such Ught

weight wear resistant parts as valve gear and other

engine components.

However, these intended appUcations primarily in the automobile industry would require some machining. Existing information on the

machinability of aluminium alloy composites

reinforced with AI2O3 fibres/particles based upon

chip forming operations like turning and drilling

(3, 4, 5 and 6) indicate that the composite materials could present machining difficulties, often

intriguing from some points of view. Tool wear is often abrasive, demanding highest tool hardness.

Another problem is the quality of machined sur- faces and sub- surface damages accentuated with

the increasing wear of the cutting tool (3) and

grinding might be the answer (2) then.

55
However, information about the grindability of

MMC is very limited. Hence as a first step it was

proposed to collect fundamental information on

the grindability of these materials such as cutting

forces and surface finish, when grinding some

standard Al alloy composites reinforced with short

alumina fibres.

The abrasive wheel types as well as the

grinding conditions were varied to cover a broad

range of parameters at this stage. Instead of the

traditional grindability studies based upon tiie grinding coefficient or G factor (relationship bet- ween the volumes of ground material to worn

abrasive wheel), the focus was directed towards

the understanding the grindability of these materials from energetic and wheel-work combinations

point of view as relevant to the surface integrity

aspects. This involves a study of the contact phenomenon between the abrasive grit and the hard

alumina fibers, as well as the soft Al alloy matrix. 1 25 fim

2. Experimental PROCEDURE

2.1 ALUMINIUM ALLOY COMPOSITES

The work material for the grinding tests was produced using the "squeeze casting", a technique in which molten Al-2Mg alloy was allowed

to infiltrate an alumina preform of required shape

under suitable temperature and pressure within a closed preheated mould (7), resulting in short

cylinders of 60 mm diameter and 40 mm high.

Accordingly Al-2Mg alloy with respectively 10

and 26 vol% alumina fibers was the principal test material. Nominal properties of the composites

used in our grinding tests are shown in Table 1.

Table 1 Properties of the Al alloy composites

No Property

10% fiber

1 . Yield stress (MPa)

2. Ultimate tensile

stress (MPa) 187

3. Vickers hardness (GPah 1.20

Material variants Al-2Mg Al-9Si-3Cu

98

26% fiber

163

263

1.40

30% fibei

200

1.52

Fig. 1 Microstructure of 26% fibre MMC

material (axial section).

Fig. 2 Microstructure of 26% fibre MMC

material (tangential section).

56
The manufacturer's specifications (7) in- dicate a mean fibre diameter of 3 mm and a length

of ~ 3.5 mm in the preforms, which in all probability gets broken during the processing of the

cylinders. The mechanical properties indicated in

Table. 1 correspond to the favourable radial direction of the cylinder. Detailed study of fibre morphology reported by van Hille et al (8) for this type of composites indicates that about 60% of the

cumulative fibre population has an orientation of

(P ~ 20° with respect to the ideal horizontal or 0°

plane, the rest occupying other positions. Normal

to the axis of the cylinder the fibers may be ex- pected to be randomly oriented although fibre

bimdling is often noticed (8). Typical microstructure (axial and radial sections)of the MMC material is shown in Figs. 1-2.

2.2 GRINDING WHEELS

From traditional grinding practice alu- minium and its alloys being soft materials are often groimd with SiC abrasives, while ceramics

(the fiber material in our MMC) are generally

ground with diamond wheels. Hence for composite material containing both hard and soft phases,

investigations involving a broad range of the abrasive types used in modem apptications was considered. Understanding the relative interaction

between the composite material and the grinding

wheels based upon different type of abrasives was

the motivating factor, not the choice of the optimal grinding wheel for the composites under in- vestigation. Accordingly representative wheels,

based upon alumina (AI2O3), silicon carbide

(SiC), cubic boron nitride (CBN) and synthetic

diamond (SD) abrasives from the same manufacturer was used. Here the grinding wheel choice

namely grain size, type of bond and the degree of

compaction was based upon manufacturers re- commendations and what was considered as comparable wheels for standard surface grinding applications to produce similar type of surface quality. All the wheels are intended for the same

peripheral grinding speed (~ 30 m/sec). The

specifications of the wheel for the proposed sur- face grinding tests are shown in Table 2.

Table 2. Grinding wheel specifications

.No Abrasive Mesh size Bond Specification

1. AI2O3 60 vitrified 43 A60 J6 VM

2. SiC 60 vitrified C60 1 VM

3. CBN

(75% cone.) 126 phenolic B126 R75B-2

4. Diamond

(75% cone.) 126 phenolic SDN126R75

B2A-3

Obviously the mesh size specification and the

number of active grains per unit area of the grinding wheel are somewhat different for these two

groups of wheels, and this should be considered

when comparing their grinding performance in

terms of material removal mechanisms. However,

from the point of view of application, these

wheels are expected to perform in a similar fashi- on.

3. GRINDING TESTS

3.1 TEST DESCRIPTION

All the grinding tests were carried out in a horizontal surface grinding machine with a constant coolant (5% emulsion in water) flow of

about -14 Ht/min. The cylindrical work pieces were held in a V block type fixture and the circular faces were ground using the down feeding of

the abrasive wheel. The dimensions of the

Al203,and SiC grinding wheels were the same,

namely 356 mm diameter and 50 mm width,

while the cubic boron nitride and synthetic diamond wheels had the same diameter, but a width

of 12 mm only. Hereafter the above four type of

abrasive wheels will be designated respectively as A60, S60, CBN and SDN. The machine spindle

speed of 1815 rpm gave a peripheral speed of 33.8

m/sec. Only plunge feeding of the wheel and upgrinding mode of the traversing table was used in

all these tests. To minimize the effects of non-uniform table speeds long stroke lengths (compared

to ground length) were used. The exact time for the test grinding strokes was used to estimate the

true table speed.

The idea behind the selection of parameters for the grinding test was the possibility of ob57
taining grinding forces and surface finish information as functions of

* the type of abrasive grit used

* the different levels of fiber content in a given Al

alloy matrix and

* the wheel depth of cut and the table speed.

In the context of the grinding parameter selection,

the relative role of the dimensions of the fiber (its diameter and length) and its concentration in a given composite and the distribution density of

the abrasive grains in the grinding wheels would

be of interest. Accordingly, wheel depth of cut was varied between 5 and 30 |im, while table

speed was varied between 7 and 25 m/min.

The grinding wheels were balanced and

trued with each change of wheel. This actually in- volves in the case of CBN and SDN, wheel balancing, followed by truing the mounted grinding

wheel on the machine using a rotating diamond

wheel dresser to eliminate radial run-out and final- ly opening of the wheel structure using AI2O3

sticks. The AI2O3 and SiC wheels were trued usi- ng multi diamond dresser stick. During the

grinding tests with A60 and S60 wheels dressing was carried out anew for each test nm.

3.2 MEASUREMENT OF GRINDING FORCES

For grinding tests with a wheel depth of

cut a =10 jxm, and a table speed of Vw = 15

m/min identified as the reference condition, the

grinding forces were measured for all the combinations of abrasive wheels and composite variants. With the MMC specimens mounted on a

three component piezo-electric force platform the

grinding forces were measured during a single upgrinding pass. The output of the dynamometer

was monitored through a charge amphfier, while

for signal storage and subsequent analysis, a multi-channel digital memory oscilloscope with

built-in hard copy out-put was used. The force

measuring system with a natural frequency of

about 4.7 KHz in the radial direction may be considered sufficient to identify phenomenon associated with the rotational frequency of the grinding

wheel (-30 Hz) such as vibrations, run-out and

non-uniformities on the wheel surface. Reliable

identification of the action of individual abrasive

grains as they move over the matrix and fibres.

would however require a measuring system with

much higher natural frequencies.

1 7 11

/

1

J'

1" mi

'A

'i

L

r

1 1 1 1-

-V

K

H-h-

"If

i

1

-

1

1 .J L-•*'*^^«»Av. 0.1 sec * * a

Fig. 3 Trace of grinding forces

(Fj and F,) A60 wheel- 10% fiber MMC

a=10 jimi; =15 m/min

The typical trace of the radial Fj. and the

tangential F^ components of the grinding forces

recorded by the oscilloscope is shown in Fig. 3. Since the true wheel depth of cut associated with a

single abrasive grain is a process variable, the

measured total grinding forces also display scatter. The difference in the wheel composition as well as the substantial difference in wheel width

(and hence damping) between the conventional

and superabrasives wheels resulted in greater

force fluctuations especially with CBN wheels.

Subsequentiy, in all measurements for any given

wheel-work combination, the maximum mean

measured forces were normahzed with respect to the maximum groimd width.

58
The consolidated results relating grinding

forces to the abrasive wheel types and grinding

conditions (wheel depth of cut and table speed)

when grinding MMC with different fiber contents

are shown in Figs.4-5. Normalized information

relating the grinding energy to the abrasive types

and the grinding conditions are shown in Fig. 6,

while the variation of grinding force ratio |J. (=Fj/ Fj.) as a function of the abrasives used, fibre content and grinding conditions are shown in

Fig. 7.

AJ'2Mg a=10 pin Vw=15 m/min

F 1. 1 0 %

S F r.l 0 %

F t, 2 6 %

Fig.4 Grinding (reference conditions)

forces for 10% and 26% fiber MMC.

10 15 20 25 Wheel depth of cut[pin]

Fig.5a Variation of grinding forces

with wheel depth of cut a for 26%

fiber MMC.

4.5

!

-

Al -2 M g

2 6 % fiber

a = 1 0 |j m

5.:

o

o -i'

u. —^ CBN

- - SDN

0

7 9 11 13 15 17 19 21 23 25 Tabl« tp««d,Vw[m/mlnl

Fig.Sb Variation of grinding forces

with table speed for for 26% fibre

MMC.

n 40 las - 30 i

25

i 20

II 10 I 5

8. „ in 0

AI-2M g a = 1 0 |jm Vw = 15 m/min

1 0 % f I b e r D 2 6 % 1 1 b e r

Fig. 6 Specific grinding energy

(reference conditions) for 10 and 26%

fiber MMC with different abrasives.

3.3 Surface FINISH IN GRINDING

Grinding is often required to fulfil tole- rance and quality requirements and thus it is essential to relate these features to the MMC

material (matrix alloy and fiber content), the abrasive mediimi and the grinding conditions used.

With this aim the ground surfaces were analysed

using a non-contact surface finish measuring system, using infrared laser beam. All the traditional s urface finish parameters such as Ra (centre line average), Rz (peak to valley depth ) and Rp(mean

peak distance from centre line) were measured.

0 70

0 60

050

I

0 40

g 0 30 o

u. 0,20

0 10

0 00

AI-2M g

a = 1 0 iJ m Vw = 15 m/min

10%fiber

26% fiber

Fig.7 Force ratio p (= Ft /Ft) when

grinding (reference conditions) 10

and 26% fibre MMC with 4 abrasives.

The probable modes of surface build-up

during grinding, the indirect influence of the

abrasives (grain sharpness) and the degree of fiber

damage/delamination could also be understood to some extent from such topographic information.

59
1.68 1.78 1.88 mm

Fig. 8 Trace of ground (reference

condition) surface of 26% fibre MMC

as recorded by non-contact laser

measuring system,

a) A60 wheel b) SDN wheel.

For each grinding condition the surface

finish across the grinding direction was measured

at three places and the average was used in subsequent evaluations. Based upon the traditional Ra

value alone, the ground MMC surfaces may be

considered to be good. Differences arise from the

MMC composition and the abrasive types used for

grinding, rather than from the grinding conditions.

Trace of a MMC surface profiles after grinding

tests is shown in Fig. 8. The consolidated results relating the measured surface finish parameters

Ra, and Rz to the abrasive types and grinding

conditions (wheel depth of cut and table speed)

when grinding MMC with 26% fiber content are shown in Figs.9-10. The influence of grinding

conditions upon Ra is marginal with some improvement with increasing wheel depth of cut

(Figs. 9a and 10a), but an exactly contrary behavi- our was observed with increasing table speed

(Figs. 9b and 10b).

0^80 |CBN

0.70 I

0.60 4

5 10 15 20 25 TabI* sp**d,Vw ((iin)

a) Variation of Ra

a) Variation of Ra

10 15 20 Table sp«*d,Vw ([n/mln)

b) Variation of Rz.

Fig.9 Variation of surface finish as a function of wheel depth of cut for

26% fibreMMC when grinding with

differem abrasives.

16 T

15 •

14

13 • I

12. 3 11 a 10-9 • 8

•

6-

b) Variation of Rz.

Fig. 10 Variation of surface finish as a function of

table speed for 26% fiber MMC using different

abrasives.

60
The deeper scratches denoted by Rz

exhibit even less sensitivity to grinding depth and

table speed in this context. The best measured

surface finish correspond to CBN and SDN

grinding and the worst to grinding with A60

wheels, while grinding with the S60 wheels often

resulted in intermediate values of finish.

a) Variation of Ra

1 10

0%

AI-?Mg I Al -9SI-3CU

b) Variation of Rz

Fig. 1 1 Influence of fibre content upon

surface finish when grinding with A60 wheels;

a=10 ]m, V^15 m/min.

The influence of increasing content of the

hard fiber in Al alloy MMC upon the surface fin- ish (Ra) is shown in Fig. 1 la corresponding to reference grinding conditions using A60 wheel.

The poor surface finish associated with the fiber

free matrix alloy improves appreciably with in- creasing fiber content, tapering off around a fibre content of about 30%. Similar influence upon Rz

values (Fig. lib) appears to be less sensitive to

fiber content.

Another general effect of the grinding

process upon the surface properties of the MMC

material is the resulting hardness. It is conceivable

that during the grinding process, due to the inter- action of the abrasive grit with the matrix and fiber the surface fiber content could be affected. To

verify such a hypothesis the micro hardness measurement of the ground surfaces was carried out using a 50 grams load. Figure 12 shows the hard- ness variation over the surface (at ten different

points) ground with different grinding wheels at the reference condition. The mean ground surface

hardness for the Al-2Mg alloy MMC was around

125 HV, and as could be seen from the figure the

influence of the abrasive type upon resulting

hardness appears to be small and was confirmed

to be statistically not significant.

Fig. 1 2 Scatter in the surface microhardness ( 1 0 measurements) of 26%

fiber Al-2Mg alloy MMC ground (reference

conditions) v^th A60 and SDN abrasives.

3.4. SEM STUDIES

A useful method forjudging the qualitative aspects of wheel-work interaction in grinding

is the direct observation of groimd surfaces in the

scanning electron microscope. Accordingly,

unetched surfaces of ground MMC were investigated in SEM both in the topographic mode (SE) as well as in the back scattered (BS) mode. The degree of surface disturbance associated with the conventional and super abrasives, when grinding

hard fiber reinforced MMC material could be seen

in Fig. 13, showing the typical surfaces ground by

A60, S60 and SDN wheels. SDN exhibits the best

quality, while S60 exhibits grooves with some

adhesion. In general, grinding with conventional

abrasives is characterized by adhered and dis- turbed surface with a few clean scratch maiks,

CBN by a number of clean scratch marks with

occasional adhesion and SDN with a large number

of clean scratch marks with hardly any adhesion.

61
The chips are hardly in the form of fine

filament like chips normally associated with steel grinding (9), due to the constant interruption by

the hard fibres and consequent grit fracture- /dislodgement.

Figures 14 to 17 show the typical appear- ance of 26% fibre Al-2Mg MMC ground with

conventional and super abrasives. At this fiber content the AI2O3 wheel has the disadvantage of

both insufficient hardness of grit (almost the same

as that of the fiber) and a rather brittle bond

(vitrified) along with contiguous grit configuration. Although there is enough chip space, this is unable to accommodate the maf'rial removed.

i

1 X500 HHIISeIIXIiHKRRT?

Fig. 13 Typical view (SEM) of the surface

of 26% fiber Al-2Mg MMC ground with,

a) A60 wheel b) S60) wheel c) SDN),

wheel.

s

Fig. 14 Typical view (SEM) of the surface

of 26% fiber Al-2Mg MMC ground with A60

wheel; a=10 pm, V^=15 m/min

a) general appearance, b) detail of a

single disturbed fiber from (a) in the

back scattered (BS) mode showing clearly

the matrix (light) and the fibers (grey).

62
The normal forces increase and the heat

concentrating on the work material (due to the

poor thermal properties of the A60 wheel), results in the physical movement of individual fiber in

the grinding direction (Fig. 14) as well as fiber

displacement, due to the inability of the grit to cut clean through the fibers.

b

Fig. 15 SEM view of surface ground with

S60 wheel; other conditions as in Fig. 14.

a) abrasive groove and exposed fiber

parallel to the grinding direction,

b) enlarged view (BS mode) of the fiber

from (a) showing fracture fiber

fracture (see arrow).

With the marginally harder SiC grit the

true interaction between the grit and fibre is more

complex as shown for a single fiber, which happens to Ue exactiy parallel to the grinding direction. Individual grit could however, cut clean

through the MMC material and get damaged mid

way (Fig. 15). The generation of the groove in turn results in the exposure of more fibres. We

further notice that the presence of even small fib- ers is accompanied by cavity formation around the

contour of the fibers. This feature appears to be a

characteristic ofgrinding MMC material as this was observed even in the case of favourable

grinding with the super abrasives.

In the case of super abrasives, the available chip space is favourable (large areas of shal- low depth), and as long as this is not filled, there

is no problem. The extent to which the harder and

sharper fine grit points of CBN could work

through the MMC material is readily seen in

Fig. 16a. Large puU out and disturbance of fibers was rarely seen here, but again around the contour

of each surface fiber worked by the grits, cavity'

formation was observed. These could be very

small (Fig. 16b).

A closer scrutiny of the bottom of individual grooves from SDN grinding (Fig. 17) displays

that the hard and sharp diamond grit could effec- tively decapitate individual fibres and the remnant

trace hardly reveals large scale cavity formation

associated with the other abrasives, although the

fiber section could be seen (Fig. 17b).

4. DISCUSSION OF TEST RESULTS

4. 1 Grinding forces

The force results show that as a function

of grinding depth of cut and table speed (Fig. 4)

these are similar to the surface grinding of metals,

namely increase in forces with wheel depth of cut. Unlike the influence of the wheel depth of cut, an

increase in table speed increases the true thickness

of cut removed by individual abrasives only marginally. Consequentiy, the influence of table speed appears to be negligible upon the smaller Fj

force, while some marginal influence could be

seen upon the larger Fj. force. This was especially

true for the CBN and SDN wheels. The measured

63
Further, the distribution of the abrasives

in the CBN and SDN wheels is characterized by

individual grains surrounded by the organic bond,

while in the case of the conventional wheels the

grits occupy contiguous regions separated by

similar regions of the vitrified bond. These

differences may be expected to affect chip packing

in grinding which in turn affects the measured

grinding forces (10). Thus the adhesive behaviour

of the wheels would be affected both due to the

intrinsic properties of the abrasives as well as by

their physical configuration. AI2O3 wheels

display the worst combination (lowest grit

hardness and maximum material pick-up) and the

diamond the best (highest grit hardness and minimal material pick-up), and the SiC and CBN

wheels come in between.

In a similar manner, the measured grinding forces as a function of the abrasive types

could be somted into two distinct groups, based

upon the relative role of the reinforcing fiber content. SiC and SDN wheels displayed an in- crease in forces with an increased fiber content,

while this was quite the contrary with AI2O3

wheel. All these effects were appreciable in the case of the larger radial force Fj.. We have to keep

in mind here that an increase in the content of

alumina fiber results in two different effects,

namely an increase in the hard phase as weU as a decrease in the matrix phase having greater influ- ence upon adhesion. Accordingly an abrasive

displaying minimal adhesive affinity to the matrix

alloy and at the same time hard enough to cut through the alumina fiber (~2(XX) HV) would be

the most successful. The basic physical properties

of the four abrasive types used in our tests is shown in Table 3.

From this table we can see that SiC, CBN and

diamond aU have greater hardness than that of

alumina fiber. Diamond with its excellent low

friction properties along with the highest hardness

thus exhibits the best capabiUty and this could be

readily seen from the smaU forces especially for

10% fiber MMC (Fig.4).

Similar behaviour could be observed in the case of S60 wheels. Both these abrasives are able to

grind and thus with increasing fiber content the

forces also increase.

Table 3. Physical properties of the abrasives used

Properties Abrasive types —

A60 S60 CBN SDN ~. Composition AI2O3 SiC BN C

Dens, p(gm/cm'') 3.92 3.21 3.48 3.52 Knoop hardness -21 26 42-46 56-102

(GPa)

Therm, cond. X 6 55 200-700 600-2000

(W/mC)

Therm, exp. a 86 45 48 15 (10^/deg.C)

Oxid. resistance 1750 1500 1400 800

(deg.C)

In chip forming operations like turning, an

increase in fibre content results in the increase of

material defects, resulting in energywise

favourable shearing in the shear plane, and

mesured forces decrease (1 1). In the case of A60

wheel, with increasing fiber content the grit loss

increases; thus material removal does not take

place, and the measured forces come down. The

role of radial force to sustain and influence actual

metal removal is well known in the case of grinding metals (10), and thus the poor grinding per- formance of A60 wheels is accompanied by a re- duction in the radial force Fj. with increasing fiber content.

4.2 Grinding energy and heat generation

The influence of the grinding conditions

could be normalized using the parameter of nominal material removal rate (a .V^ ). The grinding

forces are marginally affected by the metal re- moval rate to start with. Subsequentiy, however,

as the relative role of force ratio changes, the ef- fect of material removal rate upon forces in- creases. This effect appears to be similar for both

CBN and SDN (Fig. 5). Ultimately tiie effects of

grinding conditions influencing both the grinding

forces and the nominal rate of material removal is expressed by the concept of specific grinding

energy U (J/mm^). The value of U may be ex- pressed as U = (Ft. V )/(a.W V^) (1)

64
forces for a given grinding condition using diffe- rent types of abrasives appears to be more complex. Here, eventhough the abrasive wheels are nominally comparable in terms of performance,

both the abrasive distribution per unit area of

wheel as well as the type of bonds (vitrified for

the A60 and S60 wheels, but phenolic for CBN

and SDN) differ.

2 ^lrr\

Fig. 16 SEM view of surface ground with

CBN wheel; other conditions as in Fig. 14.

a) fine scratches in matrix and two

fibers with their contours exposed

b) region around another fibre showing

a series of micro-cavities (arrows) along

the fiber contour.

2eKu|

f

Us. H2e»eee 11* M 1

Fig. 17 SEM view of surface ground with

SDN wheel; other conditions as in Fig. 14.

a) abrasive track and exposure of fibers

(arrows)

b) detail from (a) showing the section

of a single fiber due to its decapitation by diamond grit

65
where Vg and correspond to wheel and work

velocity, while a and Wg correspond to wheel

depth and width of grinding. Of this total energy, some part is carried away by the grinding swarf,

approximated by Malkin (10) as a percentage of

the maximum possible energy associated with

chip formation namely the energy for melting a

unit volume of material U^.. The remaining energy

available for heating the woric material may

be expreesed as

U^=U-0.45Uc

= [(Ft- Vg)/(a.Wg )-0-45 Uc] (2)

The parameter U is also convenient in comparing

the grinding behaviour of materials in general. For

10% fibre Al-2Mg MMC material this energy decreases progressively (Fig. 5) in the sequence A60

S60 and SDN, indicating thereby the cutting effi- ciency of the diamond grit. With increasing fibre content S60 and diamond display a smaU increase

in the energy requirements, wlule the A60 wheel

being unable to grind, exhibits only a marginal

change. The role of wheel depth of cut in this context appears to be marginal, wMle higher table

speeds result in smaller energy requirements.

These correspond to CBN and SDN grinding of

26% fibre Al-2Mg MMC material. Thus an en- hancement of material removal rate through in- creased wheel depth of cut appears to be advantageous, if the abrasive is hard enough and the bond

could take the load.

Good correlation has been reported (10)

between the theoretical prediction and experimental data concerning the surface temperatures

associated with the grinding of carbon steels using an analysis in which the energy available for

heating (U^) is used to calculate temperature rise based upon the approximation of a moving heat source on a half-space. It is proposed to use similar approximation for estimating the surface temperature during the grinding of fibre reinforced

MMC material also. In the absence of reliable

thermal properties of the MMC material with the

two fiber content, as a first approximation we

shaU use the properties of the matrix aUoy in our

temperature computations.

Based upon the measured grinding force

(Fj ) and the velocities of the grinding wheel (Vg)

and the workmaterial (V^ ), the wheel depth of

cut a, and the width of cut Wg, and thereby the

grinding energy (Ug assumea to be imiformly

distributed over the wheel woric contact length),

we can express the maximum temperature rise (0 njax ) of ground surface (10) as

B max =

1.12-al-2.a0.75.v0^5[u_o.45U,]

Id0.25

(3)

where a and X are respectively the thermal diffusivity and thermal conductivity of the matrix material and dg is the equivalent diameter (= D, the

grinding wheel diameter in surface grinding).

Using eq.(3) the temperature rise associated with

the grinding of the Al-2Mg MMC with respectively 10% and 26% fibre content using the four

type of grinding wheels A60, S60, CBN and SDN

was calculated. The influence of the wheel depth

of cut a upon the temperature rise when grinding

26% fibre MMC material with CBN and SDN

wheels is shown in Fig. 18, whQe consohdated re- sults for aU combinations for the reference grinding conditions are shown in Fig. 18.

450

400

350

300 O

250 — 200

150

100

50

0

AI-2M g,26%tib8r.

5 10 15 20 25 30 Whttel depth of cut ,a [Mni|

Fig. 1 8a Calculated maximum surface

tempeature as a function of wheel

depth of cut during the grinding of

26% fiber MMC material.

The temperature rise for the reference

conditions appears to be nominal for most wheelwork combinations, being the order of 80-250OC.

As may be expected the lowest temperature rise corresponds to the grinding of 10% fiber MMC

with SDN wheels. With 26% fiber MMC materials S60 and CBN wheels display a rise of 180OC,

while for SDN this is somewhat less ~150°C.

66
In terais of grinding conditions the wheel depth of

cut a appears to have the greatest effect showing

up to 450°C rise during CBN grinding (Fig. 17a)

and somewhat less for diamond grinding.

250

200

150

100

: 0 % fiber 26 % fiber

A60 S60

Fig. 1 8b Calculated maximum surface

temperature in Al-2Mg MMC material

when grinding with different wheel

types at reference grinding conditions.

4.3 THERMAL EFFECTS ON FIBER

Table 4. Probable axial tensile stress in fibers due

to temperature rise in grinding.

Vol. % of

fiber

Stress in MPa

100°C 200°C emax(427°C

)

10% 325 650

26% 190 380 810

The indicated maximum fiber stress of 810 MPa

is appreciable considering the tensile strength of

alumina fibers (~20(X) MPa) reported by the

manufacturer (7). Thus there appears to be an immanent possibihty that during grinding in addition to the mechanical action of the grit, thermal

effects could affect fiber fracture. Such effects should be observable especially in the case of long

fibres lying parallel to the grinding direction ex- periencing the combined action of mechanical and

thermal stresses.

An indirect effect of heating of the MMC

material during grinding is the associated imfavourable stress build-up due to the large differ- ence in the thermal expansion behaviour of the

matrix and fiber materials (this being respectively

24 and 8 x lO'^/^C). Using the simple model for

the axial stress in the matrix (parallel to the

fibre direction) brought about by a strain mismatch due to a change in temperature A8 proposed

in (12), and the concept of force equilibriimi we

can write

^ _ f [E^Ae(a^-a,] ] [l-H(v^E^/vfEf)J

^= (Om Am + Af) (5)

where F is the resultant force, o, E, v, correspond

respectively to axial stress, elastic module and

volume fraction of phases and the suffixes m and f denote the matrix and fiber material. Substituting

the appropriate values of Ej^ =70 and Ef = 300

GPa in the above equations along with tiie other -

relevant values and solving for Of , we can esti- mate the resulting tensile stress level in the fibers

as a function of temperature rise (Table 4) esti- mated earlier.

4.4 Material removal mechanisms in

grinding

The observation of the ground surfaces

along with the evaluated surface parameters and

grinding forces could be used to get an idea about

the actual nature of material removal during the

grinding of the fiber reinforced MMC materials. If the abrasive grit is hard and sharp enough, then it can slice clean Uirough the MMC material, as exemplified by SDN and CBN wheels. Recent

analysis of energy partition in grinding using

single grit model and the appropriate thermal

properties (13) do indicate appreciable similarities

between CBN and SDN. In such a situation the

only other basic difference between these two

abrasives could be the friction between the grinding chip and the individual abrasive grit exhibited

by these two materials. Further, forces associated

with CBN grit are notoriously sensitive to the correct truing and opening of these wheels, and thus

along with the advantages of temperatures, diamond turns out to be the better among tiie two.

When the abrasive makes a clean cut the

actual separation of the material in the form of a chip is associated with cavity formation as shown

schematically in Fig. 19.

67
ABRASIVE

MMC MATERIAL

Fig. 19 Schematic view of chip separation

stage during the grinding of fiber

reinforced MMC material.

a) region with no fiber

b) fiber rich region showing the

evolution of micro-cavities for two

different fiber orientations.

This is due to the large difference in the flow

properties of the matrix and fiber material. At the

instant of chip separation, the brittie fiber frac- tures easily, while the soft and ductile matrix

continues to deform and fail in the ductile mode.

The extent and size of these cavities are dependent

upon the relation between the width of cut of individual grits and the material present here. Further,

the degree of wettability between the fibers and

the matrix has a central role in this context. In

fiber free regions cavity formation is less probable. If major part of this width is associated with

the hard fiber (single or a group) then a series of

cavities along the width corresponding to the brittle fiber will arise (Fig. 19b). If the fibers are located normal to the grinding plane, cavity formation

wQl be minimal during the grinding of fiber rein- forced MMC, while if these were located perpendicular to the grinding direction, this could result in severe cavity formation in the ground material. In a practical case like ours where a majority of

fibres are oriented at an angle (< 20° to the

grinding plane as mentioned in §2.2) the extent of

concomitant cavity formation associated with the

fibrers and grinding would be in between. While

from pure strength point of view these defects

associated with the projected length of fibers in

the grinding plane are not critical by themselves,

if corrosion and fatigue control component failure,

then the grinding mode and its optimization

would be very critical to the performance of

MMC product.

4.5. Role of fibre-abrasive interaction

The exact nature of interaction between

the abrasive grit and the hard fibers are not easily

observed in the present study. However we can

indirectiy conclude that in the absence of hard and

sharp abrasive grits as in the case of A60 wheels,

major part of the measured grinding energy is due

to the grit ploughing and bond to matrix adhesion

(as a consequence of grit fracture), rather than

chip formation. This is clear from the quality of

the ground surface, especially that of Ra value as fimction of fiber content (Fig. 1 2). Unlike Rz

(characteristic of scratch depth) value, Ra is very

much affected in this context. The situation is similar in the case of SiC grits.

In the case of CBN and SDN grinding,

evidence from single grit in-situ studies in SEM

using diamond grits (14) indicate that ft^acture of

fibres perpendicular to the scratch direction to occur when a specific normal load is reached.

Based upon the reported (14) normal forces corre- sponding to fiber fracture (0.2 N) and the geometry of the observed width of track (~ 25 |im) we

can estimate a fracture stress level of about 2036

MPa, comparable to the tensile strength of the

fiber material. In our grinding tests, the speeds are much higher and the abrasive track widths vary

from 20 to 50 |im (Figs 14 and 16). This would

mean that fiber fracture could easily take place as

is also supported by the SEM studies. Further the

68
is also supported by the SEM studies. Further the

thermal effects could also contribute appreciably

to this feature. The sum total of these factors is the high probability of severe fiber fracture in

grinding.Thus the original fiber length in the pre- form (3-4 mm long) undergoes continuous frac- turing during the different manufacturing stages of

the MMC material, including the finishing operation like grinding. An important implication of

this phenomenon is that the physical distribution

of fibres on the service surface after a manufacturing step like grinding could be different from the

original.

5. CONCLUSIONS

Based upon this very first attempt some

conclusions on the grinding of this class of materials as well as on some aspects of the surface

integrity of ground MMC materials could be

made. These will be enimierated now.

1. Super abrasive wheels are to be preferred in the

grinding of alumina fiber reinforced Al alloys. In general grinding wheel using synthetic diamond

grits results in superior surfaces irrespective of the

fiber volume content and the type of aluminium

alloy used. CBN based wheels could also be

effective, but are associated with more exacting

wheel preparation.

2. Increase in the alumina fiber content in the

MMC material results in the increase of grinding

forces when the abrasive used are hard and sharp

enough as is the case with SiC and super abrasives.

3. When the abrasive grits of sufficient hardness

and sharpness is used (as is the case with SDN

wheels), the role of the grinding parameters

(wheel depth of cut and table speed) have only a marginal influence upon the measured surface

finish parameters.

4. Both considerations of mechanical loading of

the fibers under grinding as well as thermal considerations indicate that even under successful

grinding conditions, fracture of individual alu- mina fibers in the MMC material into smaller bits to be highly probable. Diamond grits, however

often preferentially decapitate the hard fibers.

5. Surface studies also indicate that grinding of

these fiber reinforced MMC materials is characterized by the presence of a large nimiber of concomitant micro-cavities corresponding to the pro- jected contours of individual fibers in the plane of

grinding.

ACKNOWLEDGEMENTS

This research project was financed by the

National Swedish Board for Technical Development (NUTEK) as a part of MMC programme,

and thanks are due to them. Prof. B. Lindstrom,

Production Engineering Laboratory, Royal Insti- tute of Technology, Stockholm is thanked for laboratory facilities, and Sup Naxos AB,

Vastervik for the supply of grinding wheels.

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69
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Fibre Reinforced Metal Matrix Copmposite

Material", Proc of 9the RIS0 Int. Symposium on

Behaviour, of Ceramic and Metallic. Composites.,

ed. S. L Andersen, H. Lilholt, and O. B. Pedersen,

Denmark, 1988

12. G. S. Daehn, and G. Gonzalez-Doncel,

"Deformation of Whisker Reinforced Metal

Matrix Composites imder Changing Temperature

Conditions", Metallurgical Trans. A. 20a, Nov.pp2355-2368, 1989.

13. T. Matsuo, S. Toyoura, E. Oshima, E. and

Y. Ohbuchi, Y. "Effects of Grain Shape on Cutting Force in Superabrasive Single-grit Tests",

Ann.of CIRP. 38/1, pp.323-326, 1989

14. A. Alahelisten, M. Olsson and S. HOgmark, "Tribological Properties of an Aluminium Alloy Matrix Composite", Proc. of

Nordic Tribology Conference, Copenhagen, 1990
ELASTIC STRESS FIELDS CAUSED BY SLIDING

MICROINDENTATION OF BRITTLE MATERIALS

Y. AHN, T. N. FARRIS, and S. CHANDRASEKAR

Purdue University, West Lafayette, IN

An analytical model of the stress field caused by sliding microindentation of brittle materials is developed. The complete stress field is treated as the superposition of applied

normal and tangential forces with a sliding blister approximation of the localized inelastic

deformation occurring just underneath the indenter. It is shown that lateral cracking is produced by the sliding blister stress field and that median cracking is caused by the applied

contact forces. The model is combined with experimental volume change measurements

to show that the relative magnitude of tensile stresses governing lateral crack and median

crack growth varies with the magnitude of the applied load. This prediction is consistent

with the different regimes of cracking observed under a sUding pointed indenter in soda-lime

glass and other brittle solids.

1 Introduction

Microscopic observations of machined ceramic sur- faces show that brittle fracture contributes substantially to material removal in stock removal and

coarse finishing processes. A careful examination

of the fractures in machined ceramics reveals strik- ing similarities with fractures about quasi-static

and sliding indentations produced by sharp indenters in brittle solids. It was, therefore, decided

to develop a model for characterizing the stress

fields produced by sliding, sharp indenters in brittle solids. The model that has been developed is an

extension of that proposed by Yoffe [1] to describe

the stress fields and fractures generated under a

quasi-static indentation by a pointed indenter in brittle materials. We begin with a review of the

deformation produced under a pointed indentation

and a discussion of the Yoffe model.

When a sharp indenter is quasi-statically applied

onto the surface of a brittle solid, the classical theory relating the hardness to the yield stress of metals [2] no longer applies since yielding no longer occurs at a constant maximum shear stress or at a

constant volume [3]. A large body of experimental observations has shown that in brittle materials, the volume of material displaced by the penetration of the indenter is accommodated by compaction or shear deformation, or both, and this

may occur either uniformly or irregularly. In sil- icate glasses, which have a relatively open structure, the displaced material is most readily accom- modated by compaction within a zone underneath

the indenter. However, in soda-lime glasses and in most ceramics shear deformation dominates with some compaction also taking place underneath the
indenter. The experimental model used by YofFe to describe the deformation underneath a sharp conical indenter in soda-lime glass is based on the work

of Peter [4], amongst others, who showed that the

mean contact pressure under a sharp indenter re- mains constant and relatively independent of the

indenter geometry. This model assumes that as a conical indenter is applied to the solid surface

yielding initially occurs in a hemi-spherical bowl

under the indenter and continues until the yielded zone attains its stable preferred state for that pressure. As the indenter is loaded further, no additional flow occurs within this bowl, but fresh yielding takes place in an adjoining thin hemi-spherical

shell. This process continues and "it is as if a set of nested hemi-spherical bowls separated by plastic filler were forced down in turn, one within the

other, and become locked in the positions shown

in Figure 1(a) [1]." The yielding process occurs through a combination of shear deformation and

compaction within the yielded zone. Part of the

material displaced by the" indenter is accommodated by compaction within the yielded zone while

the remaining material is pushed into the sur- roundings causing a volume change. At the end

of the indentation process, the elastic half-space

is left with an over-large hemi-spheroid fixed in a

hemi-spherical cavity on its surface. This misfit

leads to a residual stress in the region outside of

the cavity. Consequently, as the material reacts to

this residual stress, cracking may occur within the

solid.

The "blister" field proposed by Yoffe [1] was for describing the development of cracks below a conical indenter acting on a brittle solid. Yoffe as- sumed that the inelastic zone underneath the in- denter is hemi-spherical in shape with a radius

equal to the contact radius (Figure la). The stress distribution in the surrounding elastic zone was obtained as the superposition of two stress fields—the

Boussinesq field for the point force which idealized

the indentation pressure distribution and a doublet

force system which idealized the localized inelastic

deformation. The doublet force system for the blis- ter field consisted of superposition of a point center

of expansion in an infinite elastic solid with an additional doublet to satisfy the free surface normal

stress boundary condition. Such a doublet force

system leads to a volume increase of the inelastic

deformation zone that must be taken up by com-

(b)

Figure 1: (a) Indentation by yielding of successive

bowls and (b) force doublet used for blister field
paction or elastic deformation of the remainder of

the solid. The volume increase is a measure of the

"strength" of the blister field. The strength of the

blister field varies with the applied force; this vari- ation is characterized by a material property which

has thus far defied a precise estimation. Qualitatively, Yolfe's model has predicted the origin and

growth of radial, median, and lateral crack systems

quite well. It's only disadvantage is from a quantitative point of view as a procedure for determining

the strength of the blister field for various indenter

geometries is not yet available.

Models to describe the sliding indentation frac- ture process in ceramics have not yet evolved to a

stage where the inelasticity under the indenter is accounted for completely [5, 6, 7]. This is in large

part due to the difficulty introduced by a lack of

information concerning the constitutive behavior

of ceramics in large hydrostatic compressive stress

fields such as that existing underneath a microindentation. It seems that a reasonable first step for modeling the sliding indentation stress fields would

be to extend Yoffe's work on static indentation to the sliding situation. This extension should include

the load-history eflfects of inelastic deformation left behind by the sliding indenter and the elastic ef- fects due to the frictional force. In this paper Yoffe's blister field is extended to model sliding microindentation stress fields. The

extension of the blister field is called the sliding

blister field and its derivation is detailed in Section

2. In Section 3, the predictions of this theory are compared to the experimental observations of Ahn

et al [8] of fractures around sliding indentations.

The paper concludes with a brief discussion of the

experimentally observed cracking that is explained

qualitatively by the sliding blister field model.

2 Derivation of Sliding Blister

Stress Fields

A schematic illustration of the sliding microindentation geometry and the location of the inelastic deformation zone is shown in Figure 2. The

pointed indenter is loaded normal to the surface

by the load P acting in the ^-direction and tangentially by the friction force Q acting in the x- direction. As we are interested in details of the

stress field outside of the plastic zone, for simplicity we use the Boussinesq and Cerruti solutions

scratch

groove \

indenter

7

plastjc; zone

Figure 2: Schematic side view of sliding microindentation and moving coordinate system

for concentrated point forces rather than the ac- tual contact pressure distribution acting between

the indenter and specimen surface. It is assumed

that the inelastic deformation occurs in a region

whose boundary in front of the sliding indentation

takes the shape of one fourth of a sphere with ra- dius a equal to the half-width of the groove. Behind the indenter this zone takes the shape of a

half cylinder of radius a and length equal to the

total sliding distance. In the following, the well

known solutions for the concentrated normal and

tangential loading of a half space is given followed

by the sliding blister field solution.

Recalling the coordinate system of Figure 2, the

Boussinesq field for the concentrated normal force

P acting normal to the surface of a semi-infinite

half-space at the origin is [9]

T — xy

73

p "1 - 2u

2^ r

2

P' J P '1 -

2^ r

2

3

2 1

P' J P '1 -

27r r

2 Sxyz

P' .

z\ x^ — zy^

p J r'^ p3

(1) — x^ zx^

z. 1 p ) t'^ P^

1 -

z \ xy xyz

p / p3

(2)

(3)
3P

~

27r ^

(4)

where = x'^-\-y^ and = x-^ + i/'^ + 2^ define the

distance from the k)a(l to the held point in the xij

plane and the total distance, respectively, and u is Poisson's ratio. The corresponding stress components due to the concentrated tangential force Q

acting at the origin in the positive x-direction are given by [9]

Q_

2-K 3

(1

+ p^{p^zY

- 2v)

3x

Q_

2tv +

xy

L p'

2 +

(1-2/.)

2xy'

(6)

Pip +

pHp +

:)2

(7)

xy Q_

2n

3xr

I P' x'y

+ (i --2//)

2xh

y pip + -]

p^{p^zY p'^{p-\-zf

Q Sxz^

2-K p^

Q Sxyz

2k p^

Q 3x'^z 2 k

(8)

(9)

(10)

The superscripts n and /. denote normal and tan- gential loading, respectively. The loads applied to the indenter are supported by these elastic stresses which decay as l/p^ as p ^ 00.

In addition to the stresses due to the loading of

the indenter, there are stresses induced in the elastic zone due to the misht of the deformation occurring in the inelastic /one. Here we focus on these

stresses which are referred to as residual stresses since they remain in the elastic zone after the in- denter slides away. Since the sliding indentation

is performed (piasi-statically, the residual stresses

ca\i be treated as a superposition of the static in- dentation residual stresses continuously formed by

distributing YofFe's blister field. That is, the in- eUvstic deformation due to Yoffe's blister field ac- cumulates behind the sliding indenter.

We shall therefore begin with the blister field derived by Yoffe [1] to describe quasi-static (normal) indentation residual stresses. Yofle's blister

field is made up of the doublet force system shown

in Figure 1(b) which consists of the superposition

of two full-space elastic stress systems. These two

systems are a center of dilatation, three orthogonal

doublets, and a doublet of twice this strength applied in opposite sense perpendicular to the surface

of the half-space. The resulting stress system, two

outward double forces in the surface plane and an

inward double force of equal magnitude perpendicular to it, satisfies the free surface condition along

the surface of the half-space. The stress components of the resulting blister field are [1]

2A

y 2A

z 2A

xy 2A

(1 - 2i/)(x-'^ 2y2) (_5 + 4i,)a;2

(1 - 2i/)(2x''^ + 37/)z2 15x22

j.2p5

+ (11)

(1 - 2/.)(2x2 + y^ {-b + 4iy)y^

f2p3

{[ -2u){3x' + 2y)^z^ Iby'z.2J2

y.2pi + (12)

1 - 2u ^ 2(4 + ^)^^ - (1 - 2vy

15; (13)

{\-2v)xy (5 - 4r/)xy

~1~ n j.2p3

{l-2y)xyz'^ Ibxyz'

f2p5

yz _

'6yz 15t/rn

2A .P' P' J 6xz 15xz^'

2 A .v~ P' \

74

(14)

(15)

(16)

where .4 is the strength of the bhster field. To superpose the blister field as the indenter

slides, the strength of the blister field per unit slid- ing length, B, is defined a,s A — B{^)d^. Thus,

defining a function / such that = Af„{x,y,z),

the residual stress due to the sliding blister field is

•b

where the indenter slides from x = a to x = b. The remaining residual stresses are calculated us-
ing the same procedure. The experimental configuration of interest corresponds to sliding the in- denter over a length, large compared to the contact size, and examination of the stresses near the

indenter's present location. This special case cor- responds to -0(0 = & constant, a = — oo, 6 = 0,

in which integration leads to the following residual

stress field: My"" - 2B

4„.2 X {2ux y

-2/ + Avy^ - 2ux'^z^

2uz^)

2B

+

(2/2 + ^2)3 (y2^22)3p5

+6a;2/ - 2z/a;2/

+4y^ - 2i// - ^x'^y'^z'^ - Ix^y'^z^

-Gux^y'^z^ - 2y^z^ - Suy^z'^

(17)

X {2xY

-Ux^y'^z'^ dux^y^z'^ ISy'z'

2B

-Uuy'^z'^ + x^z^ - 2vx^z^

-Sy'^z^ - Suy^z^ + z^ - 2uz^)

2^2(^2 _ 3^2

+

xz (^2^^2)3 (y2_^^2)3p5

-2x'^z'^ + lOx^y^z^ + I2y'^z^

(18)

X (Gx^*?/^

bx'^z^

6^^

xy 2B

yz 2B

zx 2B

y

2(1 - u)x^ + 2(1 - i/)y2 _

(19)

2uz^

- x^

+

xyz 4„.2 (y2 + ^2)3 (y2^^2)2p5

+ 10x2/ + 6/ - 4x^2 + 3t/S2

-10X^^4 _ 12^2^4 _

z(2x2 + 2y2 - z^)

9/)

X (4x y

(20)

(21)

In the above expressions for the residual stress

field, the origin is located at the present indenter

location (Figure 2).

Lastly, the complete stress field is obtained by

superposing the elastic field and the residual stress

field asa = a"" + + a'', r = r" + + r^

As is shown in the next section, qualitatively dif- ferent stress distributions are expected for different

relative values of P, Q, and B.

sliding

'direction

radial crack

scratch

groove

free surface

lateral crack

median crack

Figure 3: Schematic view of cracking induced by

sliding microindentation of brittle solids

Table 1: Classification of deformation patterns in soda-lime glass under a sliding Vickers indenter [8]

Normal

load (N)

Deformation pattern

0 ~ 0.05 No crack

0.05 ~ 0.8 Median cracking

0.8 ~ 3

Median and lateral cracking

with lateral crack growth

to the surface at higher loads

3 ~ 6 Median cracking and

crushed scratch track

75

3 Results and Discussion

With the derivation of the sliding indentation

stress field in place, it is now possible to analyze

observations of the microcracking about scratches

in brittle solids made by Swain [5], Misra and

Finnie [6], and Ahn et al [8]. Figure 3 shows a schematic view of the crack patterns which have

emerged from the studies, most of which have been

carried out at light loads in soda-lime glass, silicon,

polycrystaUine alumina, and Ni-Zn ferrite. The

typically observed crack patterns are median, lat- eral, and chevron(radial) cracks. The load regimes

at which the various cracks occur in soda-lime glass are summarized in Table 1. As noted earlier, all of these cracks are usually

observed to initiate at or close to the boundary be-
tween the inelastic deformation zone and the sur- rounding elastic solid. The subsequent growth of

these cracks occurs in the elastic region. Therefore, it will be assumed that these cracks are initi- ated and propagated by tensile stresses which oc- cur in the elastic material immediately adjoining

the inelastic zone. The stress fields are now ex- amined in detail to determine whether the peak

tensile stresses are consistent with the occurrence

of experimentally observed cracking. Under this

assumption, the occurrence of the median crack

is thus related to (7y(x,0,a), the lateral crack

to <7z{x,0,a), and the radial crack to cr^(.T,G,0).

Therefore, the magnitudes of these stresses are now

compared to predict which crack patterns are most

likely to occur at a given load. The calculations

will refer to detailed experimental observations in soda-lime glass. Hence its Young's modulus, E =

70 GPa, and Poisson's ratio, u = 0.25, are used in Equations (1-21) the ensuing calculations.

3.1 Residual Stress

The consequences of the residual stress field pro- duced by sliding indentation are examined first. Figure 4 shows the calculated residual normal

stress distributions at several locations along the

X-axis in soda-lime glass. For comparison to subsequent plotting of the complete stress field,

the residual stresses are nondimensionalized using

PI'2a? (the average contact pressure) and plotted

for the case BI P = 0.005. The stresses correspond

to the indenter having moved from x — — oo to X = 0, its present location. The residual stress

field wiU approximate the complete stress field for large negative values of x/a. The residual stress a\{x^{S^a) is tensile behind the sliding indenter.

On the other hand, the normal stresses (T^(x,0,a)

and (7^(x, a, 0), are compressive almost everywhere

behind the indenter. Furthermore, their magnitudes are much smaller than that of a^. The tensile nature of (T^(x,0,a) behind the indenter is a

possible cause of lateral cracking in a plane per- pendicular to the ^--axis in the wake of the indenter. Such lateral cracking was commonly observed

by Ahn et al [8] at normal loads in the range of

0.8 ~ 3 Newtons.

The residual stress field is in plane strain in the

y — z plane once the indenter has receded suffi- ciently far away from that point, i.e. for suffi- ciently large negative values of x/o. Figure 5 shows

0.09

^ 0.06

c\j - 0.03

CO 0.00

-0.03

-

o;(x.a,0)

1 , , , , 1

76

3.0 -2.0 -1.0 0.0 1.0 2.0 3.0 Position (x/a)

Figure 4: Residual normal stress distribution at the edge of the plastic zone (B/P = 0.005)

the residual stress distribution in this plane strain

field. The normal stress (j^(— 2a,y,a) reaches its maximum tensile value at y = 0 which is consis- tent with lateral crack initiation behind the indenter. If the growth of this crack is determined by

the maximum local tensile stress, then the direction of crack propagation may be also predicted

using the residual shear stress in Figure 5. The

sign of the shear stress for \y\ < a is consistent

with lateral crack growth that begins parallel to the X — y plane subsequently turning slightly out

the plane and growing towards the surface as indicated schematically in Figure 2. The fact that cr^ is tensile away from y — 0 is also consistent with lat- eral crack growth towards the surface. These conclusions are drawn from the assumption that the

crack grows perpendicular to the maximum principal stress. However, the stress results alone do not

explain well why the observed lateral crack is much

longer than the width of the scratch. A more detailed fracture mechanics analysis that includes the

effect of the cracking on the stress field is required

to fully explain the lateral crack growth behavior.

Such an analysis using three dimensional boundary

elements is presently underway.

3.2 Complete Stress Field

The complete stress field requires the normal (P)

and tangential forces [Q) which are related by

Q^fiP
0.100

0.075

0.050

0.025

0.000

-0.025

-n n5n

' 1 ' 1

1

1 1

'

'

Ao;(-2a,y,a)

0.25

0.00

^-0.25

CO

CO2-0.50

CO

-0.75

-1 nn

o/x,0,a)

\ o^(x,a,0) /

\_/o^(x,0,a)

1 , , 1 J —^—. —1—,—,—..—

c\3 CM

CO

CO

CD

k_ CO

-3.0 -2.0 -1.0 0.0 1.0 2.0 3.0

Position (y/a)

Figure 5: Plane strain residual normal and shear

stress distributions versus position perpendicular

to the scribing direction (B/P = 0.005)

where /x is the coefficient of friction. When slid- ing a Vickers diamond indenter against soda-lime

glass at extremely slow speeds (5 mm/min) and

at low loads (0.1~4 N), Ahn et al [8] measured ^ = 0.13 ± 0.033. Thus, to compare the stress field with crack observations the following calculations use jj. = 0.13. Figures 6 and 7 shows trends in the complete normal stress field acting in the un- cracked body at the locations where the median,

lateral, and radial cracks are observed to occur.

In particular, the stresses are given for two ratios

of B/P. The normal stress ay has its maximum at X ^ 0 so that median cracking is expected just un- derneath the indenter. The normal stress cr^ has a maximum a^t x ^ —1.3a. Thus radial cracks could

be initiated at a distance behind the moving indenter. In the range of x <~ —2a, the normal stress az maintains its maximum value, so that lateral

cracking is expected to occur in the residual stress

field. For B/P <~ 0.005 (see Figure 6), the max- ima of the tensile stress and a~ are smaller than

that of ay. Thus median cracking only is expected.

For B/P >~ 0.005, the maximum value of the nor- mal stress a^ is still small but the maximum value

of az is greater than that of ay. Therefore, in this case, it is expected that a median crack first oc- curs under the moving indenter followed by formation of a lateral crack after the sliding indenter has

passed. There could, therefore, be a time delay in emergence of lateral cracking after median cracking. Indeed, it is commonly observed in sliding

-2.0 -1.0 0.0 1.0 2.0 Position (x/a)

Figure 6: Complete normal stress distributions at the edge of the plastic zone (B/P = 0.0025)

indentation experiments that there is considerable

lateral crack growth after the indentation process

is completed.

The stress field can be observed more completely

through the contour plots of Figures 8-11. The

contours are given for the elastic zone only to emphasize that the sliding blister field does not model

the details inside the inelastic zone. Figures 8-9

show that the lateral cracks initiate at the base of

the inelastic zone behind the indenter. Figures 10- 11 indicate that the median crack is induced un- derneath or ahead of the indenter primarily by the

applied loads rather than the blister field stresses.

In conclusion, the tensile stresses which determine the various types of cracking are az(—2a, 0, a)

for lateral cracking, ay{0, 0, a) for median cracking,

and (T^( — 1.3a, a, 0) for radial cracking. Furthermore, the formation of median cracking alone is generally expected only for small values of B/P i.e. when the strength of the residual stress field, B, is small or P is large. Also it can be said that lateral

cracking will also occur when the residual stress

field becomes significant compared to the stress

field due to the applied loading.

4 Sliding Blister Field Strength

The stress fields derived from the sliding indentation model depends on the factor B which may be

termed, by analogy with YofFe's blister field con-
n3 OJ

Ui

CO

a>

0.50

0.25

0.00

-0.25

-1 .00

\ o,(x,a,0) /

o^(x,0,a)

1.—,—,—,—1—,—,—.—,—1—,—,.—.—

'-'2.0 -1.0 0.0 1.0 Position (x/a)

Position (y/a) 20 30

2.0

Figure 7: Complete normal stress distributions at the edge of the plastic zone (B/P - 0.025)

Figure 9: Complete stress contour a/la} jP^ x: 3a, B=0.025 P

Position (y/a) 20 30 Position (y/a)

-20 -30 -r

Figure 8: Complete stress contour (7^2a^/P, x=0,

B=0.025 P

Figure 10: Complete stress contour ay2o? I P, x— 0,

B=0.0025 P

78
Position (y/a)

2 0 3 0

Figure 11: Complete stress contour Gyla^ jP, x=0,

B=0.025 P

stant, A, as the strength of the sliding indentation

blister field. At present there exists no well defined

experiment which can be used to unambiguously

obtain the values oi A or B. Next B is estimated

from certain experimental observations. This es- timated value of B is then used to calculate the

stress field explicitly.

Bridgman and Simon [3] showed that most

glasses do undergo a certain amount of compaction

under stress. For calculating B, it wiU be assumed

that any compaction of the glass which occurs under the sliding indenter is completely accommodated within the plastic zone. Furthermore, the

remaining volume of material displaced by the in- denter is assumed to be pushed out of the plastic zone. For static indentation with a sharp indenter,

the increase in volume of any hemi-sphere of radius

p {p > a — plastic zone radius) is independent

p and the volume change is given by

1 - 2u AV = SttA^^ (22)

where A is the strength of the doublet producing

the static indentation blister field and G is the

shear modulus. Figure 12 shows a schematic of a typical profile taken across a scratch track in glass produced by a sliding Vickers indenter at

light loads (~ 0.05 - O.liV). See Ahn et al [8] for actual scratch profiles. Small humps appear on ei- ther side of the scratch track, which in the absence

of any cracking around the indentation, is mostly

due to upward flow of the material. Such a profile suggests that some of the material displaced

Figure 12: Sketch of surface profile across a scratch

track made by a sliding Vickers indenter in sodalime glass (load ~ 0.05-1 N). The ratio of vertical to horizontal magnification is 5:1

by the indenter is also accommodated by this upward flow. The upward plastic flow factor, Fp, is defined, using Figure 12, as

Fp = AOAD - {AABC + ADEF)

AOBE (23)

where AOAD is the area of the triangle having

vertices O, A, and D. While estimating 5, it is necessary to also account for this upward flow of

the material.

By analogy with the static indentation blis- ter fleld, we can estimate the increase in volume

around an infinitely long scratch due to the sliding

indentation blister field as

lim/_oo

lim/^oo

j'BAV{x-^,y,z)d^ =

+

2a^

tan 9

' 3 tan 6

87ra2\ 6V,

~V

(24)

79

where SVc/V is the compaction ratio for glass and

Fp is the upward plastic flow factor. The terms

in the first parenthesis of the right hand side of

Equation (24) represent the impression volume of

the sliding indentation and the terms in the second

parenthesis represent the original volume of plastic
deformation zone. The parameters a and 9 are defined in Figure 12. Combining Equations (22) and (24) gives

B

27r(l - '2v) Vtan^

1 2 V

:25:

and B can now be estimated if the compaction

ratio [bVcIV) and Fp are known.

The value of compaction ratio for soda-Hme glass can be estimated using a curve fit of the data of

Cheng and Finnie [10] for the compaction of sodalime glass under hydrostatic pressure. Their equation is

8V,

%)

0.951 - 0.729p -f 1.93V + 0-260p^

(TTpp

(26)

where p is the applied hydrostatic pressure in GPa.

In order to estimate the values of a, 6, and Fp, ex- perimental measurements of the width (2a) and

depth (a/tan^) of a scratch track in soda-lime

glass were carried out for various applied loads

(0.05 to 0.8 N). From the experimental data the

half width of the track, a (^m), can be expressed

as a function of P (N) as

a = 9.48P°-^^ (27)

The average value of 0 was obtained from Talysurf

profiles of a number of scratch tracks as 83.6" ±

1.46. Using the profilometric measurements, the

average value of Fp was estimated as 0.055.

We are now in a position to evaluate 5 as a function of normal load P. This variation is shown in Figure 13. The value of B increases monotonicaUy

with applied load P. Due to the nature of the

increase in B with P shown in Figure 13, it is ex- pected that the strength of the residual stress field eventually dominates the Boussinesq and Cerutti

stress fields at large normal loads P. This is clearly

evident in Figure 14. At low loads (<~ 0.6 N) the

stress (Ty(0,0,a) is the dominant tensile stress so that median cracking is expected in a plane per- pendicular to the y-axis. This is consistent with

the absence of lateral cracks in this regime. But

for P >~ 0.6A^ the tensile stress a^i-'Ia^O.a) be- comes more significant. This is consistent with the

observation of both median and lateral cracks in this load regime in Ahn et al [8]. It is hard to anticipate radial cracking because the tensile stress

25.0 50.0 75.0

Force, P(gm)

100.0

80

Figure 13: Strength of residual stress field as a

function of applied load

cr2;( — 1.3a, 0, a) is always much smaller than the re- maining stresses, cry(0, 0, a) and (J2(— 2a, 0, a). Per- haps this explains why the radial cracks were only

occasionally seen in soda-lime glass and, even when

they were observed, the length of the radial cracks were quite small.

The analytical predictions are in close the ex- perimental classification summarized in Table 1.

It must be noted, however, that perturbations in the stress field due to energy dissipation by one

or more crack systems have been largely ignored.

Also, it is at present difficult to explain the crushed

scratch track produced in the load range of 3 — 6A'^ in soda-lime glass using the analysis.

5 Conclusion

The stress field produced by a sliding sharp indenter in brittle solids has been derived from an extension of YofFe's model for static microindentation.

Numerical results are illustrated and compared

with experimental observations. The strength of

the residual stress field has been evaluated as a

function of the normal load has been discussed. It has been shown that the present analytical model

provides reasonably good descriptions of the crack

patterns observed around sliding indentations in brittle materials. In particular, the analysis ac- counts for the occurrence of lateral cracking under

sliding indenters above an experimentally observed

threshold load. These conclusions are presently be-
0.100

--0.075

Q- iso

(0m

s? 03

0.050

0.025

0.000

aJO.O.a)

o,(-1 .3a,a,0)

25 50 75 100

Force, P(gm)

[6] A. Misra and I. Finnie, "On the Scribing and

Subsequent Fracturing of Silicon Semiconductor Wafers," Journal of Materials Science, 14,

pp 2567-74, 1979.

[7] S.Y. Chen, T.N. Farris, and S. Chandrasekar,

"Sliding Microindentation Fracture of Brittle

Materials," STLE Tribology Transactions, 34

(2) pp 161-168, 1991.

[8] Y. Ahn, S. Chandrasekar, and T.N. Farris,

"Deformation Observations about Sliding Microindentation in Soda-Lime Glass and Ni-Zn

Ferrite," Journal of Applied Mechanics, 1993,

Submitted.

Figure 14: Tensile stress at crack location as a

function of applied load

ing further addressed through inclusion of the blis- ter field stresses in a complete fracture mechanics

analysis.

[9] K.L. Johnson, Contact Mechanics,

bridge, Cambridge, 1985.

Cam-

[10] W. Cheng, E. Ling, and L Finnie, "Median

Cracking of Brittle Solids Due to Scribing

with Sharp Indenters," Journal of the American Ceramic Society, 73 (3) pp 580-586, 1990.

Acknowledgment This work was supported in part by the National Science Foundation through

grants MSS 9057082, Dr. J. Larsen-Basse, Program Director and DDM 9057916, Dr. B. Kramer,

Program Director.

References

[1] E.H. YofFe, "Elastic Stress Fields Caused by

Indenting Brittle Materials," Philosophical

Magazine A, 46, (4) pp 617-628, 1982.

[2] D. Tabor, The Hardness of Metals, Oxford

University Press, 1951.

[3] P.W. Bridgman and I. Simon, "Effects of Very

High Pressure on Glass," Journal of Applied

Physics, 24 (4) pp 405-413, 1953.

[4] K.W. Peter, "Densification and Flow Phenomena of Glass in Indentation Experiments," Journal of Noncrystalline Solids, 5,

pp 103-115, 1970.

[5] M.V. Swain, "Microfracture about Scratches

in Brittle Solids," Proceedings of the Royal

Society A, 366, pp 575-597, 1979.
Itf)

•I !
ON-MACHINE ELECTRO-DISCHARGE TRUEING

FOR METAL BOND DIAMOND GRINDING

WHEELS FOR CERAMICS

K. SUZUKI

Nippon Institute of Technology, Saitama, Japan

T. UEMATSU and T. YANASE

Toyama Prefectural University, Toyama, Japan

T. NAKAGAWA

University of Tokyo, Tokyo, Japan

This paper describes a new method for effective and precise trueing and/or

dressing for metal-bond grinding wheels on a grinding machine by utilizing

electro-discharge machining. A simplified trueing unit with an electrode

mounted on the table of an NC grinding machine makes highly efficient and

precise trueing and/or dressing possible to the metal-bond superabrasive wheels

mounted on the spindle of the machine. The wheels trued by this method

showed a higher grinding ratio and lower grinding forces than the wheels trued

by a conventional method.

1. INTRODUCTION

Superabrasive grinding wheels, such as diamond or CBN wheels, are generally used in

the grinding of hard-to-grind materials such as cemented carbide, high speed steel, and

ceramics. Such superabrasive wheels require

trueing and dressing before use to eliminate

wheel eccentricity as a result of mounting on

the spindle, and to provide sharp cutting edges on the abrasive particles. Trueing is also

carried out to provide a desired shape on a wheel or to correct a dulled profile. But the

trueing of superabrasive wheels is rather

difficult especially in the case of metal bond

grinding wheels. Metal bond wheels have

several good features such as high grinding

ratio and ability of high efficiency grinding;

but their use requires the development of

efficient and precise trueing methods. This

paper describes a newly developed trueing

method, which uses electro-discharge

machining on a grinding machine, for metal

bond superabrasive wheels.

2. PRINCIPLE AND FEATURES OF

"ON-MACHI^^ ED TRUEING METHOD"

Trueing of conventional grinding wheels are

carried out easily with diamond dressers. But

the trueing of superabrasive wheels are rather

Spindle head

Brush for power supply

Fig.1 Schematic illustration of on-machine ED

trueing/dressing method(1).

difficult. Among three kinds of diamond

grinding wheels, i.e., resinoid, vitrified and

metal bond wheels, trueing of metal bond

wheels is the most difficult because of the

toughness of the matrix material. The authors

have developed a new trueing method named

"On-machine electro-discharge(ED) trueing

method" for metal bond diamond or CBN

wheels(l).

In this method, trueing is performed on a grinding machine by utilizing electro-discharge

machining between an electrode on the table

and a metal bond wheel mounted on the

spindle of an NC grinding machine as shown

83
in Figure 1. The expected features of the

method are as follows;

(1)Applicable to all metal bond wheels and

electrically conductive resinoid wheels.

(2)Precise trueing owing to on-machine

method.

(3)Precise trueing of small wheels and/or thin

wheels with negligibly small applied forces.

(4) Dressing of wheels by adopting finishing

ED conditions.

(5)Profile trueing by using NC movement of

a grinding machine.

(6)ED trueing with chemical solution grinding

fluids.

(7)Automatic trueing operation.

3. ON-MACHINE WIRE ED TRUEING

The most important factor in the

on-machine ED trueing is how to prevent

electrode wear, which reduces the trueing

accuracy. One solution for suppressing the

electrode wear is to adopt appropriate ED

conditions; but, this approach is difficult.

Another solution is to use a wire electrode like a wire EDG method(2).

Figure 2 shows a schematic arrangement of

the on-machine wire ED trueing method. In an actual set-up a wire electrode runs along a ceramic guide mounted on the trueing unit on

the NC table of the grinding machine as shown in Figure 3. It is possible to use conventional cutting or grinding fluids as the

dielectric fluid for the ED trueing.

Tlie main features of the on-machine wire

ED trueing are as follows;

(1)Accurate trueing/dressing is possible

because of no wire electrode wear.

(2)Trueing for a profiled wheel is possible

by utilizing NC movement of the grinding

machine, and no profiled electrode is necessary.

4. ON-MACHINE ED TRUEING WITH

A BLOCK ELECTRODE

4.1 Principle and Procedure

ED trueing with a block electrode is more

preferable than the method with a wire

electrode from the view point of the trueing

efficiency. But there exists a problem of

electrode wear in the case of the block

electrode.

To solve this problem, an on-machine

Electric power source Wheel shaft

Metal-bond grinding wheel

Wire electrode

Fig. 2 Schematic illustration of on-machine wire

ED trueing/dressing method(3).

Fig.3

[Ceramic guide

A state of on-machine wire ED trueing of a metal-bond diamond wheel(D=30mm)(3).

1 1

1 1 1 1 1

Electric power

source

Rotary

trueing

unit

[Tool magazine]

.Or

Cuner for tvletal-dond Wheel to be

j eleclroae DIa. wheel formed

ED trueing

Workplece

|NC labia of MC]

Fig. 4 Arrangement of equipments for the

on-machine ED trueing with a block

electrode on an NC grinding machine.

electrode forming method has been developed

by using NC movement and automatic tool

changing function(ATC) of a grinding

machine(3,4). Procedure of the on-machine

ED trueing method combined with the

on-machine electrode forming method is as follows(Figures 4 and 5):

84
Spindle head of grinding machine

r I I 1,111

Power source

Worn

electrode

Rotary truer NC table

(a)Fornning of electrode (b)On-machine ED trueing (c) Correction of worn electrode

(d)Finish ED

trueing(dressing)

Fig. 5 Procedure of the on-machine ED trueing method with on-machine electrode forming.

(4)ED trueing is conducted between the

formed graphite electrode and the grinding

wheel(Fig.5b).

(1)A turning or cutting tool is selected from

a tool magazine and is mounted on a spindle

of a grinding machine.

(2)Rotating graphite electrode which is set on

the NC table is machined with the selected

cutting tool to provide a desired

geometry(Fig.5a).

(3)By the help of ATC a grinding wheel is mounted on the spindle instead of the cutting

tool.

(5)When the wear of the electrode becomes

larger, the electrode is machined again to provide the accurate shape(Fig.5c).

(6)Final trueing is conducted with a finishing

ED condition(Fig.5d).

4.2 Advantages by "On-machine Electrode

Forming"

The on-machine electrode forming method

provides the following advantages:

(l)The shape of a block electrode, for a

rather complicated profiled wheel, is

Formed electrode "W' (graphite)

Wheel

shaft

Fig.6 Machining of a graphite electrode Fig. 7 Profiled metal bond diamond grinding

(D=80mm) to provide a desired shape(3). wheel(D=30mm) by the on-machine ED trueing.

85
generated accurately and efficiently with an NC program on a grinding machine.

(2)Dulled shape of the block electrode, due

to wear, is corrected with a cutting tool

mounted on a spindle of an NC grinding

machine.

4.3 Form Trueing

Figure 6 shows the machining setup for

providing a desired shape on a graphite

electrode. Figure 7 shows s profiled grinding

wheel after on-machine ED trueing. It took

only about 15 minutes to form the wheel from

the original cylindrical shape. With the

profiled grinding wheel, it is possible to carry

out profile grinding of work materials set on

the table of a grinding machine immediately.

Figure 8 shows a ceramic sample with a complicated shape ground with profiled wheels.

4.4 Trueing Efficiency

The effects of trueing conditions such as no-load voltage, settled peak current, and duty

factor on the trueing efficiency, and the

electrode wear were investigated with bronze

bond diamond grinding wheels. The diameter

of the wheel is 30mm and the width of trueing

is 5mm. The electrode was fed automatically

toward the grinding wheel so as to keep the

gap distance constant. The trueing efficiency

was expressed by the feeding speed f^ of the

electrode. The following observations were

made.

(1)fd increased with the increase of

Eo(no-load voltage), and it reached

0.18mm/min in the case of Eo=150V.

(2)fj increased with Ip(settled peak current),

and it reached 0.41mm/min in the case of

Ip=45A(Figure 9).

(3)f(, increased with the increase of pulse ontime and decrease of off-timer^ff. And

the wear of the electrode(e) decreased with

the increase of t<,„ and decrease of T(,ff(Figure

10).

(4)Positive polarity provided higher trueing

efficiency and low electrode wear.

(5)Graphite was better than copper and brass,

when chemical solution grinding fluid was

used as the dielectric fluid.

(6)Relative rotational speed of the electrode

and the grinding wheel had only a minor

Fig. 8 Silicon nitride ceramics ground with profiled

wheels made by the on-machine ED trueing.

0.6

0.5

0.4

re cc 0.2

0)

0.1

Eo=150V,Vr=Om/min,

Wheel(D=30mm,SD1 70/200)

(1.5A)

(11.5) Mean Current

J60

50

to" 40 w

CD

(U 30 §

20 Su

03 UJ 10

0

10 20 30 40 50 60 Peak Current lp(A)

Fig. 9 Relation between settled peak current and

trueing efficiency.

0.6 _ 0.5 c

E^ 0.4

E."

0.3

0.2

0.1

0

(5DCa(U

a> LL

E„=1 50V,lp=20A,Vr=0m/min, _ Wheel(D=30mm,SD1 70/200)

a(2-5A) Mean Current

(4.5)

~A- A~

(3.5) e

- 50

- 40

(7) —

20

-/h

60

OJ 30 §

20 £o

iij 10

40 60 80 100 200

Pulse On-time \„{fJis) Fig. 10 Relation between pulse on-time and

trueing efficiency.

effect. But the small relative speed

provided better trueing efficiency and less electrode wear(Figure 11).

86
4.5 Wheel Surface after ED Trueing/Dressing

Figure 12 shows the surface of a diamond

wheel after ED trueing/dressing. It is observed

that the diamond abrasives protrude sufficiently

with no deterioration and are grasped firmly by

the metal matrix. Similarly dressed surfaces were obtained for CBN wheels.

4.6 Grinding Ability of the Wheel

The grinding ability of a wheel dressed by

ED technique was compared with that of a wheel dressed with a conventional WA(white

fused alumina) stickstone(Figure 13). The

grinding force for sintered silicon nitride

ceramics was considerably lower in the case of

the ED dressed wheel. The grinding ratio GR

became higher by selecting an appropriate ED

condition(smaller pulse width in this case).

These tendencies were observed just in the

same manner in grinding cemented carbide.

4.7 Dimensional Accuracy in ED Trueing

Dimensional accuracy in ED trueing was

investigated by forming a typical shape as shown in Figure 14, which includes some

peaks with various angles, a height of 2mm

and a radius of 0.2mm. The accuracy was

expressed in the form of the error between an

electrode and a workpiece ground with a formed wheel. When the electrode was not

reformed in the course of trueing process, the error was rather large, for example the error of

the height was 0.1mm. But by reforming the

electrode once or twice, the error was reduced

to a great extent. The errors after two

reforming cycles were under 0.01mm in height,

O.OS-^-O.OSmm in radius, and -0.04~-0.38° in

angle. Figure 15 shows a formed

electrode(graphite), a trued wheel(metal-bond,

SD600) and a ground sample(alumina ceramic).

5. DEVELOPMENT OF POWER SOURCE

5.1 Servo Control of Gap Distance

The gap distance between the electrode and

the wheel has to be kept constant for stable

and efficient ED trueing. For this purpose a simple control circuit was connected between

the ED power source and the feed rate override function of the NC grinding machine

as shown in Figure 16. The electrode on the

table is fed intermittently toward the wheel

with a feed rate of 100, 20 or 0% of

^ 0.6

E. 0.4 —

mI 0. a0)

0)

Eo=150V,lp=20A,

Wheel(D=30mm,SD1 70/200)

[Vr=200m/min]

Fig.i

Vr=0 Vr=200m/min

Relative Rotational Speed Vr

1 Effect of relative rotational speed between

an electrode and a grinding wheel,

Fig. 12 Surface of a diamond grinding wheel

after ED trueing.

80

60 -

40

20

[Grinding conditions] *Wheel:D=150mm,B=10mm,

SD170.Q,Bronze

*V,=1500m/min,V,=l5m/mln,

. a=50/jm,b=4mm,Down cut *Worl<:Sl3N,{Hv1700)

1

[Dressing(D) conditions] WA(white fused aJumlna) sticl<stone : O ^Grinding ratio(GR=326) ED dressing(Eo=150V,lp=20A,T„„=5/Js) x„= 6ius^GR=423 x„=1^s--GR=326 • T„=60ps-'GR=160

0 500 1000 1500 2000 Amount of Stock Removal Vwlmm^/mm)

Fig. 13 The grinding ability of the wheel dressed

by the ED method in grinding silicon nitride ceramics.

[Top Radius Rg=0.2mm]

87

Fig. 14 Shape for investigation of trueing accuracy.
programmed feed rate depending on the

detected gap voltage. This servo control

system was confirmed to work satisfactorily.

Workpiece

(AIP3)

5.2 Power Source for ED Trueing

Conventional EDM power source can be

used for the on-machine ED trueing, but it is not always advantageous. A rather compact

and inexpensive power source was specially

designed for ED trueing. Table 1 lists the

main specifications of this power source. The

above mentioned servo control circuit for gap

distance is also included in the power source.

6. CONCLUSIONS

The authors proposed a new on-machine

ED trueing method for metal bond

superabrasive wheels, and investigated the

effect of the method. The wire ED trueing

makes precise trueing and dressing of profiled

wheels possible by using the NC movement of

a grinding machine; but the trueing efficiency

is not so high. On the other hand, the ED

trueing method with a block electrode

combined with the on-machine electrode

forming method showed superior features:

(1)Extremely high trueing efficiency

(2)Realization of profile trueing

(3) Correction of a worn grinding wheel on

the machine

(4)Automatic trueing

(5)Sufficient projection of abrasive

particles(=low grinding force)

(6)Provision of higher grinding ratio(=long

tool life)

The on-machine ED trueing method

requires a simple electrode unit and a simple

electric power source. [ACKNOWLEDGEMENTS]

The authors would like to express their

heartfelt thanks to Messrs. Sodick Co.Ltd.,

Niigata Engineering Co.Ltd., Johnson Co.Ltd.,

Yamazaki Mazak Corp. They extend sincere

thanks to Mr.S.Asano for his assistance.

[REFERENCES]

1. K.Suzuki, et al., "Proposal of ElectroDischarge Trueing with Twin-Electrode,"

Proc. of Autumn Conf. of JSPE, pp.575-578,

1985(in Japanese).

2. T.Masuzawa et al.. Annals of the CIRP

34/1, pp.431-434, 1985.

Fig. 15 Sample of formed electrode, trued wheel

and ground workpiece. (Metal-bond diamond

wheel,D=30mm,b=30mm,SD600)

Threshold volt. Eti

Threshold volt. El2(<Eti)

[ED servo control circuit] _

Fig. 16 Control circuit for on-machine ED trueing.

Table 1 Specifications of ED power source.

Capacity 3kVA

Size & Weight W600xD600xH800mm,90kgf

Monitor Volt.& Amp.-meter,

Override level meter No-load Voltage En=60 - 150V

Peak Current Ip=2 - 75A

Pulse Width On-time t„„ =1 - 200//S,

Off-time x„„=3 - lOO/us

Condition-set (1 )Manual-set,{2)Pre-set,

(3) External set Safety Device Automatic power-off device

88

3. K.Suzuki, T.Uematsu, and T.Nakagawa,

"On-machine Trueing/Dressing of Metal Bond

Grinding Wheels with Electro-Discharge

Machining," Annals of the CIRP 36/1,

pp.115-118, 1987.

4. K.Suzuki, T.Uematsu et al.," Recent

Advances in the Grinding of Ceramics and

Hard Metals," Proc. of Hard Metal Conf. , England, pp.1-16, 1988.
CERAMIC GRINDING WITH SELF-SHARPENING

DIAMOND WHEELS

R. S. HAHN

Hahn Engineering, Auburn, MA

The sharoness of resin-bonded diamond wheels was

observed to increase 20-fold when the wheels were

subjected to normal-force intensities of 30 to

50 N/mm. Once briefly subjected to those

levels, the wheel would cut 15 to 20 times faster

than the "as-dressed" wheel at moderate (lower)

force levels. The resin bond wears very rapidly

at the high-force intensities, causing wheelregenerative chatter to occur. Accordingly, the

"chatter-free-grind time" ( CFGT ) is short at the

high-force level . In order to obtain reasonably

long CFGTs it is necessary to operate the wheel at

force levels that are high enough to give good

stock-removal rates, yet not so high as to cause

frequent wheel dressing to remove the wheelregenerative chatter. The use of normal-force

grinding controls to measure the wheel work-machine

time constant, the system stiffness, the workremoval parameter , and normal-force thresholds are

descr i bed

.

ceramics and difficult to grind

materials confronts many

companies. The selection of

grinding wheels and the grinding

machine input variables to obtain

a satisfactory manufacturing

process is often difficult to

achieve

.

The development of a satisfactory

grinding manufacturing process

can be facilitated by

distinguishing grinding process

variables from grinding machine

variables and identifying the

important grinding process

INTRODUCTION!

The grinding of

variables. Grinding machine

variables, such as wheelspeed,

workspeed, feedrate, depth of

dress etc., are well understood.

Grinding process variables, such

as wheel sharpness, wheelwear

rate, threshoid normal stress,

wheel breakdown stress, workpiece

damage thresholds, surface finish

and self-excited chatter

characteristics need to be

quantitatively measured and

control led

.

89
I

WHEEL-WORK CHARACTERISTIC CHART:

Figure 1 illustrates a cylindrical grinding machine

executing a plunge grinding

operation. Upon feeding the

cross slide at the rate vf^ the

normal force between wheel and

work at the moment of contact is

zero. As the slide feeds, the

"springs" in the system compress,

building up the normal force and

causing the wheel and work to

mutually machine each other. The

time rate of change of the

workpiece radius vw, of the wheel

vs , and the machine deflection x

equals the feedrate vf , thus:

vw-«-vs + x = vf (1)

The volumetric rates

of stock removal Z'w and

wheelwear Z's per unit width of

cut are:

Z ' w=3 . 14«Dw*vw (2)

Z ' s=3. 14*Ds*vs (3)

where

Dw is the work diameter

Ds is the wheel diameter

The stock removal rate Z'w

and wheelwear rate Z's are

plotted in Figure 2 against the

normal force per unit width F'n,

for AISI 52100 steel. A rubbing

zone, a plowing zone, and a cutting zone are clearly

recognizable ( 1 ) . The rubbingplowing and the plowing-cutting

transition forces F'-j^, F'pc are

indicated. The slope of the

stock removal curve Z'w is

generally a straight line for

metals and is called the work

removal parameter WRP . That

parameter reflects the sharpness

of the grinding wheel and the

workpiece resistance to abrasion.

In the grinding of metalSj, as the

wheel becomes dull, the WRP falls

and F'no and F'pc increase.

Figure 3 shows a plot of

Z'w vs F'n for a 120 grit resin

bonded diamond wheel grinding the

ID of a Zirconia ring. The wheel

was first dressed using a diamond

dressing disc. Then grinding

tests were made at increasing

force levels, producing the

concave upward curve in Fig. 3,

followed by decreasing force | levels producing the straight ' descending line with slope of .25

cumm/s*N, The numbers at each

data point indicate the specific

energy in joules/cumm.

Figure 4 is a similar plot

for a 320 grit diamond wheel , again showing a concave upward

characteristic for increasing ». forces. However, when grinding

tests at decreasing force levels li were attempted, wheel- I regenerative chatter occurred as [1 indicated by the letter "c".

Also shown on Fig. 4 is the

wheelwear as obtained by ( micrometer measurement of the

wheel diameter before and after ; each grinding test. It is seen

that the wheelwear is negligible | except at the highest force level i where wheel -regenerative chatter

tends to occur. Wheelregenerative chatter is a phenomenon associated with ( wheelwear and is discussed in the

section below.

'

! The concave upward

characteristic for increasing \ forces may be explained in two

ways; i.e. (1) by erosion of the

resin bond to expose diamond i grits, thereby sharpening the j wheel, and/or (2) by increasing I the dominance of the brittle

fracture mode of material

removal. However, the straight

line characteristic for

decreasing force levels indicates | that a significant sharpening of i the wheel has occurred. In

addition, the used surface of the I wheel sparkles brightly as the i
result of the exposed diamonds.

The possibility of micro fissures

propagating deeply into the

surface seems unlikely since 1 mm

of stock Mas removed for each

decreasing force data point.

Therefore, it can be concluded

that brief exposure of the Mheel

to high normal stresses will

sharpen the wheel but wheelregenerative chatter will quickly

develop. (The average normal

stress can be found by pressing

the wheel statically against the

workpiece and measuring the

contact length under a given

normal force and applying a small

correction for the wheel-depthof-cut. The normal stress

corresponding to 44 N/mm in Fig.

4 was 28 N/sqmm (see Appendix).

CHATTER-FREE GRIND TIME:

Many commercial production

grinding operations operate under

conditions of dynamic

instability. That means that

self-excited vibrations will

eventually develop. Those may be

of the work-regenerative or the

wheel-regenerative type. In the

tests described here wheelregenerative chatter prevailed.

This type of chatter is analogous

to the development of a "corduroy

dirt road", wherein the pulsating

forces between an automobile

wheel and the road surface cause

a propagation of a wavy-road

surface to develop.

As pulsating forces occur between

wheel anv? workpiece, a wAvy

surface on the periphery of the

wheel tends to develop.

Figure 5 presents the

chatter block diagram for

cylindrical plunge grinding

machines (2). Associated with

the wheel -regenerative loop

involving vs, is the "wheelwear

stiffness Ks" representing the

ratio of normal force to depth of

wheelwear per wheel revolution.

Although the chatter theory is

based on linear systems and

considers Ks to be a constant,

the wheelwear stiffness for

vitreous and resin-bonded wheels

is strongly nonlinear and Ks

drops rapidly at high normal

stress levels. Since the

stability criterion in Fig 5 is

usually not satisfied in

commercial grinding machines,

the growth of vibration amplitude

is represented in Fig 6, where

the chatter—free grind time

indicates when the amplitude

reaches an unacceptable limit.

The chatter-free grind time can

readily be determined in many

cases, audibly , depending on the

natural frequencies of the

system. For the system described

in Table 1, the chatter frequency

occurs somewhat above the

wheelhead spindle natural

frequency, depending upon the

normal force, and was in the

region of 2000 Hz. For the point

at 44 N/mm in Fig 4, the chatter

free grind time was about 5 s. nee the chatter has developed it

is impossible to stop it without

retruing the wheel . For the

point at 44 N/mm in Fig 3 chatter

did not develop in the 10 s grind

time nor did it at the lower

decreasing force grinds, where

about 1 mm of stock was removed

for each grind test. The 120

grit wheel evidently is less

prone to chatter than the 320

grit wheel

.

In grinding steel the

chatter-free grind time is a very

strong function of the stock

removal rate Z'w (4) as shown in

Fig 7 for two different spindle

stiffnesses (3). A similar

behavoir is expected in grinding

ceramic materials. Although

frequent truing of aluminium

oxide wheels to remove the wheelregenerative wavy surface is

acceptible and is common

practice, the truing of diamond

wheels must be relatively

91
infrequent due to their much

higher cost. Since the dominant

parameter governing chatter

growth rate is the "wheelwear

stiffness Ks" and its functional

relationship with normal stress,

it is important to determine the

operating normal stress that

yields the desired chatter-free

grind time.

FORCE GRINDING:

On conventional feed rate

grinding machines the cross slide

is fed at given feed rates.

Normal forces are induced at the

wheel-work interface. Those

forces may be large or small

depending on the sharpness of the

grinding wheel and/or hardness of

the workpiece and are determined

by the WRP. As the wheel dulls,

the induced force increases for a given feed rate and may cause

mi crostruc tura 1 damage to the

workpiece. Those forces also

vary in fast production grinding

cycles with incoming stock

variations. The deflections in

the wheel-work system are

directly related to variations in

the normal force through the

system stiffness, and result in

size errors even though the CNC

slides are positioned to a high

degree of accuracy.

By introducing normal

force sensors into the machine

and driving the feed slide to

generate prescribed force levels,

deflections can be maintained

constant, deflection errors can

be eliminated, the sharpness of

the grinding wheel can be

monitored, and the wheel can be

operated at optimum normal stress

levels consistent with chatter

free grind time, wheelwear rate,

and workpiece damage.

A no'rma 1 -f or ce-sensing CNC

for internal grinding was used in

the grinding tests. In addition

to executing force cyc.les, the

control can also capture

grinding-process parameters. A

set of typical parameters are

shown in Table 1.

CONCLUSION

The use of normal force

sensors in combination with

computerized data acquisition and

control provide improved

performance for high precision

grinding operations. Force

control permits the wheel to be

subjected to the appropriate

normal stress level, consistent

with wheelwear rates, chatter

free grinding times, surface

integrity, and cycle time.

Force sensing can be applied

to monitor wheel sharpness to

reduce workpiece damage and can

be used to compensate for

deflections in the wheelwork

system to eliminate errors.

REFERENCES:

1. R.I.King, R.S.Hahn, "Handbook

of Modern Grinding Technology",

Chapman and Hall, New York,

N.Y. 1986.

2. R.Snoeys, D.Brown, "Dominating

Parameters in Grinding Wheel

and Workpiece Regenerative

Chatter", Proceedings of the

10th International Machine Tool

Design and Research Conference,

pp325-348, Pergamon Press,

1969.

3. R . P . Lindsay , "Chatter-Free

Grind Time of high and low

conformity Grinding Systems",

unpublished report, 1972.

4. K . Srinivasan , "Application of

the Regeneration Spectrum

Method to Wheel Regenerative

Chatter in Grinding", Trans, of

ASME, Journal of Engineering

for Industry, vol. 104, pp.46-

54, February 1982.

92
APPENDIX

:

To calculate the average

normal stress, the wheel-work

contact area is required. This

is obtained by coating the work

piece with a thin film of

Prussian Blue pigment and

pressing the wheel against the

"Prussian-blued" workpiece, and

axially rubbing it a small

distance, to determine the

contact length le. That length

may be modified because of the

"wheel depth of cut h".

The geometrical length of

contact Ig, where elastic

deflection of wheel and work are

negligible, is given by (1):

lg= \/Deq # h (4)

where

Deq =

and

Dw * Ds/(Dw +/- Ds) (5)

Dw = Work diameter

Ds = Wheel diameter

use + for OD grinding

- for ID grinding

h = vw/Workspeed (6)

If Ig <= le/2

contact length = le (7)

e 1 se

contact length = le/2 + Ig (8)

TABLE 1 COMPUTER-CAPTURED GRINDING PARAMETERS

Rough grinding force(N) ................ .300

Rough grinding velocity vw ( um/s )...,... 31 .

9

Normal force at final size(N) O

Normal threshold force Fth(N) ............. 2

Number of revolutions to round up Ill

System stiffness( N/mm ) ..........1385

Whee 1 -Work-Machine Time constant (s)... 4.5

Remaining stock (urn) .72

Final runout (urn) .0

93
TABLE 2 TEST CONDITIONS

Operation ID grind

Work speed Nw =6.7 rps

Wheel speed Ns = 173 rps

GMN high freq. motorized wheelhead 7.5 kw

Servo motor driven workhead

Workpiece Zirconia Ring 92mm 0D,76mm ID, 27mm wide

Coolant White & Bagley Heavy Duty SoluaPle #2213 5'/.sol

Dresser Wheel Truing Metal Bond Diamond 28mm OD x 3mm

Dresser Cross Axis to Wheel Axis

Grinding Wheels SD120-R100B56-1 /8

SD320-R150BXL6145-1/8

BASE

Wheel-wor k grinding system where

vf i s the cross slide feedrate

vw IS the radial stock removal rate

vs 1 s the radial wheelwear rate

Vs is the wheel peripheral speed

Vw i s the work peripheral speed

Fn is the normal component of force

Ft i s the tangential component of force

Ks i s the stiffness of wheel support

Kw I s the stiffness of work support

94
10.0

25

cuttng

H— H Wh^elwear Z s & - -o Stock RerriO^/al 2

e' WFf - Obi

Fth FfX: 10 20 60

f jcfmal Fcfce htensit/ 'T^/mm)

Fig 2. Whsel-work Characteristic Chart where

Z'w is the volumetric stock removal rate for unit width

Z's is the volumetric wheelwear rate for unit width

WRP(work removal parameter) is the slope of Z'w vs F'n

F'n is the normal force per unit width

Fth is the threshold force per unit width

F'pc is the ploughing-cutting transition

for the following conditions;

wheel = . . . A80K4

Dress lead(mm/rBv) .075

Dress deDth(um) 5

Whee 1 speed ( m/s ) 60

Workspeed ( m/s ) 1.25

Coolant Cim Cool 5 star

Work material AISI 52100 <3 60 Rc

95
MORtA^ -STRESS-SHARPENING OF RE^BTsl-BorD DIAMOND WHEELS

10

LU<

LUo

oI—

^^J^ffi^SND^cATE ^

SPEC Ef€RGr

(X!ULE/>:iimm:t --^ OECREASTK-.

FORCES/ /'

t9 /'

/ /

/ NCREASirJG

• FORCES

S012(>RK>3B56 1/8

ZIRCOr^lA

ID mi) 75mm

WHEELWEAR

10 '20 30 40

NORvlAL FORCE NTENSJTY (N/mm)

&0

to J

40

•20

5'j

3

LU

LU

UJ

Fig 3. Normal stress sharpening of resin-bonded diamond wheel

showing the stock removal rate Z w vs Normal force per

unit width for increasing force levels followed by

decreasing force levels. The cumulative wheelwear

IS also shown. The numbers over each data point

give the specific energy in joules/cumm.

Wheel 55mm OD by 21mm wide SD120-R100B56-1 /8

Workpiece .....ID Grind 75mm bore Zirconia Ring

Width of cut(mm) 10

96
MORTAL STRESS- SHARPENING OF REm BOhD CilAMOND WHEELS

1

i3

UJ

I—<DC

UJ

QC

\yOo

CO

10

C INCiOATES

CHATTER

13.'-

C

16

CECREAS^JG

FORCES

rslLtMBERSirCICATE

SPEdFIC BCRGY

gOULES/auTTfn;

ZIRCONIA IN-CREA.SING D GRiMD FORCES.-^

34'

P

80mm

50

.0

95 til 66 -JS ^

,0- —^

-• —f - —

21

•WHEaWEAjLo

Q- B'

10 20 30 40

NORMAL FORCE NTENS1TY d'lymmj

80

60 -

40 LU

Fig 4. Same conditions as Fig 3 except that chatter occurred

with the finer grit wheel as indicated for the points

marked with "C" The numbers over each

data point give the specific energy in joules/cumm.

Wheel 55mm OD by 21mm wide . . . . SD320-R100B56-1 /8

V=A exp (+at)

a1 82 83

acceptable limit

time

chatter-free grind time

Fig 6. Illustration of the chatter growth rate and the

chatter-free-grinding time for an unstable system

97
STAOSLITY CRtTERlON

Kd/Kw + Kd/Ks + 2Kd/K1 > 1

Fig 5. Chatter block diagram for cylindrical grinding machines

relating the regenerative feedback effect of wheelwear

and/or wavy work surface to the dynamic system

parameters and the stability criterion where

G(s) is the normalized dynamic compliance

Km is the static stiffness of the wheel-work system

Kd is the minimum dynamic stiffness of the system

Kw is the work cutting "stiffness"

Ks is the wheelwear wear "stiffness"

Kl is the wheel-work contact stiffness

see Snoeys and Brown ref (2)

98
1000

LU

CC<

o

100

10

»-•« static stiffness- 10.5 N/um

- static stiffness- 2a N/um

W

\

10

STOCK REMOVAL RATE {c\rYtVs"mm)

100

Fig 7. Chatter-free-grind time vs Stock-removal rate

for internal grinding with two wheel -supporting

arbors of different static stiffness.

\

99

CHEMOMECHANICAL EFFECTS OF BORIC ACID IN

CORE-DRILLING OF ALUMINA

S. JAHANMIR and H. LIANG*

National Institute of Standards and Technology, Gaithersburg, MD

Interactions between chemical compounds added to cutting fluids and the workpiece surface

in the cutting zone can have pronounced effects on the material removal process during abrasive

machining. These interactions can influence the coeflicient of friction, the wear of the abrasive

grit, and the mechanical properties of the workpiece, thus affecting the machining rate. Experiments were conducted on sapphire and polycrystalline alumina to evaluate the chemomechanical effects of boric acid mixed with distilled water. The machining tests were performed on a precision drill with metal-bonded diamond core-drills. Following the experiments, the drilled

surfaces and the diamond particles in the drill were examined by scanning electron microscopy

to elucidate the material removal process. The results indicated that addition of boric acid to distilled water increases the rate of drilling of polycrystalline alumina by a factor of two. But, boric

acid was found to be ineffective in improving the drilling rate of single crystal alumina, i.e., sapphire. Based on the results it is postulated that boric acid interacts with the amorphous oxide

grain boundary phase in polycrystalline alumina promoting intergranular fracture; thereby,

increasing the drilling rate.

Introduction

In abrasive machining processes such as grinding, material removal occurs at the contact between

the abrasive cutting points and the workpiece surface. The abrasive particle is forced to penetrate the

workpiece as the workpiece and the particle are moved relative to each other and parallel to the

surface. In this interaction, the workpiece surface

undergoes elastic and plastic deformation, followed

by fracture of small particles or chips from the

surface. Depending on the contact conditions and

the mechanical properties of the workpiece, chips are generated by plastic flow and/or brittle fracture [I]. The purpose of this paper is to elucidate the

effect of boric acid added to distilled water in abrasive machining of alumina. This compound is used

in cutting and grinding fluids for friction reduction,

corrosion protection, and as bactericides and fungi

* Research Associate, University of Maryland,

College Park, MD 20742

cides [2-9]. In these applications, boric acid is mixed with various compounds such as sodium tetra- borate, sodium tripolyphosphate, hydrofluoric acid,

triethanolamine, ammonium fluorosilicate, and

hexamethylenetetramine. It is not clear whether it is the boric acid or a compound formed by reaction of

boric acid and other chemicals that produce a positive result with these fluids. Therefore, in our study

boric acid was used as a single additive in distilled water in experiments using metal-bonded diamond

core-drills to determine the basic mechanisms by

which this compound influences the cutting process

in ceramics. In materials capable of plastic

deformation, for example, metals, chips are formed

by shear in a ductile mode. However, in ceramic

materials, which exhibit a limited ductility, chips are formed by plastic deformation and/or brittle fracture. Whether the chips are formed by brittle fracture or by plastic deformation depends on the penetration

depth and the material's susceptibility to plastic flow

as well as the environment [10-13].

Two models have been proposed to describe the

chip formation process in ceramic materials. The

101
model proposed by Evans and Marshall [10] is based on the formation and propagation of median cracks

(vertical to the surface) formed at the boundary between the elastic and plastic zones and the lateral cracks (parallel to the surface) formed as a result of

residual stresses during unloading. In the crushing

model proposed by Larchuk [14|, material removal

is described as an unstable fracture process, which

involves propagation of cracks at origins distributed

throughout the stressed region. The basic difference

between the two models is that in the latter a large

number of fragments are separated from the surface. The fluid environment used in abrasive machining can have a large influence on the removal

process [1]. These include temperature and friction control, swarf (or debris) removal, modification of

the workpiece mechanical properties, influences on

the wear and fracture of the abrasives, and reduction

of swarf transfer to the wheel.

Several publications have considered the chemomechanical effects, i.e., combined mechanical and

chemical interactions, during drilling of ceramics,

stones and glasses [15-18]. Arguments about the

effects of the fluid on diamond abrasive and the

workpiece were focused on how the chemicals affect the mechanical properties such as hardness, material

properties such as surface potential [19], and friction and wear of the diamond abrasives [20].

Generally, the chemomechanical effects of the

cutting fluids in drilling involve more than one mechanism. For instance. Swain et al. [15] studied

the effects of surface-active environments on hard- ness and drilling rate of polycrystalline alumina.

They found that the hardness of alumina was lower

in the presence of ethanol than in pentanol. The

material removal process in ethanol, therefore, in- volved substantial plastic flow preceding or accompanying intergranular fracture. Plastic flow was less evident, however, in pentanol, a more brittle behavior and large conchoidal fragments were observed.

The drilling rate of rotary diamond-studded bits were

greatest in pentanol when the material was hardest

and, presumably, most brittle. In ductile-regime grinding, Bifano et al. [13]

found an opposite result that alcohols with a small

number of carbon atoms promote embrittlement and

increase the machining rate of glass. Cuthrell [17]

postulated that the observed higher drilling rates in ceramics are primarily the result of the embrittling

effect of hydrogdn ions from liquid or gaseous

environments. Cooper [19], however, found that the

effect of liquids on drilling hard stone was that of a reduction of diamond wear rather than a direct effect

on the stone itself. Cooper and Berlie [20] reported

a significant improvement in the drilling rate upon

alternating the substitution of a series of normal

alcohols (propanol, pentanol, and octanol) for water

as the drilling fluid. The results were interpreted to reflect a reduction in wear of the diamond particles

in the drill bit. Considering that previous publications described

the role of boric acid in reduction of friction in slid- ing tests [21-23], we had postulated that boric acid

would reduce the friction coefficient at the

tool/workpiece interface. The reduction in friction and the resultant lower temperatures at the diamond/workpiece interface, would reduce the wear

rate of the diamond abrasives, and increase the

drilling rate by maintaining sharp cutting points.

However, we found that the chemomechanical effects of boric acid are complex and involve an embrittlement effect rather than affecting the wear of the

diamond abrasives.

Experimental Procedure

A high-purity polycrystalline alumina (99.8%) was used in the experiments. The properties of this material (Coors AD 998) are listed in Table 1. The

Table 1. Composition and properties of

polycrystalline alumina

Composition (%wt) <0.1% SiO^, 0.03% NajO, 0.02% TiOj, 0.01% KjO,

balance of Alfij

Grain size (\un) 3

Porosity (%) 0.2

Vicker's hardness (GPa) 15.0

Elastic modulus (GPa) 345

Compressive strength (MPa) 2071

Flexure strength (MPa) 331

Thermal conductivity

(W/mK) 29.4

Indentation fracture toughness (MPa.m"^)

I

3.2

Thermal expansion

coefficient (°C ')

1^ 6.7 X

10-*

102
specimens were in the form of circular disks with a diameter of 10 mm and a thickness of 3 mm. The

disks were cut from a cylindrical rod and were used

in the drilling experiments. In order to evaluate the

effect of grain boundaries in the process of material

removal, a limited number of tests were conducted on single crystal alumina or sapphire. In these tests the sapphire was oriented such that the (1010) plane was parallel to the flat disk surface. Pure orthoboric acid, H3BO3 (99.99% wt), was

added to distilled water at 4 %wt concentration.

Boric acid is a white and pearlescent solid which is soluble in water up to 6.35% wt. Two commercial

cutting fluids, designated A and B, were also tested

for comparison with the boric acid solution. Fluid A

is a low-foaming synthetic, whereas fluid B is a semisynthetic fluid. These two fluids were selected

for drilling because they are used in grinding of

ceramics. Table 2 lists the properties of the two

fluids and boric acid. A precision drill was used for performing the

drilling experiments. The drilling setup is shown in the schematic diagram in Figure 1. In this setup, a pump is used to circulate the cutting fluid from a 4

liter reservoir through the core-drill using a water

swivel. The fluid reservoir and all tubing were made

from plastics. A clamp was fabricated to fix the

alumina sample to the specimen holder, which was

enclosed in a fluid tank. Both the specimen holder

Table 2. Characteristics of cutting fluids

Fluid A Fluid B Boric

Acid

Solution

Concentration

in water 2% vol 1.2% vol 4% wt pH of test fluid 9 9 6

Solubility in water 100% vol 100% vol 6.35%

Boiling point

of compound

(°C) 100 100 300

Freezing

point of

compound(°C) -4 0 170

£md the tank were made from stainless steel to eliminate corrosion and other possible chemical

reactions.

The rotational speed of the drill could be adjusted to obtain three speeds (the maximum surface

speeds are listed in the parentheses): 2100 rpm (0.70

m/s), 3770 rpm (1.25 m/s), and 5500 rpm (1.83 m/s).

Metal-bonded core-drills with an outside diameter of

6.35 mm and a wall thickness of 0.13 mm were

used. A new drill had an approximately 9.5 mm

long diamond section, consisting of 30% vol natural

diamond (100 mesh or 150 iim) in a bronze matrix

bonded to a steel drill shank.

To define an optimum set of conditions for the

drilling tests and to characterize the effect of rota- tional speed and applied load, tests were first per- formed in distilled water under various loads and

speeds. The normal load on the drill was varied

from 23.0 to 46.1 N. Following these preliminary

tests, further testing with the boric acid solution and

commercial fluids were conducted at a load of 33.3

N (contact pressure of 132 KPa) and a rotational

speed of 2100 rpm (0.7 m/s) to avoid rapid wear of

the diamond particles.

The following testing sequence was used to reduce the experimental uncertainties due to the

differences in the condition of the drills, i.e., the

number and distribution as well as sharpness of the

diamond particles engaged in cutting during drilling.

First, the alumina was drilled for 15 s using distilled water, and the hole depth was measured by a dial indicator (with a precision of 0.01 mm) at the

completion of the first cycle. This was followed by

a second cycle in which the test fluid was substituted

c_r>

Water

Swivel

Metal-Bonded

Diamond

Core-Drill

Sample

(x) Pump

]

Fluid Reservlor

Stainless Steel Tank

Figure 1. Schematic diagram of the drilling setup.

103
for water in drilling the same hole for another 15 s. The hole depth was measured after each cycle. The

sequence of drilling alternating between water and

test fluid was continued for a total duration of 60-

180 s. Following these tests, the drilling rate for

each cycle was obtained by dividing the measured

hole depth by the drilling time. The relative change

in the drilling rate was obtained by comparing the

drilling rate for the test cycle in which the test fluid was used with the preceding cycle where only water was used.

During the drilling test, a 1 cm^ sample of the

cutting fluid was drawn from the stainless steel tank

for the analysis of the debris. The fluid samples were centrifuged for 45 min to separate the debris.

This step was followed by three subsequent washing

cycles in the centrifuge. The debris was then re- trieved in a distilled water solution, and the water was allowed to evaporate at room temperature. The

collected debris was coated with a thin film of gold

for examination in the SEM.

Prior to drilling a new hole, the drill was dressed

by drilling a porous silicon carbide dressing stick. Depending on the wear of the diamond particles, a dressing time of 10 to 30 s was used. The drills were then examined under an optical microscope to check the diamond particles and to ensure that the wear flat areas were removed by dressing.

A few tests were performed with a short test time to allow observation of the material removal

process. Prior to these tests, the polycrystalline

alumina samples were polished with a series of

diamond pastes of decreasing size from 30 to 0.25

^m. In these tests, the drill was engaged for only a few seconds to generate a small track on the alumina

workpiece.

After the drilling tests, the diamond drill, the

drilled surfaces and the debris were examined under

an optical and a scanning electron microscope to elucidate the material removal process.

Experiimental Results

a) Preliminary Tests in Distilled Water

A series of tests were performed in distilled water to determine the basic characteristics of the

drilling process. Each test was performed for 60 s on a new polycrystalline alumina surface, and the

hole depth was measured following the test. Figure

2 shows that the drilling rate of a new drill decreases

with the number of drilled holes. The rapid decrease

in the drilling rate is indicative of wear of the

diamond particles. The figure shows that as the drill

is dressed with a silicon carbide dressing stick for 30

s, the drilling rate increases; but it does not return to the initial high value.

Figure 3 shows the distribution of the diamond

particles in a used drill as observed in the SEM.

Figure 4 shows an optical micrograph of a section of

a used drill. In this micrograph, the diamond parti- cles that are highly reflective and appear bright have

been worn. It was determined that less than 10% of

the diamond particles were involved in cutting

during drilling. A worn diamond particle is shown

in the SEM micrographs of Figure 5, which indicates

that wear of the diamond occurs by micro-scale

fracture and attrition (i.e., tribochemical oxidation

reaction). Since the wear flat is very smooth, it is reflective in an optical microscope. No evidence for large scale fracture of the diamond particles were

found.

The main function of dressing is to sharpen the

diamond particles, to remove the bond, and increase

the clearance between the drill and the workpiece for an effective drilling. Examination of the diamond

particles in a dressed drill in the optical microscope

showed that the wear flat had disappeared after dressing. The distance between the metal bond and

the diamond particles engaged in cutting was measured in an optical microscope by focusing on the

diamond particles and the bond and measuring the

height difference. The average height of the diamond particles after a few drilling tests was 10 ± 5

|im. In a worn drill, the clearance was reduced to less than I ^im; and by dressing, the average height

value was almost restored. The drilling rate was found to be a function of

both the applied load and the rotational speed, as shown in Figures 6 and 7. In these tests, the drills were not dressed; and the test sequence is identified

by the numbers next to the data points in the figure.

The data points for the first three tests are not shown

since the condition of the drill changes rapidly at the

beginning of the test (see Fig. 2). Figures 6 and 7

show that the drilling rate increases as the speed is increased. However, the effect of load appears to depend on the speed.

At 2100 rpm, the drilling rate increases monotonically with load. However, at the higher speed of

3770 rpm the drilling rate decreases after reaching a maximum. The reduction in the drilling rate observed at the high speed and high load, must be due

to accelerated wear of the diamond particles under

these conditions.

104
1.60

\ 1-20 +

E

E

Q> "S 0.80 +

c

0.40 +

0.00

After Dressing

4 6 Number of Holes 10

Figure 2. Reduction of drilling rate by the number of

holes drilled in water (Load: 33.3 N, surface speed:

0.7 m/s, Test duration: 60s).

Figure 3. SEM micrograph showing the distribution

of diamond particles in a used drill.

Figure 4. Optical micrograph of a section of used

drill showing the worn diamond particles.

Figure 5. SEM micrographs showing the diamond

particles after drilling: (a) the diamond particles in the center of the micrograph has been subjected to wear, and (b) the worn diamond.

0.50

E 0.40

E

I

0.30 +

S. 0.20 +

c^ 0.10-

0.00

o 2100 rpm

• 3770 rpm

20 40

Load (N)

60 80

Figure 6. Effects of load and rotational speed on the

drilling rate of polycrystalline alumina in water (Test

duration: 60 s). 105
1.00

c U.ou • /mi

EJ,

0.60-- V

q: 0.40

o>

_co

0.20--

0.00 2000 4000 6000

Rotational Speed (rpm)

8000

Figure 7. Effect of rotational speed on the drilling

rate of polycrystalline alumina in water (Load: 33.3

N, Test duration: 60 s).

1.00

c 0.80-

1

0.60 -

(? 0.40 +

0.20--

0.00

o Distilled Water

• Boric Acid (4*wt)

—(

50 100

Time (S) 150 200

Figure 8. Drilling rate of polycrystalline alumina in distilled water and boric acid solution (Load: 33.3

N, Surface speed: 0.7 m/s).

b) Effect of Boric Acid on Drilling Rate

Based on the results of drilling tests in distilled water, a load of 33.3 N and a rotational speed of

2100 rpm (0.7 m/s) were used to avoid excessive wear on the diamond particles. Also, the drills were

dressed prior to starting a new hole to provide a uniform drill condition. Figure 8 shows the drilling

rate against the drilling time for a number of drilling

cycles alternating between distilled water and 4% wt

boric acid in distilled water. The figure clearly

shows that the drilling rate of polycrystalline alumina

increases on replacing the distilled water with the

boric acid solution. The results for 1 1 test cycles are tabulated in Table 3, which shows that the boric acid

solution increases the drilling rate by an average of

103% over the drilling rate measured in water.

The test results on the two commercial cutting

fluids, shown in Figures 9 and 10, clearly indicate

that the effect of these fluids on the drilling rate is small compared to the boric acid solution. The

average values for the change in drilling rate for the

commercial fluids, listed in Table 3, are 6% and

-18%. Considering the standard deviations, also

listed in the Table, it can be concluded that the

commercial fluids are ineffective in increasing the

drilling rate. The test results obtained in drilling of sapphire

in water and the boric acid solution, in Figure 11, are contrary to the results observed on polycrystalline alumina. Figure 1 1 shows that boric acid is ineffective in increasing the drilling rate of sapphire.

This provides a possible clue for the mechanism by

which boric acid solution increases the drilling rate of polycrystalline alumina. An interaction between

the grain boundaries in polycrystalline alumina and

boric acid is speculated as a factor in the material

removal process.

c) Mechanisms of Material Removal

In order to elucidate the material removal process in drilling and to discern the effect of boric acid on the process, the drilled surfaces were examined in the SEM. Observations on the sides of cylindrical

surfaces of the holes provided strong evidence for material removal by intergranular crack propagation,

as seen in Figure 12. Surfaces drilled in water and

those drilled in the boric acid solution, have the same appearance, essentially a fractured surface with

occasional plastic grooves. The bottom of the drilled holes, were also similar, showing no difference

between the surfaces drilled in water and those

drilled in boric acid solution, shown in Figure 13. The bottom surfaces, however, have a different

appearance than the cylindrical surfaces. Although,

there is ample evidence for intergranular fracture, the

bottom surfaces (Fig. 13) are covered with numerous

debris, some of which have been plastically deformation.A possible difference between the surface

drilled in water and the one drilled in the boric acid

solution is in the number of fine debris observed on

the bottom surface. Although somewhat subjective,

the number of fine debris in Figure 13 (a) for the

boric acid solution is lower than that observed in Figure 13 (b) for water. In fact, the distilled water

appeared cloudy after the tests; but became clear in a few hours. However, the boric acid solution

remained cloudy after the tests. This indicates a possible dispersion of fine debris in the boric acid

106
Table 3. Drilling rate for the three cutting fluids

1.00

TO.80

E

I

0.60

S. 0.40

c

't^ 0.20 --

0.00

Cutting Fluids % Increase in Dnlhng Rate Over

Water

Boric acid (4% wt) 103 ± 42

Fluid A (1.2% vol) 6 ± 19

Fluid B (2% vol) -18 ± 26

o Distilled Water

• Fluid A (1.2%vol)

30 60 90

Time (S) 120 150

Figure 9. Drilling rate of polycrystalline alumina for a number of drilling cycles in distilled water and

Fluid A (Load: 33.3 N, Surface speed: 0.7 mis).

1.00

E 0.80

E

I

0.60

S. 0.40

c% 0.20

0.00

o Distilled Water

• Boric Acid (45Kwt)

20 40

Time (S) 60 80

Figure 11. Drilling rate of sapphire for a number

of drilling cycles in distilled w^ater and boric acid

solution (Load: 33.3 N, Surface speed: 0.7 m/s).

1.00

E 0.80

E

I 0.60

i§ 0.40

c^ 0.20 --

0.00

o Distilled Water

• Fluid B (2*vol)

30 60 90

Time (S)

120 150

Figure 10. Drilling rate of polycrystalline alumina in distilled water and fluid B (Load; 33.3 N, Surface

speed: 0.7 m/s).

Figure 12. SEM micrographs of the cylindrical

surfaces of holes drilled: (a) in distilled water,

and (b) in boric acid solution (Load: 33.3 N,

Surface speed: 0.7 m/s). .. . .. , ^ , 107
solution. Since the unused boric acid solution is clear, and stays clear when stored, these observations

provide a second clue for the mechanism of boric

acid solution; mainly, dispersion of fme debris.

The debris, or swarf, collected after the drilling

tests, are as shown in Figure 14. The debris appear

to have a bimodal size distribution; one at 3 to 4 ^im

and the other much less than 1 |im. It is also found

from Figure 14 that the debris size in water is somewhat smaller than that collected from the boric

acid solution. The larger particle size in the boric

acid solution, suggests a more efficient material

removal process in that solution than in water.

d) Results of Short Drilling Tests Since the drilling process is complex, it is diffi- cult to identify the material removal mechanism by

examination of the drilled surfaces and the debris.

Therefore, a series of short tests were conducted on

polished surfaces to characterize the removal process

at the initiation of drilling. In these tests, the drill was engaged with polished surfaces of polycrystalline alumina and sapphire only for a few seconds;

and the surfaces were examined in the SEM. The

micrographs in Figure 15 show the typical appear- ance of the polycrystalline alumina surfaces after short tests in distilled water and in the boric acid

solution. In distilled water, the surface has been

subjected to plastic deformation by the diamond

particles, and several deep scratches are evident in Figure 15 (a). In contrast to this observation, the

surface of the sample tested in the boric acid solution in Figure 1 5 (b) shows evidence for intergranular fracture rather than plastic deformation. Although, both surfaces contained regions of plastic

deformation and intergranular fracture, the difference

Figure 13. SEM micrographs of the bottom surfaces

of the holes in polycrystalline alumina: (a) in

distilled water, and (b) in boric acid solution.

108

Figure 14. SEM micrographs of the debris collected I after drilling polycrystalline alumina: (a) from 1. distilled water, and (b) from boric acid solution.
in the two micrographs in Figure 15 are representative of the surfaces in the short tests. Based on the

short tests, it is therefore concluded that considerable

plastic deformation precedes fracture in water; but

the material removal process in the boric acid

solution primarily consists of intergranular fracture. This implies that the increase in drilling rate with

boric acid is related to an interaction between the

boric acid and the grain boundaries in the polycrystalline alumina. This conclusion is partially

supported by the fact that the boric acid solution did not increase the drilling rate of sapphire.

Two short tests were conducted on a sapphire

surface to compare the effect of boric acid solution

with water. Typical drilled surfaces are shown in Figure 16. The sapphire surfaces contain evidence

for plastic deformation, i.e., plowing marks, and

cleavage fracture. The surface appearance for both

cases are similar, which indicates that boric acid had no observable effect on the removal process in sapphire.

Discussion

The results presented in this paper clearly show

that addition of boric acid to distilled water increases

the material removal rate in drilling of polycrystalline alumina. The subsequent SEM examinations of

the drilled surfaces, in short as well as longer tests, suggested that an interaction between boric acid and

the grain boundaries in polycrystalline alumina may

contribute to the material removal process by en- hancing intergranular fracture. The implication of

Figure 15. SEM micrographs of the alumina surfaces

after short drilling tests: (a) in distilled water, and (b)

I in boric acid solution.

109

Figure 16. SEM micrographs of the sapphire surfaces

after short drilling test: (a) in distilled water, and (b)

in boric acid solution.
these observations and other possible interactions

between boric acid and alumina are described in this section.

Several publications have shown that boric acid can be used as a solid lubricant to reduce the friction coefficient of two sliding surfaces [21-23]. In these

experiments either boric acid or boric oxide powder

is used to form a coating on the sliding surfaces. In the latter case, reaction of the boric oxide coating

with the moisture in the air environment forms boric

acid on the surface [22,23]. The observed low

friction coefficient of boric acid was assumed to be

related to its layered triclinic crystal structure with

weak bonding between the layers. During sliding,

these layers are aligned parallel to the direction of

relative motion and, once so aligned, can slide over one another with relative ease to provide a low

coefficient of friction [22,23].

Boric acid is also used as an additive in metal

cutting fluids to reduce friction, increase corrosion

resistance, and control microbial degradation. In these formulations, boric acid is always used in combination with several other compounds, such as amines, fatty acids, and alcohols [2-9]. Recent

machining experiments have shown the addition of

boric acid to distilled water, without the addition of

any other compounds, can reduce the cutting force

by 20 to 50%, depending on the cutting conditions

[24]. In one test, a reduction by as much as 96%

was observed when turning a polycrystalline alumina

using a PCD (polycrystalline diamond] tool insert. Similarly, boric acid was found to be effective in milling of a glass-ceramic dental material using a high-speed steel end mill [25]. The effectiveness of

boric acid in these experiments was used as an

indication that boric acid either deposits on the

surfaces or reacts with the workpiece under the high

loads and temperatures encountered in the cutting zone. By analogy, a similar mechanism was antici- pated in drilling.

To evaluate the possibility of reaction between

alumina and boric acid, we should first examine the

state of boric acid in water. Aqueous solutions of

orthoboric acid H3BO3 are only very slightly acidic

(pH = 6 at 4% wt]. H3BO3 is considered to be a very weak monobasic acid by hydroxide ion binding;

i.e., H3BO3 + 2H2O = B(OH)/ + }i,0\

(The simple orthoborate ion, B03^ , does not exist in aqueous solutions.) [26] The following reaction may

be possible between alumina and boric acid under

the high pressures and high temperatures in the

cutting zone:

AI2O3 + H3BO3 + 3H2O = Al(B(OH)4)3 + A1(0H)3

Recent investigations have shown that aluminum

hydroxide A1(0H)3 when formed by hydrolysis on

alumina sliding surfaces can decrease the friction coefficient and the wear rate of alumina [27-29].

However, aluminum hydroxide can form in either

water or the boric acid solution. Therefore, formation of aluminum borate Al(B(OH)4)3 rather than the

aluminum hydroxide on the alumina surface during

machining may have an influence on the removal

process. Similar to sodium borate, this compound is possibly a surfactant and can be used as a deter- gent/dispersant additive in aqueous solutions [26]. It

is not clear whether aluminum borate can provide

lubrication and a lower friction, but it is expected to be softer than alumina. Therefore, by analogy with

other solid lubricants, it could act as a lubricant.

Although this reduction in friction could explain the

lower cutting forces observed in turning and milling

[24-25], it cannot be used as a direct explanation for the observed increase in the drilling rate. SEM examinations of the drilled surfaces and the

results obtained on sapphire, suggested an interaction

between boric acid and grain boundaries in polycrystalline alumina. The particular polycrystalline

alumina used for this study contains an amorphous

grain boundary phase consisting of a mixture of

SiOj, Na20, TiOj, and K2O, with the first two having

the highest concentration (see Table 1). Studies on

environmentally assisted fracture of oxide glasses

[30,31] have shown that the rate of crack growth is controlled by the chemical activity of the active

species in the environment as well as by the stress intensity at the crack tip. For example, Michalske

and Freiman[32] have described a specific chemical

mechanism by which strained Si-0 bonds in amor- phous silica react with OH" ions in water to increase

the rate of crack growth. In fact. White et al. [33]

have obtained direct evidence showing an increased

crack propagation rate in basic aqueous solutions as the pH is increased. As another example,

Weiderhom and Johnson [34] observed definite

effects of hydrogen ion activity on crack growth in

silica and other glasses, suggesting that OH' and H"^ ions may show a different stress-dependent reaction

rate than HjO. These observations on the environmentally assisted fracture of amorphous oxides that similar effects may occur during machining of

polycrystalline ceramics containing an amorphous

oxide grain boundary phase.

110
Several reactions may be possible between boric

acid and the oxide grain boundary constituents.

Identification of these reactions is beyond the scope

of this paper, and requires a detailed examination.

Nevertheless, reaction between the boric acid and the

grain boundaries in alumina that can increase the rate of crack growth is a plausible explanation for the

observed increase in the material removal rate in drilling and the decrease in cutting force in turning

and milling. In the former machining operations,

enhanced crack growth in the presence of boric acid can decrease the cutting forces since material removal may occur with less plastic deformation.

Another possible chemical effect that might en- hance drilling is the formation of surfactants such as aluminum borate, sodium borate, aluminum hydroxide, and silicon hydroxide, etc., which can be beneficial in the removal of the debris from the cutting zone during drilling. The debris must pass through

a small clearance between the bronze bond and the

exposed diamond particles. If the debris agglomerate

and form particles larger than the clearance of about

10 ^m, then the diamond particles may be pushed

away from the workpiece surface and reduce the

drilling rate. In this respect, another important

function of a drilling fluid is enhanced corrosion of

the metal bond. If bond corrosion can be controlled

to remove a sufficient amount of bond material and

increase the clearance, then drilling could be more

efficient. This effect for boric acid was not examined in the present study, but it needs to be examined in the future.

Conclusions

Based on the test results in drilling of single and

polycrystalline alumina with bronze-bonded diamond

core drills, the following conclusions can be made:

1) Addition of 4% wt boric acid to distilled water

increases the rate of drilling of polycrystalline

alumina by a factor of 2.

2) Boric acid is ineffective in improving the drilling rate of single crystal alumina.

3) The results suggest that interactions between

boric acid and the amorphous oxide grain boundary phase in .polycrystalline alumina promotes

intergranular fracture.

Acknowledgments

This project was conducted under the auspices of

the Ceramic Machining Consortium, which is a joint

research program between NIST and several companies, universities, and government agencies. The

financial support received from the U.S. Navy

Manufacturing Technology Program, is gratefully

acknowledged. The authors are grateful to Eric Ives

for his assistance in designing and fabricating the

drilling setup used in this study.

Information on product names, manufacturers or

suppliers are included in this paper for clarification. This does not imply endorsement of the product or

services by NIST.

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113

CHARACTERIZATION OF MACHINED

SURFACES

115

ULTRASONIC MEASUREMENT OF SURFACE AND

SUBSURFACE STRUCTURE IN CERAMICS

J. A. SLOTWINSKI, N. N. HSU, and G. V. BLESSING

National Institute of Standards and Technology, Gaithersburg, MD

An effort is underway to detect near-surface damage using nondestructive ultrasonic

methods. A sharply focused ultrasonic transducer was used to generate short-duration

pulses in water and to receive echoes from the surface of the submerged sample.

Direct-reflected and mode-converted interface wave echoes were separable in the time

domain echo waveforms. Amplitude scans were taken with software controlled X-Y

scanning stages capable of 0.1 |J.m resolution. The surface topography of ground

siUcon nitride (Si3N4) samples had a prominent affect on the direct-reflected echo

amplitude. It is a principal goal of this program to distinguish this topographical

structure from subsurface damage by studying the amplitude and velocity of various

echo modes. In addition to Si3N4 samples, glass samples have been investigated for

reference purposes.

INTRODUCTION

The intrinsic properties of ceramics make

them prime candidates for many applications

ranging from ball bearings to engine components.

Unfortunately, the grinding processes in

preparing ceramic materials may cause subsurface

cracks which can lead to failure when the material

is under stress. Destructive methods such as modulus of rupture tests yield quantitative

statistical information, such as the average stress at which failure occurs in specific material lots. This, however, yields limited information about

the extent and location of damage. At present

there is simply no nondestructive method for

detecting subsurface damage in a process-control

application. Ultrasonics provides a rapid,

nondestructive method of material defect

detection which may prove applicable to the online detection of subsurface damage in ceramic

materials.

In the pulse-echo ultrasonic technique, the

amplitude of the return echoes provides

information about the sample's surface roughness

[1,2] and possibly the near-surface structure [3].

It is the goal of this work to study the amplitude

behavior of certain ultrasonic echo modes from

ground ceramic samples in order to distinguish

between surface roughness and possible

subsurface damage or structure.

EXPERIMENTATION

Ultrasonic Technique

The general ultrasonic system is shown in

Figure 1. The nucleus of this system is the

ultrasonic transducer. The transducer converts

high-voltage excitations, occurring several

thousand times a second, into short (< 0.1 |J.s)

ultrasonic pulses. These pulses travel through a

fluid coupling, such as water, to the specimen of

interest, whence some of the ultrasonic energy

returns back to the transducer. The returning

117
Monitor

PiMsei^Recelver tiKlGaU

«\

t

Transducer

Sample

I Two-Axis Stage

Figure 1: General ultrasonic system

arrangement

ultrasonic energy is then converted back into an

electrical signal, which is displayed on an

oscilloscope and/or recorded for analysis.

In this work a broadband sharply focused

50 MHz F/0.8 transducer is used to transmit and

receive ultrasonic signals in conjunction with a

pulser/rcceiver unit At its focus this transducer

has a nominal half-amplitude beamwidth of 35 |j.m as determined by pulse-echo measurements while

scanning perpendicular to a razor blade edge.

This is in agreement with pertinent ultrasonic

beam-profile models [4]. In this study, a gated

peak detector acquires specific amplitude values

from the reflected ultrasonic signal as the sample

of interest is scanned. Different wave modes,

each providing possibly distinct information about

the interrogated sample, can be gated separately.

The amplitude data are then read and recorded by a personal computer through an analog to digital

converter (A/D) board

The same software which acquires the

data also controls movement of the sample by

means of X-Y scanning stages. These stages are designed for high-precision positioning, with a

resolution of 0.1 ^.m. The computer code allows

for different scanning rates and step sizes, and is

capable of raster scans. This enables efficient

examination of the samples' entire surface region.

The three ultrasonic arrangements

considered in this experimentation are illustrated

in Figure 2. The first case (a) is that of an

ultrasonic beam focused on the surface of the

target Here the ultrasonic pulses react primarily

to the surface structure, but may also be sensitive to near-surface damage. In the second case (b)

the ultrasound is focused just below the surface

of the target (i.e., defocused), where the

sensitivity to subsurface structure such as cracks

and porosity increases. The final case (c)

involves the generation of a surface wave only.

Here ultrasonic pulses strike the target surface at a critical angle 6^, as determined by Snell's law of

refraction, and propagate along the material

interface as a surface wave. This mode radiates

acoustic energy away fit)m the interface into the

coupling fluid as it propagates and is therefore

WATER/SAMPLE INTERFACE

Figure 2: Sensor/sample orientations: (a)

Normal incidence, surface focus; (b) Normal

incidence, subsiuface focus; and (c) Nonnormal incidence, surface wave.

often referred to as a 'leaky wave." (This

phenomenon is the ultrasonic equivalent of the

optical evanescent wave.) For a 50 MHz carrier

signal this wave has a penetration depth on the

order of 100 ^im in Si3N4.

118
\\\

\\\

\Vv

\\V—

direct .

reflection f/J

/// surfflfy

/// \

/ / / // ///

Figure 3: Ray diagram showing ultrasound

incident on a sample surface: orientation (b)

in Figure 2. Some of the incident energy

enters the material at 6^ where it is converted

into a surface wave. (From reference 5)

When the transducer is focused on

sample surface, only the direct-reflected

longitudinal mode is present. For the defocused

situation, however, both longitudinal and surface

modes may be observed. In this case a portion of

the wave energy is incident on the sample at 6^,

generating a surface wave by mode conversion

(Figure 3) [5].

Often it is difficult to empirically

distinguish between these two signal pulses. To

aid in this, a numerical Green's function

calculation was used for proper mode

identification [6]. For Si3N4 the numraical

results predict a leaky wave that is phase-reversed

compared to the direct reflection. This was

experimentally verified by observing the

separation of the mode-converted and directreflection waves when the transducer was

defocused in a controlled manner. Rgurc 4

shows the time separation between the direct

reflection and the leaky wave for incident

ultrasound on Si3N4. The top trace shows the

direct reflection echo waveform when the sample

surface is located at the transducer's focus. The

three subsequent traces show the echo waveforms

for various defocused positions as the transducer

is moved closer to the sample surface. As the

transducer is increasingly defocused the leaky

wave propagates over a larger surface path,

resulting in increased time separation between it and the direct reflection.

Samples

The ceramics tested were sintered

reaction-bonded silicon nitride (Si3N4) samples,

possessing a density of 3.2 g/cm-^. TTiey were

subjected to varying grinding conditions which

could in turn be expected to yield varying degrees

of subsurface damage. This report focuses on two ground samples which were prepared in highly distinct manners, in an effort to search for

contrasting features. The different grinding

parameters are shown in Table I [7]. The first sample was conventionally ground, followed by one or two spark-out passes. The second sample

was prepared by creep-feed grinding. (These two

samples are hereafter referred to as "lightly" and

"harshly" ground, respectively.)

The different grinding conditions resulted

in different surface roughnesses. The lightly and

harshly ground samples possessed average

surface roughness values (R^ equal to 0.3 and

0.5 ^im ± 0.02 ^im, respectively. (The average

surface roughness value is defined as the average

deviation of the surface contours from a plane

through their median value.) The ground surfaces

had somewhat regular surface periodicities

(grinding groove spacing) of about 40 \x.m.

Figure 4: Echo waveforms (with different

amplitude gain settings) for a transducer at its focus (top), and three defocused positions

showing the increasing time separation between

the direct-reflection (first echo) and leaky wave

(second echo).

119
A separate, lightly ground sample from

the same material batch was polished to a

sufficient depth to remove all topographical

surface features and possible near-surface

structure. TTiis sample had an Rq= 0.04 ± 0.01

jim. It was used as a material reference.

Table I: Grinding Conditions

[ Condition conventional creep-feed

Grinding

Wheel

320 grit

150 cone. 80 grit100 cone. Grinding

Ruid

Emulsion

plus 1-2

spark-out

passes

Emulsion

20:1

Table

Speed

(m/min)

13.l.DZ,

Cross Feed

(mm/pass)

4.00 4.50

Down

Feed

(mm)

0.0051 0.2540

Volume

Removal

Rate

(mm^/min)

310 1737

RESULTS

Numerous ultrasonic echo amplitude

400

i 03 0.76

Figure S: Highly polished silicon nitride echo

amplitude scan.

scans were made on the various samples, with the

transducer beam focused in some cases on the

surface, and in other cases below the surface (i.e., defocused). In all cases the center ray of the

incident ultrasound was normal to the surface.

The same transmitter energy settings and

nominally the same receiver gain were used in the

scans. This allows direct comparison between the

scans. Representative scans are presented first for the ultrasonic beam focused on the sample

surfaces.

Figure 5 is echo amplitude data taken in

10 fim steps over a 1 mm scan length on the

polished Si3N4 sample surface. While this scan was parallel to the original grinding direction, it is representative of all scans parallel and

perpendicular to that direction. The lack of

significant echo amplitude variation is to be

expected for the ultrasonic wavelength of 30 jim

interacting on this polished surface. We also note

that the observed amplitude variation in Figure 5,

with a measured standard deviation of ± 1.2 mV

about the mean value of 382.3 mV, is equivalent

to the system electrical noise level. This noise

level is principally generated by the gated peak

detector output

Figures 6 and 7 are representative scans taken parallel and perpendicular, respectively, to

the grinding direction on the lightly ground

sample. The scan length was again 1 mm, in 10

^m steps. The noticeably large loss of signal,

centered at the 0.2 num position in the parallel scan, is of uncertain origin. Possibilities are an

inclusion or a grinding wheel artifact Two

I Figure 6: Echo amplitude scan on lightly

I ground silicon nitride taken parallel to grinding

grooves.

120
features that we can explain in the parallel scan

are the same baseline system noise as observed in

Rgure 5, and the gentle rolling character of the

baseline attributed to crossing a groove or peak

along the grinding direction. In the perpendicular scan of Figure 7, the mean echo amplitude level is nominally the same as in Figure 6, but with much

greater baseline structure due to crossing the

grooves of variable height

400

1.75

Figure 7: Echo amplitude- scan on lightly

ground silicon nitride taken perpendicular to grinding grooves.

Figure 8, a data scan taken perpendicular

to the grooves of the harshly ground sample,

possesses two features of interest for comparison

to Figure 7. The first is a mean amplitude of

355.4 mV, which is lower than that of the

smoother-surface lighdy ground sample. The

second is a standard deviation of ± 9.9 mV, which

is larger than that of the lighdy ground sample.

These can be explained in terms of the different

Figure 8: Echo amplitude scan on harshly

ground silicon nitride taken perpendicular to

grinding grooves.

surface roughnesses of the two samples. More of

the incident ultrasound is scattered away and/or

extinguished due to phase cancellations in the

grooves for a rough surface than for a smooth

surface [1,2].

Longitudinal scans were also performed

with the ultrasonic beam focused just below the

surface of the ceramic samples, in an attempt to

enhance sensitivity to subsurface structure. The

echo amplitude scans with subsurface focusing at depths of 10, 20, and 30 |J,m were done on the

lightly polished sample in a direction parallel to

the grinding grooves. The recorded scans gave no indication of subsurface structure. As mentioned previously, a leaky surface

mode is also present when the transducer is defocused. Leaky wave scans on the lighdy

ground sample were done in a direction parallel to the grinding grooves. The leaky wave did not

show any signal amplitude variation. This result,

which seems rather innocent at first, is significant

for the following reason: the leaky wave did not

show any variation due to the inevitable crossing

of grinding grooves. (It is virtually impossible to

align the samples such that no grooves are

traversed, due to the beam and groove geometry.)

This suggests that the leaky wave mode is not as

sensitive to surface roughness as the focused

longitudinal wave. A glass slide, ostensibly free of subsurface

damage, was scanned to help verify this reduced

sensitivity to surface roughness. The slide had

smooth and rough (frosted) surface regions with

Ra=0.03 ± 0.01 |im and Ra=3.2 ± 0.5 ^m

respectively, as (ktermined by contacting stylus.

Neither region had a discernible surface

periodicity. There was no leaky wave amplitude

variation on the smooth region and a 10%

amplitude variation on the rough region. This is compared to the case where the ultrasonic beam

was focused on the surface, which resulted in no

amplitude variation on the smooth region, but a 60% amplitude variation on the rough region.

From this we conclude that the leaky wave is less sensitive than the reflected longitudinal wave to

surface roughness. This agrees with otiier results

[8] which show that defocused ultrasound is relatively insensitive to surface roughness.

121
Following ultrasonic testing, cross

sections of the ceramic samples were polished,

revealing comparable subsurface features at the

lightiy and harshly ground surfaces. In cross

section, the ground surfaces exhibited fracturelike structure to a depth of about 10 |im. This was observed at optical magnifications of 250 and

lOOOX.

CONCLUSIONS

A sensitivity of pulsed ultrasound to the

surface topography of various grinding^olishing

conditions has been observed on Si3N4 ceramic

samples. At an ultrasonic carrier fiequency

approaching 50 MHz, corresponding to a wavelength of about 30 (im in a water coupling

medium, wave scattering mechanisms affecting

the echo amplitude of the reflected pulse easily

discriminated amongst the ground surfaces. The

surface roughnesses ranged from 0.04 ^m on

the highly polished sample to about 0.5 ^im on

a harshly ground sample. If the surface grinding

features of these silicon nitrides could be

demonstrably related to their material quality,

ultrasonic wave scattering would provide a means

for effective quality control. Furthermore, this sensor technique, which is both rapid and

nondestructive, can be applied in-process.

An effort was also made to study

subsurface material structure by monitoring the

amplitude of surface waves generated when

defocusing the highly focused F/0.8 transducer.

No distinguishing signatures were observed

between any of the samples that could not be

attributed to surface roughness.

ACKNOWLEDGMENTS

The authors would like to thank S. Jahanmir and T. J. Strakna of the NIST Ceramics

Division for the ceramic samples used in this

research. We also thank them for technical

information regarding the samples, and for

polishing the reference sample. We also

acknowledge R. Polvani and L. Smith of the

NIST Precision Engineering Division for fruitful

discussions and for polishing the sample cross

sections. Finally, we would like to thank L. Marchetti of the NIST Automated Production

Technology Division for developing and

programming the ultrasonic scanning software.

REFERENCES

[1] G.V. Blessing and D.G. Eitzen,

"Ultrasonic Sensor for Measuring Surface

Roughness," Proceedings of the International

Congress on Optical Science and Engineering, pp. 281-289, 1989.

[2] G.V. Blessing, D.G. Eitzen, H.M. Ryan,

and J.A. Slotwinski, "Surface Micrometrology

Using Ultrasound," Proceedings of the 1990

IEEE Ultrasonics Symposium. Vol. 2, pp. 1047-

1052, 1990.

[3] K. Yamanaka, Y. Enomoto, and Y.

Tsuya, "Acoustic Microscopy of Ceramic

Surfaces," IEEE Transactions on Sonics and

Ultrasonics . Vol. SU-32, No. 2, pp. 313-319,

1985.

[4] G.S. Kino, Acoustic Waves: Devices.

&Qa£m& and Analog Signal Processing.

(Prentice-Hall, Inc. NT) p. 185, 1987,

[5] R.S. Gilmorc, K.C. Tam, J.D. Young, and

D.R. Howard, "Acoustic Microscopy from 10 to

100 MHz for Industrial Applications," Phil.

Trans. R. Soc. Lond.. Vol. 320, pp. 215-235,

1986.

[6] N.N. Hsu. S.E. Pick, and G.V. Blessing,

"Green's Function of a Liquid/Solid Interface and

Ultrasonic Materials Testing," Proceedings of the

1992 IEEE Ultrasonics Symposium. Vol. 1, pp.

395-398, 1992.

[7] S. Jahanmir, T.J. Strakna, G.D. Quinn,

R.N. Kopp, S.C. Yoon, and K.V. Kumar "Effect

of Grinding on Strength and Surface Integrity of

Silicon Nitride Part EL," Proceedings of the

International Conference on Machining of

122
Advanced Materials . Gaithersburg, MD, July 20-

22, 1993.

[8] P. Reinholdtsen and B.T. Khuri-Yakub,

"The Effects of Surface Roughness on Subsurface

! Defect Detection Using Acoustic Microscopy,"

Review of Progress in Quantitative NDE. Vol.

I 5A, pp. 485-491, 1985.

123
I

!
SURFACE CHARACTERIZATION OF ULTRASONIC

MACHINED CERAMICS WITH DIAMOND

IMPREGNATED SONOTRODE

H. DAM, S. JENSEN, and P. QUIST

Danish Technological Institute, Taastrup, Denmark

In this article a new ultrasonic machining process is presented. The tool is a diamond impregnated sonotrode, which is split at the end, so that it makes both

lateral and longitudinal vibrations. Machining experiments were performed with

glass, alumina and partially stabilized zirconia (PSZ), Material removal rates were

highest for glass and alumina, at approximately 600 mm-^/min with Rg-values

ranging from 1.6 to 2.1 ^m. The material removal mechanism for glass and alumina

was microfracture. The best surface quality was obtained with PSZ, but material

removal rates were low. The material removal mechanism was controlled by plastic

deformation.

1. INTRODUCTION

One of the major limitations in the use of

ceramics, compared to other materials, is the

relatively high cost of machining. This has

been stated in several reports [1].

To solve these problems there are generally

two ways:

1) - reduce or eliminate the need for

machining by producing near net-shape

components.

2) - develop new machining processes or improve or optimize the already existing methods.

In this paper, a new process for ultrasonic

machining is presented, and the material

removal rates are given. Further the machined

surfaces were characterized to explain the

material removal mechanisms.

2. DESCRIPTION OF PROCESS

Ultrasonic machining was developed during

the 1950's for 2 and 3 dimensional machining

of hard and brittle materials [2].

The conventional ultrasonic machining

process uses an abrasive slurry of water and

boron carbide. During processing the tool

(sonotrode) hammers the abrasive particles

against the workpiece surface; thus, creating a shape corresponding to the geometry of the

sonotrode. The main advantage of the process

is that the machining forces are low and

complex geometries are possible. However,

the material removal rates are rather low and

the tolerances are hard to control. Therefore,

attempts have been made to combine diamond grinding with ultrasonic machining. The

first known experiments with diamond impregnated tools were made by Legge [3,4], and the

process turned out to be versatile regarding

both deep hole drilling and end mill facing.

One of the latest attempts in this development

has been the use of different ultrasonic

125
vibration modes with diamond impregnated

tools. These have been reported by Kunabe

[5], and it was concluded that torsional ultrasonic vibration is suitable for high precision core drilling while longitudinal ultrasonic

vibration is suitable for core drilling with high

material removal rates.

The tool used in these experiments uses a sonotrode which converts the longitudinal

ultrasonic action into a mixed lateral and

longitudinal vibration mode, see Figure 1. The

working principle of the sonotrode is therefore

basically like a tuning fork.

Ultrasonic

Input

Ultrasonic

Output

3. DESCRIPTION OF EQUIPMENT

The equipment used in this series of experiments is described in the following sections.

3.1 Machine tool

The experiments were performed on an ultrasonic milling machine.

The main characteristics for this machine tool are ultrasonic frequency 19.5 - 20.5 KHz,

amplitudes of 10 - 23 /im, ultrasonic power of

2 KW, spindle rotation of 0-5000 rpm, full 4

axis CNC control.

Further a force measurement system using

piezoelectric sensors is built in. This system

was used for measuring the machining forces.

3.2 Sonotrode

The sonotrode is shown on figure 2.1. Its most

important dimensions and specifications are:

Length 130 mm

Top diameter 50 mm

Cutting diameter 24 mm

Diamond impregnation: Electrolytical coating with 181 as mean grain size.

3.3 Materials

The materials were commercial qualities of

glass, alumina (99.7%) and partially stabilised

zirconia (PSZ). The stabilising element in PSZ

was magnesium oxide. The material data are shown in the table below.

Fig. 1 Sonotrode for ultrasonic machining

according to the tuning fork principle

With this process, it is possible to perform a

milling process in hard and brittle materials.

126
Table 1 Material data for the ceramics

used in the experiments

Material Glass AluminaPSZ

Density

(g/cm^)

2,5 3,9 5,8

Flexural

Strength

(MPa)

50 300 500

(MPa/m)

0,7 4,5 9

Hardness

Knoop

(Mpa)

7000 23000 17000

3.3 Analysis Equipment

The surface roughness measurements were

performed using a surface profilometer with a

pick up radius of 5 fim. A complete description of the equipment and the results of the

roughness measurements are available in [6].

For the qualitative characterisation of the sur- faces a Scanning Electron Microscope, SEM

was used.

4. TEST PROCEDURE

Samples with a size of 50 x 24 mm were cut

from three materials by diamond cutting. The

samples were then glued to the base plate

with sealing wax.

Three series of experiments were performed.

1 machining of alumina to find feasible

machining parameters

2 rough machining, where the resultant

machining forces were kept at the

same level for all 3 materials.

3 finish machining, where the feed rates

were at the same level for all 3 materials.

Furthermore, a series of "high feed fi- nishing" were made with alumina and

glass with increased feed.

The general machining parameters are shown

in table 2.

Table 2 General machining parameters

Spindle speed 3000 RPM

Frequency 20,2 kHz

Amphtude (input) 19 yam

Cooling

Internally

through spindle

and externally

directly on

tool.

Coolant Water with

rust inhibitor

5. RESULTS

5.1 Selection of feasible machining parameters

In these experiments the depths of cut (vertical) and the feeds (horisontal) were varied. At

depths of cut at 2 mm there were heavy low

frequency vibrations of the tool. If the feed

was sufficiently low (5 mm/min), the low frequency vibrations disappeared, but instead the

tool made half circular machining tracks as shown on figure 2.

127
Machining tracks

Fig, 2 Principal drawing of machining tracks

appearing at to high feeds and depths

of cut.

Table 3 Machining parameters for rough

machining keeping resultant machining forces at same level

Material Feed

(mm/min)

Depth

of

cut

(mm)

F

(N)

Glass 60 0,5 24

Alumina 50 0,5 24

PSZ 7 0,5 22

With the objective to increase the material

removal rate and avoid both machining tracks

and low frequency vibrations it was found that

the highest material removal rates could be

obtained at relatively low depths of cut < 0.5

mm and high feeds > 50 mm/min for glass

and alumina. For PSZ the feeds had to be

even lower < 7 mm/min to avoid machining

tracks.

5.2 Materia! removal rates and surface

qualities for roughened and finished

samples

The machining parameters for the roughing

and finishing series are shown in Tables 3,4

and 5.

Table 4 Machining parameters for

finishing, keeping feeds at the

same level

Material Feed

(mm/min)

Depth

of

cut

(mm)

F

(N)

Glass 4 0,2 8

Alumina 4 0,2 8

PSZ 4 0,05

The forces in the tables were calculated as the

resultant of the machining forces in x, y, and Table 5 Machining parameters for high

z directions. feed finishing

Material Feed Depth F

(mm/min) of (N)

cut

(mm)

Glass 30 0,2 9,4

Alumina 30 0,2 17,3

128
The resulting material removal rates are shown on figure 3. enced by statistical deviations, because it is taken as a mean of 5 peaks.

Lagwid

I MaL Remover

AluLha Roughing High feed Rnishing

finishing

Fig. 3 Material removal rates

From figure 3 it can be seen that the material

removal rates are very high for both glass and

alumina, while PSZ has a rate 6 -7 times

smaller at the same machining force. A

problable explanation for this is that the

fracture toughness of PSZ is much higher than

alumina and it thus requires a larger amount

of energy to remove material from the surface. Because the machining force and the

spindle speed is kept constant, the only way to

obtain a larger energy input per volume of

removed material, is to slow down the feed.

The roughness measurements were performed

at the bottom surface. The two most common

roughness parameters are Rg and R^. Rg is a measure for the arithmetric mean height of

the surface profile. Let the function f(x)

describe the height of the profile at the position X. Rg can then be found by integrating

f(x) with respect to x over the measuring

length 1^ as shown below.

Ra

Fig. 4

Roughing High feed Rniahing

finishing

Roughness parameters at the

bottom of the machined surfaces.

If the surface roughness parameters are observed it can be seen that they decrease

when feed and depths of cut are lowered. This

was what could be expected. However if, the

roughness parameters for alumina are observed they are generally higher. Furthermore it can be seen that a decrease in cutting

depth and feed rate has much less effect on

the surface quality, than for glass and PSZ.

If the surface roughness are described according to the new ISO standard, which is the

same as DIN 4776 [7], a more detailled picture of the surface can be given.

In this standard three different roughness

parameters are defined; Rp^, R^,^ and R,^. They

are derived from the Abott curve, which

describes the distibution of peaks and valleys

in the roughness profile as shown on fig. 5.

R^ is the mean height of the 5 highest peaks

in the surface profile. It is comparable to R^^ax

(maximum peak height), but it is less influ129
The exact definition of the three parameters

are given in [7], but they can be interpreted as

follows:

R,5 is the depth of the roughness profile when

the highest peaks and the lowest valleys are excluded. It is a measure for the bearing ratio, so that a low means a high bearing ratio.

Rp^ (the reduced peak height) is the height of

the triangle with the shaded area, Aj. From a

tribological point of view it gives a measure

for the "running in" characteristic for the

surface.

R^^ (the reduced valley height) is the height of

the triangle with the shaded area Aj. If the

surface is lubricated, R^,. gives a measure for

the volume of lubricant on the surface.

The roughness parameters are shown in figure

6.

Fig. 6 Roughness parameters according

to DIN 4776 for machined sur- faces.

For alumina it can be seen that the reduced

valley height R^y. is almost the same for all machining rates, while both R,^ and Rp^

reduced in finishing. This indicates that there

is a constant contribution from, for instance,

porosities or grain pull-out. From a tribological point of view a low R^ has a high

priority because it gives a higher bearing ratio.

The best properties concerning friction

coefficient and wear can therefore be

expected with PSZ because it has a lower R^

and Rpk than any of the other materials.

For qualitative characterisation of the

machined surfaces a SEM-investigation was

performed. For PSZ and glass, two representative pictures are shown on figures 7 and 8.

130
Feed

Fig. 7 Glass ultrasonic machined with impregnated diamonds (finish machining).

Feed

EHT- 20.0 KV UD- 21 mm HftG- X 1.00 K PHOTO- 5 R- BSO 20.0Mm I 1 Zr02 < .

•

Fig. 8 PSZ ultrasonic machined with impregnated diamonds (finish machining).

Although the materials are machined under

practically the same conditions and their

and values are almost the same, there is a

large difference in the structure of the surfaces. Glass is characterised by having sharp

features and no machining marks. The PSZ

131
surface consits of flat plateaus with thin marks

primarily perpendicular to the direction of

feed which is indicated by the arrow. From

this it can be concluded that the material

removal mechanism for glass is brittle

fracture. For PSZ the material removal

mechanism is primarily plastic deformation

and cutting, because the single ceramic grains are cut without leaving sharp edges.

6 DISCUSSION

If this new ultrasonic machining process is compared with other ceramic machining

processes the closest "competitor" is rough

grinding. The material removal rates and the

surface qualities are practically at the same

levels. However, because the resultant machining forces are at least one order of magnitude smaller there is a potential that the risk

for machining damage and cracks is lowered.

At the same time the generation of heat

during processing is expected to be smaller.

As a temperature gradient at the surface generally gives rise to tensile residual stresses

the risk for such stresses will be lowered.

The successor for this series of experiments

will therefore include an NDT analysis and a determination of the strength of the machined

samples.

P.t several machines with a rotating ultrasonic

vibrating spindle are commercially available.

With longitudinal vibrations 'these machines

perform well in deep hole drilling (z-axis)

movement. Sonotrodes with lateral vibrations

as the one described in this article perform

well in X and y direction. A combination of

these two machining processes would therefore reveal an interesting alternative for

ceramic prototyping and complicated

geometries.

CONCLUSION

An investigation of a new ultrasonic machining process using a diamond impregnated

sonotrode with both lateral and longitudinal

vibrations has been performed.

The primary material removal mechanism

were based at microfracturing in alumina and

glass, while the material removal in PSZ was

based on plasticity.

The surface qualities obtained are comparable

to rough grinding but machining forces are an

order of magnitude lower. However, the effect

of this on damages in the subsurface zone still needs to be investigated.

In summary ultrasonic machining with a split

sonotrode combined with ultrasonic machining

with diamond impregnated sonotrodes reveals an interesting alternative for future ceramic

machining.

Acknowledgements:

The research project is financed by Nordic

Industrial Fund, the Danish Centre for

Advanced Technical Ceramics (CATC), and

the Tooling and Process Engineering Department at Danish Technological Institute (DTI).

References

[1] S. Jahanmir, L.K.Ives, A.W. Ruff and

M.B. Peterson

"Ceramic Machining: Assesment of Current Practice and Research Needs in the

United States."

NIST-Special Publication 834. June 1992

[2] R. Schmeig

"Die ultraschallbearbeitung von Hochleistungskeramik."

Keramische Zeitschrifft. No 10, P. 515-

27, 1985

[3] P. Legge

"Machining without Abrasive Slurry."

Ultrasonics. London, 4, p. 157-62, July

1966

[4] P. Legge

"Ultrasonic Drilling with a Diamond

Impregnated Probe"

Ultrasonics. London, 4, p. 1, Jan. 1964

132
[5] Junichiro Kunabe, Masahiko Jin, and

Akifumi Nishio

"Precision Drilling of Ceramics by Ultrasonic

Superposition Vibration Cutting."

Journal of the Japan Society of Precision

Engineering

Vol 55, No 2, p. 366-71 1989

[6] Zia Afreedi and Jan Efsen

"Ultrasonic Machining of Ceramic Materials"

(in danish),

Report from Institute of Manufacturing

Engineering. Technical University of Denmark November 1992

[7] Deutsche Industrie Norm 4776 (DIN

4776)

"Measurement of Surface Roughness"

May 1990

133

MEASUREMENT OF RESIDUAL STRESSES IN MACHINED

CERAMICS USING THE INDENTATION TECHNIQUE

Y. AHN, S. CHANDRASEKAR, and T. N. FARRIS

Purdue University, West Lafayette, IN

Machining produces surface residual stresses which significantly influence the strength

and fatigue resistance of ceramic components. As new methods are developed for machining ceramics, a quick and reliable technique for measurement of residual stresses

would be valuable in assessing the viability of the new machining method from a residual stress perspective. The residual stresses on ground and polished (i.e. machined)

surfaces of soda-lime glass, Ni-Zn ferrite, and silicon nitride have been measured using an indentation technique with a Vickers indenter. In this technique, the surface ex- tent of the median/radial cracks produced by the Vickers indenter in machined and in

annealed ceramics are measured. These are then combined with a fracture mechanics

model to estimate the surface residual stresses in the machined ceramics. In order to

determine the validity of the indentation technique for estimating machining residual

stresses, these stresses were also measured using an X-ray diffraction and a deflection method. The residual stress values determined using the indentation technique

in the machined ceramics were found to be reasonably close to those obtained from

the X-ray diff"raction and deflection methods.. Since it is relatively simple and easily

applied, the indentation technique appears to be a potentially promising method for

evaluating; residual stresses in machined ceramics.

1 Introduction

About seventy years ago, Griffith [1, 2] postulated

the existence of microscopic and submicroscopic

flaws (i.e. cracks) to explain the low measured

strength values of glass compared to its theoretical

strength. Since then it has become well estabUshed

that the flaws generated at locahzed contacts between hard rigid particles and solid surfaces are an important cause of strength degradation in brittle

sohds.

One method of reducing the harmful effects of

microscopic flaws is to induce compressive stresses

in the surface. This has been accomphslied to varying degrees by using different techniques: ion ex- change and thermal tempering processes in sodahme glass [3, 4]; chemical strengthening in alumina

and glass ceramics by the formation of low thermal expansion sohd solution surface layers at high

temperatures [5]; and stress-induced phase trans- formation of metastable tetragonal zirconia to the

monochnic phase in partially stabihzed zirconia

(PSZ) [6]. Not only is the load for crack initia- tion considerably increased in the presence of sur- face compression but also if the generated cracks

remain fuUy encased within the surface compressive layers, the strength of a ceramic containing

surface compressive stresses is greater than that of

the annealed ceramic [7, 8].

Recent work has shown that diamond grinding produces compressive residual stresses in ce- ramic surfaces [6, 7, 8]. However, the grinding

process also introduces flaws in the surface and

the strength of a ground ceramic is controlled by

a combination of the residual stress and flaw dis- 135
Table 1: Compressive Residual Stresses (MPa) in Soda-Lime Glass and Ni-Zn Ferrite [12]

Material Measurement Technique

Inden- Deflec- X-ray Optation tion tical

56 ±22

Toughened (Knoop)

Soda-Lime 60±5 85±11 not 85.5

Glass (Vickers) appli82±28 cable

(cone)

Lapped not

Ni-Zn 69.5±15 45±11 56±9 appliFerrite (Vickers) cable

tributions [7]. The potentially wide ranging struc- tural applications for ceramics has highlighted a need for simple measurement techniques which

could be used to estimate the residual stresses and

flaw sizes produced by machining.

While there are a number of techniques avail- able for the measurement of residual stresses, the

indentation technique offers a particularly simple

way of determining residual stresses in brittle materials [9, 10]. In this technique, the indentation

cracking produced by a spherical or pointed (e.g.

Vickers or Knoop) indenter is analyzed. Figure 1

is a schematic of the cracks typically generated by

a pointed indenter in glass. Recently, the residual

stresses in toughened glasses and lapped Ni-Zn fer- rite ceramics have been measured using the pointed

indentation technique [11, 12, 13]. The measure- ments in the glasses [11, 12, 13] were made using a

Vickers, Knoop, or 90° WC conical indenter while

with ferrite, a Vickers indenter was used [12]. In these studies, an independent estimate of the residual stress was also obtained using X-ray diffraction, a deflection method, or an optical photoelastic technique. The indentation. X-ray diffraction,

and deflection methods are described in detail in Section 2. Table 1, from Reference [12], summarizes some of the measurement results. The studies showed that the measured values of the residual stress in glass and Ni-Zn ferrite estimated us- ing the indentation technique compared well with

those obtained using the other methods (see Table 1).

It was decided, therefore, to explore the vahdity

H 2c H

Figure 1: (a) Schematic diagram of a side view

of the cracks generated in soda-Ume glass beneath

a pointed indenter. In (b) are shown the surface

traces of the half-penny median cracks: fc, full- penny median crack; he, half-penny crack; Ic, lat- eral crack; and 2c, the surface extent of the median

crack. In the figure, the shallow radial cracks are not shown. These are usually normal to the in- dented surface but they do not contain the load

axis [11].
of the pointed indentation technique using a Vickers indenter for determining residual stresses in

machined (i.e. ground and polished) ceramic surfaces. The residual stresses in machined surfaces of

hot-pressed silicon nitride (Si3N4), soda-lime glass,

and Ni-Zn ferrite determined using the pointed in- dentation technique were found to be quite close to the values obtained using other techniques such

as X-ray diffraction or a deflection method. These

results and the techniques used to obtain them are the subject of this paper.

2 Background: Residual Stress

Measurement

It is appropriate to begin with a review of the

residual stress measurement techniques used in the

present investigation.

2.1 Pointed Indentation Technique

Recently, Chaudhri and Phillips [11] and Chandrasekar and Chaudhri [12] have compared the in- dentation cracking patterns of annealed and thermally tempered soda-lime glass using pointed (i.e. Vickers, Knoop, and a 90° cone) indenters and

have estimated the residual compressive surface

stresses. In the case of Vickers indentation-induced

radial/median cracking, they estimated the residual surface stresses using the following relationship

X

-XP" - P) (1)

1.16c2

where P* and P are the loads which produce cracks

of the same radius c in the tempered glass containing a compressive residual stress of magnitude Gc

and in the annealed glass, respectively, and x is an

indenter constant. Equation (1) is derived using

linear elastic fracture mechanics. For the Vickers

indenter the value of x can be obtained from the

analysis of Lawn and FuUer [14] as

^ " 7r3/2tan680

Alternatively, an empirical estimate for x can be

obtained from a calibration test as follows

Using dimensional energy arguments, it can be

shown that in a material which is free of residual

stress an equation of the form

P

is valid for cracks formed during loading with

pointed indenters [15]. The value of x can now

be determined if Ki^ for the material is known

from an independent fracture toughness test such

as the single-edge notched beam (SENB) test and

the variation of crack radius, c, with indenter load,

P, is measured for the specific indenter and material in an indentation experiment. Previously, both

a calibrated value for x [12] and an analytical esti- mate for X (for a Vickers indenter) obtained using

Equation (2) [11] have been used for residual stress measurement.

It should be noted that the indentation technique provides measurement of the local residual

stress, that is, the stress over a small volume of size on the order of the indentation volume. Also, in the derivation of Equation (1), the residual stress

is assumed to be constant over this volume.

2.2 Deflection Method

The deflection or curvature method [16] enables

the measurement of the average in-plane residual

stresses in a solid surface and their variation within

the subsurface. The method employs a rectangular

parallelepiped specimen which contains the residual stresses to be estimated. One face of the specimen, viz. the reference surface, is polished flat and

its curvature is measured (at this stage it should

be close to zero). Then a layer of uniform thick- ness is removed from the face opposite to the ref- erence face; we denote this as the test face. The

test face contains residual stresses which are to be

determined. The layer is removed by a process

(e.g. chemical etching, electropolishing, fine polishing with soft abrasives) that does not in itself induce any significant residual stresses. The re- moval of this layer causes a change in the principal

curvatures of the reference face. These are then

determined by profilometry of the reference face,

by means of a strain-gauge rosette mounted on the

reference face, or by optical interferometry on the

reference face. The residual stress in the test face

is then given by [17]

E dCi dC2

+ v

dz dz

r.3/2 (3)

137

6(1-1/2)

+ Az[C^{z) + vC2{z)]

+ 2{wo - 2z)[Ci{wo) uC2{Wo)]

/wo "I [{C,{z)^-uC2{z)]dz\ (4)
where the test surface is at z=0, ai{z,wo) is a

principal stress at depth z below the surface, wq

is the initial plate thickness before etching, Ci{z)

and C2{z) are the principal curvatures of the refer- ence face after a depth z of the material has been

removed from the test face, Ci(wo) and €2(^0)

are the initial curvatures of the reference face before any etching, and E and w are, respectively, the

Young's modulus and Poisson's ratio of the specimen material.

The remaining principal stress ^T2(^,wo) can be

obtained by interchanging the subscripts 1 and 2

in Equation (4). A further layer is then removed

from the test surface and the residual stresses determined again as before. In this way, the residual stress distribution in the specimen is obtained.

The measured residual stress values are average

values of the stresses in planes parallel to the specimen surface.

2.3 X-ray Diffraction

In ceramic polycrystals, residual stresses can be

determined using an X-ray diffraction technique in

which the change in lattice spacing resulting from

residual stresses is measured. Comprehensive re- views of the technique are available, see for example Hilley et. al. [18] and James and Cohen [16]. The technique requires a polycrystaUine sample so that diffraction from a given set of (hkl) planes can be studied. The presence of residual stresses

in the specimen causes the diffraction angle corre- sponding to the above planes to be changed. This

is a consequence of the change in the lattice spacing of the (hkl) planes due to the residual stresses. From the measured change in the lattice spacing,

the amount of strain is calculated. Using isotropic

elasticity and knowing the X-ray elastic constant,

E/{1 -f f), for the direction normal to the (hkl)

plane (in cubic crystals this would be the [hkl]

direction), the residual stress can be determined

from the strains thus measured. The X-ray elastic constant can be calculated from a knowledge of

single crystal compliance constants [19] or obtained

from direct calibration measurements of the lattice

strain under an applied stress. If the X-ray beam

diameter is only a few millimeters across, pointto-point determination of surface residual stresses can be made in specimens of size greater than

~10 mm. Furthermore, by removing surface layers, as in the case of the deflection method, the

Table 2: Material Properties

Property Soda-Lime

Glass

Ni-Zn

Ferrite

Si3N4

E(GPa) 70 191 300

0.25 0.2 0.26

^(gm/cc) 2.5 5.3 3.22

H(GPa) 5.5 7.3 16.7

A%(MPav^) 0.75 0.8 4.0

sub-surface residual stress distribution can be obtained [6].

3 Experimental Procedure '

Three blocks each of soda-lime glass, Ni-Zn ferrite, and HP-Si3N4 (NC-132, Norton Co.) ceramics were prepared by grinding and polishing. Table 2 lists some relevant mechanical properties of

the specimen materials which will henceforth be

referred to as ceramics. Two of the blocks were of

size 25 X 12.5 x 2.5 mm while the third block was ~25 X 12.5 X 12.5 mm. The 25 x 12.5 mm faces on each of the blocks of ferrite and Si3N4 were ma- j chined by first grinding with a 600 grit size resinbonded, diamond wheel and subsequently polished

with 3 fim and 0.5 diamond paste. The grinding depth of cut was 5 //m and the wheel surface

speed was 36 m/sec. In the case of the soda-lime I glass blocks, the polishing was carried out using a

6 //m cerium oxide powder. In each case the deformed layer produced by grinding was not completely removed by the polishing process. Care was

taken to ensure that aU the three blocks of a given

material were finished under nominally identical

conditions. However, no consideration was given; to ensuring that the specimen preparation proce-j

dures conformed to those used in industry. This

latter aspect does not detract from the objectives

of the study.

The residual stresses produced by machining i

(i.e. grinding and poUshing) on the 25 X 12.5

mm face of the blocks were measured using the

pointed indentation. X-ray diffraction (only with

ferrite and Si3N4), and deflection methods. The

pointed indentation technique was used with the 25

X 12.5 X 12.5 mm blocks while the X-ray diffrac- tion and deflection methods were used to determine the residual stresses in the 25 X 12.5 x 2.5
mm blocks.

3.1 Pointed Indentation Technique

The ground and polished (25 x 12.5 mm) faces of

the soda-lime glass, Ni-Zn ferrite, and HP-Si3N4

blocks contain residual stresses induced by machining. The machined faces were indented with a

Vickers indenter on a conventional micro-hardness

testing machine. The indentation loads were typically between 10 and 300 Newtons. In a typical

test, the maximum indentation load was applied in

15 to 25 seconds and was maintained for a duration of 15 seconds before unloading. The blocks were examined using optical and scanning elec- tron microscopy after unloading to determine the

length 2c of the surface trace of the median/radial

cracks (see Figure 1) and the indentation diagonal

2a. Figure 2 is an optical micrograph showing the

surface traces of the typical median/radial cracks

in annealed Si3N4 produced by a Vickers indenter.

The lengths AA and BB of the cracks in Figure 2

correspond to the length 2c of Figure 1, while the

length CC in Figure 2 corresponds to the indentation diagonal (2a) of Figure 1. The sub-surface

shape and extent of the cracks in Si3N4 (Figure 2) were not determined. Figure 3 shows a side view

of the typical median/radial crack system in an- nealed soda-lime glass produced by a Knoop in- denter. The photograph was taken in transmitted

light. Here the length AA (which corresponds to the length 2c of Figure 1), is actually the diameter

of a half-penny median crack and not the surface

trace of of shallow radial cracks.

After several indentations had been made in

each of the blocks and the various measurements

(i.e. indentation diagonal and crack length) completed, the ferrite and silicon nitride blocks were

annealed in an inert atmosphere of argon at 800°C

for one and one-half hours to relieve the machining

induced residual stresses. In the case of the sodalime glass block, the annealing was carried out in

air at 520''C for two hours. After annealing, the

Si3N4 blocks were lightly polished using 0.5 //m

diamond paste while the soda-lime glass and fer- rite blocks were lightly polished using cerium oxide

and magnesium oxide powders, respectively. Such

polishing procedures have been found to induce

a negligible amount of residual stress. The polished surfaces of the annealed blocks were again

indented with a Vickers indenter at several different points using a range of indenter loads and the

surface traces of the cracks were measured. The

crack lengths thus obtained corresponded to sur- faces which were nominally devoid of any residual stresses. The difference in the crack lengths

measured on the ground and polished (machined)

specimen surfaces, and the annealed specimen sur- faces (free of residual stress) were used with Equation (1) to obtain the residual stress induced by

the machining (i.e. grinding and polishing) in the

ceramics.

3.2 Deflection and X-ray Diffraction

Methods

Measurements of the residual stresses were carried

out in the machined Ni-Zn ferrite and silicon ni- tride blocks using the deflection and X-ray diffraction methods. With the machined soda-lime glass

blocks, the measurements were made using the cur- vature method alone.

In the deflection method, the principal curvatures of the reference face were determined by

using a Talysurf 10 (Rank Taylor Hobson Inc.)

profilometer to trace the reference face and by

measuring the chordal rise over a fixed length.

The principal curvatures of the reference face were

also measured after layers of depth 4 fj.m and

8 /im were removed from the test face by polishing. Mechanochemical polishing with a 7 volume

percent solution of orthophosphoric acid in water

and 0.05 fim iron oxide abrasive powder was used

for layer removal in ferrite and HF-Si3N4 while a

0.05 ^m magnesium oxide powder was the polishing abrasive used on soda-lime glass. It is important that this layer removal process does not in

itself introduce substantial residual stresses compared to the stresses being measured; past unpublished results have shown that each of the polishing agents used above induces less than 8 MPa

surface residual stress in the soda-lime glass and

the Ni-Zn ferrite. From the measured curvatures

and the values of the elastic constants for sodalime glass, Ni-Zn ferrite, and Si3N4 (Table 2), the

residual stresses in the machined surface and at a depth of 4 /im from the machined surface were

obtained using Equation (4). A piecewise linear

fit of the curvature versus depth data was used in

the calculation of the derivatives and integrals in

Equation (4).
Figure 2: Optical micrograph of surface trace of a median/radial crack in HP-Si3N4 produced i'^ Vick^ is indenter (Load — 30 kg).

\

Figure 3: Photograph of a half-penny median crack in soda-lime glass produced by a pointed indenter. The

photograph was taken in transmitted light through the sides of the glass block whilst the load is applied.

140
X-ray diffraction was carried out using Crradiation (A = 0.2291 Tjm) with a spot diameter of

3 mm as it has a much smaller penetration depth

than does Cu radiation. The (119) and (323) peaks were selected for the observations in Ni-Zn ferrite and Si3N4 respectively. The diffraction angle 0 was measured to an accuracy of 0.0125 degrees using a diffractometer and the residual stress was determined by the method described in Section 2.3. By translating the ceramic specimens, the surface

residual stress was determined at three points in

the ferrite and Si3N4 specimens. Similarly, the

residual stress at a depth of 4/im from the sur- face was measured in both of the ceramics after

removing surface layers by mechanochemical polishing (as in the deflection method).

Calibration experiments were conducted to determine the X-ray elastic constants, £'/(l-f i/), cor- responding to the (119) and (323) planes in NiZn ferrite and Si3N4 respectively. Details of the

calibration method can be found in Chandrasekar

and Chaudhri [12]. The measurements gave val- ues of 150±11 GPa, (119) plane in Ni-Zn ferrite,

and 268±14 GPa, (323) plane Si3N4, for the X-ray

elastic constants.

4 Residual Stress Results

4.1 Indentation Technique

The residual stresses were obtained from indentation crack length measurements as foUows. Fig- ures 4, 5, and 6 show the variation of c'^^^ with

load, P, for annealed soda-lime glass, Ni-Zn fer- rite, and Si3N4, respectively. Here c is the radius

of the surface trace of the median/radial crack (see

Figure 1) produced by the Vickers indenter. The

solid lines in Figures 4, 5, and 6 are a least-squares

fit of the experimental data assuming that the re- lationship between P and c^/^ is linear [14]. From

the slopes of these straight lines and the values of

for the three materials (Table 2) the indenter

constant (x) was obtained using Equation (3) for each of the three materials. The values of x for soda-lime glass, Ni-Zn ferrite, and Si3N4 thus esti- mated from the calibration experiments are listed

in Table 3. The calibrated values of x for all three

ceramics (Table 3) are within 20 percent of the

analytical value for x (~ 0.0725) estimated using

Equation (2) for the Vickers indenter. The cali

Table 3: Calibrated Values for the Indenter Constant x

Material X

Soda-lime glass

Ni-Zn ferrite

Silicon nitride

0.0695

0.085

0.071

200

141

0.00 0.10 0.20 0.30 0.40 0.50 0.60

Crack Length, c'''(mm''')

Figure 4: The variation of crack length with load

in annealed soda-lime glass (maximum error bar is shown)

brated values of x are used in the estimation of

the residual stresses. Figures 7, 8, and 9 show the variation of (P* — P)

with for soda-lime glass, Ni-Zn ferrite, and

Si3N4, respectively. The figures were obtained by

combining the measurements of the crack lengths

shown in Figures 4, 5, and 6 with similar measurements carried out on the machined surfaces

of soda-lime glass, Ni-Zn ferrite, and Si3N4. The

straight lines in Figures 7, 8, and 9 are a least squares fit of the data; that the data should be on

a straight line is predicted analytically by Equation (1). The coefficients of regression for this

least-square fit are given in Table 4. Using the

slope (P* — P)/c^ of the straight lines in Fig- ures 7, 8, and 9 (see Table 4) in Equation (1) to- gether with the appropriate value of x for each of

the three materials from Table 3, we obtain the

surface compressive residual stress (<7c) in the machined surfaces of ceramics. The residual stress

(cTc) values are listed in Table 4. These residual

stresses are caused by the localized plastic defor-
400

Figure 5: The variation of crack length with load

in annealed Ni-Zn ferrite

600

0.00 0.10 0.20 0.30

Crack Length, c"'(mm''')

Figure 6: The variation of crack length with load

in annealed SiaN^ (maximum error bar is shown)

142

Table 4: Surface Residual Compressive Stress (ctc) by the Pointed Indentation Technique

Material

Slope

p*-p

(MPa)

Regression

Coefficient (MPa]

Soda-lime

glass

464.7 0.96 27.8

Ni-Zn

ferrite

1170.0 0.93 85.6

Si3N4 1951.3 0.99 119.1

I ' ' '

I I

0.00 0.10 0.20

Crack Length, c^ (nnm^)

Figure 7: The variation oi P* — P with in machined soda- lime glass {P*, P, and c are defined near Equation 1)

mation of the ceramic materials during grinding

and pohshing.

4.2 Deflection Method

The measurements of curvature showed that the

reference surfaces of the machined ceramics and

glass were ellipsoids with the two principal curvatures differing by less than 8 percent. This cor- responds to a state of equiaxial stress in planes

parallel to the surface. Here equiaxiaUty of residual stress in a plane implies that the stress is equal

in ail directions in the plane since these planes are

principal planes due to the symmetric geometry of

grinding and polishing.

The surface compressive residual stress in the

machined soda-lime glass block was determined as 21.0 ± 8 MPa; it decreased to a value of 11.0 ± 7.5
Table 5: Summary of Residual Stress Measurements in Machined Surfaces (MPa)

0' ' ' ' '

I

' ' '

I

' ' ' ' I ' ' ' ' ' 0.000 0.050 0.100 0.150 0.200

Crack Length, (mm^)

Figure 8: The variation oi P* — P with in machined Ni-Zn ferrite {P* , P, and c are defined near Equation 1)

T

0.00 0.01 0.02 0.03 0.04 0.05 0.06

Crack Length, c^ (mm^)

Figure 9: The variation of P* - P with in machined Si3N4 (P*, P, and c are defined near Equation 1)

/

Material

Technique Soda-Lime

Glass

Ni-Zn

Ferrite

Si3N4

Indentation 27.8 85.6 119.1

(surface)

Deflection 21±8 62±18 87±22

(surface)

Deflection 11±7.5 35±14.5 59±13

(subsurface)

X-ray 68±16 82±24

(surface)

X-ray 43.0±14.5 49±17.0

(subsurface)

AU residual stresses are compressive.

The subsurface values were measured at a depth

of 4 /im below the machined surface.

MPa at a distance of 4 micrometers from the surface, see Table 5. The decreases in the residual stress values were even greater with the NiZn ferrite and Si3N4. In the case of the Ni-Zn

ferrite, the surface residual compressive stress was 62.0±18 MPa while at a depth of 4 fim below the

surface the compressive residual stress had a value

of 35±14.5 MPa. The corresponding values for the

compressive residual stress in the Si3N4 block were 87.0±22 MPa at the surface and 59.0±13 MPa at a depth of A/im below the surface. The residual

stress values are summarized in Table 5.

4.3 X-ray Diffraction

X-ray diffraction measurements of the residual

stress were carried out in the machined Ni-Zn fer- rite and Si3N4 blocks. The measured values of the

residual stresses at the surface and at a depth of

4^m below the surface are given in Table 5. It is seen from Table 5 that these stress values are close to those obtained using the deflection method.

5 Discussion

The X-ray diffraction and deflection measurements

of the residual stress have shown that in all of the

machined ceramics, the surface residual stress is compressive (Table 5). The magnitude of the sur-
face residual stress produced by grinding and polishing is seen to be greatest in Si3N4 and smallest

in glass. The measured surface residual stress of

~90 MPa in Si3N4 compares well with values of

110-140 MPa reported recently for this stress in polished Si3N4 [20]. An examination of the subsurface residual stress values in Table 5 shows that

the stress gradients near the machined surface are quite steep, typically of the order of 3-8 MPa per

micrometer.

It is interesting to compare the surface residual

stress values in the machined ceramics estimated

using the indentation technique with those obtained using deflection and X-ray diffraction methods. From Table 5 it can be seen that with all three

ceramics, the indentation technique gives surface

residual stress values which are reasonably close

(within 30 percent) to those obtained using the

X-ray diffraction and deflection methods. Both

the indentation technique and the X-ray diffraction

method estimate the localized residual stresses,

while the deflection method estimates the average

in-plane residual stress in planes parallel to the

specimen surface.

It is now worthwhile to consider some of the rea- sons for the differences in the residual stress values

determined using the indentation technique and

the X-ray diffraction and deflection methods in Table 5. The analysis leading to Equations (1-3) used

with the indentation technique [9, 10, 14, 15] as- sumes that penny/half-penny median cracks form

during loading of the specimen by the indenter.

Also, it is assumed that residual stress gradients near the indented surface are not steep. The

X-ray diffraction and deflection method measure- ments have shown that with all of the machined

ceramics, the residual stresses vary rapidly near

the machined surface. There are also uncertainties pertaining to the formation and the shape

of the cracks emanating from the corners of the

Vickers-indentations (henceforth these cracks will be referred to as corner cracks). From in-situ observations of Vickers-indentation cracking, Chandrasekar and Chaudhri [12] have shown that in

soda-lime glass, the corner cracks are surface traces of half-penny median cracks. This has also been

confirmed by serial sectioning of the Vickers in- dents along planes parallel to the indented sur- faces (unpublished work). Similar serial sectioning

of the Vickers indents in Ni-Zn ferrite has shown

that for indentation loads exceeding 40 N, the cor- ner cracks in ferrite are of the half-penny, median

type. The shapes of the corner cracks in Si3N4

and in ferrite (at loads less than 40 A'^) are as yet

unresolved.

As to when the various cracks form in the difi"er- ent ceramics, median cracks have been observed to form around Vickers indentations during unloading in annealed soda-lime glass when the maximum

indentation load was less than 30-40 N [12]. In these cases there was no cracking during loading.

When the maximum indentation load exceeded 40

A, median cracks were observed to form during

loading [12]. It is quite possible that, given the

magnitude of the indentation loads used in this

study (e.g. Figure 4), some of the median cracks

formed during unloading in the soda-lime glass. It

is not yet established as to when median/radial

cracks form in Si3N4 and Ni-Zn ferrite polycrystals. However, Lankford [21, 22] has observed from

acoustic emission measurements that shallow sur- face radial (not median) cracks form during the

loading of a Vickers indenter in polycrystaUine SiC

and AI2O3 at loads as small as 0.1-0.25 A. This

could probably occur with Si3N4 and ferrite as weU.

Another point to be noted is that in the derivation

of Equations (1-3), it is assumed that only one median crack is formed. The presence of more that one median/radial crack as is commonly observed

with Vickers indentations could have an effect on

the stress-intensity factors at the median/radial

crack tips due to the interaction between various

cracks [23]. Based on the above observations, it would be un- reasonable to expect significantly better agreement

between the residual stress values determined us- ing the indentation technique and those measured

using the X-ray diffraction and deflection methods.

However, since the pointed indentation technique

is simple and can be rapidly applied, it appears to be an attractive technique for measuring residual

stresses produced by machining in brittle materials.

6 Conclusion

The residual stresses in machined soda-lime glass,

Ni-Zn ferrite, and Si3N4 are measured using the

pointed indentation technique with a Vickers in- denter. For comparison, measurements of the
residual stresses in these machined ceramics have

also been carried out using the deflection method

and an X-ray diflfraction technique (only with fer- rite and Si3N4). The measured values of the residual stress estimated using the pointed indentation technique were found to be reasonably close

(within 30 percent) to the values obtained from the

deflection and X-ray diffraction measurements.

The absence of a detailed model for the indentation crack systems, the existence of steep residual

stress gradients near the machined surface and uncertainties in the value of the indenter constant

X (Equation 2) prevent the indentation technique

measurements from agreeing more closely with the

residual stress values determined using the X-ray

diffraction and deflection methods. However, due

to the simplicity and ease of apphcation of the in- dentation technique, it appears to be a potentially

useful method for determining residual stresses in machined ceramic surfaces.

Acknowledgment The research was supported in

part by the National Science Foundation through

grants DDM 9057916, Bruce Kramer, Program Di- rector and MSS 9057082, Jorn Larsen-Basse, ProI gram Director.

References

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163-198, 1921.

I [2] A.A. Griffith, "The Theory of Rupture,"

I In Proceedings of the First International

' Congress of Applied Mechanics, pp 163-198,

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[3] O.S. Narayanswamy and R. Gardon, "Calculation of Residual Stresses in Glass," Journal

of the American Ceramic Society, 52, pp 554-

558, 1969.

' [4] M.J.C. Hill and l.W. Donald, "Stress Profile Characteristics and Mechanical Behavior

of Chemically Strengthened Lithium Magnesium Aluminosilicate Glasses," Glass Tech,

30 (4) pp 123-127, 1989.

[5] H.P. Kirchner, R.M. Gruver, and R.E.

Walker, "Chemical Strengthening of Poly145

i

crystalline Alumina," Journal of the American Ceramic Society, 51 (5) pp 251, 1968.

[6] D.J. Green, F.F. Lange, and M.R. James,

"Factors Influencing Residual Surface Stresses due to a Stress-Induced Phase Transformation," Journal of the American Ceramic Society, 66 (9) pp 623-629, 1983.

[7] R. Samuel, S. Chandrasekar, T.N. Farris, and

R.H. Licht, "The Effect of Residual Stresses on the Fracture of Ground Ceramics," Journal

of the American Ceramic Society, 72 (10) pp

1960-1966, 1989.

[8] D.P. Stinton, "Assessment of the State of

the Art in Machining and Surface Preparation

of Ceramics," Technical Report ORNL/TM10791, Oak Ridge National Laboratory, Oak

Ridge, TN, 1977.

[9] M.V. Swain, J.T. Hagan, and J.E. Field, "Determination of the Surface Residual Stresses

in Tempered Glasses by Indentation Fracture

Mechanics," Journal of Materials Science Letters, 12, pp 1914-1917, 1977.

[10] D.B. Marshall and B.R. Lawn, "An Indentation Technique for Measuring Stresses in Tempered Glass Surfaces," Journal of the American Ceramic Society, 60 (1-2) pp 86-87, 1977.

[11] M.M. Chaudhri and M.A. Phillips, "Quasistatic Indentation Cracking of Thermally

Tempered Soda-Lime Glass with Spherical

and Vickers Indenters," Philosophical Magazine A, 62 (1) pp 1-27, 1990.

[12] S. Chandrasekar and M.M. Chaudhri, "Indentation Cracking in Soda-lime Glass and Ni-Zn

Ferrite under Knoop and Conical Indenters

and Residual Stress Measurements," Philosophical Magazine A, 1993. In-Press.

[13] R. Tandon and D.J. Green, "Indentation Behavior of Ion-exchanged Glasses," Journal of

the American Ceramic Society, 73 (4) pp 970-

977, 1990.

[14] B.R. Lawn and E.R. Fuller, "EquiUbrium

Penny-like Cracks in Indentation Fracture,"

Journal of Materials Science, 10, pp 2016-

2024, 1975.
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[16] M.R. James and J.B. Cohen, Treatise on Materials Science and Technology, pp 1-62. Academic Press, New York, 1980. Ed. H. Herman.

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R.E. Ricklefs, SAE Report J7849, Society of

Automotive Engineers, Warrendale, PA, 1971.

[19] J.F. Nye, Physical Properties of Crystals,

Clarendon Press, Oxford, 1957.

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S. Kondau, "X-ray Residual Stress Measurement of Sintered Si3N4," Journal of the Society of Materials Science (Japan), 36 (407) pp

810-816, 1987.

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1182, 1981.

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(5) pp 392-396, 1986.

1
DETECTION OF SUBSURFACE DEFECTS IN MACHINED

SILICON NITRIDE CERAMICS BY OPTICAL

SCATTERING METHODS

Silicon nitride (Si3N4) materials from two manufacturers, in the form of flexure

test bars with surface finishes that were ground and polished (.03 |i.m Ra) and

unidirectionally ground (.07 |a.m Ra) and contained simulated subsurface

defects were studied by optical Fourier scattering methods. Subsurface defects

were obtained by either ultrasonically drilling 750-|jjn-diameter holes to within

75-105 |xm of the surface or by grinding slots parallel to and 75- 105 jim below

the surface. Because of the unique optical transmission properties of many

ceramic materials, including those in this study, subsurface defects, such as those likely to be caused by machining-induced damage, may be detected by

appropriate optical methods. In this study s-polarized He-Ne laser light {X = 0.6328 |mi) was used as the light source. Specular reflections were detected by

an optical Fourier method that incorporates cross-polarization. Analysis of the

scattered laser light was accomplished with a CCD-array camera, an eight-bit

A/D converter, and a PC with locally written software. Results show that

subsurface defects simulated in this study can be detected and that, for the

surface finishes studied, changes in signal-to-noise ratio are observable.

W. A. ELLINGSON, D. M. AYAZ, M. P. BRADA, and W. O'CONNELL

Argonne National Laboratory, Argonne, IL

INTRODUCTION

The life of elements that are subject to

rolling-contact fatigue, such as bearings,

can be significantly shortened if surface or subsurface defects are present. Ceramic

rolling elements, especially ball bearings,

are rapidly being developed for commercial

applications [1-4]. Currently, because of

their mechanical and physical properties,

silicon nitride (Si3N4) ceramics are the

materials of choice and offer several

advantages over steel counterparts. These

advantages include greater stiffness,

operating speed, corrosion resistance, and

thermal stabflity [4]. Further, the primary

failure mode of Si3N4 bearings is

spallation, the least harmful mode of

failure [5-6].

include a ceramic rolling element and steel races or (2) all-ceramic bearings, which

have ceramic rolling elements and ceramic

races. The machining operations are

typically different for each ceramic

component [7,8]. The bearings are usually

machined from hot-pressed "blanks,"

which are then hot-isostatically pressed

(HIP'd) and flnaUy lapped to finishes on the

order of <.05 |xm Ra. The races on the

other hand are often ground with surface

finishes <.l - .2 p,m Ra. Either of these

machining operations has the potential to

introduce surface and subsurface defects.

However, subsurface defects could also be

inherently present in the material and not

affect machining but still impact fatigue

life. Nondestructive testing methods,

which can be reliably applied to map

surface and near-surface defects would be

significantly beneficial, primarily if the

data generated could be correlated to life

Ceramic bearings are usually classified

into two types: (1) hybrid bearings that

147
prediction models. One nondestructive

characterization method that has been

extensively studied for metal surfaces

(typically optical surfaces) is optical

scattering [9-12]. Optical scattering occurs

whenever a beam of light is incident upon

a surface. Fundamental geometric optics

predicts that, for a perfect reflector, the

angle of reflection will be equal to the angle

of incidence. However, if the surface on

which the light is incident is transparent

or semi-transparent, a component of the

beam will be internally transmitted (except

at the critical angle). Certain Si3N4

ceramics are partially transparent at

selected wavelengths. Figure 1 shows a

schematic diagram of an incident optical

beam on a partially transparent material.

A reflected beam, a transmitted optical

beam, scattered light, and subsurface

optical scatterers are present. However,

many features on a scattering surface,

particularly a machined ceramic surface,

can cause scatter. These are shown

schematically in Fig. 2, which shows a

range of surface roughness, surface

particulate, cutting-oil residues, and

subsurface defects of different shape,

orientation, and location.

INCIDENT BEAM • REFLECTED BEAM

Fig. 1 . Schematic diagram of the

optical scattering components

for a semi-transparent polycrystalline ceramic.

Incident Light

Fig. 2. Schematic diagram showing

different sources of optical

scatter from a semi-transparent

or transparent material,

including a range of surface

roughness.

It is important that we develop nondestructive methods that can discriminate

between surface effects and subsurface I defects. One such method, the method

chosen for this study, involves the use of

optical scattering and selected optical

insertion and detector systems: a polarized

incident laser light and a polarization i analyzer are present in the detector. This j is often referred to in the literature as a ij cross-polarization method. Figure 3 shows | a schematic diagram relating the polarized I direction of the laser light to the surface of

a material, as well as incident and i

reflected beams. The material surface

under study is in the x-y plane. For

detection of subsurface defects surface

scatter can be considered the noise and

subsurface scatter the signal.

To improve the signal-to-noise (S/N) rati' when detecting subsurface defects, orjl

must either remove the surface effects c enhance the scatter caused by surfac

topography by removing subsurface scatt<

eff'ects. This is accomplished by selectir

appropriate source and detectc

polarization. In addition, the numeric;

aperture used in the detector system 1 important, because the reflected-light an

scattered-light vectors are slightly differei

(see Fig. 3).

148
Fig, 3. Schematic diagram showing

polarization direction relative

to the material surface.

To understand scattered light, several

detector schemes have been studied by

previous researchers [13]. Total integrated

scatter (TIS) has usually been with a

He-Ne laser, optical components to filter

the beam, a detector to collect the total

scattered light over 2 n steradians. and two

detectors specifically to detect specularly

and directly scattered light. Figure 4

shows a schematic diagram of a typical TIS

system. Note that no polariza-tion

analyzers are in place. TIS is related to

surface finish [11] as follows:

TIS =

where

Vs

Vs + As

DETECTOR

(SCATTERED

(reflectbd^'eam) /S.'^R!;^TM'^

DETECTOR

(INCIDENT BEAM)

SAMPLE

HEMISPHERICAL COLLECTING MIRROR

Fig. 4. Schematic diagram of the basic

components of an instrument for

measuring total integrated scatter.

Another method, angle-resolved scattering

(ATiS). usually includes a He-Ne laser, but

incorporates optical components in the

input optics that provide a way to set

polarization, and an analyzing polarizer in

the detector, which is mounted on a

rotating stage to detect angular-dependent

scatter. Figure 5 shows a schematic

diagram of the basic components of an

ARS instrument. According to Bennett

and Matteson [11], little data are available

to relate ARS to surface roughness.

CHOPPER

LASER

{—Jj POLARIZER SP,PATIAL ^^9^

TRR HALF-WAVE

PLATE MIRROR

NORMAL

RECEIVING TELESCOPE

LOCK-IN AMPLIFIER Vs = Voltage of detector for

scattered light only

Vs + As = Voltage of detector for

scattered light plus

specularly scattered light

5 = rms surface roughness, |im

X = Wavelength of Incident

light, ^m

Fig. 5. Schematic diagram of basic

components of instrument

for measuring angle-resolved

scattering.

Probably the most robust analytical

approach to scattering distribution

developed to date is the bidirectionalscatter distribution function (BSDF), first

proposed by Nicodemus et al., [14] in

1977. Variations of this function are BRDF, BTDF, and BVDF, for reflective,

transmissive, and volume scattering.

149
respectively. All tJiese variations, however,

are considered subsets of BSDF, which is

usually quantified in radiometric terms as the quotient of the scattered surface

radiance divided by the incident surface

radiance [9,11]. In Fig. 6. the scattered

surface radiances is the light flux scattered

per unit surface area per unit solid angle,

i.e..

S = 1^ 1 A ) projected solid angle

watts m

mm2

where

the projected solid angle = (dQs) (cosOg).

Fig. 6. Schematic diagram showing the

geometry for the definition of

the bidirectional scatter

distribution function.

The incident surface irradiance is the light

flux on the surface per unit of illuminated

surface area [11], i. e..

Incident surface irradiance I = —

Pi 9

in watts / mm-^.

A

Thus,

BRDF =

J_

Pi

S

I dP/

1 Pi /A dftscos^s

dP ^ dQg cos^s

Because the BRDF accounts for differences

in reflected light, it will be modified as a

function of the position of any analyzing

polarizer and the type of scatterer. Thus,

one should be able to discriminate

subsurface defects from surface effects if the polarization effects of subsurface

defects are large enough and the optical

properties of the material are appropriate.

Optical Fourier transform theory [15]

predicts that scatter patterns from

surfaces will not be random. This is

illustrated in Fig. 7, which shows predicted

Fourier components of light scattered from

various types of surfaces. Combining

these theories suggests that coupling

Fourier optics with an analyzing polarizer

in the detector system would make it possible to detect and characterize

subsurface defects, providing that

polarization effects attributed to the

defects are large enough and that the

material has appropriate optical

transmission characteristics.

In this study, we have found that, for two

types of Si3N4 ceramics, the optical

transmission properties are such that

depths >100 \Lm of material can be

penetrated at A, = 0.63 \ym and that, for the

two simulated types of defects, i.e., voids

(by drilled holes) and parallel cracks (bysaw cuts), the polarization effects are

sufficient to allow detection if the surface

finish is ground (<.07 ^.m Ra) or ground

and polished (<.03 Ra).

TEST PROCEDURE

Specimens

In the test procedure, flexure test bars 3 x

5 X 25 mm of two HIP'd Si3N4 materials

were produced and then machined by

either a grinding procedure or a grinding

procedure followed by polishing. Four

bars were made of each material. Two of

each were ground and polished and two

were ground only. In each group, ground

150
Fig. 7. Schematic diagram showing

predicted two-dimensional

Fourier transform plane of light

scattered off various types of

surfaces: (a) perfect flat, (b)

periodic (sinewave), (c) periodic

(d) periodic (saw tooth), (e) random, and (f) periodic plus

random.

and polished or ground only, ultrasonically

drilled non-through holes. 750 \im in

diameter, were drilled perpendicular to the

machined surface from the side opposite to

the machined surfaces. Thus, the

machined surface appeared defect free. A

series of two to four holes, =2-4 mm apart,

were drilled so that the defect depth below

the machined surface ranged from 50 to

200 ^m. Similarly, slots were cut in bars

so that 50-200 \Lm of material remained.

Test equipment

The hybrid optical detector system used in

this study is schematically illustrated in

Fig. 8, which shows that a horizontally

polarized 35 mw He-Ne laser is followed by

a V2 plate to allow polarization of the light

to selected angles. A spatial filter with a

6.3 liui pinhole and 24 mm collimator is

used together with a 95-cm focal length

plano-convex lens to focus the light on a

spot. The spot size was =1 mm in

diameter. The detector system consisted 'of a polarization analyzer, a 110-mm

diameter. 155-mm focal length double

convex lens, an image display plate and a

512 X 512 charge-coupled-device (CCD)-

array TV camera coupled to a lens with

remotely controllable zoom and iris. The

output of the CCD camera was fed to an

eight-bit A/D board of a PC that was

loaded with image software locally modified

to perform digital image subtraction.

EXPERIMENTAL SET-UP

Sin SPARC n/OS

wiib 24 bil pbne

Mimr

Fig. 8. Schematic diagram of

experimental optical system

used to study surface and

subsurface defects.

The test specimens were mounted on a

rotate-translate stage equipped with

stepper motors. Thus, a specimen could

be translated at constant incidence angle

or rotated to change the angle of incidence

(hence reflection). Although the entire

optical system was mounted on a 1.3 x 2.7

X .3 m solid granite vibration-isolated

151
optical bench, we found that it was

unnecessary to isolate the optical table.

Procedure

After each specimen was carefully

mounted, the polarization analyser was set so that almost total extinction of light

occurred. This angle changed slightly with

each finish type. Four to six optical

Fourier transform images were taken at

locations on the samples away from the

intentional defects. These images were

then averaged in the computer. This

averaged data set was stored in the

computer memory. By virtue of the

modified software, we could set the

computer so that image subtraction from

"live video" from the CCD camera occurred.

What this implies is that as the specimen

was translated and subsequent images

were displayed on the computer monitor, if these images did not substantially deviate

from the "average image", the result would

be a blank (or totally dark) image.

However, if any subsequent image would

deviate from the "average image" within the

limits of the detection system, then only

that "difference image" would appear.

Such "difference images" were obtained by

careful step-wise scanning of each ceramic

test piece. The size of the step could be

varied from 50 jim to over 1 mm by the

translation stage. By using this method,

and digitally storing scanned images,

analysis of the difference images by digital

image processing methods could be

performed. At the present time, this

analysis consists primarily of gray-scale

histogram analysis to establish differences

in the images that can be used to establish

defect detection sensitivity.

RESULTS AND DISCUSSION

At the time we began these studies, we did

not know if the optical transmission

properties of the selected polycrystalline

ceramics would allow sufficient penetration to allow detection of subsurface

defects. Thus, we determined the initial

optical transmission properties of the

selected materials.

To obtain initial optical transmission

properties, step wedges of both ground and

polished and unidirectionally ground

materials with step thicknesses of 100,

200. 400 and 600 \im were used with laser

light of X = 0.6328 ^m. The results

demonstrated sufficient optical transmission to begin an initial subsurface

detection experiment. Previously available

optical transmission data were for singlecrystal Si3N4 and these data show that

there is a wide window of optical

transmission. However, no data were

available for optical transmission

properties of polycrystalline SislSL^..

Although the microstructure of each

material is quite different, as shown in

Fig. 9, these microstructures (obtained by

polishing and etching) are typical for

Si3N4. Notably, each material has

elongated 3-5-(mi grains and grains with

diameters up to 1 |xm. Although quoted as having a grain size <0.5 }im by some

producers, the Si3N4 grains have high

aspect ratios and these microstructures

will need to be well defined for light

scattering studies, especially for detection

of subsurface defects.

All experimental data reported here were

obtained by using an incidence angle (off

surface normal) of 20°. This was the

smallest angle achievable with the test

equipment. Optical Fourier scatter

patterns were obtained on all specimens

and are similar to theoretically predicted

patterns for randomly rough or unidirectionally roughened surfaces. Typical

patterns obtained are shown in Fig. 10 for

both a ground and polished specimen (Ra = 0.03 |J.m) and a unidirectionally ground

specimen (Ra = 0.07 fxm). These images

can be compared to those shown concepn

tually in Fig. 7. With the experimental

equipment used, no significant differences

between ceramic materials specimens were

verified. However, we observed experimental suggestions of differences that we

have not had time to thoroughly study at

this time.

For artifical subsurface defects (drilled

holes), using the "difference image"

method, we could easily detect 75- and

105-{jm-deep defects on either the ground

and polished or ground specimens. This

was observable in real-time during live

video subtraction and was verified by

152
histogram analysis. A typical example of

our results is shown in Fig. 11, which

shows optical Fourier transform difference

images of specimens with no defects, a 75-

jim-deep defect, and a 102-|jm-deep defect

I on one of the ground and polished

' specimens. What is notable in these

difference images is that the "no-defect

image" is not totally black. This is thought

to be caused by local differences in

microstructure, as noted in Fig. 9.

I Microstructural differences may give rise to

I changes in polarization of scattered light,

j which would then pass the analyzing

polarizer and be detected.

' Fig. 9. Microstructure of two Si3N4

I ceramics, used in this study:

I (a) material A, and (b) material B.

Fig. 10. Typical experimentally obtained

optical Fourier transform scatter

patterns obtained away from

subsurface defect for (a) ground

and polished and (b) unidirectionally ground only specimens of

polyciystalline Si3N4. Note that

the images show the significant

differences in surface theoretically

predicted in Fig. 7.

Fig. 1 1 . Typical optical Fourier transform

difference images (100 x 100

pixel) experimentally obtained on

ground and polished (0.03 ^im Ra)

specimens with drilled-hole

simulated defects: (a) No defect,

(b) 75-Min-deep defect, and

(c) 102-)j.m-deep defect.

For synthesized defects, a different scatter

pattern is seen. The differences are manifested in both a large number of highamplitude components (intensity of white

spots) and Fourier component distribution

as shown in Figs, lib and c. These

153
differences are quantified in Fig. 12, which no Sub-Surface Defects

shows typical histograms of the gray-scale , , , , _ values of the resulting total difference ^ " ' ^ * ' ' ^ ' ^ ' ' *

^ images. i - "•- - - •

To understand the histograms in Fig. 12,

we study the number of pixels with

intensity values less than 138 (0 is white,

255 is black). The reason for this is that

lower gray-scale values indicate larger

quantities of light passing through the

analyzing polarizer than in the no-defect

average image. Thus, we see significant

differences between the 75-^m-deep defect

and the 102-^im-deep defect. This is as

expected, because of the more significant

attenuation (scatter) for deeper defects.

Further, if one carefully examines the

different images of Fig. 1 1 , it is clear that

the higher spatial frequency components of

the 102-|a.m-deep defect have less intensity

than the high spatial frequency components of the 75-|xm-deep defect. Thus,

analysis of these components, unless

totally defect-size dependent, may provide

a method to establish depth sensitivity.

Similar results were obtained on other

ground and polished specimens.

Figure 13 shows the experimentally

obtained optical Fourier transform difference images for a ground specimen (Ra = .07 |xm). As with the ground and

polished specimens, we see distinct highfrequency components with much greater

intensity (brightness), for regions of the bar

with a defect than for regions Avithout a

defect, and further, the intensity is greater

for the 75-)im-deep defect than for the

102-|j.m-deep defect.

This observation is quantified by

examination of the histograms shown in

Fig. 14. Once again, the highest brightness is observed for the specimen with the

75-)jjn defect.

For the synthesized crack defect specimens, only the edge of the cut could be

detected by the scatter. This is currently

thought to be caused by the repeated

internal reflections within the cut cavity

itself, which reduces the power of any light

scattered back to the detector.

6

Pixel Values

75|iinHistograin

1

mi

i

ii m

1 -4hl-i—

40 72.5 105 137.5 170

Pixel Values

102|ini Histogram

40 72.5 105 137.5 170

Pixel Values

Fig. 12. Histograms of the gray scales t of the optical Fourier difference ( images shown in Fig. 1 1 for j ground and polished specimens I (0.03 M-m Ra) with drilled-hole i simulated defects: (a) No defects, ii s[ (b) 75-^m-deep defect, (c) 102- i j [Lui deep defect i
Machined No Defect Histogram

-

• * *

>

-

Fig. 13. Typical optical Fourier transform

difference images experimentally

obtained on ground specimens

(0.07 |j.m Ra) with drilled-hole

simulated defects: (a) no defect,

(b) 75-|j.m-deep defect, (c) 102-

M,m-deep defect.

DISCUSSION AND SUMMARY

Optical scattering by a hybrid crosspolarization, Fourier-optics method has

been used to detect subsurface defects, in

the form of simulated voids, in ground and

polished and unidirectionally ground

Si3N4 specimens. Detection of subsurface

defects is possible because of the optical

transmission properties of polycrystalline

Si3N4. Specimens from two sources were

studied. The grain sizes of these

specimens were found to be quite large,

with diameters of 1 [xm and aspect ratios of

3-5. For He-Ne lasers, X = 0.63 ^im, this

means that the scatterer is as much as 10

times the wavelength of the light used. We

have not yet studied grain size effects on

scatter and only limited machined surface

variations have been attempted. These

defect and grain-size effects must be

separated if possible. The impact of grain

size will need to be carefully studied

because it is likely that subsurface,

machining-induced damage will be grainsize dependent. Further, it is not clear

exactly what effect angle of incidence has

on detection sensitivity, even though Fan

and Huynh [16] and Cuthbert and Huyeh

[17] have recently studied some of these

aspects. Our current work studied

scattering with a 20° angle of incidence

and this also must be optimized.

o

o

J2

o

Machined Sub-Suface Defect at 75^m Histogram

210 215 220 225 230

Pixel Values

235

Machined Defect at 102|im Histogram

35-

3

25

15-

220 225

Pixel Values

Fig. 14. Histograms of gray scales of the

optical Fourier difference images

for drUled-hole simulated defects

shown in Fig. 13 (0.07 |im Ra):

(a) No defects, (b) 75-|im-deep

defect, (c) 102-|im-deep defect

155
In addition, the effect on polarization of

surface characteristics of ceramics

developed by different machining methods

must be established. Bennett and

Mattesan [11] and Stover [9] have

indicated that greater penetration of

semi-transparent materials will be possible

if other than s-polarized light is used,

because s-polarization maximizes reflection and minimizes transmission. Clearly,

this needs to be examined in more detail.

Because of their complex curvature, ball

bearings will have to be rotated rather

than translated for surface studies. For

full 47c-steradian coverage, methods to

spin the balls must be developed; this is

considered possible at the present time.

For most ceramic races, the fatigue life will

mainly be influenced by the contact points

below the balls. Because bearing races are

cylindrical, optical scatter details are not

currently thought to present significant

problems, and will be addressed in the

future.

This method appears capable of detecting

subsurface microcracking as well.

ACKNOWLEDGMENT

The authors thank Jack Lucek of Cerbec

for providing the NBD 200 SisNi material.

Bill Mandler of Enceratec for providing

Toshiba Si3N4 material, and Mike Ressl of

AMOCO Laser Company for helpful

discussions.

REFERENCES

1. Y. Nishihara, H. Nakashima.

N. Tsushima, and S. Ito, "Factors that

Affect Rolling contact Fatigue Life of

Ceramics and Rolling Contact Fatigue

Life of Ceramic Balls and Rollers."

ASME paper 90-GT-377, Brussels.

Belgium, 1990.

2. H. Takebayashi, T. M. Johns,

K. Rukkaku, and K. Tanimoto,

"Performance of Ceramic Bearings in

High-Speed Turbine Application,"

Society of Automotive Engineers,

paper 901629, Milwaukee. WI, Sept.

1990.

3. E. V. Zaretsky, "Ceramic Bearings for * Use in Gas Turbine Engines," ASME

J. of Engineering for Gas Turbines

and Power. Vol. 111. No. 1. pp. 146-157, 1989.

4. J. F. Chudecki, "Ceramic Bearings -

Applications and Performance Advantages in Industrial Applications," SAE . Technical Paper Series 891904, 1989. i

5. S. A. Horton, "Detection of Surface , Defects in Ceramic Rolling Elements,"

proceedings of the 4th International

Congress on Materials for Advanced

Heat Engines. Goteborg, Sweden,

1991, Elsevier App. Sci., pp. 897-904. »

6. RT. Cundill, "Material Selection and

Quality for Ceramic Rolling

Elements," ibid., pp. 905-912. 3

7. S. Jahanmir, L. K. Ives, A. W. Ruff,

and M. B. Peterson, "Ceramic

Machining: Assessment of Current

Practice and Research Needs in the

United States." National Institute for

Science and Technology, Special ii Publication 834, June 1992.

8. S. Jahanmir (editor), Cersimic Bearing

Technology, National Institute for

Science and Technology Special

Publication 824, November 1991.

9. J. C. Stover, Optical Scattering

Measurement and Analysis, McGrawHill. New York, 1990.

10. T. V. Vorburger and E. C. Teaque.

"Optical Techniques for On-line

Measurement of Surface Topography,"

Precision Engineering, Vol. 3, No. 2,

pp. 61-83, 1981.

11. J. M. Bennett and L. Mattson,

Introduction to Surface Roughness

and Scattering, Optical Society of

America, Washington, DC. 1989.

12. B. Ruffing and J. Anschutz. "Surface

Roughness Measurement by 2-D

Digital Correlation of Speckle

Images," Society of Photo-Optical

Instrumentation Engineers, Vol. 814,

pp. 105-112, 1987.
13. M. Shiraishi, "Scope of In-process

Measurement, Monitoring and

Control Techniques in Machining

Process: Part 1. In-process

Techniques for Tools," Precision

Engineering. Vol. 10. No. 4.

pp. 179-189, 1988.

14. F. E. Nicodemus, J. C. Richmond,

J. J. Hsia, I. W. Ginsberg, and

T. Limperis, Geometric Considerations and Nomenclature for

Reflectance, U.S. National Bureau of

Standards (NBS) Monograph 160,

U.S. Department of Commerce, 1977.

15. J. W. Goodman, Introduction to

Fourier Optics, McGraw-Hill, New

York, 1968.

16. Y. Fan and V. M. Huynh,

"Investigation of Light Scattering from

Rough Periodic Surfaces: Numerical

Solutions," Optics and Laser

Technology, Vol. 24, No. 3. pp. MSISO, 1992.

17. L. Cuthbert and V. M. Huynh,

"Statistical Analysis of Optical Fourier

Transform Patterns for Surface

Texture Assessment," Measurement

Science and Technology, Vol. 3. No. 8,

pp. 740-745. 1992.
I

I

i (Jl
CHATTER AND SURFACE PATTERN DETECTION FOR

CYLINDRICAL GRINDING USING A FLUID COUPLED

ACOUSTIC EMISSION SENSOR

Y. P. CHANG and D. A. DORNFELD

University of California, Berkeley, CA

An acoustic emission (AE) sensor setup has been designed to monitor the cyhndrical grinding process using the coolant fluid as the signal transmission media. Modifications have been made in the coolant flow path to compensate for the problem

of coolant flow instability inherent in a fluid coupling setup. Experiments have been

conducted to evaluate the eff'ectiveness of this setup on monitoring grinding chatter

and grinding patterns. The root mean square energy of AE is used as the primary sensor signal characteristic. Successful detection of chatter vibration has been

achieved through this setup. Defection of workpiece surface patterns existing prior

to grinding has been investigated.

INTRODUCTION

Grinding is usually used as a finishing process in manufacturing. Therefore, the workpiece quality (surface pattern, roughness, and

roundness) produced by the grinding process

is critical to the quality of the product, and

damage to the workpiece caused by abnormalities in the grinding process is very often irreversible. Since it is often difficult or practically

impossible to predict these abnormahties much

attention has been paid to the use of sensors

for process feedback[l-4] and control [5]. Originally applied to the destructive test- ing of structures, acoustic emission (AE) has

became a monitoring technique in manufacI turing due to its sensitivity to process parameters. Active research in AE has shown the

strong potential of this technique in manufacturing process monitoring [6]. Successful applications of AE for grinding processes have been

developed in monitoring contact and sparkout, wheel dressing, wheel wear, chatter detection, workpiece surface roughness, and process parameter variations [7-13]. A fundamental study of AE in grinding was conducted by

Dornfeld and Cai [7], in which the sensitivity of AE for detecting contact/noncontact between wheel/ workpiece was established. Research has been performed on AE for monitoring wheel loading conditions [7-8], and

wheel dressing conditions [10-11], and to predict wheel life in different grinding processes.

With respect to monitoring of surface properties, Inasaki has investigated the influence on AE amplitude due to the change in workpiece

surface roughness [8]. And several studies were done on correlating the AE signal to the work

surface finish in turning process [14-17].

This study diff"ers from earlier research in

accessing surface roughness in grinding with

AE [12] in that it determines the feasibility of

using AE to observe the repetitive 2-D patterns generated on the work surface during

cylindrical grinding. The attempt here is to

establish a relationship between the in-process

AE signal with those patterns on the work surface. Surface roughness values will be used to

quantify the work surface characteristics.

159
EXPERIMENTAL SETUP

The design of the AE instrumentation in

the present setup, which is different from that

used in previous work [12], is shown in Fig- ure 1, and is aimed at solving the problem

of coolant flow noise. The instability of the

output flow rate from the coolant pump, and

the interaction between grinding wheel and

coolant fluid have been the major contributors to coolant flow noise generation. In this

work, a fluid reservoir has been incorporated

into the coolant flow path, as indicated in Fig- ure 1. This reservoir acts as two filters removing the flow rate instability (since the output

flow rate is determined by the inlet diameter,

outlet diameter, and fluid height in the reser- voir and not directly related to the output of

coolant pump) and minimizing the bubbles in

the output flow of the coolant pump.

To avoid interaction between the coolant

and the grinding wheel, the coolant stream is directed just adjacent to the grinding wheel. In

this way the stream has minimum interaction

with the grinding wheel (preferred from the

sensor point of view) while still giving a satisfactory cooling efi"ect for the process. The cutting signal generated at the cutting point travels through the workpiece and coolant to the

AE transducer mounted on the coolant supply

tube.

Figure 1 illustrates the setup on a cylindrical grinding machine. The diameter of

the aluminum oxide grinding wheel is 9.6

in. (243.8mm); the workpiece is made of

1045 steel with the size of 72.1mm dia. and

304.8mm length. The grinding wheel speed is 1720rpm, and workpiece speed is 200rpm. The

feed rate (grinding table speed) is 20in/min

(508mm/min). And the typical value of depth

of cut is 0.0005in. (0.0127mm). A broadband

transducer is used to detect the AE signals,

which pass through a cascade of preamplifier

and main amplifier providing a total gain of 50

dB. The ampfified AE signals were recorded by

a modified video recorder. A proximity transducer was installed on the head stock spindle

to measure the rotational speed of the workcooiant pump

video recorder

80386 computer

(digitizalion, FFT, topograph.—)

Fig.l The Set-Up for Grinding Experiments

with AE Sensor

piece. The signal from the proximity transducer also acts as the base line for positioning

the AE signal on the workpiece surface. This

signal was recorded on the audio channel of the

video recorder. Both the AE root mean square

(rms) energy signal and the proximity sensor

signal were monitored on a digital oscilloscope.

Digitization of the signals for analysis (calculation of Fast Fourier Transfer and rms) were performed on a personal computer. , ; The surface profile measurements were done by a stylus surface measurement instrument connected to a personal computer. Measurement of surface roughness along the axial direction of the cylindrical workpiece was

straightforward since the instrument was designed to measure "flat" surfaces. For measurement along the periphery of the workpiece,

the surface profile at 15 degrees to the center

axis was measured and the peripheral profile

approximated by geometrical calculations.

CHATTER DETECTION

The chattering phenomenon investigated

in this study can be classified as self-excited

grinding vibration. The vibration can be a 160
result of many factors including flexibility of

workpiece-wheel structure, repetitive roundness error of the wheel/workpiece, which gives

periodic variation on the depth of cut, and

wheel wear. In this study, the wheel wear condition is the major cause of chatter.

At the beginning of each chatter experiment, the grinding wheel is dressed. Several

cutting passes are conducted with one sparkout pass following each cutting pass. As the

cutting passes continue, the wheel becomes

progressively more loaded. The loaded wheel

contributes additional thrust force components

to the wheel/workpiece interaction in addition

to the cutting force from the fresh wheel, which

acts on the workpiece tangential to the workpiece circumference. This thrust is repetitive

since the grinding wheel is rotating at a constant speed. As the wheel loading increases,

the repetitive thrust will have an increasing

magnitude, and the magnitude of the machine

resonant frequency excited by the thrust will

also increase. This phenomenon has been re- ferred to as "chatter growth". The chattering

noise is audible when it reaches a certain magnitude. The experiment is halted when the

chatter noise is first detected by the technician

operating the experiment to prevent machine

damage.

Figure 2 shows the frequency spectrum

of the root-mean-square energy, AErms-, for

five consecutive cutting passes. The sparkout passes conducted after each cutting pass

have not been included in the figure. These

sparkout passes are used to clean up the workpiece surface without any infeed of the grinding wheel. The chattering disappears in these

sparkout passes.

The "chatter growth" phenomenon can be

easily observed in the AE^rns signal at the fre- quency of 417.5Hz (Fig.2). According to other

research results, an increase in depth of cut

(radial engagement) gives a larger AErms signal level [7] and more severe chatter vibration.

The exponential trend of the signal growth

would have been .more dramatic had the depth

of cut not been reduced for different cutting

passes.

Fig.2 FFT of AErms for consecutive

Cutting Cycles

The workpiece surface profile after the fifth

pass in Figure 2 has been preserved and measured. There are chatter marks along the cir- cumferential direction, which can be easily observed. Visually counting the marks (approximate 126 chatter marks can be spotted on

the workpiece) with a workpiece rotation fre- quency of 3.15Hz the chatter frequency is calculated to be at 396.9 Hz, which is close to the

resonant frequency of the AErms signal.

15

-15 -

0 0.5 1 1.5 2 2.5 (< ^ Length (mm'

one period

Fig. 3 Surface Profile along the Workpiece

Circumference

Figure 3 is the measurement result of the

surface profile in the circumferential direction.

The length for one period of chatter mark is calculated to be 1.8 mm. The average peakto- valley height for the chatter marks is 30

//m. Since generating this surface resulted in
a significant AE signal level, seen in Figure 2,

and from the signal growth in the same figure,

chatter vibrations with smaller peak-to-valley

heights are detectable by AE. Figure 4 shows

the AErms for a less severe chatter vibration

and the generated surface profile with peakto- valley height at an average of 20 //m.

To examine the similarity between AErms

and chatter marks, the AE signal has been fil- tered by a band pass finite impulse response

(FIR) filter at 417.5 Hz. The filtered signal

is then synchronized with work rotation us- ing the proximity sensor signal. Finally this

synchronized signal is mapped to a two dimensional contour plot to show the AE signal patterns. Figure 5a shows the 3-D plot of

AErms after filtering and synchronization. Fig- ure 5b illustrates the blow up of figure 5a in

2-D contour plot, which clearly shows the patterns contained in the signal. It is observed

that the AE signal pattern is similar to the

chatter marks on the work surface.

DETECTION OF WORKPIECE SURFACE PATTERNS

This is a preliminary approach to detecting workpiece surface patterns generated during grinding by using the AE signal. In this approach, surface patterns have been artificially

generated on the workpiece before the grinding process. The workpiece is then ground and

AE signals recorded to determine the sensitivity of the AE signal to the existing patterns of

the surface profile. The first group of experiments qualitatively tested the feasibility of this

approach, while in the second group of experiments a quantitative test was attempted.

Two or three strips with a coarse surface

roughness were generated on the workpiece

surface along the axial direction by manual fil- ing, as indicated in Figure 6. Figures 7- 9 show

the 1-D, 2-D, and 3-D views of the AErms after

synchronization to work rotation for the case with two strips on the workpiece. As illus- trated in Figure 7 there are two signal spikes

in each period of workpiece rotation. When all

Fig.4 EFT of AErms for Less Severe Chatter

3-D Diagram

2 4 6 8 10" 12 14 16 18 Number of Workpiece Rotation

Fig.5 AErms Signal at the Chatter Frequency
the workpiece rotation periods for the cutting

pass were joined in the 2-D contour plot, as in

Figure 9, it can be observed that the AErms

forms two strips along the axial direction of

the workpiece. These strips are at the equivalent location to those physically on the workpiece. Figure 9 gives a three dimensional view

i of the AE signal, in which the strips contained

j in the signal can be clearly observed. Similar results were obtained for the case of three

j surface strips.

Fig. 6 Simulated Grinding Patterns on

the Workpiece

50 - f

0' •

. . One Workpiece Rotation

Fig.7 1-D Diagram of AE^ms for Workpiece

with 2 Strip Patterns

Number ot Workpiece Rotation Fig. 8 2-D Diagram of AErms for Workpiece

with 2 Strip Patterns

Fig. 9 3-D Diagram of AErms for Workpiece

with 2 Strip Patterns

In another set of experiments, quantification of the signal sensitivity was attempted.

Three slots were machined on the workpiece

surface by slitting saws. The width of the slots are 0.008 in. (0.203mm), 0.030 in. (0.762mm),

and 0.040 in. (1.016mm). The slots lay along

the axial direction on the workpiece surface at constant increments (120 degrees apart on the

circumference). Figure 10 is a typical 1-D signal representation from a cutting pass for this

workpiece. Signal spikes similar to previous ex- periment results were observed. The 2-D con- tour plot of the signal, as depicted in Figure

1 1, reflects the pattern of the slots with equivalent locations. Consecutive cutting passes were performed on the workpiece, with depth of cut

varying at some passes, to gradually remove
3100 10 030- slot (0.762mm)

0.040" slot (1.016mm)

0.008" slot (0.2032mm)

One Workpiece Rotation end

Fig. 10 1-D Diagram of AEr

with 3 Slots

for Workpiece

10 15 Number of Workpiece Rotation Fig. 11 2-D Diagram of AE^

with 3 Slots

for Workpiece

the slots in the workpiece. The signal patterns

for different passes were investigated.

The average mean and standard deviation

(STD) of AErms were plotted in Figures 12

and 13. The average AErms rnean, Fig. 12,

reflects the variation in depth of cut but not

the dimensions of the slots. This results from

the fact that the change in metal removal rate

introduced by the slots is so small that it does

not affect the average value for the whole cutting pass. From Figure 13, two levels of standard deviation value were observed. For the

first 1 1 cutting passes, when all the slots were

present existing, the standard deviation val- ues were at a higher level. During the 12th

and 13th cutting passes where part of the slots

were removed, the level for standard deviation

started to drop. Finally, for the 14th to 16th

cutting passes where no slots existing on the

workpiece, the standard deviation stayed at a

lower level which was approximately 50 mV

below the higher level. This figure, however,

only gives information about pattern existence

in a macroscopic view. To understand the ex- istence of individual slots, the AErms signal

corresponding to each slot was investigated.

3200

3000

2800

2600

2400

2200

2000

Fig.

D.O.C=0.001in. (0.0254mm)

,^--D.O.C..0.00075in.

(0. 0127mm)

D.O.C=0.0005in. (0.01905mm)

—I—'—r—'—I—'—1—'—(—'—I——I—'—I—'—I

0 2 4 6 8 1012141618 20

Cutting Passes

12 Ave. AErms Mean for Different

Cutting Passes

300

260

100

patterns existing

^t*—Q—

->• no patterns

2 4 6 8 1012141618 20

Cutting Passes Fig. 13 Ave. AErms Standard Deviation for

Different Cutting Passes

Figure 14 is the 1-D signal diagram from

the last cutting pass where no slots existed on

the workpiece before grinding. The signal positions corresponding to the three slots are indicated on the diagram. It can be observed

that the signals at these positions are the same

as others, which is very different compared to

the spikes they generated in Figure 10. It is, therefore, reasonable to assume that the signal generated by the slots are among the most

spiky for the whole cutting pass. The "spikiness" of the AErms signal generated by the

164
slots were characterized by the deviation of

the spike from the mean of the cutting pass

signal. If the assumption is correct, the deviations should be bounded by a certain value

related to the standard deviation of the cutting

pass signal.

(0.2032mm) (0.762mm)

0.008- slot 0.030- slot

One Workpiece Rotation

Fig. 14 1-D Diagram of AErms for Workpiece

with no Slots (previous location of

slots shown by arrow)

Slot Widlti B— 0 040in. (l.Oiemm)

0.030 in. (0.762mm)

0 008 in. (0.2032mm)

10 12 14 16 18 20

Cutting Passes Fig. 15 Ration of Spike Level over Standard

Deviation for Different Cutting Passes

repetitive patterns. The 2.2 standard deviation threshold implies that the signal spikes

generated by the slots on the workpiece will

always fall among the highest 2.3 percent of

the signals for the cutting pass.

CONCLUSIONS

To verify the assumption, the deviations were compared with the average standard deviation for each slot in every cutting pass. As

illustrated in Figure 15, there were two levels

of deviation values for each slot - the higher one for slot existence and the lower one for

slot when removed. Through visual inspection between each cutting pass, it was recorded

that after the 11th cutting pass, the 0.030

in. (0.762mm) slot was completely removed.

And after the r2th cutting pass, the 0.008 in. (0.203mm) slot was also completely removed.

At the end of the 13th cutting pass, all three

slots were completely removed. From the cor- relation of this observation and the signal levels in Figure 15, it was concluded that the as- sumption was correct. To utilize the assumption for slot existence

detection, a threshold value is necessary to divide the spikes from other signals in the cutting

pass. The threshold value of signal deviation was set to be 2.2 standard deviations based on

the experiment results. It is assumed that a normal distribution describes the AEyms from

regular cutting passes on workpieces with no

The detection of grinding chatter and

workpiece surface patterns has been studied

using a refined fluid coupled acoustic emission

sensor setup. It was shown that the AErms signal was able to pick up the chatter vibration

generated through the grinding process. The

chatter growth phenomenon was captured in

the frequency spectrum of AErms during several cutting passes. Simulated grinding patterns on the workpiece were detected by the

variation of the multi-dimensional AErms signal spikes. A suitable threshold of the variation for indicating surface defects has been determined to be 2.2 times standard deviation.

This technique is presently under evaluation in an industrial cyhndrical grinding operation. The objective of that evaluation is to

determine whether or not process instabilities

(chatter ) and the resulting surface degradation can be detected in process using acoustic emission sensing. Signal levels exceeding a

threshold would indicate generation of surface

defects and the grinding machine controller

could respond by a change in feed rate or cycle stop to alert operators.

165
ACKNOWLEDGEMENT

This work is supported in part by Alcoa

Co. and the industrial affiliates of the Laboratory for Manufacturing Automation at the

University of California at Berkeley.

References

[1] Gao, Y. and .Jones, B. 1992, "An In- tegrated Sensor System for Size and

Roundness Control in Plunge Grinding",

Int. J. Mach. Tools Manufact. . Vol. 32,

No. 3, pp. 281-290.

[2] Shiraishi, M. 1988, "Scope of In-Process

Measurement, Monitoring and Control

Techniques in Machining Processes- Part

1: In Process Techniques for Tools",

Precision Engineering, Vol. 10, No. 4,

pp. 179-189.

[3] Shiraishi, M. 1989, "Scope of In-Process

Measurement, Monitoring and Control

Techniques in Machining Processes- Part

2: In Process Techniques for Workpiece",

Precision Engineering, Vol. 11, No. 1,

pp.27-37.

[4] Shiraishi, M. 1989, "Scope of In-Process

Measurement, Monitoring and Control

Techniques in Machining Processes- Part

3: In Process Techniques for Cutting Processes and Machine Tools",

Precision Engineering, Vol. 11, No. 1,

pp.39-47.

[5] Gao, Y. Jones, B. and Webster, J. 1992, "An Integrated Size and Roundness Adaptive Control System for the

Plunge Grinding Process", Int. J. Mach.

Tools Manufact. . Vol. 32, No. 3, pp.291-

303.

[6] Dornfeld, D. A. 1992, "Application of

Acoustic Emission Techniques in Manufacturing" , Proc. nth Int'l Acoustic

Emission Symposium, JSNDI, Fukuoka,

.Japan, pp. 1-15.

[7] Dornfeld, D. A. and Cai, H. G. 1984,

"An Investigation of Grinding Processes

with Acoustic Emission", Trans. ASME,

J. Eng. Industry, Vol. 106, no. 1, pp. 28-

33^

[8] Wakuda, M. and Inasaki, I. 1991, "Detection

of Malfunctions in Grinding Processes",

Proc. 4th World Meeting of Acoustic

Emission . ASNT, Boston, MA, pp. 494-

501.

[9] Inasaki, I. 1991, "Monitoring and Optimization of Internal Grinding Process",

Annals CIRP . Vol. 40, no. 1, pp. 359-362.

[10] Inasaki, I. and Okamura, K. 1985, "Monitoring of Dressing and Grinding Processes with Acoustic Emission Signals",

Annals CIRP. Vol. 34, no. 1, pp. 277-280.

[11] Konig, W. and Meyen, H. P. 1990, "AE

in Grinding and Dressing: Accuracy and

Process Rehability" , Proc. 4th International Grinding Conference, Technical

Paper MR90-526, SME, Detroit, Michigan.

[12] Chang, Y. P. Diniz, A. E. and Dornfeld, D. A. 1992, "Monitoring the Grinding Process through Acoustic Emission Using a Squirter", Proc. Japan-USA

Symposium on Flexible Automation, ASME, San Francisco, CA, pp. 335-340.

[13] Blum, T. and Dornfeld, D. A. 1990,

"Grinding Process Feedback using Acoustic Emission", Proc. 4th International

Grinding Conference , Technical Paper

MR90-525, SME, Detroit, Michigan.

[14] Diniz, A. E. Liu, J. J. and Dornfeld, D. A.

1992, "Correlating Tool Life, Tool Wear,

and Surface Finish Roughness by Monitoring Acoustic Emission in finish Turning" , Wear.

[15] Blum, T. Suzuki, I. and Inasaki, I. 1988, "Development of a Condition
Monitoring System for Cutting Tools

using an Acoustic Emission Sensor",

Bulletin Japan Soc. Prec. Eng'g, 22, 4,

pp. 304-308.

[16] Dalpiaz, G. 1991, "Monitoring Workpiece

Quality Deterioration in Turning by AE

Analysis", Proc. 4th World Meeting on

Acoustic Emission , ASNT, Boston, MA.,

pp. 479-486.

[17] Diniz, A. E. Liu, J. J. and Dornfeld, D.

A. 1991, "Monitoring the Surface Roughness Through AE in Finish Turning",

Proc. 4th Meeting on Acoustic Emission,

ASNT, Boston, MA., pp. 487-493.

167

EFFECT OF GRINDING ON

STRENGTH

169

fflGH EFFICIENCY GRINDING OF STRUCTURAL

CERAMICS

H. G. WOBKER and H. K. TONSHOFF

University of Hannover, Hannover, Germany

Advanced structural ceramics, such as silicon nitride or alumina, are attractive

for many advanced applications due to their high strength at elevated

temperatures, wear resistance and chemical stability. During the sintering

process ceramics undergo a volumetric shrinkage and thus are far from

meeting high demands on their form and size accuracy. To meet these demand

in most cases grinding with diamond tools is necessary as a finishing process.

The cost associated with this manufacturing step may reach up to 60 % of the

total cost thereby diminishing the competiveness of ceramic parts. The paper

presents an approach in grinding of these materials, based on the model of the

equivalent chip thickness in grinding. The correlation between the chip thickness and the subsurface damage is used for the process layout. Thereby it is possible to reduce machining time effectively and to avoid effects of the

grinding process on materials strength.

0 Introduction

Advanced structural ceramics such as

silicon nitride or alumina are attractive for

many advanced applications due to their

high strength at elevated temperatures,

resistance to chemical degradation, wear

resistance, and low density. Despite these

advantages there are considerable

impediments to the introduction of advanced ceramics. With current technology

fabrication costs are high, compared to

other materials and component reliability is uncertain. The cost of machining can be as

high as 90 percent of the total cost of

some high precision components; and

damage produced during machining can

produce premature failure due to crack

growth under stress. Thus apart from

suitable material properties safe and eco- nomic machining is another prerequisite

for a wider use of structural components

(1-4).

1 Aim of the study and experimental

details

Although advances have been made in near-net shape technology, machining in the fired condition remains to be the

predominant process for the manufacturing

of structural ceramics. In order to achieve

high quality surfaces and high geometric

accuracy the last step in manufacturing of

ceramic components is grinding with

diamond wheels. However, cutting,

grinding, and polishing operations were

found to leave subsurface cracks and

171
residual stresses in the surface layer of,

brittle materials (Fig. 1) (1-3). The strength

degration effect of microcracks in the

surface of ceramics can easily be

understood by equations of fracture

mechanics. The stress necessary to

propogate a microcrack is given in

equation 1 (5),

<^B=K,cY.-l 71/

While K|c is the fracture toughness of the

material, Y refers to the crack geometry

and c gives the crack length. If we assume

a semicircle-shaped crack with a radius c = 50 |im, the bending strength of silicon

nitride ceramic will be reduced by at least

30 percent.

It is obvious that compressive stresses

acting at the flaw crack tips can increase

strength by suppressing crack growth

radial crack —

elongated

radial crack ~

diamond groin -

surface flow - >^

longitudinol crock

residual stresses

6282

residual flexural strength

K Ic Y

1

a, 99.9 (-

O 0/ O

O 10

1 o

cup whee

0,1

surface grinding

I I I I I I I I 100 200 too MPa 1000 bending strength

663^

3i/6681 © IFW 8i59

Fig. 1Influence of machining on the strength of brittle materials

under stress. The interaction of flaws of

different size, orientation and length and

residual stresses is the origin for the effect

of grinding on bending strength and

fracture probability of ceramics. Fig. 1 shows strength values obtained from

conventional surface and cup wheel

ground silicon nitride (6).

The present stody intends to investigate

the interaction between machining conditions in grinding and the induced stresses

and microcracks in order to achieve an

optimized process design in terms of

productivity and product quality.

This paper reports on investigations

conducted on alumina reinforced with

dispersed zirconia and on hot isostatic

pressed silicon nitride. The materials

properties are given in Table I. The

ceramics were ground with diamond cup

wheels of different grit size and

concentration in order to achieve flat 172
surfaces of different roughness and states

of surface integrity. Residual machining

stresses were evaluated by using x-ray

diffractometry. For the variation of the

penetration depth of the x-rays copper,

chromium and titanium anode x-ray tubes

were used (7). The bending strength of the

ground samples were evaluated by fourpoint bending tests. For each grinding

condition a set of 40 samples was tested.

Table I: Material Properties

Al203/Zr02 Si3N4-Y203

Vickers hardness (HV 0.5) 2000

Flexural strength (04^5) MPa 600

Youngs modulus (E) GPa 380

Poissions ratio 0.24

Weibulls Modulus (m) 10

Fracture toughness

(K|J MPa • m°-5 5.8

1500

750

280

0.26

15

6.0

The grinding experiments were conducted

with resin bonded diamond cup wheels of

350 mm diameter and D15 - D181 (15 |im181 Jim) grit size. Tangential, normal and

radial forces were measured and recorded.

Castrol "Honilo" grinding oil with a viscosity

of 4.15 mm2/s (at 40 °C) containing phosphorus additives was supplied to workpiece/wheel interface as coolant using a

single high pressure jet.

2 Results

2.1 Forces and wheel wear i n grinding

Grinding damage in brittle materials

orginating from mechanical loads is mainly

due to the grinding normal force Fp. Initial grinding test were therefore carried out to

identify the influence of machine settings

on grinding forces and wheel wear. As the

true grit load can not be measured, the

average normal grinding force per unit

wheel width Fp' is used to compare the

different grinding conditions. For the

comparision and explanation of results a

term called grit depth of cut is introduced

which refers to a theoretically value the

average depth of cut of a cutting point

during machining. The grit depth of cut, or undeformed chip thickness hcu, in face

grinding is given by equation 2 (8).

C-r

hcu - 1 TTT I IT' (2)

Where C is the number of cutting edges

per unit area, r the chip width to thickness

ratio, Vc the cutting speed and Vfa the axial

feed speed.

For the test described in this part forces

and wheel wear were measured after approximately V'v^, = 600 mm^/mm of stock

had been removed. The wheels were

dressed and profiled after each grinding

cycle with SiC-wheels mounted on a rotary

brake truer.

From Fig. 2 the influence of different wheel

speeds on the grinding forces and the

surface roughness can be observed. With

increasing wheel speed the grinding forces

gradually decrease. Higher wheel speeds

at a constant material removal rate lead to

a reduction of the grit depth of cut resulting

in lower grinding forces.

From Fig. 3 the effect of the axial table

feed speed in grinding of silicon nitride on

grinding forces and surface roughness can

be derived. In general, higher feed speeds

should lead to higher grinding forces due

to an increase in grit depth of cut. However

an increase of the material removal rate by

means of the feed speed Vfg in grinding of

silicon nitride results in lower grinding

forces. This indicates that the diamond

grits are prevented from blunting at higher

depth of cut and short grinding cycles.

173
C71

Q.

If)

O if) O o

in CDD

O

250

N

mm

150

100

50

0

5

(im

3

2

10

Ft

\ •1— \ .1 '——

-^s:^

**•-»— . •

0 20 /O m/s 60

cutting speed

o

o

CDc

_C

"i_cn

300

mm'^

mrn^

100

0 1

10 20

cutting

35 m/s 55

speed

material SijN^ wheel 5A2-D91K+8821JYC100

truing SiC-wheel

dressing alumina stick coolant oil

v,jj = i mm/min

V;, = 600 mmVmm

3/./6585 © IFW

Fig. 2: Grinding forces at different cutting speeds

Consequently it is possible to increase

efficiency of the grinding operation without

any loss in accuracy.

Following general grinding technology, the

use of "coarser" grit diamond wheels

should result in lower grinding forces and a

detonation of the surface roughness of the

ground parts due to the increased depth of

cut. Fig. 4 shows results obtained from

grinding using wheels of grit size D15-

D181 at a constant material removal rate

Q'w. For the wheels with grit size D15-D91

the tendency is as expected, the forces

decrease, simultaneously the surface

roughness increases. However, comparing

the results of the D91 wheel with those

obtained with the D181 wheel the

tendency is reversed. It appears to be

possible to achieve surface roughness of

the ground parts as good as those with the

D46 wheel. This tendency can be ex- plained by the number of cutting points per

unit wheel surface. A coarser grit often

show more than one cutting point (8),

therefore the total number of cutting points

per unit wheel surface is higher compared

to fine meshed wheels. The increase of the

grinding forces observed with the coarse

wheel supports this assumption.

All in all it can be stated that it is possible

to maintain surface quality even with coar- ser wheels provided adapted truing and

dressing cycles are applied.

Apart from grinding forces and surface

roughness wheel wear in grinding of

ceramics is another important factor. A

comparison of G-ratios (i.e. stock removal/wheel wear) obtained in grinding of

silicon nitride is given in Figure 4. From

Figure 4 it appears that finer grit wheels

wear relatively rapidly. Especially the DI81

wheel shows a good performance, due to

its higher mecanical and thermal wear

resistance.

174
250

N

cnc mm

c

O 1/1 m a;

in

O d;

o

150

100

50

0

5

|im

3

2

1

0.

i

i

1 1

1 1

»

c3 300

L 8 mm/min

feed speed

16

o mm3

mm3

o

1_ 200

cnc

ind 100

CD

8 mm/min 16

feed speed fa

material Si3N, wheel 5A2-D91KH-8821 JYC100

truing SiC-wheel

dressing alumina stick coolant oil

= 20 m/s

V;, = 500 mmVmm

3Z,/668^ ©IFW8A65

Fig. 3: Grinding forces versus axial feed speed

D15 Di.6 D91

grit size D

D181

500

o mm

E 300

c 200

X)z 100

0 D15 D/.6 D91 ^m D181

grit size D

material wheel 5A2-K+8821 JY C100

truing SiC-wheel

dressing alumina stick coolant oil

= 20 m/s

v,g = i. mm/min

V;^ : 600 mmVmm

3i/5583 IFW 8i67

Fig. 4 : Grinding forces and G-ratios for different grit sizes

175
The various grinding tests have

demonstrated so far some potentials for an

increase of efficiency in grinding of

ceramics. However, the effect of the

different machine settings and the wheel

specifications on the workpiece quality i.e. residual stresses and subsurface damage,

needs to be evaluated in order to maintain

workpiece quality.

3.2 Residual stresses

For the exact description of the state of

residual stresses in ground surfaces it is necessary to determine the stress depth

profile and the orientation of the main

stress direction. From metallic materials it

is known that non-directional processes

like a thermal influence produce no

preferred orientation of the residual

stresses. In contrast processes with a

defined direction, like cutting or grinding are known to produce a distinctive

orientation between residual stresses and

the direction of cut. This is also true for

residual stresses in ceramics as Fig. 5

demonstrates. The measurements re- vealed that the main stress direction is oriented perpendicular to the direction of

cut, whereas a minimum is found in the

direction of cut. The measurements were

realized on the (030) lattice plane of the

AI2O3/IO % Zr02 ceramic.

The directional dependence of stresses in the surface layer correlates with the

mechanical effects of the cutting edges.

The largest amount of plastic deformation

Fig. 5: Influence of the machining direction on residual stresses

is found perpendicular to the cutting

direction. The material remaining in the

cutting path is deformed plastically to a

smaller extent due to crushing effects

during material removal.

The penetration depth of x=rays depends

on their wave length. It is defined as the

value where the primary intensity is reduced to 1/e, or 37 percent of the initial value. In ceramics the layer of residual

stresses extend well below the surface and

176
there are high stress gradients within the

penetration depth of the radiation. Three

different radiations where used to measure

the stress versus depth profile in ground

material. A typical profile found is shown in Fig. 6. The penetration depth of copper-Ka

radiation in the Al203/Zr02 ceramic is about 1 7 ^im at the (226) lattice plane and

21 ^im at the (416) plane. Due to their

longer wave length the use of chromium or

titanium radiation reduces the penetration

depth to 8 [im at the (030) plane and 4 |im

at (113) plane respectively. Directly

underneath the surface, high compressive

stresses are found.

Figure 7 shows the effect of different

cutting speeds on residual stresses in

ground silicon nitride. The stresses

decrease with increasing wheel speed in grinding. This corresponds to the alteration

of the grinding forces which can be

explained by the reduced grit depth of cut

at high wheel speeds (4). Again, a

comparison of stress values measured

with chromium and copper radiation

indicates the existance of stress gradients

in the subsurface layer.

To evaluate the influence of different wheel

specifications on residual stresses, wheels

of different grit size and concentration

were compared (Rg. 8). All tests were

conducted with diamond cup wheels at

different cutting speeds. All wheel

specifications tested showed the same

tendency of lower compressive stresses at

higher cutting speeds. Furthermore, at a

defined cutting speed a coarser wheel with

fewer cutting edges (i.e. coarser grit size or lower concentration) showed higher

compressive stresses.

The reduction of cutting edges causes an

increase of the grit depth of cut. With a

higher grit depth of cut the forces acting at

a single cutting point increase and thereby

causing higher mechanical loading of the

subsurface. With higher mechanical

loading an increased plastic deformation

and higher compressive stresses are li- kely.

0

MPa

u) -500

Q)i_ too -1000

13 g

'(/) <u--1500

-2000

^ (^16)

)

(226)

^(hkl)

(116)

13)

0 5 1 0 15 20 ^im

depth beneath surface x

material

g rindin g conditions: peripheral surface grinding

= 25 m/s

v„ = 167 mm/s

a, = 0.05 mm

wheel : D91 Kt888 NY C50

used X-radiation CuK,

CuK„

o <s„ TiK. A a„

= 0.15418 nm

a

= 0.22911 nm

• cr, =0.27508 nm A a.

326/6713 ©IFW

Fig. 6: Stress versus depth profile in ground Al203/Zr02 ceramic

177
0 10 20 30 1,0 cutting speed

material wheel

truing

dressing

coolant

= 600 mm^/mnn

= ^ mm/min

SLD9r-6A2K+882 JY C100

SiC - wtieel alumina stick

oil

wave length X bragg' arlgle 29

lattice plane

penetration depth

CrKa

0,22911nm

131.84°

(212)

12 \im

CuKa

0,15418 nm

141. 5*

(323) 37 um m/s 60

34/2856 'IFW 6369

Fig 7: Influence of cutting speed on residual stresses

1/1 in

o

20 30 ^0

cutting speed

material wheel

truing

dressing

coolant

V'^ = 600 mmVmm

= 4 mm/min

radiation wave length

bragg angle 29

lattice plane

penetration depth

SijN,

6A2-K*8821 JY C 100 SiC-wheel

alumina stick

oil

CrK.

0,22911 nm

131,84

( 212 ) 12 \im

34/2855 IFW 7900

Fig. 8: Residual stresses in silicon nitride for different wheel specifications

178
0.2 O.Z, 0,6

chip thickness h^^j

material SijN, wheel 6A2Kt8821 JY ClOO

truing SiC-wheel

dressing alumina stick coolant oil V, = U mm/min. v = 10-55 m/s

V;, = 0 - 600 mmVmm

radiation CrK, wave length 0,22911 nm

bragg'angle 20 131.84°

lattice plane (030)

penetration depth 8.6 (im

as fired O

polished D15 C100 D46 C100 •

D91 C100 B

D91 C50 A

D181 C100 O

34/ 5694 IFW 8477

Fig. 9 : Residual stresses in silicon nitride versus the chip thickness hcu

In Fig. 9 the residual stresses found in grinding silicon nitride are plotted versus

the undeformed chip thickness h^u ^or

different wheel specifications; cutting and

feed speeds. Residual stresses are compressive in all cases and show an

almost linear correlation with the chip

thickness hey. Starting from a low state of

stress for the polished material the compressive stresses increase at higher h^u

values. For the coarsest wheel, at a wheel

speed of 10 m/s, the maximum stress was

observed at -400 MPa. The good correlation between the residual stresses and

the chip thickness supports the assumption that residual stresses in ceramics are mainly caused by the plastic deformation

of the material by the cutting edges.

3.3 Grinding and strength of ceramics

In addition to residual stresses depend on

the machining process subsurface

microcracks may be generated in grinding

of ceramics . Flaws are responsible for the

premature failure of ceramics. For an

appropriate process layout in view of the

workpiece quality knowledge about the

influence of the machining process on

materials strength is a decisive factor.

Fracture of ceramics originate from

machining flaws as well as from material

intrinsic failures. The determined strength

values are generally depend on the

specimen size and number. In this

investigation the bending strength of

silicon nitride samples with dimensions

25.4 X 8 X 4 mm was tested via 4-point

bending test. The measured strength and

Weibull modulus have been corrected to

give values independent of the specimen

number by means of the Maximum- Likelihood-method. The grit size of the diamond

wheels, the cutting speed and the feed

speed were changed to produce groups of

samples with different states of surface

integrity. The achieved variation in the grit

depth of cut is again characterized by the

parameter hcu- influence of grinding

conditions on bending strength and Weibull's modulus is shown in Fig. 10.

179
00 VD

c

OJ

1_

1/1

1 cn

_c

T3C

cu x>

1 c

'o

Q.

1 -Jchip thickness

grit size cutting speed

feed speed

material SijN, wheel 6A2K+8821 JYClOO

truing SiC-whieel dressing alumina steel coolant oil

34/57/.2 IFW 8475

Fig. 10: Influence of grinding conditions on bending strength and fracture

probability of silicon nitride

21. /5947 © IFW 8 A 7 6

Fig. 1 1

Intrinsical material flaw in silicon nitride

180
material wheel

truing dressing

coolant

L point bending - strength ||

Si3N,

6A2-D181K-f 8821 JYC100

SiC - wheel

olumina stick

oil = 35 m Is

V, = 12 mm/min

12U MPa

3i /59A6 IFW 8 i 7 9

Fig. 12: Machining damage in silicon nitride

Strength values remain on the same level

almost independent of the parameter

variation. At high values of hcu a decrease

of the Weibull's modulus and successively

a slight decrease of bending strength

indicates an influence of machining. The

reason for the almost unimpared strength

of the material at low grit depth of cut is the

presence of volume defects. The defects

act as crack sources thereby limiting

materials strength. These material intrinsic

failures lead to a bending strength

unaffected by the machining conditions,

Fig. 11.

The decrease of Weibull modulus and of

bending strength at large values of hcu

points to an influence of machining. At high

grit depth of cut flaws (Fig. 12) are generated in the surface near region which

exceed the dimensions of intrinsic failures

causing an increase in fracture probability.

This grit depth of cut should not be

exceeded in grinding to avoid any

decrease fracture strength or component

reliability. Furthermore these results

indicate that the grinding seems to flaw

size exceed the thickness of the layer

under compressive stresses. This may

explain why no strength increasing effect

of compressive stresses were found.

3.4 Optimization of the process

From this study it has become apparent

that it is possible to optimize conventional

ceramic grinding processes which are

characterized by low cutting speeds, low

infeed speeds and the use of fine-grained

diamond wheels. For efficient grinding of

ceramics the workpiece material properties

have to be considered for process layout.

Because of the brittleness of ceramics, parameters must be chosen that avoid

grinding damage. This can be accieved by

evaluating the critical depth of cut for the

181
material. Based on this parameters the

process design towards an high efficiency

can be made considering three basic

process parameters. To reduce wheel

wear in grinding a coarse type of grit

should be applied (D 151 -D 181). High

wheel speeds are capable of limiting

grinding forces. High feed speeds can be

used to reduce machining time and

attritious grit wear.

material

grinding wheel

grinding wheel

dressing

sharpening

coolant

D91 C100 VTA

= 20 m/s

= L, mm/min

V;, =290 mmVmm

D181C100 r>x?xl

= 55 m/s

v,g= 12 mm/min

V^=600 mmVmm

SiC - wheel corundum block

mineral oil

. 80

N

: mm

60

iO

-r

20 -

_ 30

E

cn 20

c

,C^ 10

2i

A

0

it)0 mm'

mm'

300

200

100

3i/5938 ' IFW 8478

Fig. 13: Optimizing the grinding process

Grinding of silicon nitride cutting tool

inserts may serve as an example for a

comparision of a conventional process and

one optimized with respect to the achieved

results (Fig. 13). A conventional process in industry was substituted by an adapted

process with optimized process

parameters. The adaptation of the process

parameters wheel speed, feed speed, and

wheel specification were done based on

the critical grit depth of cut of the material.

The application of this technology reduces

cycle time drastically by more than 60%. In addition the wheel wear is reduced to a

third.

Grinding forces and surface roughness are maintained. If such process design is applied a considerable increase in

efficiency in machining of advanced

ceramics is possible.

Acknowledgement

The authors wish to thank the German

Ministry for Research and Technology and

the industrial partners Blohm Machine

Tools, Winter & Sohn, and Feldmuhle for

their support.

References

(1) Tonshoff, H.K.; Brinksmeier, E.; Wobker, H.-G.: "Evaluation of

Surface Layers of Machined

Ceramics, Annals of CIRP Vol. 38/2,

pp. 699-708, 1989

182
(2) Marshall, D.B.; Evans, A.G.; el. al.: "The Nature of Machining Damage in Brittle Materials", Proc. of Roy Soc.A.

385, pp. 461-475, 1983

(3) Hakulinen, M.: "Residual Strength of

Ground Hot Isostatically Pressed

Silicon Nitride", Journal of nat. science. 20, pp, 1049-1060, 1985

(4) Tonshoff, H.K.;Roth, P.:"Chip

Formation and Material Removal in Grinding of Ceramics" Proc. 4th. Int. Grinding Conf. Dearborn,

9.-11.10.90, MR 90-539, pp. 1-18,

1990

(5) Rosmanith, H.P.: "Fracture

Mechanics", Springer Verlag.

Heidelberg, 1982

(6) Allor, R.L.; Baker. R.R.: Effect of

grinding Variables on Strength of Hot

Pressed Silicon Nitride, ASME Paper

83-GT-203, 1983

(7) Brinksmeier, E.; Siemer, H.; Wobker,

H.-G.: Requirements on Nondestructive Testing Methods after

Machining of Ceramics", 3rd. Int. SympQSiQm en Nondestructive

Testing . DG28 P, Saarbrucken,

3.-6.10.1988

J (8) Malkin, S., "Grinding technology",

Ellis Horwood Limited City, 1 989

II (9) Wobker, H.-G., "Schleifen

keramischer Schneidstoffe", Dr.-lng.

Diss. Universitat Hannover, 1992

EFFECTS OF CERAMIC MICROSTRUCTURAL

CHARACTER ON MACHINING DIRECTION - STRENGTH ANISOTROPY

A literature review as well as substantial new experimental data are used to better

define important microstructural parameters impacting the effects of machining direction on room temperature uniaxial flexural strength of ceramics. It is shown that grain size has

a pronounced effect, with fine grain bodies showing the greatest strength anisotropy; i.e. strengths are commonly -50% for flexure bars ground perpendicular versus bars ground

parallel with the tensile axis. However, as grain size increases, this strength anisotropy

first decreases, approaching zero at intermediate grain size (e.g., 20-50 |im) then

increases again as grain size further increases, extrapolating to the range found for single

crystals (which vary due to the effects of crystal orientation and machining direction relative to the tensile axis and preferred fracture planes). These observations are consistent with

machining flaw sizes being approximately constant as a function of microstructure for a

given machining condition, such that at intermediate grain size, flaw sizes approach the

grain size, which then becomes a constraint on the ultimate flaw size until grain sizes are

sufficiently large that failure occurs before flaws grow to the size of the grains. It is also

shown that relatively fine homogeneous porosity has no significant effect on strength

anisotropy from machining direction. However, increasing the pore, pore cluster or other

defect sizes reduces strength anisotropy as a result of machining direction, with the

anisotropy becoming zero when the scale of such defects is > that of machining flaws so such defects dominate failure. Grain size effects on the anisotropy are nearly universal,

i.e. most ceramics fall within a limited band as a function of grain size when effects of other sources of failure, e.g. pore clusters, are eliminated. Some ceramic particulate composites,

whisker toughened AI2O3 and microcracking materials showed possible evidence of intrinsic reduction of strength anisotropy as a function of machining direction, mainly in toughened

ZrOg materials.

R. W. RICE

W. R. Grace & Co., Columbia, MD

This author accidentally discovered the effect

of machining direction on uniaxial flexure (a)

strength while studying surface finish effects on

ceramics. Unusually wide strength scatter of machined polycrystalline AI2O3 bars was found to separate into two groups showing normal scatter, with the average strengths of the two groups

differing by nearly a factor of two. More detailed

evaluation revealed that the two groups correlated

exactly with which direction the tensile surface of test bars had been machined relative to the bar axis. This had come about as a result of the machinist

accidentally rotating the rectangular plate from which

the bars were subsequently cut 90° when he turned

1. Introduction it over to grind the second side after grinding the first side so one side of the test bars were ground

perpendicular, the other parallel with the bar axis. This strength anisotropy with machining direction

discovery led to a detailed study of strength

anisotropy as a function of machining direction

relative to the tensile axis for a wide variety of ceramic materials ranging from glassy carbon and

soda lime glass to single crystals (AI2O3, MgAlgOJ,

polycrystalline oxides (AI2O3, MgO, Zr02, and

MgAlgOJ and non-oxides (B^C, SiC and Si3N4).'^'

Limited tests also revealed similar, but somewhat

less effects as a result of sanding with SiC abrasive

paper. Examination of the literature revealed a

precursor observation by Mould and Southwick,*^'

who noted (as an aside to static fatigue studies) that 185
abrading glass specimens with 150 grit, emery cloth perpendicular to the tensile axis, resulted in strengths averaging -50% lower than for abrasion

parallel with the axis of subsequent flexure testing.

Also, in the same time frame as the present author's

initial work, Kirshner et al'^' independently showed

strengths of small (-3.6 mm dia.) extruded rods of a

commercial, sintered, 96% AI2O3 (G ~5nm) were

lower when the rods were circumferentially scratched

with a single point diamond versus axial scratching

(Fig. 1A). The strength difference

increased with increasing load on the diamond, due

mainly to an initially faster strength decrease of circumferentially scratched specimens. An

approximately constant strength differential was

reached at higher loads where strengths of circumferentially scratched rods were -45% of those

for axially scratched a. Greater depth of cut (i.e. of infeed) of the diamond point rapidly decreased

strengths (Fig. 1 B).

Q.

400 -

350

5 300 h

liJ cc

cn

_i

<

ECZ3X

m

250 -

200 -

150 -

100

200 400 600 800 0

LOAD ON DIAMOND POINT (grams)

0.001 0.002 0.003 0.004 0.005 U- INFEED DEPTH IN SINGLE POINT

^'^^^ DIAMOND CIRCUMFERENTIAL SCRATCHING

Fig. 1 Strength data of Kirchner et al'^' for extruded, 96% AI2O3 rods tested at 22°C as a function of scratching

with a single point diamond. A) Data for scratching circumferentially (open squares) or axially (solid

squares) as a function of load on the single point diamond, B) Data for circumferential scratching as a

function of depth of diamond cut (infeed).

Subsequently, this author extended his work

to include another silicate glass as well as a sulfide

glass, stoichiometric MgAl204 crystals (versus

previous AljOg-rich MgAl204 crystals) and a TiC

crystal.''" Several investigators subsequently

confirmed substantial strength anisotropy as a result of machining direction in various high-strength SigN^

bodies over a limited range of typical machining

parameters (e.g. see refs. 5 and 6 as examples).

The present author and colleagues showed this strength anisotropy arose primarily, if not exclusively,

from differences in the geometry of flaws generated

parallel to the machining direction versus that generated perpendicular with the machining

direction. The depth of both flaw populations

tends to be similar and does not change significantly

with grain size (G) (including single crystals, i.e. G=oo). However, while flaws generated

perpendicular to the abrasive direction (which

dominate failure for stressing parallel to the

machining direction) typically have an approximately

half-penny shape, those generated parallel to the

abrasive direction (which dominate failure for stressing perpendicular to the machining direction)

typically have a substantially elongated shape.

Thus, flaw depth to half length ratios were less for

failure initiating flaws from machining perpendicular versus parallel with the bar axis.

Much more recently, some work has been

done to address important parameters affecting this strength anisotropy. Thus, Ota and Miyahara'^^'^^'

report greater strength reductions for dense SigN^ as

a result of changing diamond grinding wheel grit size, due mainly to greater strength reductions for grinding perpendicular to the bar (and hence tensile)

axis than for grinding parallel with it. Their work

differs some from, but is basically similar to, work of 186
Kirschner et al'^' (Fig. 1) and Anderson and Bratton*^'

(Fig. 2). Rice and Mecholsky showed diamond

sawing resulted in similar but somewhat more

severe strength anisotropy effects (while earlier studies have shown similar though possibly

somewhat less effects of sanding versus grinding).'^'

It was also clearly shown that failure of diamond sawn specimens occurred from flaws associated with more serious sawing scratches. More recently, Rice

has shown that: 1) the strength anisotropy as a

result of machining direction is a significant function

of grain size, with a definite anisotropy minimum at intermediate grain size*^^' (e.g., see Figs. 3-5); and

2) relatively fine homogeneous porosity by itself has

no significant effect on strength anisotropy as a

result of machining direction.'^*'

ABRASIVE IVIESH SIZE

200 120

s

X 600

1-z

lUfE 500

U)

111 eci

400

1 1 1 Grinding

II i X Anderson & Bratton + 0 • Ota & Miyahara

° °' o —0 ^"^^ -

1 1 J*-—1 1 1 1

90

50 75 100 125 150 ABRASIVE GRIT DIAMETER (urn)

composites are evaluated. Ramification and

application of these results are also discussed.

2. Experimental Procedure

A variety of experimental and commercial

materials (Table 1) were cut into bars ground parallel

(on one side) and perpendicular (on the other side)

relative to the bar length and subsequent tensile

axis. Following sawing, all grinding was done with a 400 grit resin bonded diamond wheel with 1 00%

concentration, (15.2 cm. dia. 1.3 cm in width)

operated at 3,000 rpm (i.e., a surface speed of 143

m/s) with an in-feed rate -7.6 cm/sec and 8|im depth

of cut per pass, using a synthetic coolant.* Such

grinding typically involved -30 passes with a cross feed of - 8mm at the end of each pass. Because of the variable size samples available, resultant bar

dimensions were also varied, but typically falling in the range of width to thickness ratios of 1.5-2. Three-point flexure testing in air at 22°C was utilized with fixed steel supports and spans of 1 .25-2.0 cm

and typical span to thickness ratios of 5-7. Subsequent to testing, extensive fractography was

carried out to determine fracture initiation location,

flaw character and microstructure. Keeping the key

bar test ratios - constant as well as the extensive

fractography minimizes limited effects of varying bar

and test dimensions. Further, the primary focus of

this paper is the relative difference in strength** for the same bars ground parallel with the tensile axis on one side and perpendicular on the other side.

3. Experimental Results and Discussion

jii Fig. 2 Room temperature strengths of dense SigN^

' samples for machining perpendicular and

parallel with the tensile axis and the ratio of

Ij the strength-ground perpendicular to that ground parallel with the tensile axis as a

function of abrasive grit size. Data after Anderson and Bratton'^* and Ota and

Miyahara.*"'^"'

This paper further extends data on the

effects of grain size and porosity on the strength

anisotropy as a result of machining direction,

I including a broader range of material compositions

and microstructures. Further, significant extension

I occurs as a result of evaluating effects of ceramic

toughening failure mechanisms on the strength

anisotropy from machining. Thus, in contrast to

I previous studies being limited to "monolithic"

( ceramics, a number of ceramic composites including

various particulate composites with different

; toughening mechanisms as well as some whisker

Both the strength and fractography results are summarized in Table 1.** The present AljOg,

MgO, ZrOj, MgAl204, and B^C*** results further document and extend previously shown grain size

effects on strength anisotropy** as a function of

* Doall Kleen Kool synthetic #01958.

** Throughout the rest of this paper, the measure of strength anisotropy will be the variation of the ratio of the strength tested perpendicular to the machining

direction (al) to that tested parallel to the machining

direction (ct||) from a value of 1 or 100%. This is similar to using the difference in strengths, e.g. Gli- al.

*** Note that greater scatter in the CeO^ and

especially the B4C data is due to grain size

variations in some of these bodies as well as less

definition of grains in the 100% transgranular

fracture (e.g. see Fig. 8).
machining direction. Thus, the greatest strength

anisotropy (aJ_/a|| of ratios) of 40-60% occurs at fine grain size, with less anisotropy as grain size

increases, approaching zero (i.e. cUg\\ ratios of -100%) at intermediate grain size, then increasing to the single crystal values as grain size further increases (Figs. 3-5)). That a minimum anisotropy occurs at intermediate grain size is shown clearly by

the AI2O3, MgAl204 (Fig. 3), MgO, and cubic zirconia

(CZ, Fig. 4) and B4C (Fig. 5) data, i.e. where there

is more data at both moderate and larger grain sizes

and for single crystal orientations relative to stressing

and machining directions (their average is appropriate for comparison to polycrystalline

behavior.) Similarly limited results for the other

normally single phase materials support this trend,

e.g., the boron'^^' results (where SEM examination

showed no grain structure, but it is reported to be

~2|im'^^'). Thus, except for some materials showing

less anisotropy, most materials follow a common

trend, and, in fact, fall within a limited band (e.g.

Figs. 3-5). The possibly somewhat lower strength

anisotropy observed in this study (e.g. Fig. 3) versus

the previous studies of this author is consistent with

reduced strength anisotropy (i.e. higher aiVall ratios)

with less severe machining conditions, mainly due to less reduction in strength for machining

perpendicular to the tensile axis (e.g. this author's

previous studies utilized a removal depth of ~50|j.m

per pass vs. 8|im here).

The rationale for plotting the ratio of the

strength (a) for grinding perpendicular versus that

parallel with the bar axis as a function of G"'''' is to emphasize the relationship of the grain size

dependence of this machining effect to the basic ct- G

'''' behavior of ceramic materials. It has been

extensively shown elsewhere that ceramics failing

from flaws follow a two-branch behavior on a a-G

''''

plQ^ (13.16) ^j^g \arge grain size regime, there is a

substantial grain size dependence of strength, but at intermediate grain size this switches to either a zero

or, much more commonly, a modest grain size

dependence of strength. The underlying mechanism

of this behavior arises from flaw sizes from

machining typically not changing much with grain

size. Thus, in the fine grain size regime, one has

flaws which are substantially larger than the grain

size, while in the large grain size regime flaws reach

catastrophic failure before they grow to the grain

size. Thus, the two branches of the ct-G"'''' behavior

intersect when the flaw size is approximately equal

to the grain size. The fact that the strength ratio for grinding perpendicular versus parallel with the tensile

axis reaches a maximum (i.e. minimum anisotropy)

at intermediate grain size (e.g. 20-50 |im) is

consistent with this mechanism. This implies that as

the flaw size and grain size approach one another as the grain size becomes a constraint to flaw

dimensions for the more elongated flaws associated

with specimens failing from machining flaws as a

result of machining perpendicular to the tensile axis. Some fractographic support for this observation has

been previously presented, and is provided by this study as discussed further later.

'^'^^^

This study further shows that relatively fine,

homogeneous porosity has no obvious effect on

strength anisotropy as a function of machining

direction. However, increasing heterogeneity of porosity increasingly reduces intrinsic machiningstrength anisotropy (e.g. from finer grain sizes).

Thus, comparison of various bodies of B, AI2O3,

MgO and MgAlgO^, with generally fine (e.g. <1nm)

homogeneous pores showed no significant strength

anisotropy differences from other bodies of the same

material with much less porosity (Table 1 , Figs. 3, 4). For examples of somewhat homogeneous porosity see Figs. 6 & 7. As the size of the pores or pore

clusters increased, the strength anisotropy as a

function of machining direction decreased, becoming

zero (i.e. gUoW = 1) when the pores or clusters

became the flaws, e.g. CeOj (Fig. 8), ZrOj (Fig. 9),B

(Figs. 10, 11), and aluminum titanate-zirconium

titanate (AT-ZT) composites. These trends are

further illustrated by the actual strength-porosity data

(Figs. 12, 13). Materials with more (but not

perfectly) homogeneous porosity show nearly (but not completely) parallel trends for strengths from

parallel (a||) and perpendicular (al) grinding vs. porosity (e.g. AI2O3, Fig. 12), those with intermediate

homogeneity of porosity less parallelness of all and

al vs. porosity (e.g. MgAl204, Fig. 12), and those

with low homogeneity of porosity still less

parallelness of the two strength-porosity curves e.g. B, Fig. 13. In all cases, the reduced parallelness is due to a reduced slope for the strengths of perpendicular bars vs. porosity.

A key question is whether the reduction in strength anisotropy from machining direction is due

to interaction of the machining flaws with the pores

to cause failure or due to the pores acting

independently as sources of failure, hence overriding

the machining effect. The above strength-porosity

results (Figs. 12,13) can be interpreted as implying

that strength anisotropy decreases due to increasing

machining flaw-porosity intersections as pore size

and heterogeneity increases. This is supported by

the reduction in parallelness of the parallel and

perpendicular ground strength vs. porosity curves as
Table 1

Flexure Strengths of Ceramics Ground Perpendicular

vs. Parallel to the Tensile Axis

Material

Procedure/

Source'" G (nm)''' P(%)'^'

w, t, Span'"' (mm)

al'='

all

01'='

Fracture Origins'^'

1. Oxides

AI2U3 HP (LA)/B

(A- 1-39)

4 0.5 5.0, 2.5, 15.0

(6) 330120 (5)

(8) 370130 (7)

90 5SLG, 1CLG

7SLG, 1CLG

AI2U3 HP (LA)/B

(A-3-30)

1 1.5 (5) 310140 (11)

(5) 6201100 (16)

51 2MF, 2MF?

3C, 1ILG &P's

AI2U3 AI n HP (LA)/B

(A-3-1)

0.7 6.0 (5) 250130 (14)

(6) 5851105 (18)

42 4MF

20, 11 PA, 11 PALO, 1SPA

AI n HP (LA)/B

(A-3-39)

1 16.5 - (5) 230175 (32)

(4) 3351145 (44)

69 (51) 3LSP, 2MF

2 LSP, 1IPA

MI2U3 AI O HP (LB)/B

(A-3-1 1

1.5 1.0

- (7) 320130 (10)

(5) 6401100 (16)

50 3MF, 2MF?, 11-SP 2MF?, 10, 10?

AI n HP (LB)/B

(A-3-42)

2 0.8 (7) 415140 (10)

(7) 6401105 (16)

65 (59) 5MF, 2MF?

30, 3MF?

AI2U3 AI n HP (LB)/B

(A-3-44)

1 1.2 (7) 295170 (23)

(6) 4651120 (26)

63 (63) 4MF, 2MF?,

11 PA?

10, 11-SPA, 1ILG

AI r\ AI2U3 (P-88) S/Diamonlte

5 5.2, 2.5, 19.0

(5) 310130 (10)

(4) 380130 (9)

81 IMF?

none identified

A 1 r\ AI2U3 (P-99) S/Diamonite

5 5.2, 2.4, 19.0

(5) 305135 (12)

(5) 315+35 (12)

98 porous areas?

AI2O3 S (GE (Lucalox) 30 -0 5.2, 2.5, 12.7

(6) 260135 (14)

(6) 275130 (11)

95 none identified

Ce02 S/Boeing -30 6.6 5.1, 2.5, 15.2

(5) 80110 (10)

(5) 80110 (9)

95 (99) 5SPA

3SPA, 1SPA+LG,

10

oeu2 S/Boeing -30 6.6 2.4, 1.5, 12.7

(4) 95110 (10)

(4) 120145 (39)

80 All SPA

MgO HP/Boeing 0.6 -0 4.0, 1.8, 12.7

(5) 185110 (12)

(5) 340155 (16)

54 (51) 4MF?

40, 1MF?

MgO HP/Boeing

(M-5-14)

0.5 4.5 5.5, 2.6, 12.7

(3) 105115 (15)

(3) 165130 (18)

62 none found none found

MgAljO, S/Boeing

(M-5-13)

8 26 5.3, 2.5, 15.2

(6) 120120 (18)

(6) 130110 (9)

89 none identified

MgAl20, S/Ohio State -10 -3 4.9, 2.4, 12.7

(4) 225125 (11)

(4) 265130 (11)

85 2SP, 2SPA

1SP, 1IPA?

Zr02+5 w/o

Y2O3

Crystal from

skull melting

00 0 3.5, 1.5, 15.2

(6) 17101235 (14)

(5) 17401255 (15)

98 IMF, 10

2MF

ZrO2+20

w/o Y2O3

0 4.2, 2.2, 15.2

(5) 195150 (26)

(5) 235135 (15)

83 4MF

5MF

189
Material

Procedure/

Source'" G (nm)"' P(%)<''

w, t, Span'"

(mm) oil

01'=' o||% Fracture Origins'®'

Zr02+11.2 W/O Y2O3

HP/Zyttrite

(WPAFB)

20 -1 5.3, 2.6, 15.2

(5) 230±20 (80)

(5) 275±50 (18)

84 2PD, 1MF, IMF?

2SPA?, IMF

Zr02+2.9 w/o CaOS/Zircoa-C 30 (15-60) -7 5.1, 2.6, 15.2

(6) 170±15 (8) (6) 165±10 (6)

101 none detected

Zr02+2.8 w/o MgO

S/Zircoa-1027 30 ~5 5.8, 2.9, 20.3

(5) 300±25 (9) (5) 320+25 (8)95 none detected

Zr02+6 w/o Y2O3

S/Union Carbide -0.5 -2 5.8, 1.7, 15.2

(4) 365185 (25)

(3) 545±65 (12)

67 (73) ISA, ISP, 1SPA

11-SP, 1SPA

Zr02+-3

m/oYjOa

S/Kyocera 3 -5 5.2, 2.6, 15.2

(6) 60±5 (8)

(6) 60±5 (15)

104 none detected

ZrO2+-10

w/o mixed

stabilizers

S/Diamonlte

Zirmonite 1000

2 ~5 5.2, 2.5, 19.0

(6) 710165 (9) (5) 6701110 (16)

105 3SPA, 1CPA, 1C

3SP, IIP, IMF?

ZrO2+~10

w/o mixed

stabilizers

S-HIP/

Diamonite

Zirmonite 1000

2 1-2 4.7, 2.3, 12.7

(6) 850135 (4) (6) 835150 (6)

102 1C, ISPA, ISP?,

11-SP, IMF?.

1C, ISPA, 1IS- SPA

II. Non Oxides

• B, Borides

Boron HP (383) 3.2 5.1, 2.3, 12.7

(4) 4601140 (30)

(4) 5401130 (25)

88 (67) 3MF, 1C

2MF, IMF?,

1MF-P

Boron HP (560) 4.5 (5) 340120 (6)

(4) 440185 (19)

77 3MF, 2MF?

2MF, 1MF-P, 1C

Boron HP (566) 9.8 (5) 320130 (10)

(5) 3901100 (26)

82 (68) 3MF, ISP, 1C

2MF, ISPA, 1C

Boron HP (594) 21 (5) 245135 (13)

(3) 240130 (12)

101 ISPA, 1CVC?,

CPA, 1IJPA,1iPA,

11-SPA, 1C-IPA

Boron HP (584) 29 (5) 1901100 (26)

(3) 190170 (38)

98 1C?, ISP?

ISP. 11-SP?

TiBj S/PPG (1900) 3 2 4.8, 2.3, 15.2

(1) 343

(1) 249

73 1C?

IMF?

TiBj S/PPG (1900) 3 2 3.4, 1.3, 12.7

(7) 430160 (14)

(7) 385165 (16)

112 1PD, ISPA, 1IPA, 1MF?,1C?

31-SP, 1C

TiB2 HP Ceradyne 15 -0 2.5, 5.3, 15.2

(5) 150185 (58)

(4) 135165 (48)

115 None found None found

TiB2 S/PPG (2100) 5 -0 5.2, 3.2, 17.8

(5) 355130 (8)

(4) 340120 (6)

105 None found None found

TiB2 S/PPG (2250) 150 -0 5.2, 2.8, 15.2

(3) 4717 (15)

(3) 4913 (6)

96 None found None found

• C, Carbides

Graphite /Unknown - Reactor Grade

20 21 5.0, 2.5, 15.2

(5) 3013 (10)

(4) 3311 (40)

90 None found None found

190
Material

Procedure/

Source'" G (nm)<=' P(%)">

w, t, Span'"' (mm)

ol'^'

oil

ol'^' o||% Fracture Origins'®'

Graphite /POGO 1-2 16 5.2, 2.6, 15.2

(8) 130±4 (3)

(8) 133±4 (3)

99 3C, 2C?

1C, 3C?

B,C HP/Union

Carbide

5-40 ~0 12.7, 5.3, 1.8

(8) 225±85 (38)

(8) 345±135 (39)

66 (54) 2MF, 1PD, 1C

1SPA, IMF, 1C,

1C-LG

B,C HP/Union

Carbide HT

20-100 -1 12.7, 4.3, 1.8

(8) 305±80 (27)

(8) 340±55 (17)

84 (80) 3MF, IMF?, 1SPA

1SPA, IMF?

HP/Carborundum

5-10 -2 N (8) 360120 (6) (7) 450180 (18)

80 2MF, 1C?

2C, 1C? IMF?

B,C HP/ESK 10 -0 - (8) 275+25 (9)

(8) 300135 (11)

93 (93) 1SPA, 1CPA,

IMF

1SPA, 1IPA, 1LG

B,C HP/ESK, HT 10 -0

;

(6) 3901100 (25)

(7) 435+80 (18)

90 (90) 1C, IMF,

1MF+LG, 1LG

1LG, 1IPA

B4C HP/FMI 2 -0 4.2, 1.5, 12.7

(3) 305150 (16)

(3) 5001120 (24)

61 IMF, ISP, 1C?,

15LG? 2C

B,C S/FMI -50 -7 3.1, 1.2, 12.7

(6) 190115 (8)

(6) 190+10 (5)

100 ISP?, 11-SP? ISP?, 1SPA?

Bfi HP/Borid Prod. -15 -0 4.5, 2.1, 1.3 (4) 415120 (5) (4) 570150 (8)

73 2MF, 1C?

2MF, 1PD, 1C

SiC HP/Cercom 5 -0 5.1, 2.4, 12.7

(6) 495155 (11)

(6) 530145 (9)

94 3C?, IIP?, IMF?

4C?, IMF?

TIC HP/Ceradyne 20 -1 5.2, 2.5, 12.7

(5) 365180 (22)

(4) 405175 (18)

90 None found None found

ZrC HP/Ceradyne 15 -2 5.3, 2.5, 15.2

(4) 260125 (10)

(3) 265120 (8)

98 IMF, 1C?,

1SPD?, 1CPA?

1CP?, 1IPA?, 1IPD?

• Nitrides

AIN S/Dow 4 SI 6.1, 3.1, 20.3

(7) 320130 (10)

(8) 335135 (11)

96 none detected

AIN HP/Grace 15 -0 4.2, 1.7, 12.7

(6) 235120 (8) (7) 260115 (5)

90 none detected

BN HP/Union

Carbide

0.5 thick 3 dia.

~1 4.9, 2.3, 12.7

(6) 7016 (8)

(6) 85115 (17)

82 none detected

SI3N, S/Diamonite -0.5 -2 4.9, 2.5, 19.0

(5) 640140 (6)

(4) 914165 (6)

70 4MF?

1SPA, 11-SPA; 1IPA

SiaN, -4 -7 4.6, 2.7, 12.7

(5) 340120 (6)

(5) 335145 (14)

102 2CP, 11-SLG's, ISP, 1PD

2C, 1C?, IIP, 1SPA, 1C-LG-P?

HP/Diamonite

(SN2)

4 -2 4.6, 2.8, 12.7(6) 370155 (15)

(6) 425125 (6)

88 (92) Specific origins not detected

III. Composites

i;
Material

Procedure/

Source'" G (^m)'^' P(%)'^'

w, t. Span*'' (mm) oil o||% Fracture Origins'®'

AI2O3+ 27V/0 TiBjHP/Grace ~0 11.5, 5.7, 12.7

(5) 600130 (6)

(5) 555±60 (11)

108 LAI2O3 G clusters?

AI2O3+

42v/o TiC

HP/Grace -0 (4) 360±30 (8) (4) 365±40 (12) 99 All PA'S?

AljOj+TiC

-20v/o

HP/Valeron 1-2 ~0 5.8, 1.8, 12.7

(5) 390±50 (12)

(5) 560+60 (11)

70 3MF?

2C?, 1PD?

MgAljO^+a

1 v/o SiC

HP/Grace ~0 (4) 275±20 (7)

(4) 300+40 (13)

91 All PA'S?

Ti02+25%

AI2O3 HP/Ceradyne 20 5.0, 2.3, 15.2

(4) 125+5 (4)

(4) 140+15 (9)

91 All probable PA's

Pyroceram F/Corning 0.5-3 -0 5.2, 2.6, 15.2

(5) 165±5 (3)

(5) 220+40 (17)

75 IMF?

2C, IMF?

AI2O3+26 w/o ZrOjS/Diamonite

Zirmonite 200

2-3 (Zr02

0.5-1)

-6 5.2, 2.4, 19.0

(3) 450+20 (4)

(4) 390+145 (37)

115 (90) 1TSP?, 1SP,

1SPA

10

AI2O3+26 w/o ZrOjS/Diamonite

Zirmonite 200

- 6 5.2, 2.6, 19.0

(5) 510+45 (9)

(5) 555+55 (10)

92 2SA, 2SP, 1SPA

3SA, IMF?

AI2O3+26 w/o ZrOjS/Diamonite

Zirmonite 200

1 (ZrO^ -

0.5)

u (2) 605±10 (2) (6) 625±105 (16)

96 (92) 2MF?

2SP, 2MF?, 11-CP

AI2O3+26 w/o ZrOjS-HIP/Diamonite

Zirmonite 200

1 -1 5.2, 2.5, 12.7

(6) 830195 (11)

(6) 9001200 (22)

92 (87) 2SP, IIP, IMF?

3IP, ISP, IMF?

AI2O3+7 w/o SiCwHP/Cercom

(Disl<)

-2 -0 5.1, 1.9, 12.7

(4) 610150 (8)

(4) 7151125 (18)

85 1ILG?

20?, 1SLG?

AI2O3+25

v/o SiCw

HP/Cercom

(Plate)

-1 -0 5.1, 1.7, 12.7

(5) 420140 (10)

(5) 560155 (10)

75 11-SLG, 10?,

1PD?

20?, 1SLG?

70AT30ZT'" (A)S/Grace 4 11 5.2, 2.4, 12.7

(6) 75+6 (7)

(6) 75+6 (8)

96 none found

60AT40ZT"'' (B) S/Grace 2-10 7

12.7

(D) OOlO (0)

(6) 9517 (8)

by none found

40AT60ZT'" (D)S/Grace 4 23 5.2, 2.5, 12.7

(6) 6516 (9)

(6) 6516 (9)

96 none found

Footnotes:

(1) F = fused (i.e., melted), HP = hot pressed, S = sintered, HIP = hot isostatically pressed.

(2) Average diameter of grains measured on a fracture surface.

(3) P = porosity.

(4) w = bar width, t = bar thickness, span = test span in 3-point flexure.

(5) al = average strength for grinding perpendicular to bar axis shown over average strength ground parallel to bar axis (a||). First number (in ()) is the number of tests averaged, next is the average strength and

standard deviation, both rounded to the nearest 5 MPa (unless <10 MPa), and the last number (in ()) is the coefficient of variation. Note that the data were measured in psi units, so conversion to MPa and

192
rounding to the nearest 5 MPa may give slightly different coefficients of variations than shown (from the

original psi), and from the aJVall ratio. a±\G\\ ratios in ( ) are corrected values due to unusually low

strengths or failure from gross defects.

(6) Fracture origin designations: Locations: C=from (rounded) corner of fracture, i.e., rounded bar edge; I = interior of bar, i.e. not at surface, S = bar surface, f-S = source not at, but sufficiently close to the

surface to be uncertain whether it was an internal or surface (connected) origin. Types: LG = large grain;

MF = machining flaw; P = Pore (P's = a few pores); PA = porous area (i.e., cluster of pores). Possible

combinations shown by - or &, and uncertainty is designated by ?.

(7) AT = aluminum titanate, ZT = zirconium titanate.

GRAIN SIZE (G, [im)

OO 50 20 10 4 2 1.5 1.0 0.7

30 ^

1 \ I I \ 1 \ I I I I i I 10 20 30 40 50 60 70 80 90 100 110 120

0

-1/2 r- 1/0 ^^^^^ (GRAIN SIZE) (1/n/G , cm'^^^

)

Fig. 3 Strength anisotropy, i.e. the ratio of the strength (at 22°C) ground perpendicular versus parallel with the

tensile axis, as a function, the inverse square root of grain size for AlgOg and MgAl204. Note strengths

for several orientations for crystals of each of the two materials, specific orientations are shown for stoichiometric Czochralski-grown MgAl204; all other crystal data points are for different orientations cut from Vernuil-grown boules of each material. Data for both the present study and previous studies*^ "' are shown. Numbers adjacent to some data points are the volume % porosity where >0.5%. Note the trend

lines shown for the materials, including a possible single trend for both materials (e.g. the dashed line), which are consistent with those of most other materials (Figs. 4 & 5). -193
Fig. 4 Ratio of the strength (at 22°C) of other oxides ground perpendicular versus parallel with the tensile axis as a function of the inverse square root of grain size. Both current and previous data shown, including

new data for both fully and partially stabilized ZrOg single crystals. Data for both the present study and

previous studies*^ "' are shown. Numbers adjacent to some data points are the volume % porosity. Trend

lines shown for some of these materials. Note the peak of the MgO data at intermediate grain size being

>100% is attributed to surface work hardening effects.'^^'

porosity heterogeneity increases. While the

fractography results are clearly not unequivocal, they

both directly and indirectly indicate that at least part

of the reduction of strength anisotropy as a function

of machining direction results from interaction of the

machining flaws with the larger pores or pore

clusters available for failure. First, where machining

flaws were found in AI2O3 of different levels of porosity they were similar (Figs. 14, 15). Second,

some strength anisotropy is typically observed, even

in sample populations in which failure is predominantly associated with isolated larger pores

or pore clusters. Although machining flaws generally

could not be specifically identified in most porous

samples, they must be interacting with the pores, since the pores themselves show no basis for providing strength anisotropy in contrast to

machining flaws clearly having this effect. Further,

previous cases of specific interactions of machining

flaws with larger pores have been shown in glass*^^'

and polycrystalline specimens.'®' It is also true that some of the materials of this study having pore

clusters as a major factor in their fracture initiation had large grain size, e.g. particularly clear for the case of CeOg (e.g.. Fig. 8). This may also have

been a factor in the PSZ material (Zircoa 1027, while

specific fracture origins were not identified in the

samples of this study, previous studies of this material have shown fairly frequent initiation from

grain boundaries with pores). In these cases where pore clusters dominated failure initiation, it is

still quite probable that the grain structure was

limiting the ultimate size of the flaws even where the

pore clusters were the initial flaws or a major portion

194
100 50 20 10

GRAIN SIZE (G, iim)

4 2

110

<

-J3O

o

z

UJ

Q.

OC

LJ

Q.

Q

Z

O

OCo

X

H

o

z

LU

QC

I-0)

100

ill _l

_I

<

OC<

Q.

o

z

Z)

o

OCo

X

I-o

z

LU

DC

h90

80

70

60

50

Ot

1.5 1.0 0.5

At

O

TiBg B4C SiC TiC ZrC Si3N4

0 A 8 0 Present • Previous

{100} <100>

=.0

o

Range for

Si3N4 of

Anderson

Bratton

10 20 30 40 50 60 70 80 90

(GRAIN SIZE)"^^^ , cm'^'^)

-1/2

Fig. 5 Room temperature strength anisotropy of non-oxide specimens ground perpendicular versus parallel with

their tensile axis as a function of the inverse square root of grain size. Data for both the present study

and previous studies*^ are shown Note the TiC single- and poly-crystal data and that the B4C trend line would also be consistent with the other non-oxide materials (as well as most oxides, Figs. 3,4).

of them (i.e., in combination with machining flaws).

Thus, these materials further support the concepts of intermediate grain sizes limiting flaw sizes.

Next, consider the strength anisotropy due to

different machining directions in composite materials

with differing toughening mechanisms. All of these

appear to suggest less strength anisotropy than

would typically be expected for their average grain or

particle sizes. However, in the case of most hot- pressed oxide-non oxide composites (Table 1), fracture origins appear to be predominantly either

larger AI2O3 grain clusters (see Fig. 21) or porous

areas (which, in the case of AlgOg-TiC were often

associated with isolated graphite flakes or clusters of these), in both cases consistent with previous

development and evaluation of these materials.*^^'

Fig. 6 Example of reasonably homogeneous

porosity in fine grain hot pressed AI2O3. SEM of specimen A-3-39 of Table 1 having

-16.5% porosity.

195
Fig. 7 Examples of somewhat homogeneous

porosity in MgAl204. A) lower magnification

SEM photo showing a somewhat

homogeneous distribution of larger pores, many of which are not sufficiently large in size to be failure-causing flaws themselves

and they are sufficiently sparse to not necessarily be frequent sources of failure. B) higher magnification SEM showing finer- scale pores, i.e. approximately the same or

less than the grain size. These photos are

of the MgAip^ (SMA1-7) with -26%

porosity.

Fig. 8 Example of heterogeneous porosity at the

fracture origin of a ceria sample. Fracture

initiated from the collection of larger pores in the bottom center of a photograph. This

specimen failed at 93 MPa. Total porosity in these samples was approximately 30%.

Clusters of porosity dominated failure of this material as shown in Table 1 . However, this particular sample shows a somewhat greater

degree of concentration and consolidation of pores than was true in some other samples.

Fig. 9 Sample of heterogeneous porosity in CaOPSZ. Failure initiated in this Zircoa-C

specimen from the porosity near the tensile

surface (bottom of the photograph) at a

fracture stress of 165 MPa. Failure initiation from such clustered pores (which also

typically limited grain sizes in the immediate

region, as in this photo) was common for

this material. Note larger grains from which

failure did not initiate, e.g., in the bottom

right area of the photo).

196
Fig. 10 Example of extreme porosity heterogeneity

causing failure in a highly porous hot

pressed boron. This sample (584 of Table

1) failed at 215 MPa from the large pore and

surrounding finer porosity in approximately

the center of the photo. While this set of samples contained -21% porosity, typically

much of it was quite heterogeneous such

that larger pores or pore clusters dominated

failure initiation in this material, as was the case in this specimen as well as more

generally shown in Table 1

Fig. 11 Example of substantial porosity

heterogeneity in a hot pressed boron sample

of intermediate porosity (9.8%). This

sample (566 of Table 1) failed at 353 MPa

from the pore cluster shown.

100

22X, Ground: 1| 1

O • AI2O3 MgO

A A MgAl204

MgAl204, 1

b~ 2.8

I 5 10 15 20 PERCENT POROSITY (P,%)

25

G20e

Fig. 12 Strengths for parallel and perpendicular

ground AI2O3 (G~0.7-4|j.m) and MgAl204

(G~6-8|im) samples vs. porosity. Note

these two materials show respectively

nearly, and less, parallelness of the two sets

of curves consistent with less porosity

homogeneity in MgAl204 than AI2O3.

1 r

1 r

22°C, Ground: II L

O • B

10 15 20 PERCENT POROSITY (P,%)

30

G209a

Fig. 13 Strengths for parallel and perpendicular

ground boron vs. porosity. Note substantial

lack of parallelness of the two curves

consistent with the substantial heterogeneity

of the porosity.

197
Fig. 14 Example of the failure initiation from a machining flaw in fine grain, nearly dense

AI2O3. A) and B) show respectively finer and higher magnifications of the failure

initiating flaw from machining perpendicular

to the tensile axis (specimen A-3-42, 0.8%

porosity). C) shows higher magnification of the demarcation between the original flaw

(bottom half, all intergranular fracture) and

the area of failure (top half, mixed inter- and

transgranular fracture). The demarcation

between total intergranular failure in the

original flaw versus substantial transgranular

failure beyond this was characteristic of most of the AI2O3 studied. A primary

exception was extremely fine grain size

materials where there was little or no

transgranular failure. The failure stress was

387 MPa. K|c calculated from the failure stress and the flaw dimensions was -3.0

MPam'"', which corrected to P=0 via e"'*" gives 3.1 MPam''*.

Fig. 15 Example of failure initiation from a machining flaw in a more porous (P=16.5%)

AI2O3 body. A) shows a lower, and B) a

higher magnification of the failure initiation flaw. This sample failed at 280 MPa,

yielding a K,c of 2.2 MPam'^, which corrected

to P=0 via e"^^ gives 4.2 MPam'-^.

An important exception is the commercial hot

pressed (and more homogeneous) Al203-TiC

showing substantial anisotropy. Porous areas also

appear to be the source of failure in the finer,* and

especially the coarser grain graphite and the TiOaAI2O3 composites and hence a possible explanation

for the limited strength anisotropy, as could possible

microcracking due to the larger Ti02 and graphite

grains. Similarly, the AT-ZT composites which have

low thermal expansion due to microcracking, exhibit

quite limited strength anisotropy as a result of

* Note, however, that heterogeneity in the porosity of the finer grain graphite clearly did not appear to be

extreme, questioning whether it was the total reason

for very little strength anisotropy. Whether the soft character, or other factors such as twinning or microcracking play a role in its reduced anisotropy is unknown. Some measurable anisotropy in BN

question this, but tests on dense materials, e.g. CVD

graphite and BN (though complicated by preferred

orientation and colony structures) could be useful.

198
machining direction, especially In view of their fine grain size, as noted earlier. While It may be

tempting to attribute the limited strength anisotropy

due to microcracking, It is equal or, quite probably more likely, that this Is due to the substantial porosity

in view of considerable heterogeneity in both pore

spatial distribution and size. However, the limited strength anisotropy in Pyroceram may be somewhat

less than might be expected for the grain size(s) of

this material and no processing defects were evident nor expected in this material. Finally, the two hot

pressed AljOg-SiC whisker composites may also

indicate some modest reduction in strength

anisotropy as a function of machining relative to pure

AI2O3 of the same grain size of the alumina matrix.

However, note the trend for the strength anisotropy

in these two bodies (and the commercial hot pressed

AlgOg-TiC) to follow trends expected from the AI2O3

matrix grain size.

Trends are similar for various Zr02-

toughened materials. Processing defects (mainly

pores) were the common source of failure in most of the AljOg-ZrOg composites which might explain their very limited a anisotropy as a function of machining

direction. HIPing these Increased strengths and

reduced the severity, but not the frequency, of such

processing defects as fracture origins, modestly

increasing the strength anisotropy. Similar results were obtained with various partially stabilized

polycrystalline ZrOg materials. However, these

showed more evidence of a possible intrinsic reduction in strength anisotropy. There are clearly

other factors limiting strength anisotropy in some of these materials. Thus, the larger grain size and

grain boundary pores of the sintered PSZ are also

likely an important factor in the limited strength

anisotropy as noted earlier. Also, one TZP material was extremely weak (apparently due to moisture

degradation during actual use of the component from

which the material was obtained). More generally,

pores were a common factor In failure, including in the HIPed TZP, thus limiting strength anisotropy as

l[ in other materials. However, three factors suggest

I some intrinsic strength anisotropy reduction In

! zirconia toughened zirconia materials. First and

most significant are the single crystal results, where

j partially stabilized ZrOj shows higher strength and -

, ~0 strength anisotropy versus lower strength and

|i higher strength anisotropy in fully stabilized ZrOj

crystals. Porosity and grain size are not factors in these crystals (both sets of which are believed to have ~<1 10> tensile axes). Second, there is clearly more evidence of strength anisotropy in the TZP

materials, especially at finest grain size (though less

I than in other materials at the same grain size).

Third, HIPing of TZP, though reducing the severity of pore origins, did not significantly increase strength

anisotropy. Surface transformation and resultant

surface compressive stresses In those materials'^"' would be a likely explanation of such effects in ZrOg

toughened materials since such surface stresses would likely inhibit flaw elongation. Greater effects indicated in TZP versus AlgOg-ZrOg bodies would be

consistent with this, since such surface stresses should be greater In the former.

While the materials of this study did not yield

nearly as extensive clear fractography results as some of the materials In earlier study,'®' these results nonetheless are consistent with the flaws from

machining causing failure due to machining

perpendicular to the tensile axis being more

elongated relative to those causing failure as a result of machining parallel with the tensile axis of the test bars (Table 2). Thus, the flaw depth to half length

ratios were typically less for grinding perpendicular versus parallel. The one possible exception was

B4C, where less differentiation In flaw anisotropy is expected because of the larger average grain size

(and less definitive determination of grain size, as

noted earlier). Examples of machining flaws are shown in Figs. 14-18, corner flaws (Fig. 19), and

failure from large grains (Figs. 20, 21). Where flaws were definitive enough for calculating a (e.g.

Figs. 14-18), values were consistent with previous

data.'^^'^^'

Two issues regarding parameters effecting

the strength anisotropy as a function of machining

direction require further consideration. The first is the effect of machining parameters such as abrasive

grit size (Fig. 2). Results of both Anderson and

Bratton'^' and Ota and Miyahara'^^'^^' provide strength

anisotropy data for machining parallel and

perpendicular as a function of abrasive grit diameter.

However, there Is some disagreement in the trends.

This may reflect differences in rate of material

removal, as well as possibly of sample

microstructure. Thus, further work is needed to better define how the anisotropy changes as a

function of both machining and material parameters.

Some work has been conducted to assess

intermediate angles of machining relative to the

tensile axis (showing a nearly linear transition

between al and a||).'^^' The second aspect is the

effect of specimen size relative to grinding wheel

width. With typical strength test specimens, as used

in this and most other studies, specimen width is often substantially less than that of the grinding

wheel, such that in grinding parallel with the tensile

axis the grinding wheel will typically more than cover
TABLE 2

Machining Flaw Depth to Half Length Ratio

Material'" G

(|im)

Flaw

Ratio

al<''

Flaw

Ratio

oil'" B 0.5±0.3

(15)

0.6±0.2

(15)

AI2O3 <2 0.5±0.3

(21)

0.810.3

(1)

B4C 5-50 0.6±0.3

(13)

0.5±0.2

(8) CZ xl 00 0.1

(1)

0.3

(1)

1 ) B = boron, CZ xl - cubic zirconia crystal.

2) G = grain size. Although the G of the B was not seen in examinations of this study, it has been

reported to be ~ 2|im.'^^'

3) Gl and a|| are respectively for specimens tested

perpendicular and parallel to the machining direction. Number of values averaged shown in ( ).

the complete bar width. On the other hand, in grinding perpendicular to the bar axis, multiple

adjacent passes of the wheel over the bar are

required. This thus results in a number of regions in which the machining of the specimen occurred at or near the edge of the grinding wheel, where abrasive

interactions with the workpiece may be different than

under the main part of the grinding wheel. A few

trials have previously been made with substantially

larger test specimens, such that possible edge

effects of grinding wheels would be experienced in both grinding parallel and perpendicular to the

tensile axis.*®' However, these were not definitive since other (e.g., processing) defects became an

increasing factor in the failure as specimen size was

increased. Thus, it is still uncertain at this stage

whether or not machining flaws associated with

abrasive-workpiece interactions near the edge of the

wheel may be different from those occurring from

abrasive particles well away from the wheel edges,

hence another area for further investigation.

Finally, three points should be noted. First, the strength anisotropy is quite important in the

200

design and manufacture of many components.

Second, the flaw shape factors that underly this anisotropy are an important probe to help identify

mechanism of failure, e.g. strength-grain size

relations as emphasized here. However, other uses

should be feasible, e.g. probing effects of surface

stresses, e.g. as indicated for MgO and TZP here.

Third, the grain size trends appear to be very

common, if not universal, when effects of other sources of failure are removed.

Fig. 16 Samples of machining flaws in B4C.

Representative SEM photos of machining

flaws causing failure in test bars (typically

approximately the bottom center of the

photograph). The samples, their failure stresses and the fracture toughnesses

calculated from their failure stress and flaws are as follows: A) Carborundum, 370 MPa,

-1.8 MPa m^; B) Union Carbide heat

treated, 266 MPa, K,c -1.7 MPa m^; C)

ESK, 354 MPa, K,c -4 MPa m"". All samples represent failure from flaws from

machining perpendicular to the tensile axis. Note the essentially exclusive transgranular

failure and that the grain structure is not

especially well defined.
100 Mm

and B) higher magnification of the failure

initiating flaw. Note the less-elongated

character of this flaw relative to that shown

in Fig. 17. Specimen was from the batch

designated 560, with a failure stress of 437

MPa. The calculated from the failure stress and flaw character is -1.7 MPam''^

(correction for the -4.5% porosity versus e"*^ would raise this to -2 MPa m^ at P=0.

25 jUfT

Fig. 17 Example of failure initiation from a

machining flaw from grinding perpendicular

to the tensile axis in hot pressed boron. A)

and B) are respectively lower and higher

magnifications of the failure initiation, with B)

showing details of the elongated machining

flaw from grinding perpendicular to the

tensile axis. Sample 383, failure stress 340

MPa. calculated from the failure stress and flaw character -1 .9 MPa m'"^ (correcting

for the 3.2% porosity using e"'' would

increase the fracture toughness to -2.2 MPa

m'^ for P=0. Note the absence of any

indication of grain structure in this material.

B

10 urn

Fig. 19 Samples of corner failures. A) a ceria

sample (failure stress 74 MPa). Note that while the failure markings clearly show

failure initiation from the rounded corner,

specific size and character of the flaw

causing failure cannot be delineated. B) a

hot pressed boron (560) failing at 460 MPa.

Note a reasonably well-defined flaw

. extending in from the corner of this material

which was generally much better for showing specifics of flaw character.

Fig. 18 Example of a failure initiating flaw for machining parallel with the tensile axis of a

hot pressed boron. A) lower nriagnification

201
Fig. 20 Examples of larger grain clusters at probable fracture origins in AI2O3+27 v/o

TiBg sannples. A) and B) represent two

probable fracture origin areas in this composite material. Note the heterogeneous distribution of larger AI2O3 grain

clusters, some of which are often associated

with one or two moderately sized pores,

e.g., the lower left-hand area of photo B). ^ ^

- < ^ ^

100 urn

Example of failure initiation from larger

grains in hot pressed AI2O3. A) and B)

represent respectively lower and higher

magnifications of the large grain fracture

origin in a sample (A-1-39) failing at 396

MPa. Note that using the large grain as the

failure flaw size yields a K^^ of -2.7 MPa m"'."

4. Summary and Conclusions

Substantial further data on the strength

anisotropy as a function of grain size at room

temperature for grinding parallel or perpendicular

with the tensile axis of specimens has been

provided. This includes a much broader range of materials in terms of composition, crystal structures,

microstructures, and fracture mechanisms. This

clearly reinforces and extends previous findings

showing that grain size is a major variable in this

effect. The greatest anisotropy occurs at fine grain

size (or in amorphous materials) and decreases as

grain size increases, rising to ~ zero anisotropy, i.e. limited or no difference, for strength ground

perpendicular versus parallel with the tensile axis of bars at intermediate grain sizes (e.g., 20 to 80 ^m).

It then increases again, extrapolating to the range of single crystal values (which vary with crystal

orientation relative to machining direction due, for example, to orientation of preferred fracture planes

relative to abrasive motions and resultant stress

directions). Fractography corroborates previous

findings that a major factor in the machining strength

anisotropy, i.e. lower strengths from machining

perpendicular to the tensile axis versus machining

parallel with the tensile axis, is greater elongation of flaws formed parallel vs. perpendicular to the

machining direction (which dominate failure respectively for stressing perpendicular and parallel to the machining direction). The concept of the

minimum strength anisotropy being due to the grain

size constraining the flaws, hence first and more

extensively constraining elongated flaws from

perpendicular machining, is further supported. Thus,

data shows a minimum in anisotropy at intermediate

grain size unless greater anisotropy at larger, and

especially finer, grain size are reduced or overridden

by heterogeneous larger grains, porosity or other

processing defects (e.g. as shown by fractography).

This study also further reinforces results showing that fine homogeneous porosity has little or no effect on strength anisotropy as a function of machining direction. However, the presence of larger pores or pore clusters which are frequent

origins of failure can significantly reduce the

machining anisotropy by being alternate sources of

failure to the machining flaws, or combining with

machining flaws, thus also limiting anisotropy.

Finally, evaluations of the machining anisotropy in various ceramic composites indicate that some of these may provide some intrinsic limitations on

strength anisotropy as a function of machining

direction, with the most probable of these being TZP

materials. However, even in these materials, both

202
grain size and porosity also have some impact on

strength anisotropy similar to other materials.

Acknowledgement

The author gratefully acknowledges those

who have generously provided samples for this 6. testing, in particular Dr. T. Quadirfor providing many

of the Diamonite materials and especially those that have been hot pressed or HIPed; Dr. R. Brezny for the AT-ZT samples, and Dr. C. Cameron for the

three hot pressed oxide/non-oxide composites, and

Drs. J. Enloe and L. Dolhert for AIN samples (all of W. R. Grace's Washington Research Center); f^r. R. Palicka (of CERCOM) for SiC and AiPg-SiC whisker

composite samples; Dr. R. Ruh (of Wright-Patterson

Air Force Base), for both Boron and cubic stabilized 7.

(Zittrite) ZrOj; and Dr. R. Ingel (of the U.S. Naval

Research Laboratory) for partially and fully stabilized

ZrOj crystal samples. It is also a pleasure to acknowledge the assistance of Mr. C. Golden for 8. machining samples, Mr. L. Peters for mechanical

testing, and Mr. R. Milton for SEM analysis of fracture surfaces, all of W. R. Grace's Washington

Research Center.

REFERENCES

9.

1. R.W. Rice, "The effect of grinding direction on the strength of ceramics," NBS Special

Pub. 348, The Science of Ceramic 10. Machining and Surface Finishing, Eds. S.J. Schneider Jr. and R.W. Rice, pp. 365-76,

U.S. Govt. Printing Office, Washington, D.C.

(Issued May 1972).

2. R.E. Mould, and R.D. Southwick, "Strength 11,

and static fatigue of abraded glass under

controlled ambient conditions: II, Effect of various abrasions and the universal fatigue

curve," J. Am. Cer. Soc. [42], 582-92 (1 959).

3. H.P. Kirchner, R.M. Gruver, and R.E. Walker, "Strength effects resulting from 12. simple surface treatments," NBS Special

Pub. 348, The Science of Ceramic

Machining and Surface Finishing, Eds. S.J. Schneider Jr. and R.W. Rice, pp. 353-63,

U.S. Govt. Printing Office, Washington, D.C. 13,

(Issued May 1972).

4. R.W. Rice, "Machining of Ceramics" in "Ceramics for high-performance

applications," Proc. of the Second Army 14,

Materials Technology Conference, Eds. J.J. Burke, A.E. Garum, and R.N. Katz, pp. 287-

343, Metals & Ceramics Info. Center, 15,

Columbus, Ohio (1978).

5. C.A Anderson and R.J. Bratton, "Effect of 203

surface finish on the strength of hot pressed

silicon nitride," NBS Special Publication 562,

The Science of Ceramic Machining and

Surface Finishing II, B.J. Hockey and R.W.

Rice, Eds., pp. 463-474 U.S. Government

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

D.W. Richerson, J.J. Schuldies, T.M.

Yonushonis, and K.M. Johansen, "ARPA/

Navy Ceramic Engine Materials and Process

Development Summary," Ceramics for High

Performance Applications II, Proc. of the

Fifth Army Materials Technology

Conference, Eds. J.J. Burke, E.N. Lenoe

and R.N. Katz., pp. 625-650, Metals &

Ceramics Info. Center, Columbia, Ohio

(1978).

J.J. Mecholsky, S.W. Freiman, and R.W.

Rice, "Effect of grinding on flaw geometry

and fracture of glass, J. Am. Cer. Soc. [60]

3-4 (1977).

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

nature of strength controlling machining

flaws in ceramics," NBS Special Publication

562, The Science of Ceramic Machining and

Surface Finishing II, B.J. Hockey and R.W.

Rice, Eds., pp. 351-78, U.S. Government

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

R.W. Rice, "Machining flaws and the

strength grain size behavior of ceramics,"

Ibid ., pp. 429-452.

R.W. Rice, J.J. Mecholsky, Jr., and P.F. Becher, "The effect of grinding direction on

flaw character and strength of single crystal

and polycrystalline ceramics," J. Mat. Sci. 16

(1981) 853-862.

Ota, Minoru and Miyahara, Katsutoshi, "The

influence of grinding on the flexural strength

of ceramics," Soc. of Manufacturing

Engineers, MR90-538, paper presented at 4th International Grinding Conference,

Dearborn, Ml (1990).

K. Subramanian, and S. Ramanath,

"Principles of abrasive machining,"

Engineered Materials Handbook Vol. 4, pp. 315-28, The Materials Information Society

(1991).

R.W. Rice, "Correlation of machining-grain

size effects on tensile strength with tensile

strength-grain size behavior," J. Am. Cer. Soc. in press, (1992).

R.W. Rice, "Effects of porosity on strength

anisotropy due to machining direction," (to be published). D.R. Petrak, R. Ruh, and G.R. Atkins,

"Mechanical properties of hot-pressed boron

suboxide and boron. Am. Cer. Soc. Bull.
53[8], pp. 569-573 (1974).

16. R.W. Rice, "Microstructure dependence of mechanical behavior of cerannics," Treatise on Materials and Technology, Vol. II, Ed. R.K. McCrane, pp. 199-381, Academic

Press, Inc., New York (1977).

17. R.W. Rice, "Pores as fracture origins in ceramics," J. Mat. Sci. 19 pp. 895-914

(1984).

18. R.W. Rice, K.R. McKinney, and R.P. Ingel,

"Grain boundaries, fracture, and heat

treatment of commercial partially stabilized

zirconia," J. Am. Cer. Soc. 64 [12] C-175 - C-177 (1981).

19. CP. Cameron, J.H. Enloe, L.E. Dolhert, and

R.W. Rice, "A comparison of reaction vs. conventionally hot-pressed ceramic

composites," Ceram. Eng. Sci. Proc. 11 [9- 10], pp. 1190-1202 (1990).

20. J.S. Reed, and A-M. Lejus, "Affect of Grinding and Polishing on Near-Surface

Phase Transformations in Zirconia," Mat. Res. Bull. [12], pp. 949-54 (1977).

21. R.W. Rice, R.C. Pohanka, and W.J.

McDonough, "Effect of stresses from thermal

expansion anisotropy, phase

transformations, and second phases on the

strength of ceramics," J. Am. Cer. Soc, [63],

11-12 (Nov.-Dec. 1980).

22. R.W. Rice, "Effects of thermal expansion

mismatch stresses on the room-temperature

fracture of boron carbide," J. Am. Cer. Soc.

[10] 3116-18 (1990).

23. I. Tomayasu "Grinding characteristics and

material strength after grinding of fine ceramics," Toyoda Machine Works, Ltd.

204
DIAMOND GRINDING OF SILICON NITRIDE CERAMIC

J. E. MAYER, JR. and G. P. FANG

Texas A&M University, College Station, TX

Diamond grinding of hot pressed silicon nitride (HPSN) ceramic material is the main

focus of this paper. The object of this work is to achieve high overall removal rates while at the same time providing high strength in the final ground product. Grinding wheels with

different diamond grit sizes have been explored. Results reported include surface

characterization of the ground surface, transverse rupture strength data, and grinding flaw

detection results. Key findings are: All finish ground surfaces showed surface damage (pitting) over the range of wheel grit sizes employed (5.1 to 249 microns). For longitudinally ground

test bars, there is no apparent loss of strength as the grit depth of cut increases for the range

investigated (0.085 to 0.281 microns). For transversely ground test bars, there is a decrease

in strength as the grit depth of cut increases beyond a critical value. For diamond grinding of

HPSN ceramic this critical value appears to be about 0.160 microns.

1. Introduction

I Advanced ceramics, such as silicon nitride

' and silicon carbide, have the potential for broader use

{ in structural applications such as heat engines which

I would greatly increase their efficiency; however, the

strength variability and machining cost must be

j improved before this can take place (1-4). This

I paper is directed at providing ceramics which are ground with improved strength while at the same

' time achieving efficient grinding rates.

The conventional ceramic machining

operation for final shaping and finishing is diamond

' grinding. Material removed by this process is

ij usually by a microfracturing process (5-10). This

i fracturing process induces microcracks in the finished

I surface. The induced microcracks range down to 5-

i| 20 microns (11). These flaws result in variable

strength of the produced components which causes

the unreliable strength reputation. Grinding

II parameter investigations can improve the induced

flaw situation. Parameters can be selected which

result in less microfracturing, and thus can lead to higher strength of the ground ceramic material.

The main focus of this paper is diamond

grinding of hot pressed silicon nitride (HPSN)

ceramic material. Grinding wheels with different

diamond grit sizes were used to grind transverse rupture strength bars. Grinding direction was in both the longitudinal direction and transverse to the

bar length. These specimens were employed to obtain transverse rupture strength data (12-13) and

to characterize the ground surfaces by use of

scanning electron microscope (SEM) and

profilometer. The grit surface density of the

grinding wheels was measured by means of an

optical microscope sighting on the grit flats; and

this data was employed to compute the grit depth of

cut. In this paper, the relationships of grit depth of

cut and grind direction with strength and surface

characteristics of the ground specimen for HPSN

are discussed. Grit depth of cut to achieve high

strength in the ceramic during grinding (10,14-19)

is also discussed.

An investigation into methods to detect

flaws (microcracks) induced by grinding of

ceramics is being conducted utilizing surface

monitoring techniques. Flaws (microcracks) in the

range of 5-20 microns, which are of typical interest 205
regarding grinding craclcs in ceramics (11), are the

target size for these techniques. The flaw detection

methods and progress in this investigation are discussed in this paper.

2. Experimental Procedure

2.1 Machining

HPSN bars were obtained with approximately

0.508 mm (0.020 in) oversize all around to allow for

the grind stock required in MIL-STD-1942(MR) (12-

13). MIL-STD-1942(MR) was followed because it is a tighter specification than ASTM-Cl 161-91. These

bars were all cut from one hot pressed billet. In this research, the bars were rough ground for the first 0.457 mm (0.018 in) and then finish ground for the

final 0.0508 mm (0.002 in) in accordance with MILSTD-1942(MR). The final bar dimensions were

3.00 mm x 4.00 mm x 45.0-50.0 mm (0.118 in x 0.158 in X 1.75-2.00 in), which are the dimensions

required for the transverse rupture strength tests.

Resinoid-bonded diamond wheels with

different grit sizes which were employed in this research are shown in Table 1. These wheels have

150 concentration and uncoated diamond. The outer diameter of the wheel is 152.4 mm (6 in) and the

width and thickness of the diamond layer are 3.175

mm (0.125 in). The grit surface density was

measured in this research by means of an optical

microscope sighting on the grit flats.

A commercial surface grinder with microprocessor

control on the wheel depth of cut was used in this research. The wheel depth of cut could be

controlled at a unit of 1.27 microns (0.00005 in). The following grinding parameters were held

constant at the indicated values: wheel speed =

24.69 m/s (4860 ft/min), table feed for roughing =

11.18 mm/s (26.4 in/min), table feed for finishing = 7.62 mm/s (18.0 in/min), and crossfeed = 1.59

mm (0.0625 in). The wheel depth of cut for

roughing was 0.0254 mm (0.001 in) and for

finishing 0.00254 mm (0.0001 in), except for the

5.1 micron grit wheel which used a 0.00127 mm

(0.00005 in) wheel depth of cut for finishing. The

same wheel was employed for roughing and

finishing in the case of each grit size, except for the

5.1 micron grit wheel where the 50.8 micron grit

wheel was used for roughing. The specific material

removal rates for roughing and finishing conditions

are shown in Table 2. For strength tests, five bars were prepared at each grinding condition with the

grind direction longitudinal to the length of the bar,

and another set of five bars were ground transverse

to the bar length in order to determine the effect of

grind direction on the strength. A water miscible

synthetic grinding fluid was used in these

experiments. The grinding wheels were trued with

a commercial diamond roll truing device and were

dressed with aluminum oxide dressing sticks. Truing and dressing procedures were standardized

and were those recommended by the diamond

grinding wheel supplier.

Table 1. Grit Size of Diamond Grinding Wheels

Wheel Mean Grit Size Grit Surface Densitv

microns inches No. /mm' No./in^

D4/8MIC-R 1 50BXL6 1 45- 1 /8 5.1 0.0002 49.91 32,200

SDG320-R 1 50BXL6 1 45- 1 /8 50.8 0.0020 27.59 17,800

SDG 1 80-R 1 50BXL6 1 45- 1 /8 87.5 0.0035 11.66 7,520

SDGl 10-R150BXL6145-1/8 152 0.0060 8.324 5,370

SDG 1 00-R 1 50BXL6 1 45- 1 /8 249 0.0098 6.402 4,130

Table 2. Specific Material Removal Rates

Grind Type

Roughing

Finishing

Finishing (5. 1 micron

grit wheel only)

mm^/s • mm

0.284

0.0194

0.00968

in /mm * in 0.0264

0.0018

0.0009

206
2.2 Surface Analysis 3. Experimental Results and Discussion

A scanning electron microscope (SEM) was

utilized to characterize the ground surfaces for all of

the grinding wheels. Magnifications employed were

either 1500X or 3500X. Magnification of 7500X was

used to investigate grinding swarf, and 10,000X was

employed to illustrate plastic deformation on the

ground surface. A contacting surface profilometer was utilized to measure the surface finish of the

ground specimens.

2.3 Strength Testing

Transverse rupture strength tests were

performed in accordance with MIL-STD-1942(MR)

(12-13) on all bars ground in this research. The test fixture shown in Figure 1 was employed in these

strength tests and the fixture was loaded by means of

a universal test machine. A total of five bars were

tested for each grinding condition to establish the transverse rupture strength. The range of strengths

and the Weibull characteristic strength were utilized to present the results in this investigation.

2.4 Flaw Detection

The overall research plan includes attempts to detect flaws (microcracks) induced by grinding via

various non-destructive surface monitoring

techniques. The first two methods being investigated

are the nuclear magnetic resonance imaging (NMRI)

technique (20) and a thermal acoustic surface wave

(Thermal Acousto Photonic NDE (TAP-NDE))

technique (21-25).

Loading

n

Member

3.1 Grit Depth of Cut

The controlling factor in determining the

material removal mechanism is believed to be the

grit depth of cut, t, which is illustrated in Figure 2

(16-17,26-27).

Path of grit tip

Figure 2. Illustration of the Surface Grinding

Process Showing Various Process

Variables.

Test Specimen

1.5d^^^^

Support Member

L »j Note: Bearing Cylinders are held in

I place by low stiffness springs

' or rubber bands. Figure 1 . Four-Point Transverse Rupture Strength Test Fixture.

207
The value of t depends on both machine and wheel

parameters and can be computed from the following

expression (28):

t = 4v_

N vcA

d_

D

(1)

where t

V

V

d

D

is grit depth of cut (maximum

undeformed chip thickness)

is the number of active cutting points per

unit area of the wheel periphery (grit

surface density)

is the ratio of chip width to average

undeformed chip thickness

is wheel peripheral speed

is workpiece speed (table speed)

is wheel depth of cut

is wheel diameter

The value of r is reported to be in the range of 10-20

(28); r is assumed to be equal to ten in this paper

which should give a conservative value for t, i.e. the

computed value of t may be larger than the actual value. The value of C was determined by means of

an optical microscope sighting the grit flats after grinding with the wheel. The values of C so determined are shown in Table 1 . The values of t for

the grinding conditions employed in this research were computed using Equation (1) and are reported

in Table 3, listed by diamond grit size. The values

of grit depth of cut, t, shown in Table 3 will be

utilized to help explain the surface characteristics and

strength results.

for each of the grinding conditions, i.e. for each

wheel grit size. These micrographs are shown in Figure 3. All of the SEM micrographs in Figure

3(a) - 3(f) show surface damage (pitting) and a combined ductile and fracture mode grinding

(16,19,29); however, the trend is toward less damage and more ductile mode appearance as the

diamond grit size becomes smaller, i.e. as the grit

depth of cut becomes smaller. In this paper, the

pitting is assumed to be caused by a fracturing

process; and surface areas which are smooth and

without pits are assumed to have been produced by

a ductile material removal process, i.e. by plastic

deformation. None of the micrographs, not even

for the largest grit size Figure 3(e)-3(f), show

evidence of extreme brittle fracture as observed

when grinding glass as illustrated in Figure 4. This

can be explained by the much higher fracture

strength of HPSN compared to glass.

For the 5.1 micron grit wheel, a SEM

micrograph was also made after the first 25.4

microns (0.001 inch) depth was removed during

finish grinding. Figure 5 shows this 25.4 microns

(0.001 inch) finish depth SEM micrograph. All

other SEM micrographs of the ground surface were made after completion of finish grinding, i.e. after 50.8 microns (0.002 inches) depth was removed by

finish grinding. By comparing Figure 5 with

Figure 3 (a), one can observe the surface at the

halfway point in finish grinding relative to that at the endpoint of finish grinding. Figure 5 exhibits

greater surface damaged area than Figure 3 (a). The surface damage in Figure 5 was probably

Table 3. Grit Depth of Cut, t Mean Grit Size Computed t Values

microns inches

Roughing Finishing

microns micro-inches microns micro-inches

5.1 0.0002 0.0845 3.33

50.8 0.0020 0.2910 11.46 0.1351 5.32

87.5 0.0035 0.4478 17.63 0.2079 8.19

152 0.0060 0.5299 20.86 0.2460 9.69

249 0.0098 0.6041 23.78 0.2805 11.04

3.2 Surface Characteristics

Scanning electron microscope (SEM)

micrographs were made of the finish ground surface

induced initially by the roughing grind wheel (50.8

micron grit) and is being removed by the 5.1

micron grit wheel during finish grinding.

208
1

i

1 © m 0 19 3 1 3j

3 (a) 5.1 Micron Diamond Grit Size.

f 209
3 (c) 87.5 Micron Diamond Grit Size.

3 (d) 152 Micron Diamond Grit Size.

210

A
3 (e) 249 Micron Diamond Grit Size.

3 (f) 249 Micron Diamond Grit Size (Higher Magnification).

Figure 3. SEM Micrograph of Finish Ground Surfaces.

211
Figure 4. SEM Micrograph of Ground Glass Surface (102 Micron Diamond Grit Size).

2 0KU X1,S00 10>*m 12923-

Figure 5. 5.1 Micron Diamond Grit Size (at halfway point in finish grinding).

212
The percentage of surface damaged area (total pitted area relative to total surface area) was

measured by means of optical image processing of

the SEM micrographs for the ground surfaces

produced with the 5.1 micron grit wheel up through

the 249 micron grit wheel. SEM micrographs with

750X magnification were employed for these

measurements. The results are listed in Table 4. The percent area values in Table 4 are the average of

five different locations on the ground surface. Table

4 indicates that the percent surface damaged area ranges from about 3% to 25% as the wheel grit size

ranges from 5.1 microns to 249 microns.

Table 4. Percentage of Surface Damaged Area

Mean Grit Size Percent Damaged

microns inches (Pitted) Area

5.1 0.0002 2.83

50.8 0.0020 11.91

87.5 0.0035 16.56

152 0.0060 23.85

249 0.0098 24.91

Figure 6 shows SEM micrographs of

grinding swarf and Figure 7 a grinding groove

under higher magnification. The swarf was

collected by employing fresh coolant, collecting

coolant coming from the wheel during grinding,

depositing a coolant drop on a SEM stud, and

letting air dry. In Figure 6, Sample A and Sample

B are SEM micrographs taken at different locations

in the grinding swarf. Figure 8 shows the bulk

structure of the HPSN which was obtained by

etching the ground surface with 49% hydrofluoric

acid for 2 hours at room temperature. From a comparison of the swarf in Figure 6, the surface area region next to the grinding groove in Figure 7,

and the damaged areas in Figure 3 with the bulk

structure in Figure 8, it appears that the material

removal mechanism in the damaged areas is by

fracturing out HPSN grains or clumps of grains.
6 (b) Sample B.

Figure 6. Grinding Swarf (102 Micron Diamond Grit Size).

214
Figure 8. Etched Surface.

A contacting surface profilometer was used to measure the surface finish of the finish ground

specimens. The surface finish results are recorded in Table 5, listed by diamond grit size. Ra is the

arithmetic mean deviation of the roughness profile

from the mean line. All measurements were made

perpendicular to the grinding direction. The surface

finish is plotted versus grit depth of cut in Figure 9. A plot was also made versus wheel grit size and is shown in Figure 10. These results show, as expected, that the surface finish improves as the

wheel grit size and the grit depth of cut become

smaller.

Table 5. Surface Finish, Ra

Mean Grit Size Ra

microns inches microns ^-in

5.1 0.0002 0.0508 2.0

50.8 0.0020 0.1594 6.28

87.5 0.0035 0.2985 11.75

152 0.0060 0.2864 11.28

249 0.0098 0.4090 16.10

0 0.1 0.2 0.3

t, microns

Figure 9. Surface Finish, Ra, versus Grit Depth of

Cut, t.

215
0.5

0-^ \ \ \ 1 \

1 0 50 100 150 200 250 300 Wheel Grit Size, microns

Figure 10. Surface Finish, Ra, versus Wheel Grit

Size.

3.3 Strength

Strength data obtained by means of the transverse rupture strength tests on a!l bars ground in this research is given in Table 6 and is summarized

in Figure 1 1 with strength plotted against wheel grit

size, and in Figure 12 plotted versus grit depth of

cut, t. The range of strengths obtained for the five

test bars at each grinding condition and the WeibuU

characteristic strength for the five bars are both

shown in these figures. Separate plottings are made

for the longitudinally ground specimens and those

transversely ground. From Figures 11 and 12 it can be seen that for grinding in the longitudinal

direction there is no loss of strength as the grit

depth of cut, t, increases (or the grit size increases)

for the range investigated. Statistical analysis of the

data for all longitudinal grinding showed no

significant difference in strength. For transverse

grinding, there is a decrease in strength as t increases (or the grit size increases) beyond a

critical value of t of about 0.160 microns.

Statistical analysis of the data for transverse grinding showed that the, strengths for the 152 and

249 micron grit sizes are significantly different

compared to the strengths for the other grit sizes

employed for transverse grinding; they are also significantly different compared to the strengths for

all grit sizes employed for longitudinal grinding.

Figure 11 agrees with general trends reported

previously by Ota (30); however, the strength levels are lower in Ota's work. Ota's maximum strength

levels were about 600 to 700 MPa. Additionally,

Ota's strengths for transverse grinding never reached those for longitudinal grinding over the 30-

160 micron wheel grit sizes explored. Ota's

ceramic workpiece material was sintered silicon

Table 6. Transverse Rupture Strength Test Results

Wheel Grit Microns 5.1 50.8 87.5 152 249

Size Mesh No. 1200 320 180 110 100

Grind Direction T L T L T L T L T

Breaking

Stress

for Individual Bars

MPa 868.6

835.9

731.8

869.7

635.0

759.2

856.3

910.8

725.1 688.0

962.3

794.0

823.7

812.6

804.4

839.6

668.7

880.8

702.1 703.6

850.4

545.7

715.1 641.3

691.0

857.1

699.1 788.9

881.5

848.2

566.8

602.8

507.1 572.0

503.8

802.2

667.3

752.9

917.9

884.1

604.2

556.8

597.9

683.6

458.9

Ksi 126.0

121.2

106.1 126.1

92.1

110.1 124.2

132.1 105.2

99.8

139.6

115.2

119.5

117.9

116.7

121.8

97.0 127.7

101.8

102.0

123.3

79.1 103.7

93.0 100.2

124.3 101.4

114.4

127.9

123.0

82.2 87.4

73.6

83.0

73.1

116.3

96.8 109.2

133.1 128.2

87.6 80.8

86.7

99.1 66.6

Weibull MPa 827.0 826.3 870.9 798.1 732.5 842.8 568.4 844.8 611.9

Characteristic Strength Ksi 119.9 119.8 126.3 115.8 106.2 122.2 82.4 122.5 88.8

Weibull Modulus 8.22 7.21 8.48 6.79 5.20 12.27 11.68 7.48 6.57

* Note: L = Longitudinal Grinding, T = Transverse Grinding

216

1
1200

1000 -

(T3 Q. 800 H

^ 600 - c

i: 400

200 -

50

I

\ 1

1 00 1 50 200

Wheel Grit Size, Microns

T

250 300

Figure 11. Transverse Rupture Strength as a Function of Wheel Grit Size in Longitudinal (L) and

Transverse (T) Grinding.

1200

t. Microns

Figure 12. Transverse Rupture Strength as a Function of Grit Depth of Cut, t, in Longitudinal (L)

and Transverse (T) Grinding.

217
nitride (SSN) which can have lower strength and

toughness than hot pressed silicon nitride (HPSN),

which possibly explains the differences between the

present research and the results of Ota (31).

The work of Bifano (16) suggests that the grit

depth of cut, t, for ductile mode grinding of silicon

carbide ceramic should be about 0.200 microns or

less; above this value the grinding mode changes

from a ductile to a brittle mechanism. Additionally,

since silicon nitride ceramic has higher toughness

than silicon carbide ceramic, it is expected that the

critical t value for ductile mode grinding of silicon

nitride ceramic will be larger than 0.200 microns.

Bifano's definition of ductile mode grinding is based on the relative amount of fractured surface area on

the ground surface. He defines ductile mode

grinding as grinding which produces 10% or less fractured (damaged) surface area. Looking at Figure

3 and Table 4, it appears that 10% damaged surface area occurs between the 5.1 micron and 50.8 micron

diamond grit size wheels (i.e. between Figures 3(a)

and 3(b)). Table 3 indicates that this corresponds to a t value of 0.085-0.135 microns, which appears to disagree with the findings of Bifano.

If one defines ductile mode grinding as those

grinding conditions where maximum strength is achieved, then from Figure 12 it is seen that the

critical t value under this definition is about 0.160

microns. This critical t value is larger than that determined by applying Bifano's definition to this present research, which possibly indicates that the

arbitrary percent damaged area in his definition

should be larger than 10%. From Table 4, it appears

that about 15% might be a more realistic value to employ. Since the real goal is to produce high

strength ceramic components, it appears that a ductile

mode grinding definition based on strength is a more appropriate criterion.

During the course of this research, it was

observed that the strength of the HPSN employed

can vary from one hot pressed billet to another.

Strength results for two different billets ground with

50.8 and 87.5 micron grit wheels are shown in Table 7. The results in Table 7 are for test bars

ground longitudinally. Statistical analysis of the

strength data for Billet AA and Billet BB showed

that their strengths are significantly different.

Previous results did not show a significant

difference in strength from billet to billet (31).

This result means that a control set of transverse rupture strength test bars must be ground and tested

for each billet in order to make meaningful

conclusions when more than one billet is employed

in a test series.

3.4 Flaw Detection

The brittle removal process produces

microcracks on the surface of the ground ceramic;

whereas the surface produced by the ductile

removal process should be free of microcracks and

thus have higher strength. Thus if cracks on the

ground surface can be detected by some nondestructive technique, then it might be possible to employ this technique to determine the extent to which the ductile process exists and the level of

strength of the ground surface.

To pursue this possibility, attempts are being made to detect flaws (microcracks) induced

by grinding via surface monitoring techniques.

Flaws (microcracks) perpendicular to the surface

Table 7. Strength from Billet to Billet

HPSN Material

Transverse Rupture Strength

(Weibull Characteristic Strength)

50.8 Micron Grit Wheel

MPa Ksi

87.5 Micron Grit Wheel

MPa Ksi

Billet AA

Billet BB

703.6

826.3

102.0

119.8

653.4

798.1

94.8

115.8

218
Figure 13. NMRI (Nuclear Magnetic Resonance Imaging) Image of Cross-Section of Ground Silicon

Nitride Specimen.

and in the range of 5-20 microns, which are of

typical interest regarding grinding cracks in ceramics

(11), are the target size for these techniques.

Investigation of the nuclear magnetic resonance imaging (NMRI) technique (31) has shown promise.

J Figure 13 shows the NMRI image for a ground

silicon nitride specimen. The "plus sign" in this figure is an image indicating a flaw. The dots

leading down to the "plus sign" are near to the 20

; micron range.

ij Additionally, a technique called Thermal

Acousto Photonic NDE (TAP-NDE) is being

employed to investigate detection of grinding cracks

in ceramics utilizing Rayleigh surface wave and

Lamb wave techniques (21-25). This research is in the early stages of investigation. The techniques of

[ thermal wave analysis (imaging) and photothermal

I microscopy (32) are being studied for possible

j investigation.

ji

4. Conclusions

; Hot pressed silicon nitride (HPSN) ceramic

t was ground on a surface grinder with various grit

size diamond wheels, and the finished surfaces were

observed with SEM and profilometer. Grit depths of

1 ^ 219

cut were estimated. Transverse rupture strength

data was obtained. Additionally, flaw detection

methods are being investigated. The following

conclusions may be made on the results of this research work.

1. All finish ground surfaces showed surface

damage (pitting) over the range of 5.1 to 249 micron diamond grit sizes investigated

(i.e. grit depths of cut, t, ranging from

0. 0845 to 0.281 microns). The amount of

surface damaged area ranged from 3% to 25% as the wheel grit size ranged from 5.

microns to 249 microns. None of the

surfaces showed evidence of severe

fracturing or cracks. The trend was toward

less damage and more ductile mode

appearance as the grit size became smaller,

1. e. as the grit depth of cut became smaller.

2. It appears that the material removal

mechanism in the damaged areas is by

fracturing out HPSN grains or clumps of

grains.

3. For grinding in the longitudinal direction

along the test bar, there is no loss of

strength as the grit depth of cut, t, increases
(or the grit size increases) for the range

investigated. For transverse grinding, there

is a decrease in strength as t increases (or the

grit size increases) beyond a critical value of

t. This result is in general agreement with

previously published work. The present

research provides new information in that it estimates the critical value of t at about 0. 160

microns for diamond grinding of hot pressed

silicon nitride ceramic. For the grinding

conditions employed in this research, this can be achieved with a 50.8 micron grit size (320

mesh number) wheel.

4. It appears that high strength grinding of

HPSN occurs when the damaged surface area

is less than about 15%.

5. It has been observed that the strength of

HPSN can vary from one hot pressed billet to another. This means that a control set of

transverse rupture strength test bars must be

ground and tested for each billet in order to make meaningful conclusions when more

than one billet is employed in a test series.

Acknowledgements

The authors would like to express sincere

thanks to those who have supported this research

work as follows: Texas Higher Education

Coordinating Board Advanced Technology Program

(Project No. 999903-172), National Science

Foundation (Grant No. DDM-9020981), Ford Motor

Company, Norton Company, Allen-Bradley

Company, Texas A&M University, and Ashland Oil

Company.

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222
MACHINING OF SILICON NITRIDE: EXPERIMENTAL

DETERMINATION OF PROCESS/PROPERTY

RELATIONSHIPS

R. L. ALLOR, T. J. WHALEN, and J. R. BAER

Ford Motor Company, Dearborn, MI

K. V. KUMAR

GE Superabrasives, Worthington, OH . .

Strength controlling flaws in ceramics such as silicon nitride can be introduced during the

grinding process. An experimental program was conducted to investigate and improve the

efficiency of machining a structural ceramic material, silicon nitride, with diamond abrasives.

This paper analyzes the data from two groups of experiments and determines the significant

results. The first group of designed experiments, half-factorial, investigated the effects of

diamond mesh size and concentration in the grinding wheel, the rate of removal of the

ceramic, wheel speed and type of grinding on the strength, surface roughness, power

consumption, vertical force and wheel wear. The grinding machine spindle contained air bearings. Two or three significant factors were found to affect each result. For example, the

strength of the ceramic was affected by the diamond mesh size when creep-feed grinding was

used, whereas diamond concentration and speed of the wheel were found to be significant

factors when pendulum-feed grinding was done. The second group of experiments were run on the same grinding machine with a spindle containing steel ball bearings. Results from the

second experiment showed that creep or pendulum modes of machining had no effect on

strength of machined silicon nitride, nor did removal rates.

iNTRODUCnON

Machining of ceramics usually introduces

strength controlling features at the surface or in the

I subsurface region. Allor, Govila and Whalen(l)

I presented data on the machining of silicon nitride

i (Si3N4) and silicon carbide (SiC) and showed the

: influence of machining variables on the strength of

these advanced ceramic materials. Many other

authors have also presented data on machining

these advanced ceramics and compilations of papers

1 presented at conferences are available(2-9).

j Recently, an experimental program was established

J; to investigate and improve the efficiency of

machining hot-pressed silicon nitride with

|i manufactured diamond abrasives. An approach

based on statistical design of experiments was used

to analyze the effect of grinding parameters on

I strength of one silicon nitride material.

! Experimental Design

The original design was a half factorial ly ^

j design with five variables at two levels each and two

replicates at each experimental condition, totaling

32 experimental runs performed in random order.

The five variables were mesh size of diamond

(80/100 and 120/140), concentration of diamond in the wheel (75 conc.=3.3cts/cm^ and 100

conc.=4.4cts/cm^), specific material removal rate (6

and 12 mm^/mm-sec), wheel speed (30 and 45 m/s),

and type of grinding action (pendulum and creep

feed). After noting an error in the rate of removal

after several experiments, the 2v'^ design was

redone to form two 2fv^ blocks, one for pendulum

feed and one for creep feed. The block for

pendulum feed was performed at rates of 6.0 and

12.0 mmVmm-s, while the block for creep feed was

performed at 0.6 and 1.2 mmVmm-s. Six bars were

ground per test condition. The results of the

experiments were transverse rupture strength

(MPa), average surface roughness (Ra, /xm), average

power consumption (watt/mm), average vertical

force (N/mm) and radial wear on grinding wheel

The spindle of the grinding machine for the

223
Table I: Grinding Conditions

Work Material Hot pressed silicon nitride (NC-1 32) - Norton

Wheel Specification

Diamond Type RVG-880 Diamond

Mesh Size 80/100 and 120/140

Concentration 75 and 1 00

Bond Type Phenolic resin

Wheel Size 1753 x 6.3W x 31 .75B, mm

Grind Type Plunge Grind - pendulum and creep feed up cut

Wheel Speed 30 and 45 m/s

Depth of Cut 0.025 and 0.050 mm for pendulum

0.25 mm for creep feed

Specific Removal Rate 6 and 1 2 mmVmm-s, pendulum (first experiment)

0.6 and 1 .2 mm^/mm-s, creep feed (first experiment)

1 .2 and 1 2 mm^/mm-s (second experiment)

Coolant HD-90 water soluble, 20:1

above experiment was equipped with air bearings.

A second experiment was conducted in which the

grinding machine spindle was changed to steel ball bearings. Based on the results from the first experiment, a simple 2pu,| design was used for the

second experiment to investigate differences in

resulting strength of the material, if any, between

samples ground by creep feed mode and pendulum

mode at removal rates of 1.2 and 12 mmVmm-s. In

this case, 8 bar samples were ground per test condition.

Two general references on the design and

analysis of these experiments are given(10,ll).

Table II: Results for 1st Experiment/Creep Feed (2^^;

Experimental Procedure

A hot pressed silicon nitride (NC132) billet was

cut into blank test bars. The dimensions prior to grinding were 60L x 4.5W x 3.5T mm. During the

test, approximately equal amounts of material were removed from each side to obtain the finished

dimensions. Table I gives the grinding conditions

used for the experiment. A phenolic resin bonded

wheel was used for all the tests and the choice of

RVG*-880 diamond, was based from an earlier

study(12). In creep-feed grinding the process used

is a large depth of cut with a very slow table feed, as compared to pendulum feed grinding that uses a shallow depth of cut with a fast table feed cutting in

' Matrix with 2 Replicates)

Exp

#

Mesh

Size

Cone.

Removal

Rate

(mm^/mm-2)

Wheel

Speed

(m/s)

Strength

(MPa)

Surface

Roughness

Power

(Watt/mm)

Vertical

Force

(N/mm)

Wear

ijjm)

1 80 75 0.6 30 848/869 0.50/0.50 58/58 56/103 7.6/7.6

2 120 75 0.6 45 662/717 0.90/0.55 55/55 32/27 25.4/20.3

3 80 100 0.6 45 848/834 0.35/0.45 158/123 88/85 7.6/8.9

4 120 100 0.6 30 696/738 0.35/0.45 78/112 65/49 7.6/8.9

5 80 75 1.2 45 896/800 0.45/0.50 110/100 56/46 6.4/7.6

6 120 75 1.2 30 751/800 0.90/0.70 45/33 15/14 25.4/16.5

7 80 100 1.2 30 814/841 0.50/0.55 78/78 49/49 5.1/6.4

8 120 100 1.2 45 779/676 0.55/0.35 88/110 35/54 7.6/6.4

224
Table III: Average Effect on Results of Factors in 1st Experiment/Creep Feed

Strenath (MPa) Surface Roughness Power

(Watt/mm)

Vertical Force

(N/mm)

Wear

(pm)

Mesh Size (M) -116.4 0.12 -23.4 -29.6 7.6

Conc.(C) -14.7 -0.18 38.9 15.1 -7.3

Removal Rate (R) 18.1 0.06 -6.9 -23.9 -1.6

Wheel Speed (S) -18.1 -0.04 32.4 3.4 0.6

MC + RS 4.3 -0.16 11.1 13.6 -7.0

MR + CS 30.2 0.01 0.9 10.1 -0.0

MS + CR -19.9 0.03 -22.4 -2.1 -0.3

Average 786.6 0.53 83.7 51 .2 11.0

both directions. In the case of creep feed grind

tests, only up cut was used and all of the stock per

side was removed in a single pass with the work

speed adjusted to obtain the high and low material

rates. For pendulum grind tests, the work speed was kept constant and the depth of cut was adjusted

to obtain the high and low material removal rates. Following grinding, all the samples were chamfered

in the lengthwise direction.

After machining, the bars were broken in fourpoint bend test to obtain the transverse rupture

strength. The baseline strength for the hot-pressed

silicon nitride billet used in this study was obtained

from bar samples ground according to Table IV: Results for 1st Experiment/Pendulum Feed

reference(13).

First Experiment: Creep Feed Grinding

The data determined in the grinding

experiments with creep feed rates of 0.6 and 1.2 mm^/mm-s are listed in Table 11 and IIL

Strength

The strength data listed in Table 11 and in

Table IV are shown as two values or replicates at each experimental condition. Each value of

strength is the average value of six bars.

Only the mesh size variable, of the four

variables evaluated, had a significant effect on the

(2^v^ Matrix with 2 Replicates)

Exp

#

Mesh

Size

Cone.

Removal Rate

(mmVmm-s)

Wheel

Speed

(m/s)

Strength

(MPa)

Surface

Roughness

Power

(Watt/mm)

Vertical

Force

(N/mm)

Wear

ipm)

1 80 75 6 30 786/676 0.55/0.55 213/190 30/28 3.8/3.8

2 120 75 6 45 696/662 0.75/0.80 113/113 20/19 27.9/25.4

3 80 100 6 45 710/724 0.45/0.60 203/215 37/40 3.8/2.5

4 120 100 6 30 758/765 0.45/0.50 203/203 29/27 2.5/3.8

5 80 75 12 45 662/717 0.65/0.60 333/303 58/52 6.4/8.9

6 120 75 12 30 738/752 0.70/0.80 223/223 28/24 20.3/22.9

7 80 100 12 30 800/738 0.50/0.55 41 5/460 56/59 3.8/3.8

8 120 100 12 45 772/717 0.60/0.55 370/303 52/46 5.1/8.9

225
strength of the machined silicon nitride, and these

data are shown in Figures 1 and 2. The effect of a variable or factor is the change in response or result

produced by a change in the level of the factor.

When the level is examined at two levels only, as is the case in these experiments, the effect is simply

the difference between the average response of all trials or experiments carried out at the first level of

the factor and that of all trials or experiments at the

second level. An assumption is made that the

distribution of experimental errors is Normal or

Gaussian, and this assumption has often been found

to hold in practice. Therefore, if the effects are

plotted on Normal probability paper, as is done in

Figure 2, the effects that fall within experimental

error will show a linear relationship and the

statistical effects will lie some distance from the line. As in the case shown in Figure 2, the significant

effect of mesh size is some distance from the line

representing the effects which are included in

experimental error. The greater strength was

attained in samples machined with 80 mesh diamond

abrasive. Diagnostic tools used to determine the

significance of the data are the plot of the scaled

t-statistic (Figure 1) and the probability plot of the

effects (Figure 2). A plot of the probability versus

the residuals of a model employing only the effect

of the mesh size showed a linear relationship, which

indicates that the mesh size variable adequately

describes the data.

Surface Roughness

Two significant effects were noted which

influenced the surface roughness: the concentration

of the diamond in the grinding wheel and the mesh

size - concentration two-factor interaction. When

an effect of one factor or variable depends on the

level of another factor, the two factors are said to interact and these interactions of two factors are

called two-factor interactions. The data and plots

showing these effects are given in Figures 3, 4 and

5. Since a two-factor interaction is present, one

• tM • • •

I

' ' '

I

' ' '

I

' ' '

I

I 'I

I

' ' '

I

' ' ' I' ' '

I

'

I '

I -lao -1O0 -M -ae -40 -ao o ao 40 EWact

^ ^ ^ ... Figure 1: Plot of strength effects and scaled

t-statistic shows the region of experimental error, and effect of mesh size is well beyond

experimental error.

90.0

99

96

!5

1 S 60 o y

o o/ /

M

6

0

1

0.1

-120 -80 -40 0 Flexure Strength (MPsO

40

Figure 2: Probability plot of strength effects

shows significant mesh size factor separate from

other effects which are not significant.

C MC • • • • • • •

I

I 'I

' M ^

I I

I

I I I

'

I

I I I M I

'

i I

I

I I I I M I I

'

I (u .^.15 -aos 0 o« 0.1 ai9 EffMt

=8gure 3: Plot of surface roughness effects and

scaled t-statistic shows that concentration and

mesh-concentration two-factor interaction are important factors for surface roughness.

99.9

99

1

-0.18 -0.13 -0.08 -0.03 0.02 0.07 0.12 Surface Roughness Oim)

Figure 4: Probability plot of effects of machining

variables on surface roughness.

226
sees that the level of the factors influence the

results of the experiment. In this case both levels of

concentration result in a low surface roughness

when 80 mesh diamonds are used in the wheel, but

only at the high concentration level does one

observe a low surface roughness when a 120 mesh

diamond wheel is used. Therefore, low surface

roughness was generated with a 80 mesh diamond

wheel with either 75 or 100 concentration diamond,

and also with 120 mesh diamond wheel with only

the 100 concentration (Figure 5).

Average Power

Average power used during grinding was

significantly affected by speed, mesh size and

concentration. The greater amount of power was

used with 80 mesh diamond wheels of 100

concentration at the higher speed of 45 m/s.

Figure 6 contains the data for the power results. The best conditions fi^om a power standpoint

(lowest power needed) were 120 mesh diamond

wheel of 75 concentration at 30 m/s.

Average Vertical Force

I The variables of removal rate and diamond

i mesh size significantly influenced the average

; vertical force. The effects of removal rate and mesh

' size on vertical force in Figure 7 are both negative.

Therefore, the greater vertical force is found with

the lower level of removal rate (0.6 mmVmm-s) and

ij mesh size (80 mesh size). Since lower vertical force

is desired,

preferred.

120 mesh size and 1.2 mm7mm-s is

Wheel Radial Wear

Diamond mesh size, diamond concentration,

Surface Finish (^m)

0.8

0.7

0.6

0.5,

0.4

0.3

Conc.75 Cone.100

80 120 Mesh

Figure 5: Plot of effects of mesh-concentration

two-factor interaction on surface finish.

and a mesh-concentration two-factor interaction

influence the radial wear of the diamond grinding

wheel. Figures 8, 9 and 10 contain data which

clearly show these effects. The presence of a twofactor interaction requires one to consider the level

of each of the significant factors, mesh size and

concentration, when one predicts the better

operating conditions. In this case, the lower wear

rate is obtained for 80 mesh diamond wheels with

either 75 or 100 concentration, whereas for 120

mesh wheels, the low wear rate is observed at only

100 concentration.

First Experiment: Pendulum Feed Grinding

The data generated in these experiments with

pendulum feed rates of 6 and 12 mm^mm-s are given in Tables IV and V.

Figure 6: Plot of effects of speed, mesh size and

concentration on average power, with scaled t plot to show extent of experimental error.

Figure 7: Plot of significant effects of removal rate and diamond mesh size on average vertical force.

Figure 8: Plot of significant effects of mesh,

concentration and a mesh-concentration twofactor interaction on wheel radial wear, with a

scaled t plot to show extent of experimental error. 227
Strength

Both the grinding wheel speed and the

diamond concentration were significant factors in

influencing the strength of the silicon nitride, as shown in Figures 11 and 12. The higher strength

test bars were those machined at the higher speed,

45 m/s, and higher diamond concentration of 100.

All other factors fell within the experimental error. The plot of probability versus effects indicates

clearly the significance of the two factors, speed and

concentration. The linear relationship shown in the

plot of probability versus residuals. Figure 13, when

the model includes these two factors, enhances our

confidence in the interpretation. (See Reference

10, p.333 for discussion of residual analysis.)

Wear (fim)

99.9

99 M 95

/

O

o

CB 50 n

o

Q. 20 MO y

c /

5

1

0.1

1 , , , 1 , , , 1 1 ... 1

0 4 Wear (^m)

120 Mesh

Figure 10: The mesh-concentration two-factor

interaction is shown for wheel radial wear.

Figure 9: Probability plot of factors showing

significant factors for wheel radial wear some

distance from line representing experimental error.

Figure 11: Plot of significant effects of whee

speed and concentration on strength of machined

samples, along with a scaled t plot to show extent

of experimental error.

Table V: Average Effect on Results of Factors in 1st Experiment/Pendulum Feed

Factor

Strength

(MPa)

Surface Roughness

(^m)

Power

(Watt/mm)

Vertical Force

(N/mm)

Wear (^m)

Mesh Size (M) 6.1 1 0.09 -72.6 -14.4 10.0

Conc.(C) 37.1 -0.15 82.6 10.9 -10.7

Removal Rate (R) 14.7 0.04 147.1 18.1 0.8

Wheel Speed (S) 44.0 -0.05 22.1 -5.4 -3.0

MC -1- RS 4.3 -0.09 19.1 4.9 -8.4

MR -1- CS 9.5 0.00 -25.4 -4.4 -1 .4 MS CR -2.6 -0.01 -33.9 -1.9 -1.4

Average 729.5 0.60 255.2 37.8 9.6

228
Surface Roughness

Identical to the effects observed for the creepfeed grinding, the surface roughness is influenced

significantly by the mesh size, the diamond

concentration and the two-factor interaction

between mesh size and concentration. These results

for pendulum feed machining are seen in Figures 14

and 15, which is similar to Figures 3, 4 and 5 for

creep feed grinding. The lower surface roughness

is observed with either an 80 or 120 mesh diamond

wheel with 100 concentration.

7 17 27 37 Flexure Strength (MPa)

47

Figure 12: Probability plot of effects showing

significant factors of wheel speed and

concentration on strength of machining samples.

99.9

99

95

>. 80

o !5 o^ « 50 O n

o^

o£ 20 0

5

o^

1

0.1

, 1 , , 1 , . 1 , . 1 , , 1 , , 1

-9 -6-3 0 3 Residuals

6 9

Figure 13: Probability plot of residuals showing

linear relationship.

Average Pdwer

Average power recorded during grinding was

significantly influenced by removal rate, diamond

concentration, and mesh size (Figure 16). A larger

amount of power (12 watts/mm) was used for

grinding at a high rate, the larger diamond

concentration of 100, and the smaller diamond size

of 120 mesh. This result is similar to the creep feed

grinding in which concentration and mesh size are

significant, but wheel speed, and not removal rate,

is also influencing the power. The lowest power

condition is predicted to be at a material removal

Figure 14: Plot of effects on surface roughness,

showing concentration, mesh size of diamond

and a mesh-concentration two-factor interaction.

Figure 15: Plot of the two-factor interaction of

mesh and concentration for surface finish.

I

I

I M II

i M I

I I I

I

I M II I

I

I I I

I

I I I

I

I I I

I M I I

I M I

I M I «> «) .40 .20 0 20 40 ao to 100 120 140 ieo

Effect

Figure 16: Plot of effects on average power with

scaled t distribution to show extent of

experimental error. 229
rate of 6 mmVmm-s, a 120 mesh wheel with 75

concentration diamond.

Average Vertical Force

The variables of machining rate, diamond

concentration in the wheel, and diamond mesh size

all influence significantly the vertical force

measurement. The data presented in Figure 17

show that a greater force is measured when the rate

is 12, the concentration is 100 and the mesh size is 80.

Wheel Radial Wear

The mesh size and the concentration of the

diamond, along with the mesh size-concentration

two-factor interaction significantly influence the wear of the wheel (Figures 18, 19 and 20). These

factors are identical to those observed for the creep

feed grinding (Figures 8, 9 and 10). The better

R

16 20

Figure 17: Plot of effects on vertical force with a

scaled t distribution plot.

MC

I M I

I M I

I M I I M I

I

-11 -3 1 6 Wear Qjtn)

13

Figure 19: Probability plot shows significant

effects of mesh size, concentration and meshconcentration two-factor interaction on wear.

Wear(Mm)

120 Mesh

Figure 18: Plot of effects on wheel radial wear

with scaled t distribution plot.

Table VI: Comparison of Pendulum & Creep Feed Results (2nd Experiment)

Figure 20: Plot of the two-factor interaction on

wear.

Grind

Mode

Cutting

Rate

(mmVmm-sec)

Average

Strength

(MPa)

Surface

Roughness

ijjm)

Average

Power

(watt/mm)

Vertical

Force

(N/mm)

Wear

(^m)

Pendulum 1.2 800 0.6 48 13 6

Pendulum 12.0 772 0.6 422 66 4

Creep 1.2 827 0.5 74 85 8

Creep 12.0 1 786 0.5 564 87 3

230
grinding condition (lower wear) predicted from this

study would be either 80 or 120 mesh wheel at 100

concentration level of diamond (Figures 18, 19 and

20).

Second Experiment

The grinding machine dynamics for this experiment were different from the first experiment

in that the wheel spindle bearings were changed

from air bearings to steel ball bearings. Based on

the results of the first experiment, the wheel

specification was fixed to a mesh size of 80/100 and

100 concentration. The wheel speed was set to 45

m/s. The results of this 2|u„ design are listed in

Table VI. The analysis of variance of the means

table is shown in Table VII and Figure 21 and the

result is that neither of the factors, mode of

machining nor removal rate, had a significant effect on the strength of the machined samples.

Discussion of Results

A summary of the significant effects of the

machining variables on the results is shown for the

first experiment in Table VIII. One sees that for

strength, the mesh size was important for creep feed

at the lower removal rate (0.6 and 1.2) whereas the

wheel speed and diamond concentration were

significant effects for the pendulum feed at the

higher removal rate (6 and 12). Roughness is seen

to be affected by diamond concentration and mesh

size for both types of feed. Power is influenced by

concentration and mesh size for both types of feed,

Table VII: Analysis of variance table (second

experiment) for comparison of samples machined

witln ball bearings.

Source DF SS MS F

Factor 3 222 74 0.79

Error 25 2331 93

Total 28 2553

0.509

and by speed of the wheel for creep feed, and

removal rate for pendulum feed. The average force

is affected by rate and mesh size for both types of

feed, and the pendulum feed is also influenced by

diamond concentration. Wear of the wheel is influenced significantly by mesh size, diamond

concentration, and a two-factor interaction of mesh

size and concentration for both types of feed.

In the second experiment in which the wheel

spindle bearings were changed to steel ball bearings,

it was found that the mode of grinding (creep or pendulum) and the machining rate had no

significant effect on the strength of the machined

samples. This is seen in the analysis of variance of

the means table shown in Table VII and Figure 21.

The F test listed in the figure indicates that the

results could be expected by chance about 50% of

the time. Therefore, it is concluded that there are no significant differences in the groups of samples

tested.

A summary of the preferred conditions for

machining NCI32 Silicon Nitride ceramic for

improved properties and machining efficiency are

Pandulum Low Rate aoOMPa

Cro«p Low Rate 824MPa

Pendulum

High Rats 774MPa

Creep

High Rate 789MPa

700 750 800

Flexure Strength (MPa)

860 900

Figure 21: Plot of the 95% confidence intervals

for the averages of the four groups.

Table VIII: Significant Machining Factors (1st Experiment)

Results: Strength (MPa)

Surface Roughness Power

(Watt/mm)

Vertical Force (N/mm)

Wear (pm)

Creep Feed

(0.6 & 1.2 removal rate) Mesh Size

Cone.

Mesh Size-Conc.

Speed

Mesh Size

Cone.

Removal Rate Mesh Size

Mesh Size

Cone.

Mesh Size-Cone.

Pendulum Feed

(6 & 12 removal rate)

Wheel Speed

Cone.

Cone.

Mesh Size

Mesh Size-Conc.

Removal Rate Mesh Size

Cone.

Removal Rate Cone.

Mesh Size

Mesh Size

Cone.

Mesh Size-Cone.

231
listed in Table IX. The table lists only the

significant factors and the other factors can be

chosen for convenience or economy. One notices

that the conditions leading to higher strength will not lead to lower power usage.

A comparison of the strength of samples from

all the second experiment (ball bearings) and the

baseline data for the billet of silicon nitride from

which samples were cut shows that there is no

significant difference between these two groups of

Pendulum -i- Creep

Steel Ball Bearings

(29 Samples)

Baseline

(10 Samples)

750 775 800 825 353

Flexure Strength (MPa)

875 900

Figure 22: Plot of the 95% confidence intervals

for strength averages of pendulum and creep

feed data (29 samples) with baseline strength

data (10 samples).

Table X: Analysis of variance table for

comparison of samples machined with air bearings and steel ball bearings.

Source DF SS MS F p

Factor 1 319 319 2.81 0.102

Error 37 4200 114

Total 38 4519

samples. Table X and Figure 22 have an analysis of

variance of the means chart which indicates no

difference. However, if we compare the averages

from the first and second experiments v^th the

baseline data for the billet material we see in

Table XI and Figure 23 that there are significant

differences in these comparisons. The data on

strength of samples machined with an air bearing

spindle are significantly lower than the strengths of

samples machined with the ball bearing spindle.

However, the averages reflect the diverse machining

PwKfukjm + CrMp

Sto*) Ball BMrings

(29 StunplM)

B&irine Str»ngtti Stsel B«l! Bearings

(10 SamplM)

Cre«p

Air Baving*

(16 Sampiss)

Pandulum

Air Bearings

(16 Samples)

750 800

Flexure Strength (MPa)

850 900

Figure 23: Plot of 95% confidence intervals for

strengths of several groups.

Table X!: Analysis of variance table for

comparison of all strength data with baseline

strength data.

Source DF SS MS E E

Factor 3 3069 1023 11.55 0.000

Error 67 5937 89

Total 70 9007

Table IX: Best Factor Levels for Machining of NCI 32 SigN^ Ceramic (1st Experiment)

Results: Higher

Strength

Lower

Roughness

Lower

Power

Lower

Force LowerWear

Creep Feed

(0.6 & 1.2) 80 Mesh Size

... 1 80 Mesh Size

75 or 100 Cone.

[OR]

120 Mesh Size

100 Cone.

120 Mesh Size

75 Cone.

30 m/s

120 Mesh Size

1 .2 Removal Rate

80 Mesh Size

75 or 100 Cone.

[OR]

120 Mesh Size

100 Cone.

Pendulum Feed

(6 & 12)

100 Cone.

45 m/s

80 Mesh Size

100 Cone.

[OR]

120 Mesh Size

100 Cone.

120 Mesh Size

75 Cone.

6 Rate

120 Mesh Size

75 Cone.

6 Removal Rate

80 Mesh Size

100 Cone.

[OR]

120 Mesh Size

100 Cone.

232
conditions included in the experimental design. Six

of the sixteen experiments in the creep feed group were equal in strength to the baseline data, and one

of the sixteen experiments in the pendulum feed

group was equal to the baseline data.

Machining with the air bearing spindle appears

to have a deleterious effect on strength compared

to the machine with the steel ball bearing and this may reflect the greater stability available in the steel

bearing. It is generally recognized that the grinding

of ceramics requires very stable machines.

Conclusions

1. Designed experiments with a machine using an

air bearing spindle showed that there were

significant effects of machining variables on the

strength and surface roughness of the machined

silicon nitride, as well as the power and force

observed and the wear on the grinding wheel.

2. Two or three significant variables were observed

for each result, with a two-factor interaction of

mesh size and diamond concentration observed

for surface roughness of the workpiece and wear of the grinding wheel.

3. The second experiment with a machine using

ball bearings in the spindle showed that the

mode of machining (creep and pendulum) and

removal rates (1.2 and 12 mm^/mm-sec) are not

significant factors in the strength of the

machined silicon nitride.

4. Baseline strength data generated from samples

cut from the NC132 Silicon Nitride billet were

the same as strength measured from samples

machined in the second experiment using the

ball bearings.

5. Averages of strength from the first experiment

using the grinding machine with the air bearings were significantly lower than strengths measured

in the second experiment using the machine

with the ball bearings. Creep or pendulum

modes of machining had no effect on strength

of machined silicon nitride, nor did removal

rates.

References

1. R. L. Allor, R. JL Govila, and T. J. Whalen,

"Influence of Machining on Strength Properties

of Turbine Materials", Ceramic Engineering and

Science Proceedings, Vol.3, pp.392, 1982.

2. The Science of Ceramic Machining and Surface

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

(Eds.), Nat.Bur.Standards Special Publication

No.348, pp.416, 1972.

3. The Science of Ceramic Machining and

Finishing 11, R. W. Rice and B. J. Hockey

(Eds.), Nat.Bur.Standards Special Publication

No.562, pp.518, 1979.

4. W. Konig and J. Wemhoner, "Optimizing

Grinding of SiC", Ceramic Bulletin, V0L68,

N0.3, pp.545, 1989.

5. W. H. Daniels, A. C. Hawkins, et al, "Superabrasives for Ceramic Grinding and

Finishing", Advanced Ceramics Conference,

Society of Manufacturing Engineers Paper EM

89-125, Feb.1989.

6. T N. Farris and S. Chandrasekar, "On the

Characterization and Control of Surface

Finishing Damage in Ceramics", Journal of

Mechanical Working Technology, No.20, pp.69,

1989.

7. K. Subramanian and P. P. Keat, "Parametric

Study on Grindability of Structural and

Electronic Ceramics-Part I", Machining of

Ceramic Materials and Components, American

Society of Mechanical Engineers, PED-Vol.l7,

pp.25, Nov. 1985.

8. P M. Stephan and S. Tuffs, "Machining of High

Performance Silicon Nitride Mechanical Face

Seal Rings", Machining of Ceramic Materials

and Components, American Society of

Mechanical Engineers, PED-Vol.l7, pp.75,

Nov. 1985.

9. H. Kawamura, "Study of Grinding Process and

Strength for Ceramic Heat Insulated Engine",

Society of Manufacturing Engineers

Superabrasives '91 Conference Proceedings,

June 1991.

10. G. E. P Box, W. G. Hunter, and J. S. Hunter,

Statistics For Experimenters, John Wiley and

Sons, New York, New York, 1984.

233
11. D. C. Montgomery, Design and Analysis of

Experiments, 2nd Edition, John Wiley and Sons,

New York, New York, 1984.

12. K, Unno and T. Imai, "Performance of

Diamond Wheel in Grinding Ceramics",

Proceedings of 6th International Conference on

Production Engineering, Osaka, pp.26, 1987.

13. G. D. Quinn, E I. Baratta and J. A. Conway,

"Commentary on U.S.Army Standard Test Method for Flexural Strength of High

Performance Ceramics at Ambient

Temperature", Army Materials and Mechanics

Research Center Report AMMRC 85-21,

AD A160873, August 1985.

234
DAMAGE DETERMINATION AND STRENGTH

PREDICTION OF MACHINED CERAMICS BY X-RAY

DIFFRACTION TECHNIQUES

W. PFEIFFER and T. HOLLSTEIN

Fraunhofer Institute fur Werkstoffmechanik, Freiburg, Germany

The complex relationship between machining parameters and strength-controlling surface

properties of high-strength ceramics has been investigated using newly developed x-ray

diffraction techniques for the determination of profiles of depth of machining induced

plastic deformation and residual stress. The resultmg mformation is compared with

results of strength tests performed at different loading conditions and SEM observations.

A fracture mechanics approach, which correlates fracture stress and nondestructively

determined near-surface properties (residual stress and plastic deformation) is used to evaluate defect size in the surface layer and - as a consequence - to predict fracture

strength of machined ceramics. Typical cases of machining like grinding and lapping

applied to silicon nitride and alumina are discussed and examples of successfully

predicted strength data of machined alumina bars are presented.

Introduction

The high surface quality required for components

of structural ceramics can often be achieved only

by an additionally applied machining process.

This strongly influences the strength and service

behaviour of the components and, in addition,

t causes a considerable increase in manufacturing

I cost. Therefore, there is a need for nonI destructive methods to determine the surface properties of machined components and for models

to predict the fatigue and fracture behaviour in

order to minimize cost of machining procedures

and optimize the quality of components.

Common and easily determinable characteristics

of machined surfaces of ceramics like SEM

micrographs and roughness values may lead to erroneous assessments of the strength behaviour

of the machined parts. This is demonstrated in Fig. 1, where SEM micrographs, roughness values

(average peak-to-valley height, R^) and

characteristic bending strength data for silicon

nitride samples lapped with two different grain

sizes of boron carbide are compiled: In spite of

the 'rougher' and obviously more damaged

surface of samples lapped with the coarse grained abrasives (Fig. la) these samples show

a higher bending strength than those samples

which have a smoother surface due to lapping

with a fine grained abrasive (Fig. lb). Similar

results were obtained for lapped partially

stabilized zirconia [1].

More meaningful methods like fractographic

investigations on critical surface flaws or nondestructive observations using, e.g., acoustic or thermal microscopy [2, 3] suffer from

conflicting demands for both an utmost global

investigation of the surface region relevant for

surface behaviour and highest local resolution in

order to detect the small machining induced

flaws.

235
Fig. 1: SEM micrographs and roughness values

of lapped silicon nitride surfaces lead to erroneous assessments of the strength:

(a) : 100|am abrasive, = 11.5}im,

characteristic bending strength =

587 MPa

(b) : 16|am abrasive, = 7.9|im,

characteristic bending strength

552 MPa

The aim of the presented investigation was,

therefore, to

i) evaluate and improve nondestructive methods

for determination of machining induced

residual stresses and damage,

ii) quantify the effect of these near-surface

characteristics on the strength, and to

iii) validate these methods by comparison of

experimentally determined strength with

predicted strength data.

Materials and Machining

Flat samples of sintered pure (99.7%) alumina

and Y-;03/Al203 (5%, each) doped sintered

silicon nitride with the dimensions 50x50x3mm

were investigated in lapped and ground conditions. Lapping was performed at different normal

pressures F"^ ranging from 0.02 MPa up to 0.06 MPa and at average surface speeds Vl of

0.9 m/s and 1.8 m/s using boron carbide abrasi- ves (30 Vol%) with average grain sizes dj.; from

9|im up to 100|am. Grinding (up-cut) was performed using a diamond grinding wheel type

D151K+C100 (151 = average grain size in jam)

and infeeds a^ from 0.03mm up to 0.5mm, tan- gential feed rates Vf; from lOmm/s up to 167mm/s and cutting speeds v^ of 25m/s and

50m/s, respectively. The specific material removal rates Q\ during grinding cover the range

from 10mm' /(mm-s) to 40mm' /(mm-s).

Additionally, alumina rods (length = 165mm,

diameter = 6mm) were prepared for tensile tests by cylindrical grinding using three different grinding conditions:

(1) Wheel type=D76K+C75, Q\ = 0.001mm'

/(mm-s), a^ = 0.015mm, Vft = 0.08mm/s.

(2) Wheel type=D151V+C100, Q',=0. 165mm'

/(mm-s), a^ = 0.025mm, Vf; - 6.6mm/s.

(3) Wheel type=D151V+C100, Q'^ = 0.5mm'

/(mm-s), a^ = 0.025mm, Vft = 20 mm/s.

Strength Tests

Bending strengths of the machined flat samples were determined using the concentric-ring test with a loading ring diameter of 8mm and a supporting ring diameter of 40mm. According to the

linear theory of plate bending, a constant maximum biaxial stress state prevails in the surface

236
within the loading ring. The advantage is, that

effects of unproperly prepared edges of the

samples can be neglected. Due to the symmetry

of the stress state the 'weakest' direction of the

surface is expected to lead to failure. Nine

samples with identically machmed surfaces were

broken at stress rates of about 90 MPa/s.

Tension strengths of the alumina rods were

determined using a universal testing machine

with the so called 'supergrips' for momentum

free loading of ceramic specimens with

shoulders. Seven to ten samples with identically

machined surfaces were broken at stress rates of

90 MPa/s.

The statistical distribution of the measured

strength values was evaluated by means of the

two-parameter Weibull distribution with the

parameters characteristic strength o„ and Weibull

modulus m.

Nondestructive Determination of Machining

Induced Residual Sti^sses

The assessment of defects within the machining

influenced surface layer during loading needs the

knowledge of the contribution of residual

stresses to the overall stress state acting at flaws.

Mean values of near-surface macroscopic resi- dual stresses were determined by x-ray diffraction using the well-established sin^V|/-technique

[4]. The distribution of residual stresses versus depth was evaluated using a newly developed

application of grazing incident x-ray diffraction

techniques [5].

Figure 2 shows the nondestructively determined

distributions of machining induced residual

stresses for alumina and silicon nitride samples.

They were lapped using different grain sizes of

the abrasives: Fine grained abrasives lead to high

compressive residual stresses at the surface.

Coarse grain abrasives lead to smaller residual

stresses at the surface but to thicker surface

layers containing compressive residual stresses (especially for silicon nitride).

Alumina

(0 Q. -500

(0

0)2 -1000

"aI -1500

(0

0)

-2000

f \\\

' \V ^= (To: \ \ \ \ \ 16 //m 352 MPa

\ 64 325 MPa

\ 100 t"^ 318 MPa

1 5 10

depth in

Silicon Nitride

15

n

Q. -500 c

(0

(09 -1000

0)«3

-1500

"35 0)

-2000

/ —

\/)\\

- Jo:

552 MPa

J \ 64 f/m 563 MPa

1 \ 100 /im 587 MPa

5 10

depth In

15

Fig. 2: Lapping of alumina and silicon nitride

using abrasives with a large average

grain size dj^ leads to deep-reaching

compressive residual stresses which

overcompensate the effect of machining

induced damage in the case of silicon

nitride (see bending strength data a^).

The experimentally determined bending strength

data for alumina qualitatively show the expected

strength reducing-effect of coarse grained abrasives. In contrast to alumina the compressive

residual stresses in silicon nitride obviously over- compensate the effect of machining induced

damage on the bending strength.

237
Evaluation of Defect Sizes

The quantitative effect of machining induced

residual stress states and the amount of

machinmg induced damage is evaluated by

means of fracture mechanics calculations;

The size, distribution and position of strength

controlling defects are normally not accessible to direct measurements. Since they are important

for the strength, 'effective' flaw sizes were

calculated by means of fracture mechanics

methods from the bending strength data and the

experimentally determined residual stress distri- butions. These 'flaw sizes' are, of course, affected by the used simplifications but nevertheless

they characterize the amount of damage.

The determination of the 'depth' of flaws is based on the equation describing the fracture

toughness of brittle materials:

The critical fracture stress c^. is evaluated from

the characteristic fracture stress obtained

through the bending test and from the residual

stress Gres determined by the x-ray diffraction

method. For values of fracture toughness Kj^, of

the materials tested those values are used, that

take into account the absence of 'R-curve effects' in the case of small flaws (AI2O3:

Kj =2.4 MPa-m"' [6], S13N4: Kj, 1.9 MPa m"'

[7]). The form factor f takes into account the

ratio of length/depth of flaws and is 1.12 in case of a surface scratch and 2/71; in case of a half-penny shaped surface flaw. The critical depths of flaws are calculated by numerical integration of the nonlinear loading and residual

stress distributions using the weight function

method [8].

Figure 3 shows an example of calculated depths

for half-penny shaped flaws of lapped alumina

and silicon nitride, for two cases:

i) effect of machining induced residual stresses

is taken into account, and

ii) effect of machining induced residual stresses

is not taken into account.

Especially in the case of silicon nitride the

results point out that calculated flaw sizes are

larger if the strengthening effect of compressive

residual stresses is taken into account. The

expected larger flaw sizes (when using coarse grained abrasives) are confirmed by those calculations that include the effect of residual stresses.

Alumina

100 64 16

grain size of abrasive in

Silicon Nitride

100 64 16

grain size of abrasive in ^ivci

Fig. 3: Lapping induced damage is quantified

by calculated depths of flaws only if the

influence of residual stresses is taken

into account.

238
The comparison of the depths of flaws detected

by SEM fractography (see Fig. 4) with the calculated values additionally shows, that the calculated depths of flaws are of similar magnitude.

' Fig. 4: Micrographs of fracture surfaces with

machining induced flaws of alumina (a)

and silicon nitride (b). The flaw sizes

f are estimated on the basis of

half-penny shaped cracks.

I Nondestructive Determination of Machining

I Induced Damage

I Indentation experiments performed on different

1 brittle materials and theoretical investigations

' indicate the existence of a (material dependent)

relationship between the plastically deformed

surface layer and the flaw size (e.g. [9, 10]). The

presented method for the nondestructive

determination of machining damage is based on

the determination of the thickness' a.^ of the

micro-plastically deformed surface layers and the

establishment of correlations between these

thickness' and calculated flaw depths a^.

The thickness of the plastically deformed surface

layer can be estimated by a newly developed

application of grazing incident x-ray diffraction

techniques [5] using the effect of machining

induced lattice defects (e.g. dislocations) on

microscopic lattice strains.

Figure 5 shows such distributions of microscopic

strain within the surface layers of alumina and

silicon nitride samples which were ground using

different infeeds a„ different specific material

removal rates Q'^ and different cutting speeds Vg.

The effect of different grmding conditions on

micro-plastic deformation can clearly be separated by the x-ray diffraction method. The obtained thicknesses of deformed surface layers

correspond to the expectations regarding the

influence of different grinding conditions on the

elementary mechanisms of material removal.

Creep feed grinding conditions (large infeed aj

as well as high cutting speeds and small

tangential feed rates reduce the undeformed chip

thickness at the single cutting edge [11]. This,

combined with reduced local forces, results in

reduced depths of plastic deformation. The

bending strength data indicate that the depth of

plastic deformation is related to the amount of

machining induced damage and that grinding

compressive residual stresses (see below) do not

significantly influence the bending strength.

In order to validate the relationship between the

thickness of plastically deformed surface layers

and machining induced damage, the depths of

experimentally determined plastic deformation

and the calculated flaw depths are compared for

alumina and silicon nitride samples in various

machined conditions in Fig. 6.

239
Alumina

a.: Q' : v»: (To:

0.03 5 25 253

0.5 20 25 278

0.5 5 25 341

0.5 10 50 338

5 10

depth in pm

Siiicon Nitride

15

5 10

depth in ym

15

Fig. 5: Grinding using large infeeds

(mm), small material removal rates (mmV(mms)) and high cutting

speeds Vg (m/s) lead to small

micro-plastically deformed surface

layers and obviously small damage (see

characteristic bendmg strength data

(MPa)).

The depth of plastic deformation was determined

atEp=10-', the fl aw depth a, was evaluated using

characteristic bending strength data,

experimentally determined residual stress distributions and assuming half-penny shaped

flaws.

For alumina, a unique Imear dependence

between deformation depths and calculated flaw

depths can be established for both lapped and

ground samples. Few data points (marked by ^)

do not fit to the regression line: These data

points belong to those grmding conditions, where

fracture during the strength tests resulted in preferred orientation parallel to the grinding direction.

Alumina

30 E

a.c

J= 20

a.

o> o

c

o

'S 10 E

^

£ )K , effect of crack geometry

machining procedure:

/ ^ grinding ^ lapping

30 50 70 calculated flaw depth a^ in [im

Silicon Nitride

90

40 E

a.c

a 30

Q.«

C

o

flE

\-o«

20

10

I

In

f machining procedure:

13 grinding ^ lapping

10 20 30 calculated flaw depth a^, in40

Fig. 6: Correlation between nondestructively

determined deformation depths ap and calculated flaw depth a^. of

differently lapped and ground alumina

and silicon nitride, half-penny shaped

flaws are assumed in all cases.

The more realistic assumption of scratch-like

flaws (form factor f 1 . 1 2) in these cases would

lead to smaller calculated flaw depths and by

this improve the quality of correlation. For

machined silicon nitride, different (linear)

240
dependencies exist for ground and lapped

samples: Obviously the ratio of crack depth/

deformation depth is larger m the case of

grinding than in the case of lapping or the

grinding induced microcracks are more serious

(which means more scratch-like). In spite of

these uncertainties about the geometry of flaws

(which could not be clarified by SEM

investigations), the established relationships

between deformation depth and crack depth may

be used to determine machining damage

quantitatively by the advanced nondestructive

x-ray diffraction technique.

Alumina

O

E 6

machining procedure: B grinding ^ lapping

30 50 70 90 calculated flaw depth a^ in

Silicon Nitride

20 o

£ 16

12

8

4

machining procedure: B grinding ^ lapping

0 10 20 30 40

calculated flaw depth a^ in ym

Fig. 7: Correlation between nondestructively

determined mean values of deformation

and calculated flaw depths a,, of

differently lapped and ground alumina

and silicon nitride (half-penny shaped

flaws are assumed).

The determination of the distribution of

microscopic strain is a time-consuming

experiment, which (due to defocussing effects) is additionally limited to samples with simple

geometries. Therefore, a similar relationship

between (more easily determinable) mean values

of microscopic strains of the surface layer and

the calculated flaw depths is established in

Fig. 7. This is possible because the mean

microscopic strain of surface layers is dominated

by the depth of plastic deformation.

Effect of Residual Sti«sses and Damage on

Sti^ngth

On the basis of calculated flaw sizes a, (taking

the effect of residual stresses into account) and

the appropriate K,^ values, the critical fracture

stress values o, were calculated. Due to the

theoretically given dependence between and

a,, (see eq. 1) the critical fracture stress values

are plotted versus l/aj " and compared with the

experimentally determined characteristic bending

strength data a^, (Fig. 8).

The results show that the effects of machining

induced damage and residual stresses depend on

the type of machining and material:

- For lapped and ground alumina, the

contribution of machining induced

compressive residual stresses to the bending

strength is small, the bending strength is dominated by machining induced damage.

- For ground silicon nitride increasing

machining damage is combined with

increasing strengthening effects due to compressive residual stresses but these

residual stresses do not overcompensate the

effect of damage on strength.

- For lapped silicon nitride, the machining

induced compressive residual stresses compensate the strength reducing effect of

machining induced damage.

- With increasing amount of damage also

increasing compressive residual stresses are created due to machining in all cases

241
investigated. Thus an additional material and

machining dependent relationship between

micro-plastic deformation and residual stress states can be established [5, 8]. This allows

to quantify both the effect of damage and the

residual stresses on strength by determination

of just one of these surface properties.

Prediction of Strength

It was shown, that the strength dominating

properties (residual stresses and flaw sizes) of

surface layers of machined ceramics can be

evaluated quantitatively by nondestructive x-ray

diffraction methods.

Knowing the residual stress distribution within

the machining influenced surface layer and using

the established relationship between deformation

depth and 'crack sizes', it is possible to predict

the surface-strength of machined components.

The experimental validation of this thesis is performed using the relation between more easily

determinable mean values of microscopic strain

and damage (see Fig. 7) and the material and

machining dependent relationships between

fracture stress, characteristic bending strength

and amount of damage (see Fig. 8).

Lapping of Alumina

450

0.14 0.16 0.2

700

Q.£ 500 «

«)9

o

IB

300

100

0.1

Lapping of Silicon Nitride

resid. stress \ ^^^^

effect V*^^^

'B- bending strength —critic, tract, stress

0.15 0.2 0.25

a,(-^/2) ,n ^^(-1/2)

0.3 0.35

Grinding of Alumina

450

0.12 0.14 0.16 0.18 0.2

Grinding of Silicon Nitride

700

a.S 500

9

9tH- 300 3

100

resid. stress B' effect , >-TM~^

>- ^ ^ a, ^

"B- bending strength — critic, tract, stress

0.1 0.15 0.2 0.25

ac<

^'2> in /im<-^/2)

0.3 0.35

Fig. 8: Strength - damage diagrams showing the effect of machining induced compressive residual

stresses on the bending strength of alumina and silicon nitride, denotes the critical fracture

stress which equals the strength in the absence of residual stresses.

242
The following procedure for the prediction of

tension strength of alumina rods which were ground using three different grinding parameters,

therefore, consists of two steps:

i) From the mean values of experimentally

determined microscopic plastic deformation

of the tension rods the corresponding crack

depths are derived (Fig. 9a).

ii) From the estimated crack sizes the critical fracture strength and the additional

contribution of compressive residual stresses

to the overall characteristic fracture strength are predicted (Fig. 9b).

The results show that grinding with increasing

material removal rates (see 1^2^3, Figs. 9, 10)

leads to increasing flaw depths and reduced

strength. The small flaw depths evaluated for the

tension rods ground with the smallest material

removal rate, (1), are comparable with the

calculated flaw depth of bend samples which

were lapped with a fine grained abrasive (see

Fig. 3).

Fractographic investigations showed, that the

fracture causing defects of these samples are

'natural' defects. This means, that the effect of

machining induced damage is negligible,

whereas grinding at higher material removal

rates (2, 3) leads to failure due to machining

induced flaws.

The experimentally determined Weibull-distributions of the tensions rods are shown in Fig. 10

and the evaluated characteristic strengths are compared with the predicted values in Fig. 9. The predicted and measured strengths are in

good agreement. The small discrepancies

between predicted and measured strength data

may be due to the effect of the fracture relevant

volume on strength. The prediction is based on bend samples with a relatively small fracture

relevant volume but applied to tension rods with

a larger volume. This leads to a slight

overestimation of strength.

(a) Determination of Damage (b) Prediction of Strength

0.2 ~cracir depth in /im

Fig. 9: Prediction of the strength of alumina bars in differently ground conditions:

(a) : Evaluation of crack depths by the nondestructively determined plastic deformation.

(b) : Assessment of the strength dominating effects of damage and residual stress and

comparison with experimentally determined strength data.

243
1 0151 V+ C100, V - 35 m/»

1 3 2 1 0,: 0.025 mm <y f

/» - 3 : m - 14. a, - 247 MPa 7 1 L'

2 : m - 21. o, - 281 MPa r T ^' /o 4

076(91) K+ C75(100). 20 m/« o/ r *

0 I 0-015 I ' w ' Cf /• A

1 • m H m m M UPrt /

/o

/ / * # /

/o

/

*

/ / / /

/ '/ A V, - 0.08 mm/s / / • V, - 6.6 mm/i / , / O V. - 20 mm/s

100 200 300 400

Frocture stress in MPa

Fig. 10: Weibull distributions of fracture

stresses of alumina rods ground with

different material removal rates (due to different tangential feed rates VfJ. Characteristic tensile strength data (o^)

and Weibull moduli (m) are indicated.

Using the established relationships between

damage and residual stresses, respectively, and

strength, the tension strength of differently

ground alumina bars has been predicted with

high accuracy. J

Acknowledgement

% Machining of the ceramic materials was

performed in the Fraunhofer-Institut fiir Produktionstechnologie by J. Wemhoner, M. Popp and

E. Verlemann. Part of the investigations has been

sponsored by the Bundesministerium fiir Forschung und Technologie and the companies

M. Braun GmbH, Friedrichsfeld AG, Hutschenreuther AG, Instron Ltd. and Siemens AG.

References

Summary

The complex relationships between machining

parameters and strength-controlling damage and

residual stresses of alumina and silicon nitride

have been investigated using advanced x-ray

diffraction techniques, strength tests and fracture

mechanics calculations. Correlations between

micro-plastic deformation and amount of damage

can be established. They allow to determine

nondestructively and quantitatively the amount of

machining induced damage. The contribution of

machining induced residual stresses to the

strength of machined ceramics can be quantified

both by calculation of flaw sizes taking the

residual stresses into account and by establishing

material- and machining-dependent strength - damage - diagrams.

The investigations show that in the case of

lapped and ground alumina and of ground silicon

nitride bending strength is dominated by

machining induced damage. In the case of

lapped silicon nitride the effect of damage is compensated by machining induced compressive

residual stresses.

[1] W. Pfeiffer, "Influence of Machining-Induced Residual Stresses on the

Strength of Ceramics," Proc. of Conf. on j Experimental Mechanics, Copenhagen,

Denmark, pp. 1237- 1245, 1990.

[2] H. Reiter, M. Maisl, U. Netztelmann and

W. Arnold, "Nondestructive Testing of

Ceramics by Computed Tomography and ^ High-Frequency Ultrasonics," 4th

International Symposium on Ceramic

Materials and Components for Engines. Elsevier Science Publ. LTD, London, New | York, pp. 880-887, 1992.

[3] H.K. Tonshoff, E. Brinksmeier, and H.G.

Wobker, "Randzonenanalyse als Werkzeug , zur keramikgerechten Prozefigestaltung,"

DKG-Symposium "Bearbeiten von Keramik" . Bayreuth, 13.12.1992.

[4] P. Muller and E. Macherauch "Das

sinV-Verfahren der rontgenographischen

Spannungsmessung," Z. ang. Phys. , Vol.

13, pp. 305-312, 1961.

244
[5] W. Pfeiffer, "Characterization of NearSurface Conditions of Machined Ceramics

by use of X-ray Residual Stress Measurements," Residual Stress III. Elsevier

Science Publ. LTD, London, New York,

Vol. 2, pp. 607-612, 1992.

[6] T. Fett and D. Munz, "Evaluation of R-Curves in Ceramic Materials based on Bridging

Interactions", KFK-Report 4940. 1991,

ISSN 0303-4003.

[7] E. Sommer, R.Mohrmann and M.Rombach,

"Lifetime Prediction for Components of

Ceramics under Quasistatic and Cyclic

Loading," Xlth Intern. Colloquium on

Mechanical Fatigue of Metals. Kiev, USSR,

May 13-17, 1991.

[8] W. Pfeiffer and E. Sommer, "Bewertung der

Festigkeitseigenschaften Oberflachenbearbeiteter Keramikbauteile," Werkstoffe. DGM-Informationsgqsellschaft mbH, Oberursel, pp. 575-584, 1991.

[9] R.W. Rice, "The Effect of Grinding Direction on the Strength of Ceramics," The

Science of Ceramic Machining and Surface

Finishing II , B.J. Hockey and R.W. Rice

(Eds), National Bureau of Standards Spec.

Publ. 348, U.S. Government Printing

Office, Washington D.C., pp. 365-376,

1979.

[10] K. Tanaka, "Elastic/Plastic Indentation

Hardness and Indentation Fracture Toughness: The Inclusion Core Model," J. of

Mater. Sci. . Vol 22, pp. 1501-1508, 1987.

[11] W. Konig and J. Wemhoner, "Optimizing

Grinding of SiSiC," Ceram Bull. . Vol. 68,

No.3, pp. 545-548, 1989.

245

\

INFLUENCE OF MICROSTRUCTURE ON GRINDABILITY

OF ALUMINA CERAMICS

p. ROTH and H. K. TONSHOFF

University of Hannover, Hannover, Germany

This paper presents the results of grinding experiments with various grades of alumina ceramics.

When machined identically, the materials reveal systematical differences in grindability.

Grinding forces and energies are controlled by material grain size and toughness. Brittle as well

as ductile behaviour of alumina ceramics are identified in grinding. Material removal

mechanisms depend strongly on the microstructure of the materials. Grain size is identified as one dominant parameter. With decreasing grain size, plastic deformation becomes more

prominent over micro-brittle material removal. With a given material, two parameters determine

grit-work interaction: Micro-brittle fracture, ploughing, and cutting action of the grits are observed as a function of (1) grit depth of cut and (2) cutting point density. A thin layer of

heavily deformed material is frequently found to cover ground surfaces. Compressive residual

stress originates from the plastically deformed sub-surface zone. The magnitude of residual

stress is a function of the grit depth of cut. Over a wide range of conventional machining

conditions, grain size limits the obtainable surface finish.

Introduction

Modem structural ceramics consist of chemical

compoimds from the groups of oxides, nitrides, and

carbides. In engineering, they are valued for their

superior stif^hess/weight ratio, high hardness and wear resistance even at elevated temperatures. Yet, it

is these very qualities that limit the use of ceramic

components in machine design. Prepared from

powders by various compaction and sintering

techniques, ceramic parts can vary considerably in

size and shape. To meet close dimensionaJ.

tolerances, the more intricate parts have to be

finished by grinding, lapping and sometimes

polishing with diamond grit. Such abrasive

machining is critical in respect of cost and quality

control. The surface quality attained by finishing

processes can determine the performance of ceramic

components. This is a result of the low fracture

toughness of ceramics and their sensitivity to surface

damage. As modem ceramics are constantly

improving in quality, volume defects are reduced in

size and number. Thus, it can now be the surface

flaws that determine the mechanical strength of a finished part. In order to determine favourable

grinding conditions which retain unimpaired surface

integrity, the process of material removal must be

investigated systematically.

While in grinding of ceramics the aspects of

machining parameters and tool specifications have

been examined closely [1-6], it is of equal interest to study the influence of varying microstmcture with

ceramic materials. As of today, industrial standards

such as those for metals do not exist for this rapidly

developing group of materials. It is therefore

desirable to predict the machinability of a given

ceramic from its mechanical properties such as hardness, bending strength, etc. These qualities of a

material depend on its chemical composition as well

as on its microstructure, i.e. the size and distributions

of grains and pores. With multi-phase materials, the

interaction between the different phases has to be

considered as a further influence.
Experimental procedure

The investigations described in this paper have been

carried out to determine the influence of ceramic

microstructure on grindability. A single-phase

alumina ceramic with varied grain size has been

chosen for investigation. For the grinding tests, a

series of high-purity alumina ceramics were

produced under controlled conditions (99.98%

AI2O3). No additives were used for sintering of five

grades at temperatures ranging from 1300°C to 1700°C. In all cases, almost identical densities of at least 98.5% were achieved. Equal in chemical

composition, these ceramics differed in

microstructure and properties due to the varied

sintering cycle. In accordance with the sintering

temperatures the alumina grades were designed as TM 30.. .70. Their mechanical values are given in

Table 1 and Figure 1 below.

With rising sintering temperature, average grain size

increases from 0.8 to 12 (im. Correspondingly, the

hardness of the five alumina grades is found to decrease from 19.4 GPa to 15.8 GPa while fracture toughness rises slightly to remain at a constant level with the larger grained varieties. Bending strength

decreases from a maximum at the second finest alumina grade. The important differences in hardness

and strength with varying microstructure can be

expected to notably influence the grindability of the

different aliunina variefies.

Table 1 : Material properties

Material Sintering Grain Hardness Fracture Bending

grade temp. Tg size dg HV 10 toughness Kj^, strength o^^

[°C] [jim] [GPa] [MPa-mi/2] [MPa]

TM30 1300 0.83 19.4 3.68 320

TM 40 1400 2.2 19.2 3.80 380

TM 50 1500 4.6 17.7 4.10 340

TM 60 1600 7.8 16.5 4.15 300

TM 70 1700 12.2 15.8 4.07 300

Surface grinding tests were performed using phenolic

resin bond diamond wheels with grit sizes ranging

from D7 (5-10 ^im) to D151 (125-150 ^im). Truing was performed by a brake-type rotary truer with

silicon carbide wheels. The grinding wheels were

then dressed with a porous alumina block at conditions optimized separately for each wheel.

Earlier testing had shown that the grinding forces

stabiUzed after an initial phase of rapid wear after dressing. Consequently, each test was carried out

until a specific material volume of 3(X) mm^/mm had

been removed. Throughout the tests, low viscosity

mineral oil was used as a coolant.

In existing studies of brittle workpiece materials, the

depth of cut of an individual grit, h^^, has been

demonstrated to determine the transition from microbrittle to ductile material removal [5]. To study the onset of the different mechanisms, the grit depth of

cut was varied via the wheel depth of cut, a^, the feed

speed, Vft, and the grinding speed, v^,. The strongest

influence on hj^^^ results from the density of cutting

edges C in the abrasive process as demonstrated by

equation (1) below. Variation of this parameter has

been accounted for by the use of different grit size

grinding wheels of uniform diamond content.

h - M ^ft fC

"^^"^C-r'v, 'Vd [7] (1)

The tests encompassed a range from conventional to creep feed grinding conditions. The parameters used

in the grinding experiments are summarized in

Table 2 below.

248
g/cm

3.9

c- 3.8

C

o

^ 3.7

CD "O 3.6

3.5

2100

HVio

1900

1800

°

1700

1600

1500

1300 1500°C1700

sintering temp, tg

V

1300 1500°C1700

sintering temperature t^ source R.Telle.MPI/PML Stuttgart

.9 jc

c ^

-o 2!

. -4-' O

Q.

1

500

MPa

AOO

350

300

250

200

1 /

1 1

1 1 \1 1 1

1

1300 1500°C1700

sintering temperature t^

(ISB method, n=16-18)

(/) <uc - Jl

D)

O

0)

l_Z3o

a

A.6

MPa/m

^.2

A.O

3.8

3,6

3,2

1

1

1

1300 1500°C1700

sintering temperature t^

321/^836 ©IFW

Fig. 1: Mechanical properties of workpiece materials

Table 2: Process settings for the grinding tests

Grinding wheels lAl-300-10-3 D7...D151 C50,

V^' = 300 mm3/mm (50 mm^/mm for D20B, 10 mmVmm for D7)

Truing rotary brake type truer, silicon carbide wheels 37 C 60.. 80 N5V

Dressing porous alumina block, = 0.5... 10 mmVmm-s

Grinding speed V(, = 10.. .25 m/s, up-grinding

Work feed speed Vft= 10...10'000mm/s

Wheel depth of cut ae = 0.001. ..2.0 mm

Coolant straight mineral oil (v = 4. 15 mmVs at 40°C),

q = 30 1/min

249
In order to analyse the interaction of the grit with the

workpiece surface more closely, plane alumina

specimens were polished, sputtered with gold, and

then ground in a slightly tilted position. By this, the

wheel depth of cut increased gradually from zero to a preset maximum and the average grit depth of cut along with it.

During grinding, horizontal and vertical grinding

forces were constantly monitored by a piezoelectric

dynamometer and recorded by computer.

Subsequently, the specific normal and tangential

grinding forces were computed off-line. From the

tangential grinding force, the specific grinding energy

Cc was calculated. It denotes the energy necessary to remove a unit volume of material. It is derived from

the tangential grinding force and the settings of the grinding process by

= (Ft Vc) / (ae • Vft). (2)

Surface finish (average peak-to-vaUey roughness) was measured by stylus tracing to characterize the

ground ceramic surfaces. Measurements were

performed in several directions relative to the

grinding direction. Micrographs of the groimd

surfaces were prepared by SEM.

To determine residual grinding stress x-ray stress analysis was carried out with Cr-K„ radiation on the

(030) plane of alumina. Residual stress was

measured across the grinding direction.

qnndinq : conditioning ; D91K + 888 NY C50 SiC - wheel

- 35 m/s V(.(j = 35 m/s

Qg = 10 nm Qed - 10 fim

Vf, = 2.5 mm/s Qgj, = 5 mm /mms

V;^ = 0 mmVmm alumina block

coolant : mineral oil. v = ^.15 mmVs at40°C

= 1500°C

sintering temp.tg = 1300°C = 1600°C

= uoo°c t. = 1700°C

321/U2Z,® IFW 7682

Fig. 2: Surfaces of various alumina ceramics ground at identical conditions

Results

As expected in view of their mechanical values, the

five alumina varieties behave very differently in

grinding [8]. Figure 2 shows SEM micrographs of

ground surfaces. These surfaces were ground

identically at a low grit depth of cut in order to

facilitate ductile material behaviour. Most

importantly, the ground surfaces exhibit different

mechanisms of surface formation.

250
With the fine-grained varieties TM 30 and 40, the

ground surfaces appear covered with smooth grooves.

Some lateral pile-up of displaced material can be

observed. The regular groove pattern is broken by

two different types of cavities. Small pores form as grains are broken from the surface imder the action of

a grit. Larger, flake-like particles are removed by

chipping. Some of these have only partly separated

from the substrate and remain on the surface close to their original position.

With the increase in grain size, the surfaces become

continuosly rougher. The removal of single grains

dominates over groove formation. The surfaces of the coarser alumina varieties TM 50 and TM 60 reveal

Fig. 3: Ground surface of fine-grained sample

Surface roughness was measured to characterize the

ground ceramic surfaces. Figure 4 gives the roughness of the various materials in a polar

diagram relative to the grinding direction.

The roughness measurements confirm the visual

impression obtained from the SEM pictures in Fig. 2. Roughness increases monotonously with grain size. In TM 30 and TM 40, roughness values are significantly different along and across the grinding

direction. This reflects the observed formation of

grooves. The directionality of surface roughness

many near-perfect grain facets. Finally with TM 70,

transgranular fracture is frequently observed. Some of

the remaining plateaus bear traces of plastic

deformation as with the fine-grained materials.

In a detail SEM micrograph of a ground TM 30

sample (Fig. 3) the sub-surface structure becomes

visible at the edge of a cavity. The top layer is heavily

deformed by ploughing and appears unstructured at

its broken rim. Its thickness can be estimated at approximately 0.5 |im. Below this layer, the original

material microstructure made up of regular submicron alimiina grains is retained. Visible grain facets suggest that fracture has proceeded predominantly by

an intergranular process.

88 NYB C50

D m/s

0

5 mm/s

el

I m/s

I \im

) mm^/mm-s

1 block )

jteriol :

ity alumina.

321/2065 © IFW 7923

gradually disappears with larger material grain sizes

to approach isotropic roughness as in a lapped surface

with TM 70.

Figure 5 represents the roughness values determined on samples of TM 40 ground under various creep

feed and conventional surface grinding conditions.

With alumina, unlike ductile woric materials, no direct

connection exists between the roughness and the

grinding conditions. Roughness rises only

insignificantiy over a wide range of chip cross

sections.

251
work material:

high purity alumina

(TMDR 30 ...70)

g rindin g o peration:

surface grinding

Vg = 30 m/s

Vjj= 100 mm/s

a,= 10 ^m

coolant:

mineral oil tool: D91K+888NYC50

average grain

size dg |im 0

angle if relative to

grinding direction

10 ^im 8 0 6 8 |im 10

321/Z,831 ©IFW

IgQo

Fig. 4: Directional surface roughness

N

2 ^

CD

CD 2

O)

O >N

I- Q)

Q) =

> O

O >

workpiece material

grinding

tool

high-purity alumina peripheral surface grinding D91 C50

d =2.1|im V(. = 30 m/s 5^,= 5 mm^/mm s

vt, = 2.5/100 mm/s =300mm3/mm

Oe = var.

8

6

0

creep feed grinding surface grinding

0,5 1 2 5

non - dim. chip cross section

10

heq

/10

0

1*0

321/A832 ©IFW

Fig. 5: Roughness as a function of the grinding conditions

252
The grit depth of cut h^^^x found useful in

metal grinding to predict the roughness of ground

surfaces [9, 10]. The parameter of hg^/lg on the x-axis

of Fig. 5 denotes the dimensionless theoretical chip

cross section which is closely related to h^^. While

hj^ax is the truly significant quantity in a kinematic

sense, its calculation demands detailed knowledge

about the spacial distribution of the cutting edges and

about their profiles. To facilitate the comparison of

grinding conditions, here the parameter of hg^/lg is used instead which represents the dimensionless cross

section of cut of the individual grit or cutting point.

It is readily derived from the settings of the grinding

process by eq. (3) and reqires no experimental input.

'eq _

(3)

The quantity of hgq/lg can be used to characterize a

set of grinding parameters provided that the spacing

of the cutting edges remains constant. For a given

grinding wheel and identical dressing conditions, this approximation appears justified. However, different

grinding wheels cannot be compared using hg^/lg.

The specific grinding forces F^ and Ft are important

quantities to characterize a grinding process. The

grinding forces vary strongly between the five

alumina grades. The largest forces are found with the

smallest grain size materials (Fig. 6).

US

O

100

N

mm

10

c

C

o

o

0)Q.

1.0

0.1

creep feed grinding

Vjt = 2.5 mm/s

conventional surface

grinding v,.=100mm/s

0.^ 0.63 1. non - dim. cross section of chip -J^-^/ 10

Fig. 6: Normal and tangential grinding forces

work material : high purity alumina

(TMDR)

g rindin g_operation:

surface grinding

= 30(35) m/s

v,, = 2.5/100 mm/s

Oe = var. coolant;

mineral oil tool : D91K + 888NYC50

Vw = 300 mm^/mm

321/57^5 ©IFW

In creep feed grinding, the specific normal force F^' increases only slightly over a wide range. A material

influence is hardly detectable. In ordinary surface

grinding, the normal forces of the different alumina

grades separate systematically with increasing chip

cross sections. The larger grain sizes produce lower

grinding forces. At the upper end of the range of the

tests, the normal forces of the fmest and the coarsest alumina grade reach a ratio of three.

Quite similar at low depths of cut, tangential forces Fj' of the various work materials rise at different rates in

creep feed grinding. Again, the forces are correlated

inversely with grain size. Alumina with a grain size

of 0.8 ixm requires four to five times the tangential

253
grinding force of the 12 |im variety. This trend can

equally be found in surface grinding over the whole

range of cutting depths.

The specific grinding energy e^ derived from the

tangential grinding force by eq. (2) depends of the

grain size in the same characteristic fashion. In Fig. 7, Cc is plotted versus the dimensionless chip cross section with the grain size included as an additional

parameter.

cross section of chip Hfl / io'

work material

high purity alumina

(TMDR)

grinding operation:

surface grinding

v^ = 30(35) m/s

Vj, = 2,5/100 mm/s

g, = var. coolant. minerol oil tool : D91K + 888NYC50

Vw = 300 mm^/mm

321/^868 ©IFW

Fig. 7: Specific grinding energy

As in metal cutting, the energy demand for the

removal of a unit volume of material decreases with

increasing chip thickness. Grain size, however, exerts an even stronger influence on the specific grinding

energy than the input parameters of the grinding

process. Except for very small material removal rates

in creep feed grinding where the curves for the

different materials approach an identical value on a high level, specific energy demand varies by as much

as a factor of four among the alumina grades with the

finest and the coarsest microstructures.

To determine the causes for the strong influence of

grain size on grindability, an attempt was made to determine the modes of material removal dominating

with each kind of material. For this purpose, grinding

tests were performed on tilted ceramic specimens.

Figure 8 presents the surface structures obtained in these experiments.

The micrographs demonstrate the results on two

different workpiece materials with grain sizes of 0.8

and 4.6 |im respectively for grinding with a constantly increasing grit depth of cut. As the grits contact the polished surface, sharp cutting edges

immediately start removing material in the ducfile

mode (cutting). Duller cutting points first remove the

sputtered gold coating in elastic contact (rubbing) and

then displace material sideways in plastic

deformafion (ploughing). With increasing grit depth

of cut, cracks begin to form, and material is removed

by micro-brittle fracture. The two woricpiece

materials are clearly distinguished by their different

degree of ducfility. In the 0.8 |im variety, heavily

deformed layers of material are smeared along

grooves formed in the surface. Grain breakout and

fracture start much earlier in the 4.6 |im alumina

grade and thus little deformation can be detected on

the the remaining surface.

254
a) TM30 b) TM50

. 4 =N—

1

spot of first contact

in up - grinding of polished samples

of alumina

material AI2O3

dg =0.8 [Jim la) dg = i*M |im (b)

321/3303 ©IFW

Fig. 8: Traces of rubbing, cutting and ploughing actions of the grit

Figure 9 demonstrates the important influence tool specifications have on the obtainable surface finish. Samples of 0.8 )xm alumina were ground with three

grinding wheels of grit sizes 65/80 |xm (E>91),

30/40 ^rni (D20B) and 5/10 |xm (D7) respectively.

With the coarsest grit size wheel part of the sample

surface reveals plastic flow but equally a large

amount of grain break-out. Flakes of partly-removed

material are observed. A D20B wheel at the same

settings produces a corrugated surface covered with

grinding grooves. A number of smaller cavities from

fallen-out grains have not been completely filled up

by deformed material. With an extremely fine D7

wheel no fracture is visible at aU; the surface is very

smooth and completely free of any irregularities.

The measurement of residual grinding stress reflects the amount of sub-surface plastic deformation.

Deformation is a function of the load on a single grit as long as no fracture occurs. Due to the difficulty in

determining the number of cutting points on a wheel

surface, average contact pressures Fjj" obtained by

dividing the normal grinding force by the contact area

are comparable only among results from a single

wheel. Figure 10 demonstrates that a correlation

exists between the normal grinding force per unit area and residual stress in the workpiece sub-surface.

255
DS

workpiece materiai

high-purity alumina dg= 0 8 ^jm grindin g

peripheral surface grinding

= 25 m /s

- 5 jjm v„= 157 mm/s

tools

091 C50

020 8 C50

D7 C50

resin bond

coolont low viscosity mineral oil

Fig. 9: Tool influence on surface topography

workpiece moteriol

high - purity alumina,

dg = 0.8|im

grindin g : peripheral surface grinding

D91C50 (US170 / 200),resin bond

coolant :

1 2 3 4 N/mm2 5

normal grinding force per unit area F^' 321/ 890A ©IFW

Fig. 10: Residual stress vs. average contact pressure F^" in grinding

256
In Fig. 10 the creep feed grinding tests appear at low

Fji" while surface grinding produces much larger contact pressures. Up to a limit around 2 NAnm^,

increased grit loads produce increasing values of

compressive residual stress. If this limit is exceeded,

residual stress remains constant in the range between

-350 MPa and -400 MPa. Scatter is somewhat smaller

in grinding with small grit loads than at large .

Discussion

The present study has been conducted on a series of

alumina ceramics with varied hardness, toughness,

and transverse rupture strength. These aspects of

material behaviour originate directly from the grain

and pore sizes of the materials. In a brittle material

like alumina, the strength of a specimen containing

defects is inversely proportional to the square root of

the flaw size (Griffith's equation). Most obviously,

the grain boundaries themselves act as low-strength

imperfections. Their sizes correlate with the

respective grain size. Furthermore, the noted

segregation of pores in the large grain varieties

weaken the material by promoting intergranular crack

propagation. Hardness of a ductile material is determined by the resistance toward plastic

deformation. As long as no fracture occurs in

hardness testing, even brittle ceramics deform

plastically in the hydrostatic stress field of the

indenter. Deformations are the result of dislocation

movement which in a polycrystaUine material is obstructed by the grain boundaries. Thus, the

hardness wiU increase with reduced grain size, i.e. higher density of grain boundaries.

Roughness of groimd surfaces provides a first indicator of the prevailing mechanism of material

removal. Generally, very little influence of machining

conditions is found on roughness if machining is carried out in the brittle-fracture domain. Surface

finish is only about 20% better in creep feed than in

typical surface grinding. Instead, material grain size

dominates the roughness of ground surfaces. While

the fine-grained aluminas exhibit mariced differences

in roughness along and across feed direction, the

influence of feed direction is almost lost with coarse

grain sizes. Grain faces are uncovered by

intergranular fracture in micro-brittle material

behaviour. As can be gathered from SEM

observation, fracture dominates with coarse

microstructures while the surfaces of fine-grained

aluminas appear covered with a layer of deformed

material less than 1 |iin thick. In the latter case, the

roughness is determined by the traces of the grit.

Secondary chip formation with lateral cracking

behind the passing grit under the action of residual

sub-surface stress is the origin of comparatively large

surface defects in surfaces ground with large grit. As

grit depths of cut are reduced either via the process

input parameters or by the use of finer grit size

wheels, the transition from micro-brittle to ductile

material removal is passed. Cracking can be avoided

completely thus producing surfaces of almost mirror

quality. Apart from the grit depth of cut, the density

of cutting points plays an important role. To perform

a cutting action of the grits and avoid ploughing

phenomena as with coarse wheels set at low grit

depths of cut, the stress fields of the neighboring grits

must overlap. The state of stress created in grinding

then forces the workpiece material to be ejected

forwards and prevents any groove formation with

pile-up of deformed material.

It has been foimd that both components of the

grinding force are linearly dependent of a power

function including the equivalent chip thickness h^q

and the working engagement a^ at different powers

(comp. Fig. 1 1).

By plotting the specific forces versus this quantity

(heq'""ae"), it is possible with a given ceramic to directiy predict both force components from the

process settings. The exponents determined here are not specific to the grinding of ceramics but turn out

identical to those found e.g. in high-efficiency deepgrinding of case hardened steel [11]. At a constant

material removal rate, the woricing engagement ag determines the contact length and thus the grit depth

of cut. Specific forces are higher in creep feed

grinding with smaU grit depths of cut than in

conventional surface grinding with much larger grit

depths of cut. This principle is a parallel to the decrease of tiie cutting forces per unit area with

increased depth of cut in metal cutting [12].

257
O

c

C

i—cno

"o

0)Q.

if)

30

N

mm

20

15

10

5

0

0 0.1 0.2 0.3 OX 0.5 0.6 0.7 |im0-8^0.9

.0.7 O.U

work material : high purity alumina

dg = 0.8 ^m

g rindin g : peripheral surface grinding

V, = 21.5 / 25 m/s

Q;,= 0.33 / 1.0 /3.3mmVmm.s

tool : D91 C50, resin bond

Q'sb= 5 mmVmms

V;^ = 300 mmVmm

coolant : low viscosity mineral oil

321/963/. ©IFW

Fig. 1 1 : Specific grinding forces as a function of the process settings

i_

0)c

c

T3

c

cno

o

(UQ.

(/)

100

J/mm

60

Z.0

20

10

0

surface grinding,

-Q' = 0.1 mmVmm-s

Q' = 1.0 mmVmm-s

creep feed grinding.

Q' = 0.1 mm'^/mms

0' = 1.0 mm^/mm-s

1

workpiece materials

high purity

alumina ceramics

g rinding : surface grinding

= 30 m/s

v„ . a, = var.

coolant : min. oil

tool : 1A1 D91 C50

diamond wheel

grain size d

1 1/2

g

'

J L

^ Hmi/2 5

0 2 A 6 8 10

grain size dg

15 ^im 25

321/ 5132 ©IFW

Fig. 12: Specific grinding energy e^, vs. workpiece grain size d^

258
The specific grinding energy e^, is useful in the

assessment of the effectivity of a grinding process

[13-15]. It is composed of energy consumed for the

formation of new surface in crack propagation,

energy necessary for plastic deformation of the work

material, and energy dissipated in friction. The

specific grinding energy has been shown to vary

inversely with the undeformed chip thickness h^^^x hcu [7, 8, 16]. hi grinding of ceramics, indicates the

relative importances of brittle and energy-consuming

ductile material behaviour. Compared to metals,

plasticity is much less important with ceramics. Thus,

the grinding energy is composed mainly of the

frictional and surface energy terms. The considerable

deformation of fine-grained alumina ceramics does

not contribute significantly to the measured grinding

energy. The results represented in Figure 12

demonstrate the close connection of the grinding

energy with workpiece grain size. The two sets of

graphs refer to different grinding conditions typical

for surface and creep feed grinding. All cases exhibit

the same trend, namely an inverse correlation

between energy demand and grain size. Between the

finest-grained and the coarsest alumina varieties,

ratios of up to four in e^ are found, hi particular, the

specific grinding energy is close to inversely

proportional to the square root of the grain size dg in brittle-domain grinding.

This indicates a significant amount of micro-brittle

behaviour in material removal. Fracture mechanics

can be used to explain the dependence of grinding

force on grain size. The grain boundaries act as defects in ceramic microstructure, especially if, as in the case of coarse grain, rows of pores collect along

the grain boundaries. The linear dimension of the

individual grain boundaries is a fraction of the grain

size. With the boundaries acting as crack origins, the

strength of the material in resisting the tangential

grinding force is thus directiy controlled by grain

size. The ratios of measured specific grinding energy

among the different workpiece materials agree

reasonably well with their grain (flaw) sizes and

fracture toughnesses (see Table 3). Kj^*, dg*, and e^.* refer to the values for the 0.8 |iim grain size alumina.

With small grit depths of cut, fracture is avoided and

plastic deformation together with friction determines

the energy demand for material removal. Again, grain

size influences the grinding forces via the grain

boundaries' resistance to dislocation movement.

Table 3: Grinding energy demand related to the material properties

Woricpiece material

(grain size)

at

^"^^ =3-10-6

* 1

0.8 ^im 1.00 1.00

2.2 ^im 0.63 0.78

4.6 ixm 0.48 0.45

7.8 \im 0.37 0.33

12.2 \im 0.29 0.25

Residual stress correlates well with the normal grit

load up to a limit where brittie fracture begins to dominate over plastic removal. Residual stress is partiy relieved by cracks. Another cause for stagnation at -380 MPa combined with increased

scatter is the loss of deformed material from part of

the surface which thus no longer contributes to the

value of residual stress averaged over the x-ray spot.

Further research will have to be directed at measuring

the stress levels in surfaces ground in the completely

ductile regime at extremely small grit depth of cut.

259
Conclusions References

From grinding tests done on several alumina grades

of different grain and pore sizes it is concluded that grain size plays a significant role in material removal.

Over a wide range of machining conditions, grinding

forces and energies vary inversely with the average

grain size of the ceramic. Surface structures give

evidence of shear deformation and plastic flow. These

are found most maricedly with fine-grained material

and at low grit depths of cut. While the major part of

material removal is probably accomplished by brittle fracture, plastic deformation is held responsible for the final appearance of the ground surfaces. The

extent of plastic deformation is determined by

material grain size and grain retention in the matrix as machining starts with extremely low grit depths of cut

at the beginning of the arc of contact in up-grinding.

The brittle/ductile transition takes place at very low

grit depths of cut. To produce ground surfaces

completely free of chipping, both small grit depths of

cut and a high density of cutting points are required.

Residual grinding stress in alumina is found to be

compressive in all cases. It is an indicator of the

average mechanical loading of each individual grit in

the grinding process. With a given wheel, the

magnitude of compressive residual stress correlates

with the grit depth of cut and with the normal

grinding force per unit area. A reduction of load per

grit via the use of a finer grit size grinding wheel with

a larger number of cutting points will produce

reduced compressive stress in the sub-surface zone.

Acknowledgements

The authors gratefuUy acknowledge the support of

Dr. R. Telle (MPI fuer Metallforschung, Stuttgart,

Germany) who processed and tested the work

materials. Many new ideas have sprung from the

enlightening discussions with him. Thanks are equally due to S.-H. Park of the same institution for characterising the different material grades. The

authors would like to thank the German Research

Foundation (DFG) for funding the research over a period of four years.

[I] H.K. TOnshoff, G. Wamecke, H.-Z. Choi, "Der

Beitrag der Forschung zum SSgen von Stein",

Internationale Diamanten Rundschau. Vol.11,

No.l,pp.5-ll, 1977

[2] H.P. Kirchner, J.C. Conway, "Mechanisms of

Material Removal and Damage Penetration

during Single Point Grinding of Ceramics",

Machining of Ceram. Mat.ls and Components. Vol. 17, ASME, New York, pp. 53-61, 1985

[3] H.K. TOnshoff, E. Brinksmeier, "Abrasives and

their Influence on Force, Temperature and

Surface", Proc. 2nd Intematl. Grinding

Conference. Philadelpliia, June 10-12, 1988

[4] E. Saljd, H. MOhlen, "ProzeBoptimierung beim

Schleifen keramischer Werkstoffe", Internationale Diamanten Rundschau. Vol.21, No. 4, pp. 243-247, 1987

[5] E. Saljd, H. Mohlen, "Material Removal in

Grinding of Advanced Ceramics", SME Techn.

Paper MR88-498. pp. 1-19, 1988

[6] P. Blake, T. Bifano, T. Dow, R.O. Scattergood,

Precision Machining of Ceramic Materials,

Ceramic Bulletin. Vol. 67, No. 6, pp. 1038-

1044, 1988

[7] G.S. Reichenbach, J.E. Mayer, S. Kalpakcioglu,

M.C. Shaw, "The Role of Chip Thickness in Grinding", Transactions of the ASME. Vol. 78,

No. 5, pp. 847-859, 1956

[8] H.K. Tonshoff, R. Telle, P. Roth, "Chip

Formation and Material Removal in Grinding

of Ceramics", Proc. 4th Intematl. Grinding

Conference. Dearborn, October 9-11, 1990,

pp. 539-1/539-18

[9] G. Werner, "Kinematik und Mechanik des

Schleifvorgangs", Dr.-Ing. Diss. . TH Aachen,

Germany, 1971

[10] G. Kassen, "Beschreibung der elementaren

Kinematik des Schleifvorgangs", Dr.-Ing.

Diss. . TH Aachen, Germany, 1969

[II] T. Tawakoli, "Hochleistungs-Flachschleifen",

Dr.-Ing. Diss. . VDI-Verlag, Duesseldorf,

Germany, 1990

[12] 0. Kienzle, H.Victor, "Spezifische

Schnittkrafte bei der Metallbearbeitung",

Werkstofftechnik und Maschinenbau. Vol. 47,

No. 5, pp. 224/225, 1957

260
[13] H.W.Zheng, G.Q. Cai, S.X. Yuan, "An

Experimental Study on the Mechanism of

Cemiet Grinding", Annals of the QRP. Vol. 38, No. 1, pp. 335-338, 1989

[14] T.W. Liao, G. Sathyanarayanan, "An In- vestigation into Creep Feed Grinding of

Alumina Ceramic", Transactions North Am.

Manuf. Res. Inst, of SME (NAMRC XVID. Vol. 17, No. l,pp. 175-184, 1989

[15] J. Akbari, Y. Ichida, K. Kishi, T. Machida;

"Grinding Energy of Fine Ceramics", Proc.

First Intematl. Conference on New Manufacturing Technol. . Chiba, 1990, pp. 323-328

[16] I. Inasaki, "High Efficiency Grinding of

Advanced Ceramics", Annals of the CIRP. Vol. 35, No.l , pp. 211-214, 1986

261

EFFECT OF GRINDING ON STRENGTH AND SURFACE

INTEGRITY OF SILICON NITRIDE: PART I

S. JAHANMIR, T. J. STRAKNA", G. D. QUINN, and H. LIANG"

National Institute of Standards and Technology, Gaithersburg, MD

R. L. ALLOR

Ford Motor Co., Dearborn, MI

R. D. WEST

Therm Advanced Ceramics, Ithaca, NY

Grinding is often the method of choice for machining structural ceramics and has a high

potential for large-scale production and automation. Because there are few guidelines available

for grinding ceramics, a conservative approach is usually used to minimize the potential grindinginduced damage. With this approach, machining time is high, which increases the cost of

machining. The purpose of this paper is to report data and information on the effect of grinding on strength of two types of silicon nitride. The work is part of a study to develop a database that can be used to assist in the optimization of the grinding process to achieve low cost, high

reliability, and eventual automation. Standard flexure specimens, made from a reaction-bonded

and a sintered reaction-bonded silicon nitride (RBSN and SRBSN), were surface ground in the

longitudinal direction under three different conditions. The resulting ground surfaces were

characterized with a three-dimensional profiling instrument and a scanning electron microscope.

Four-point flexure tests were performed on the specimens and WeibuU parameters were calculated

to determine the effect of the different grinding conditions on flexure strength. Fractography was

performed on the specimens after testing to determine the location and to characterize the fracture origin in each specimen. The results showed that varying the material removal rate by a factor

of 20, from 30 to 590 mmVmin, had no significant effect on the characteristic strength or the

Weibull modulus of four-point flexure specimens. This implies that material removal rates higher

than those currently used for grinding these two types of silicon nitride materials may be possible

without strength degradation.

Introduction

Structural ceramics, such as silicon nitride, are

attractive for many applications due to their high

strength at elevated temperatures, resistance to chemical degradation, wear resistance, and low

density. Industry analysts forecast that the market

for structural ceramics will grow from $350 million

in 1990 to $865 million by 1995 [1,2]. Major

market segments for these materials are cutting tools, wear resistant parts, heat engines, and aerospace and

defence related applications [1-3]. The high cost of

ceramic components versus metallic components,

however, has been cited as one of the main

* Graduate Research Assistant, University of

Maryland, College Park, MD 20742

** Research Associate, University of Maryland,

College Park, MD 20742

263

problems preventing the automobile manufacturers

from employing ceramics in automotive applications

[1]. Some studies estimate that of the total production cost machining can account for 30 - 60% and

sometimes even up to 90% for high-precision components [4,5]. Therefore, recent reports recommend

a coordinated research effort in machining of advanced ceramics [4,6] to reduce machining costs. Grinding is often the method of choice for machining structural ceramics [4,5] and has a high

potential for large-scale production and automation.

Because there are few guidelines available for grinding ceramics, a conservative approach is usually used

to minimize the potential damage induced by grinding. This involves very low depths of cut and many

passes to obtain the desired tolerances, thus increasing the cost of machining. The purpose of this paper

is to report data and information on the effect of

grinding on the strength of two types of silicon
nitride. The work is part of a study to develop a database that can be used to assist in the optimization of the grinding process to achieve low cost, high

reliability, and eventual automation.

Although many recent papers [6-13] and confer- ence proceedings [14,15] have addressed the issue of

niachinability for advanced ceramics and provided

data, much more data and information are needed on

the machinability of specific materials. In Particular,

it is important to achieve the maximum possible

material removal rate, while minimizing the effect of

machining on performance and properties. Previous

studies have shown that grinding-induced damage

can be detrimental to the strength and performance

of ceramic parts [16-22]. Since the process of

grinding involves removal of material by deformation and fracture, the machined components often

contain residual stress, a deformed layer, surface and

subsurface microcracks, phase transformation re- gions, and other types of damage [23-25]. The

extent and the specific nature of the machininginduced damage depend on the type of machining

process and the machining parameters. Also, it should be recognized that formation and the influ- ence of machining-induced damage depend on the

material type and microstructure, for example, the

grain size [23,24].

Flexure testing [26] has proven to be a reasonable, low-cost method for evaluating the effect of

machining on strength of ceramic materials [16,17,1

9,22,23,27-30]. AUor and Baker [27] studied the

effect of several grinding parameters on the flexure

strength of a hot-pressed silicon nitride. Grinding in the same direction as the maximum tensile stress was

found to increase flexure strength, and down feeds

up to 0.254 mm/pass had no significant effect on

flexure strength. Grinding wheels of grit size between 220 and 600 increased flexure strength, while

varying the diamond concentration of the wheel

between 75 and 100 on 320 grit wheels had no

significant effect on flexure strength. Thomas and

West [28,29] studied the effect of grinding parameters on grinding forces, surface finish, and strength

of a hot-pressed silicon nitride. The flexure strength

of the specimens was found to decrease as the down

feed was increased from 0.0254 mm to 0.1778 mm.

Higher strengths were obtained, as well as better surface finish, for a 320 grit wheel than for a 150

grit wheel. Miyasato, et al. [30] explored the effect of down feed, different wheel types, and grinding

direction on strength, surface finish, and residual

stress of a hot-pressed silicon nitride. No significant

strength differences were found between different

wheel types and down feeds, while grinding normal

to the maximum tensile stress direction decreased

strength by 40% compared to grinding parallel to the

tensile stress direction.

In the present study, four-point flexure testing was used to assess the effect of grinding on the

strength of two commercially available silicon nitride materials, one reaction-bonded and the other sintered

reaction-bonded. Weibull statistical analysis [31] was performed on the data. Three grinding conditions were used to achieve a wide range of material

removal rates. The surface integrity resulting from

the grinding conditions was investigated with a three- dimensional profilometer and SEM. Fractography was performed to identify the location and source of

fracture in each specimen.

Experimental Procedure

a) Materials

Two types of silicon nitride materials were used

in this study; reaction bonded (RBSN), produced by

Eaton Corporation, and sintered-reaction bonded

(SRBSN), produced by Ceradyne, Inc. The RBSN

was prepared from a mixture of silicon powder with

iron oxide, as a nitriding agent, and water [32]. The

SRBSN was also processed with silicon powder, but

it contained yttria and alumina sintering aides. The bulk density, crystal structure, and grain size

of both materials were determined. Bulk density was

measured via fluid displacement. The density of the

RBSN was 2.36 g/cm^ and the density of the SRBSN

was 3.22 g/cm^. The crystal structure of the materials was determined by x-ray diffraction. Both alpha

and beta phases were present in the RBSN material

with an alpha to beta ratio of 7.7 : 1, whereas the

SRBSN primarily contained the beta phase. Grain

sizes were measured by direct viewing of polished

and etched surfaces of the materials. For the RBSN,

the alpha phase was 1-2 p,m in length and 0.5 - 1 jim in width, while the beta phase was 1 - 3 ^m in length and less than 1 }im in width. For the SRBSN,

the beta phase was 6-8 )Xm in length and 1-3 ^m

in width. The manufacturers gave the porosity of the

materials at approximately 20% for the RBSN and

less than 1% for the SRBSN. Both materials contained small inclusions ranging in diameter from 0.1

to 2 |im.

The micrographs in Figure 1 show the microstructure of the RBSN and SRBSN materials as observed in an optical microscope. A distinguishing

feature of the RBSN material is the characteristic

mottled appearance on a size scale of approximately

50 |im.
The materials were received in billets. The

RBSN billets were approximately 86 x 10 x 3.2 mm

and were received in three batches. One or two

flexure specimens were made from each RBSN

billet. The SRBSN billets were 76 x 76 x 13 mm.

These billets were sliced into 4 equal size pieces

with approximate dimensions of 50 x 25 x 13 mm

and the location in the billet of each sliced piece was

tracked. A total of 8 to 12 flexure specimens were made from each SRBSN sliced portion. The billet number, location, and batch number from the billets were tracked for each specimen to check the possibility of billet to billet variation and the influence of

specimen location within each billet on strength.

b) Grinding Conditions

A total of 25 to 30 flexure specimens from each

material were machined under each set of grinding

conditions that will be described shortly. These

specimens had dimensions of 3 x 4 x 50 mm, as specified by the ASTM "Standard Test Method for

Flexural Strength of Advanced Ceramics at Ambient

RBSN

Figure 1. Optical micrographs of polished surfaces

of RBSN and SRBSN.

Temperature" [33] . In order to reduce the possibility of fracture initiation from the edges, all the bars were chamfered as required by the ASTM standard.

A sufficient amount of material was removed from

the test bars by grinding to reduce effects of surface

layers that are often formed during processing and

damage caused by slicing. Three sets of grinding

parameters, shown in Table 1, were used to finish

grind the specimens. In all cases, grinding was

performed on a horizontal-spindle surface grinder

with resin bonded diamond wheels in the longitudinal direction (parallel to the specimen tensile axis).

No spark-out or further surface treatment was per- formed.

Grinding condition A used a 320 grit wheel,

0.0025 mm/pass down feed, 1.58 mm/pass cross

feed, and 7.62 m/min table speed. This resulted in a material removal rate of 30 mmVmin. During

rough grinding, the same parameters were used with

the exception of the down feed, which was decreased

in the following steps: 0.0127 mm to net dimension

+0.2032 mm, 0.0254 mm to net dimension +0.1270

mm, 0.0127 mm to net dimension +0.0508 mm. A

total of 0.0508 mm was removed in finish grinding.

Grinding condition B also used a 320 grit wheel.

Under this grinding condition, the down feed and the

table speed were larger than that of condition A, but

the cross feed was smaller. This condition is slightly more aggressive than A, resulting in a material

removal rate of 59 mm'/min. Larger grit wheels

Table 1. Grinding parameters

A B C

Griffiding Wheel Grit 1 Concentration

320 / 100 320 / 75 120 / 63

Grinding

Fluid 20:1

Emulsion

50:1

Emulsion

50:1

Wlieel Surface Speed

(m/s)

28 28 28

Tabk Speed

(m/min)

7.62 12.2 6.10

Cross Feed

(nun/pass)

1.58 0.381 0.381

Down Feed (mm)

0.0025 0.0127 0.2540

Volumetric Removal

liaie (mmVmin)

30 59 590

265
(120 grit for RBSN and 150 grit for SRBSN) were

used during rough grinding, with a table speed of

6.10 m/min and a down feed of 0.381 mm. A

total of 0.254 mm was removed during fmish grinding. Conditions A and B represent bench-mark

conditions usually used for ceramic materials and

recommended as a conservative procedure in the

ASTM standard.

Grinding condition C used a 120 grit wheel,

0.2540 mm/pass down feed, 0.381 mm/pass cross

feed, and 6.10 m/min table speed. The same set of

parameters was used for rough grinding. Condition

C is far more aggressive than either A or B, and

resulted in a material removal rate of 590 mmVmin.

Other parameters such as diamond concentration

and grinding fluid are also listed in Table 1. In addition, each grinding condition was performed on

a different machine tool by a different operator. Although these factors may influence the results, this study was mainly concerned with the volumetric

removal rate.

c) Surface Integrity Evaluation

The surface topography of the ground samples was evaluated with a three-dimensional stylus

profiling system, an optical microscope and a scan- ning electron microscope (SEM). The profiling

system consisted of a stylus profilometer, a moving

stage powered by a steppermotor, a personal computer and data acquisition and processing software,

and electronics to run the system [34]. Three-dimensional topographical maps were produced on selected

test bars of each material and each grinding condition by compiling multiple profile traces taken at fixed intervals on the specimen surfaces. The test bars were mounted in the system such that each trace ran perpendicular to the grinding direction. A total of 200 traces, 3 mm in length and set 10 jim apart, were taken. This corresponds to an area 3x2 mm.

A subset of the traced area, approximately 1x1

mm, was used to calculate the topographical parameters: root mean square roughness (R^), average peak

to valley roughness (R), and centerline average

roughness (RJ. Final leveling of the surface was

performed in the software, as well as filtering to remove lower frequency form deviations and system

error. The repeatability of the roughness measure- ments was determined by performing 4 three-dimensional profiles, as described above, on one specimen

from each material. The standard deviation about

the mean R^ value was 0.02 [im for RBSN and 0.01

\lm for SRBSN, representing 7% and 6% of the

mean, respectively.

d) Flexure Strength Tests The flexure strength of the ground test bars was

measured in accordance with the ASTM Standard C

1161 for four-point flexure testing [33]. A universal

tension/compression testing machine was used at a cross-head speed of 0.5 mm/min to apply a steadily

increasing load to each specimen. A load cell was

used to measure the load, and an integrated sl^rip- chart recorder was used to record the applied load. A fully articulating steel fixture [35] was used to mount the specimens for testing in four-point flexure

position. This fixture has the advantage of articulat- ing during loading to correct for any misalignments

due to dimensional errors in the specimens. This

ensures a good line-contact between each load-pin

and the specimen. The system was calibrated each

day before beginning the flexure tests, and checked

periodically during the testing sequence.

To perform a test, the width and thickness of

each test bar was first measured with a micrometer

within 0.0025 mm. It was then placed in the fixture. After mounting, a preload of 5 - 10% of the expected failure load was applied. The specimen and

fixture were then inspected and adjusted, if needed,

to ensure line-contact between the load-pins and the

surface of the specimen. When the mounting was

determined to be satisfactory, the contacts of the

load-pins were marked on the specimen and the

application of load was continued until failure occurred. The fractured specimen was then stored

for future examination, and the load at fracture was

recorded from the strip-chart.

Following the flexure tests, the strength data were analyzed to determine possible effects of

grinding on strength and the variation of strength. In order to identify the location and source of fracture

in each test bar, the fracture surfaces were analyzed

by stereo optical microscopy. Selected specimens were examined with an SEM. All fractography was

performed in accordance with procedures specified

in Reference [36].

Experimental Results

a) Surface Integrity

Topographical maps of selected samples ground

under the three different conditions are shown in Figures 2 and 3 for RBSN and SRBSN. Differences

between the grinding conditions can be seen in these

maps as a variation in severity of the grinding

striations. It is evident that grinding condition A

imparted more severe or deeper striations in both

materials. This is more clearly evident in Figure 4,

Figure 3. Topography of ground surfaces of SRBSN for grinding conditions A, B, and C. 268
which shows the effect of grinding condition on the

Rq roughness of both materials. This figure also

shows that the roughness of RBSN is greater than

SRBSN for each grinding condition. Grinding

condition A produced the largest R^ roughness for

both RBSN (0.41 \im) and SRBSN (0.32 ^m),

followed by C and then B for RBSN (0.38 \im and

0.27 [Lm), while the R^ roughness of SRBSN surfaces ground under conditions B and C were nearly

equal (0.20 [im). The effect of grinding conditions on the other roughness parameters R and R^ were

similar to that described for R^. As seen in Figures 2 and 3, the two materials

display distinct surface characteristics regardless of

the grinding conditions used. The RBSN ground

surfaces show somewhat random peaks and valleys

and relatively obscured grinding striations, though

directionality is evident. The ground surfaces of

SRBSN, however, display sharp and distinct grinding

striations with obvious directionality. These observations exemplify the differences in the mechanical

properties and/or microstructure of the two silicon

nitrides and may indicate different material removal

mechanisms for each in the grinding process. The SEM micrographs in Figures 5 and 6 are consistent with the topographical maps of the ground

surfaces. The SEM micrographs suggest two possible modes, brittle and ductile, for material removal

in grinding silicon nitride. The striations in RBSN

are not as continuous as those observed for SRBSN.

Also, larger pits or pores are seen on the surface of

RBSN (Fig. 5) than those on the SRBSN (Fig. 6). The surface structure of these two materials is shown

in Figures 7 and 8 at a higher magnification. These

figures clearly show that while plastic flow and

plowing were the primary material removal mechanisms in SRBSN, these features are less evident in RBSN. This suggests that a brittle fracture process

may be the dominant mechanism for grinding RBSN

because of its lower strength and higher porosity

compared to SRBSN. The difference between the

material removal processes of the two materials is a

critical issue in determining the extent and nature of

possible machining-induced damage.

b) Flexure Strength

Due to the dependance of strength on the flaw

population within the specimens, the strength data

obtained from the flexure tests is well-fitted by the

two-parameter Weibull statistical distribution. The

Weibull distribution allows for the fact that failure in a specimen occurs at the largest flaw among a population of flaws existing in the specimen. From

this analysis two parameters are obtained: the charac0.4

^ 0.2

i

^ResN

SRBSN i

B C

Grinding Condition

Figure 4. Rq roughness of RBSN and SRBSN for

each grinding condition.

teristic strength (scale parameter) and the Weibull

modulus (shape parameter). The characteristic

strength is the stress at which 63.2% of the speci- mens failed. The Weibull modulus is an indication

of the scatter in the strength data, where the scatter decreases as the Weibull modulus increases. By

plotting the probability of failure on Weibull probability axes and the flexure strength on natural log axes (termed Weibull plots) the data can be fitted by

a straight line. The least squares fit method was

used to fit a line to the data whose slope is the

estimate for the Weibull modulus. The stress corr- esponding to the point on the line where 63.2% of

the specimens will fail is the estimate for the characteristic strength.

The 95% confidence intervals for the Weibull

modulus and characteristic strength were calculated

using the method developed by Ritter, et al. [37].

This analysis is based on the law of propagation of

errors and assumes that the scatter in the characteristic strength and Weibull modulus is normally distrib- uted and the coefficient of variation is less than or equal to 10%. This analysis method has been

successfully applied to flexure data of ceramics [38].

The Weibull plots for each material and grinding

condition are shown in Figures 9 and 10. The data

is fitted by a straight line satisfactorily in each case, and thus is well represented by the two-parameter

Weibull distribution. The characteristic strength and

the Weibull modulus are displayed for each material

and grinding condition in Figures 11 and 12 along

with the associated 95% confidence intervals. The

269
Figure 5. SEM micrographs of the ground surfaces

of RBSN for grinding conditons A, B, and C.

Figure 6. SEM micrographs of the ground surfaces

of SRBSN for grinding conditons A, B, and C.

270
Figure 7. SEM micrographs of the ground surfaces

of RBSN for grinding conditions A, B, and C (at a higher magnification than in Figure 5).

Figure 8. SEM micrographs of the ground surfaces

of SRBSN for grinding conditions A, B, and C (at a higher magnification than in Figure 6).

271
0 999 099

0.95

0.90

0.80 0 60

> 0.40

; 0.20

i. 0.10

! 0.05

>

0.02

0.01

0 005

0.0025

o Condition A

Welbull Modulus: I I Chorocterislic Strengih: 225 MPo

Condition B

Wcibull Modulus: 15 Chotoclcnslic SIrength: 208 MPo

« Condition C

Weibull Modulus: 18

Chofocterislic SIrenglh: 214 MPa

200 300

Froclure Stress (MPo)

Figure 9 Weibull plot of RBSN samples for each

grinding condition.

0.999

0 99

0 95

0.90

0 80

0.60

0 40

0 20

0.10

0 05

0 02

0 01

0.005

B

o Condition A

Weibull Modulus, 12

Chorocterlstic strength 723 MPo

• Condition B

Weibull Modulus: 14 Characteristic Strength. 688 MPo

* Condition c

Weibull Modulus: 1

Chorocteristlc Strength: 700 MPo

600 700 800 900 1000 1200

Frocture Stress (MPo)

Figure 10. Weibull plot of SRBSN samples for

each grinding condition.

no machining effect on the flexure strength, it can be

concluded that the testing sequence had no deleter- ious effect on the results, as well.

c) Fractography

In order to identify the source of fracture initia- tion and determine the role of machining damage in the fracture, fractography was performed on each

specimen using an optical microscope. Fractography

provides a practical means for confirming the influ- ence, or lack of influence, of machining procedures.

In many advanced ceramics, it is possible to detect

and fully characterize the strength limiting flaws, and

determine whether the flaws are processing, microstructural, handling, or machining related. Although

fractography can be subjective, a recent publication

[36] makes the practice much more objective and

serves to minimize the subjective factors. The

guidelines suggested by this publication were used in the present study to assess the effects of the machining procedures used.

Crinding ABCCondition Grinding Condition

characteristic strength varies from 208 to 225 MPa

for RBSN and from 688 to 723 for SRBSN. The

difference in strength between the two materials is expected due to the different material properties.

The characteristic strength for each material, however, does not vary significantly at the 95% confidence

level between the grinding conditions. The Weibull

modulus varies from 11 to 18 for RBSN and from 1

to 14 for SRBSN. This difference is not statistically

significant at the 95% confidence level. The results,

therefore, indicate that the three grinding conditions

used in this study do not have a significant effect on

the characteristic strength and the Weibull modulus

of the two silicon nitride materials. Statistical analysis on the characteristic strength and Weibull modulus confirmed that billet to billet variation in SRBSN

and batch to batch variation in RBSN had no influ- ence on the fracture results. Since the results are conclusive at the 95% confidence level in showing

Figure 11. Characteristic strength and Weibull

modulus of RBSN for each grinding condition.

Grinding Condition ABC Grinding Condition

272

Figure 12. Characteristic strength and Weibull

modulus of SRBSN for each grinding condition.
Fracture of ceramic materials typically induces a

set of characteristic features on the fracture surface

known as the fracture mirror, mist and hackle. The

fracture mirror is a smooth, flat, circular or semicircular region surrounding the fracture origin. Mist

is a border region found along the edge of the mirror

and is less smooth than the mirror. Hackle is larger

grooves and crack lines radiating out from the

fracture origin. The approximate origin of fracture can be easily determined by identifying these features.Often the terms flaw, irregularity, discontinuity, or defect are used to identify the source of fracture

initiation. In the present paper, the term flaw is used

to identify the fracture origin, and should not be construed to mean that the material has been improperly

prepared. This usage is consistent with the recommendation given in Reference [36]. Fracture origins were characterized by identity

and location, as recommended by Quinn et al. [36].

Identity characterizes the phenomenological aspects

of the flaw and identifies the type of flaws. For

example, agglomerates, large grains, pores, inclusions, and other microstructural features are volumedistributed flaws, while machining damage, handling

damage, and pits from oxidation or corrosion are surface-distributed flaws.

Location of the flaw identifies where the flaw is located in a specific specimen using the three

classifications: volume, near-surface or surface, and

edge. Classification of the location of the flaws is subjective, and depends on the researchers definition

of bulk vs. near-surface. In the present paper, volume-located flaws refer to those fracture origins

(fracture mirror centers) observed approximately 100

|im or more below the surface, which correspond to being visibly below the surface when viewed at 5 Ox

in an optical microscope. Any fracture origin that appeared closer to the surface was identified as surface or near-surface. Selected specimens were examined in an SEM since, in general, machining

damage is very difficult to discern with optical

microscopy [23,36,39]. During examination of the

fracture surfaces, a simple drawing was made of

each fracture surface and notes were taken on points

of interest. Examples are shown in Figure 13. Fractography revealed that the majority of

fractures were initiated from near-surface or surface

flaws. Fracture origins in the RBSN included 15-40

\im sized pores, agglomerates, and inclusions. The

SRBSN origins also were mainly pores and inclusions in the 20-30 ^im size range. In several instances, however, the fracture origins were zones of

clearly transgranular fracture, whereas the fracture

a ;

3

Figure 13. Drawings of fracture surfaces of speci- mens observed with optical microscopy.

273

mode of most specimens was highly intergranular.

Examples of fracture initiation from volume-distributed and volume-located microstructural features are shown in Figure 14. The fracture origin in the

RBSN specimen is identified with a large inclusion,

suspected to be iron rich, while a large pore initiated fracture in the SRBSN flexure bar. The results of the fractographic study are given

in Table 2. In some cases, fracture was initiated from the edge and chamfer, and in only a few

specimens was machining damage identified as a possible source of fracture initiation. Among the

RBSN flexure bars, only those finished under grinding condition A contained fracture origins that were

identified as being associated with machining damage. For each grinding condition used on SRBSN,

a similar number of fracture origins (9 - 18%) could

be associated with possible machining damage.

However, the SRBSN specimens had more edge and

chamfer fracture origins than the RBSN specimens.

The larger number of possible machining damage

and edge and chamfer fractures in SRBSN than in RBSN was not unexpected. Since SRBSN is far more dense than RBSN, the strength of this material

is expected to be more sensitive to the effects of

machining damage. Machining damage, however, was never absolutely determined to be the cause of

failure in any of the specimens. Two examples

showing fracture initiation from possible machining

damage are shown in Figures 15 for RBSN and

SRBSN. In the RBSN specimen, two semi-circular

regions extending from the surface can be seen, but

they are not sufficiently clear to confirm machining

damage as the fracture initiation source. There are other features, as well, such as substantial porosity in the immediate area, leading to uncertainty as to the
SRBSN

Figure 14. Fracture surfaces of RBSN and SRBSN

specimens whose fracture origins are both volumelocated and volume-distributed microstructural features.

source of fracture. In the SRBSN specimen, a semicircular region extending from the machined surface can readily be seen, but it is certainly interacting

with a pore. Thus, machining damage is probably

not the sole cause of fracture in this specimen,

either.

Discussion

As compared to machining metals, ceramic

machining technology is relatively undeveloped. The

large foundation of machining data and research that exists for metals does not exist for ceramics. This

implies that research in ceramic machining my be

Table 2. Summary of fractography

RBSN

A B C

Volume 32% 4% 7%

Edge and Chamfer

7% 0% 3%

Surface or

Near-Surface 61% 96% 90%

Possible Machining

Damage

13% 0% 0%

SRBSN

A B C

Volume 10% 3% 9%

Edge and Chamfer27% 19% 18%

Surface or

Near-Surface 63% 78% 73%

Possible Machining

Damage

17% 9% 18%

fruitful in improving current methods of machining.

The general findings of previous research and

experience on the effect of grinding parameters on

the strength of ceramics is that a smaller depth of cut may impart less damage, a finer grit size may do the

same, and that grinding in the direction of the

maximum tensile stress maximizes strength. However, the current procedures may be too conservative in terms of material removal rate; and therefore, the

lack of knowledge tends to push machining practice

to the "safer" end of the spectrum, i.e., very low

material removal rates, and usually, higher cost. Many of the earlier studies evaluated strength as a function of machining parameters, such as down

feed or table speed. The influence of the individual

parameters had to be inferred because of the lack of

understanding of the grinding process from first principles. In addition, the inherent scatter in the

strength results hampered unequivocal interpretation.

With the advent of standardized test procedures

[33,36] that measure strength accurately to about 1%,

and analyses that permit comparison of strength

samples, it is now possible to analyze quantitatively

the machining influences on strength. In the interest of obtaining narrow confidence intervals in the

strength parameters, however, large sample sizes are necessary. The current test standards recommend a
RBSN

Figure 15. Fracture surfaces of RBSN and SRBSN

specimens with possible machining damage as the cause of failure.

sample size of 30 as a compromise between cost and

statistical considerations.

The results of the present study clearly indicate

that the grinding conditions used in the study did not have a deleterious effect on the flexure strength and

Weibull modulus of the two silicon nitride materials.

The implications are that the material removal rate may be increased up to and possibly beyond the

values used in this study without strength degradation for RBSN and SRBSN. Because no strength

degradation was found when higher removal rates were used and the dominance of pre-existing fracture-initiating flaws in the specimens, it may be

stated that the optimal grinding condition in terms of

removal rate for these two materials occurs at the

maximum removal rate in which the influence of the

machining induced flaws is less than the influence of

pre-existing flaws in the material. The influence of

the pre-existing flaws is, of course, different for each

275

material and may change as the process used to manufacture the material is improved.

It is thus important to recognize that machinability data is material-specific. The great difference

in surface integrity between RBSN and SRBSN

resulting from the grinding conditions used in this study demonstrates this. In addition, if it can be

accepted that grinding will not degrade strength until the influence of the flaws induced by grinding are more severe than the influence of pre-existing flaws,

then it would be expected that the strength of

SRBSN would be more sensitive to grinding-induced

damage than RBSN because of its lower porosity and

greater homogeneity (hence fewer pre-existing

flaws). This approach to grinding optimization must

be used until the mechanisms of material removal in grinding ceramics are better understood. Trends in the effect of grinding parameters on strength and

surface integrity may eventually appear as more

machinability data is obtained, thus pointing to underlying mechanisms of material removal.

The material removal rates used in this study are by no means the optimum values. To define the best

set of grinding parameters, one must increase the

removal rate further and analyze for the effect of

grinding on strength and other performance characteristics. This approach has been adopted for the

second part of this paper, which follows this paper in the proceedings.

Conclusions

1. Varying the material removal rate by a factor of

20, from 30 to 590 mmVmin, in the grinding of

RBSN and SRBSN did not have a significant

effect on the characteristic strength or the Weibull modulus of four-point flexure specimens.

2. Surface roughness and topography were affected

by the grinding conditions.

3. The results presented in this paper indicate that higher removal rates than those currently used

for the grinding of these two types of silicon

nitride may be possible without strength degradation.

Acknowledgments

This project was conducted under the auspices of

the Ceramic Machining Consortium, which is a joint

research program between NIST and several compa-
nies, universities, and government agencies. The

authors are grateful to James Edler of Eaton Corporation and John Mangels of Ceradyne, Inc. for providing the silicon nitride materials used in this study. The financial support received from the

Department of Energy — Office of Transportation

Materials, is gratefully acknowledged. We thank

Peter Blau, Ceramic Machining Project Coordinator,

Oak Ridge National Laboratory, for his encouragement and his support.

Information on product names, manufacturers, or

suppliers are included in this report for clarification. This does not imply endorsement of the products or

services by NIST.

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Surface Finish and Strength," Therm, Inc., Ithaca,

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July 24, 1990.

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(Eds.), Vol. 22, 1976.

277
iI
EFFECT OF GRINDING ON STRENGTH AND SURFACE

INTEGRITY OF SILICON NITRIDE: PART II

S. JAHANMIR, T. J. STRAKNA*, and G. D. QUINN

National Institute of Standards and Technology, Gaithersburg, MD

R. N. KOPP

Norton Co., Worcester, MA

S. C. Yoon

Cincinnati Milacron, Cincinnati, OH

K. V. KUMAR

GE Superabrasive, Worthington, OH

Standard flexure specimens, made from a reaction-bonded and a sintered reaction-bonded silicon

nitride (RBSN and SRBSN), were surface ground in the longitudinal direction under three different

conditions involving high material removal rates. The resulting ground surfaces were investigated with

a three-dimensional profiling instrument and a scanning electron microscope. Four-point flexure tests were performed on the specimens and Weibull parameters were calculated to determine the effect of the

different grinding conditions on flexure strength. Fractography was performed on the specimens after testing to characterize and determine the location of the fracture-initiating flaw in each specimen. Varying

the volumetric removal rate in surface grinding from 310 to 516 mmVmin for RBSN and from 310 to 1737 mmVmin for SRBSN by employing creep feed grinding conditions was found to have no significant

effect on the characteristic strength of these materials. However, one of the grinding conditions, which

used spark-out, reduced the Weibull modulus, i.e., increased the variation in flexure strength. The results also showed that surface roughness and topography were affected by the grinding conditions.

Introduction

Grinding is an extensively used technique for

machining structural ceramics [1]. There is, however, very little data available on grinding ceramics and

cost-effective grinding techniques have yet to be

developed [2]. The purpose of the present investigation was to collect data and information on the effect of grinding on strength of two types of silicon

nitride. The data will be part of a database being

developed to assist in the optimization of the grinding process for low cost, high reliability, and eventual automation. The database will serve to efficiently transfer this information to manufacturers in an effort

to make advanced ceramics a practical engineering

material.

In Part I of this investigation [3], the effect of

i three grinding conditions on the strength of a reac- tion-bonded and a sintered reaction-bonded silicon

nitride was investigated. The grinding conditions

used encompassed a wide range of material removal

* Graduate Research Assistant, University of

Maryland, College Park, MD 20742

rates. Two of the conditions were bench mark

conditions which represented a conservative procedure in terms of material removal rate (30 and 59

mmVmin) and the amount of damage they may

induce, while the third condition was significantly more aggressive, with a removal rate of 590 mmV

min. Because strength was not affected by these

three grinding conditions, the tests were extended

using creep feed grinding at higher removal rates. The results of these tests are presented in this paper. The grinding conditions used in this paper and in Part I serve to cover the operational envelope of a typical surface grinder.

Creep feed grinding is characterized by low table speeds and large down feeds [4]. Creep feed grinding of a part is typically completed in one or a few

passes and the machining time can be significantly

lower than in conventional techniques.

In one of the grinding conditions, spark-out was

used for comparison with the others. Spark-out is a technique that is often used in grinding metals to improve surface finish [4]. In this technique, the

workpiece is subjected to several passes under the

grinding wheel without incrementing the down feed.
This will allow for material removal at smaller

depths of cut, until no further removal takes place

(i.e. sparks are no longer generated). Although

spark-out is not a high removal rate technique (in

fact, it removes very little material), it may have an advantage in obtaining a good surface finish. Four-point flexure testing was used to determine

the effect of grinding conditions on the strength of

two commercially available silicon nitride materials.

Weibull statistical analysis was performed on the

flexure strength data and fractography was performed

to identify the location and source of fracture in each

specimen. To characterize the surface integrity,

surface roughness was measured with a three-dimensional profiling instrument and the ground surfaces were examined in a scanning electron microscope.

Experimental Procedure

a) Materials

In this study, two types of silicon nitride were

used — reaction-bonded (RBSN, produced by Eaton

Corporation) and sintered reaction-bonded (SRBSN,

produced by Ceradyne, Inc.). Some important

material properties of each are shown in Table 1. The composition and niicrostructure of these

materials are detailed in Part I of this investigation

[3]. The materials were received in billets. One or two flexure specimens were made from each 86 x 10

X 3.2 mm RBSN billet, which was received in three

batches. The SRBSN billets were 76 x 76 x 13 mm

and were sliced into 4 equal pieces with approximate

dimensions of 50 X 25 X 13 mm. A total of 8 to 12

flexure specimens were made from each SRBSN

sliced portion. The billet number, location in the

billet, and batch number of the billet were tracked

for each specimen to check the possibility of billet to

billet variation and the influence of specimen

location within each billet on strength.

b) Grinding Conditions

A total of 25 to 30 specimens from each material were machined under each set of grinding conditions

used. The specimens were ground to dimensions and

tolerances required by ASTM Standard C 1 161

Standard Test Method for Flexural Strength of

Advanced Ceramics [5] size B specimens (3 x 4 x 50

mm). All edges along the length of the specimens were chamfered to reduce the possibility of fracture

initiation from the edges, as prescribed in the

standard.

280

Table 1. Material properties of RBSN and

SRBSN

Material RBSN SRBSN

Density

(g/cm"

2.36 3.22

Porosity 20% < 1%

Alpha to Beta

Ratio

7.7 : 1 All Beta

Grain Size Length x Width (m)

1-2 X 0.5-1 6-8 X 1-3

Three sets of grinding parameters were used to finish grind specimens from each material. The

grinding parameters are shown in Table 2 (condition

D was used on both RBSN and SRBSN). In all cases, horizontal spindle surface grinding was

performed in the longitudinal direction of the flexure

bars with resin-bonded diamond grinding wheels.

These conditions were chosen to be rather aggressive

as a continuation of the first part of this investigation

[3]. A 320 grit wheel was used for grinding

condition D with 15.2 m/min table speed and 0.0051

mm down feed; the cross feed of 4.00 mm/pass

represents grinding the full width of a specimen in each pass. In condition D, 1-2 spark-out passes were

performed as the final grinding step. Rough grinding was performed at the same conditions with the exception of the down feed, which was incremented

in the following steps: 0.0254 mm to net dimension

+0.1524 mm and 0.0127 mm to net dimension

+0.0635 mm. Conditions E through H represent

creep-feed grinding, where typically the table speed

is lower and the down feed is larger than in conventional grinding. In conditions E and G, a cross-feed of 1.91 mm/pass was used, whereas, the

full specimen width was used in conditions F and H. The resulting volumetric removal rates ranged from

310 mmVmin to 516 mmVmin for the RBSN and

310 mmVmin to 1737 mmVmin for the SRBSN

material. In addition, each grinding condition was

performed on a different machine tool by a different

operator.

In terms of material removal rate, there was some overlap between this study and Part I. However, the high removal rates were obtained by

different means than those in Part I. Conditions E

through H approach a creep feed grinding mode with

high down feeds and low table speeds.
Table 2. Grinding parameters

RBSN D E F

Grinding

Wheel Grit / Concentration

320 / 150 150 / 75 80 / 100

Wlieel Surface Speed

(m/s)

28 33 28

Grinding

Huid

Emulsion

14:1

Emulsion

20:1

Emulsion

20:1

Tallin ^n^<w1 (m/min)

15.2 1.52

Cross Feed (mm/pass)4.00 1.91 4.50

Down Feed (mm)0.0051 0.1778 0.0508

Volumetric Removal

Rate (mmVmin)

310 516 347

SRBSN D G H

Grinding

Wheel Grit / Concentration

320 / 150 150 / 75 80 / 100

Wheel

Surface Speed

(m/s)

28 33 28

Grinding

Fluid

Emulsion

14:1

Emulsion

20:1

Emulsion

20:1

Table Speed

(m/min)

15.2 1.52 1.52

Cross Feed (mm/pass)

4.00 1.91 4.50

Down Feed (mm)0.0051 0.2540 0.2540

Volumetric Removal

Rate (mmVmin)

310 737 1737

Notes: Condition D used 1-2 spark-out passes as a final

step. Conditions D, F, and H used Cimperial

HD90, whereas E and G used Cimperial 1010

soluble oil grinding fluids.

c) Surface Integrity Evaluation

The surface integrity of the ground specimens was evaluated with a three-dimensional stylus

profiling system [6] and scanning electron

microscopy. Three-dimensional topographical maps

were made of the surfaces of selected test bars of

each material for each grinding condition by

compiling multiple profile traces taken at fixed inter- vals on the sample surfaces. Each trace ran perpendicular to the grinding direction. A total of

200 traces, 3 mm in length and set 10 |im apart, were taken. This corresponds to an area of 3 x 2 mm. A subset of the traced area, approximately

1 X 1 mm, was used to calculate the topographical

parameters: average roughness (RJ, root mean

square roughness (R^), and average peak to valley

height (R).

d) Flexure Strength Tests The flexure strength of the ground test bars was

measured in accordance with the ASTM Standard

[5]. A universal testing machine was used at a cross-head speed of 0.5 mm/min to apply a steadily

increasing load to each specimen. A load cell was

used to measure the load, and an integrated stripchart recorder was used to record the applied load. A fully articulating steel fixture [7] was used to mount the specimens for testing in four-point bend

position. This fixture has the advantage of

articulating during loading to correct for any

misalignments due to dimensional errors in the

specimens. This ensures a good line-contact between

each load-pin and the specimen. The system was

calibrated before beginning the flexure tests, and

checked periodically during the testing sequence. A

more detailed description of the test procedure is given in Part I [3].

Experimental Results

a) Surface Integrity

The Rq roughness values corresponding to each

grinding condition are shown in Figure 1 for both

materials. The observed dependance of Rq roughness on grinding condition and material type was also evident in other roughness parameters such as R and

Rg. Figure 1 shows that RBSN has a larger surface

roughness than SRBSN for each comparable grinding

condition. Even for grinding condition F, where the

down feed is lower than condition H by a factor of

5, RBSN exhibits a much higher roughness than

SRBSN. The figure also shows that the R,

roughness values for conditions D and E are nearly
0,6

D EG F H

Grinding Condition

Figure 1. roughness resulting from each

grinding condition for each material.

equal. This indicates that spark-out with a 320 grit

wheel did not produce a finer surface finish than that obtained with a 150 grit wheel at a much higher

down feed. Although it is not possible to establish a

correlation between the wheel grit size and surface

roughness, because of the variations in grinding

parameters, a certain trend is clear. Grinding

conditions F and H, which used an 80 grit wheel,

also produced the largest surface roughness. It is particularly interesting to note the difference between

conditions E and F, where the down feed is larger in the former condition by a factor of 3. Nevertheless,

the Rq roughness produced by condition F is approximately twice that produced by condition E. Similarly, the roughness produced by condition H is larger than that produced by condition G. This

shows the influence of grit size on roughness.

The topographical maps resulting from the three- dimensional profiling are shown in Figures 2 and 3

for RBSN and SRBSN, respectively. It is important

to note the difference in surface characteristics

between the two materials. In general, the RBSN

surfaces are more random and grinding striations are not so clearly evident. The SRBSN, however,

displays distinct and regular grinding striations. In addition, for each material, differences in the severity

of grinding striations and the character of the

surfaces between each grinding condition are evident.

The micrographs of the surfaces in Figures 4-7

confirm the above observations. For RBSN, in Figures 4 and 6, condition D displays semi-flat

regions with distinct pits on the surface. The surface

generated under condition E is similar to the surface

generated by condition D, but lacks much of the

smooth regions. The condition F surface is more

broken up or crushed, on a small scale, but the peaks

and valleys of the grinding striations are revealed as dark horizontal bands across the surface.

In the SRBSN material, there is little difference

in the surface characteristics observed in the

micrographs in Figures 5 and 7. Under grinding

condition H, however, deeper grinding striations are observed, and the bottom of the deep grooves appear

smooth and somewhat featureless. These observations suggest that considerable

plastic flow occurs during material removal in SRBSN, whereas for RBSN, brittle fracture is the

dominating process. The spark-out process, however, appears to be associated with plastic

deformation or polishing of the surface, especially in RBSN. Although the polishing action produced

locally smooth areas, it did not reduce the overall

roughness.

b) Flexure Strength

The two-parameter Weibull distribution was used

to statistically evaluate the strength data. Weibull

plots were used, where the probability of failure, on

Weibull probability axis, is plotted against the

fracture stress on natural log axis. The slope of the

line obtained by least squares fit is the estimate of

the Weibull modulus and the strength at which the

probability of failure is 0.632 is the estimate of the

characteristic strength. The 95% confidence intervals were calculated using the method based on the law

of propagation of errors developed by Ritter, et al.

[8]. This analysis has been applied successfully to flexure data of ceramics [9]. The Weibull plots of flexure data for each

grinding condition are shown in Figure 8 for RBSN

and Figure 9 for SRBSN. The data is well fitted by

a straight line in each case, thus indicating good

representation of the data by the Weibull distribution.

The Weibull parameters were calculated without the

3 lowest strength data points in condition E and the

lowest data point in condition F. These data were

not representative of the population as failure was

determined to be caused by the presence of

unusually large iron-rich inclusions resulting in extremely low flexure strengths. In the Weibull

plots, it is apparent that for both materials condition

D stands out with lower slopes.

The Weibull parameters are displayed for each

grinding condition in Figure 10 for RBSN and Figure

11 for SRBSN, along with the 95% confidence

intervals. As shown by the Weibull plots, the

characteristic strengths do not vary significantly

among the grinding conditions, ranging from 206

MPa to 216 MPa for the RBSN and from 690 MPa

to 703 MPa for the SRBSN. The difference in


Figure 4. SEM micrographs of the ground surfaces Figure 5. SEM micrographs of the ground surfaces

of RBSN for grinding conditons D, E, and F. of SRBSN for grinding conditons D, G, and H.

285

i
D

II

D

H

Figure 6. SEM micrographs of the ground surfaces

of RBSN for grinding conditons D, E, and F (at a higher magnification than in Figure 4).

Figure 7. SEM micrographs of the ground surfaces

of SRBSN for grinding conditons D, G, and H (at a higher magnification than in Figure 5).

286
0.999

0.99 0.95

0.90

0.80

1 0.40

E 0.20 >

I

0.05

0.02

0.01

0.005 0.0025

o Condition 0

Weibull Uodulus: 9

Chorocterislic Strength; 206 UPo

° Condition E

Welbull Uodulus 8

Chcjrocteristic Slrenqlh: 210 MPo

' Condition F

Weibull Modulus: 12

Choroctcristic Strength; 216 UPa

200 300

fracture Stress (UPo)

Figure 8. Weibull plot of RBSN samples for each

grinding condition.

0.999

0.99 0.95

0.90

0.80 0.60

E 0.40

° C.20

J- 0.10

I

""^

0.02

0.01

0.005 0.0025

o Condition 0

Weibull Uodulus; 7

Chorocterislic Strength: 690 "Po

* Condition G

Weibull Modulus: 1 1 IhoroCleristic Strength; 692 MPo

• Condition H

Weibull Uodulus: 1 1 Chorocterislic Strength 70.3 UPo

D >/

/ H

400 500 600 700 800 900 1000

Froctur* Strasf (UPg)

Figure 9. Weibull plot of SRBSN samples for each grinding condition.

strength between the two materials is expected due

to their different processing methods and

microstructures. In terms of the 95% confidence

intervals, a statistically significant difference is observed only for the Weibull modulus between

conditions D and E in RBSN in Figure 10. The data

in Figures 8 and 9, however, indicate that the

variation in strength for both materials is increased

by grinding condition D, since the Weibull modulus

is lower for this grinding condition.

c) Fractography

Fractography was performed, according to the

published guidelines in Reference [10], on each

specimen to identify the location and source of

fracture initiation and to determine if machining

damage played a role in fracture. Fracture origins were characterized by identity and location, as recommended by Quinn et al. [11]. The procedure

used in this study was identical to that used in Part

I.

Identity characterizes the phenomenological

aspects of the flaw and identifies the type of flaws.

For example, agglomerates, large grains, pores,

inclusions, and other microstructural features are usually volume-distributed flaws; machining damage,

handling damage, and pits from oxidation or corrosion are usually surface-distributed flaws.

Location identifies where the flaw is located in a specific specimen. The three classifications used

are volume, near-surface or surface, and edge.

Classification of the location of the flaws is subjective and depends on the researchers definition

of bulk vs. near-surface. In the present paper, volume-located flaws refer to those fracture origins

observed in the examination of the fracture surfaces

by optical microscopy to be approximately 100 pjn or more below the surface. Any origin that appeared

closer to the surface was identified as surface or

near-surface. The latter specimens were considered

to be candidates having possible machining related fracture origins if no obvious flaw, such as an

inclusion, was observed at the fracture origin.

5 200

D E F Grinding Condition

0 E F Grinding Condition

Figure 10. Characteristic strength and Weibull

modulus of RBSN for each grinding condition.

nding Conditioi

0 G

Grinding Condition

287

Figure 11. Characteristic strength and Weibull

modulus of SRBSN for each grinding condition.
Table 3. Summary of fractography

RBSN

D E F

Volume 4% 17% 13%

Edge and Chamfer11% 13% 13%

Surface or

Near-Surface 85% 70% 74%

Possible Machining

Damage

0% 0% 0%

SRBSN

D G H

Volume 10% 29% 21%

Edge and Chamfer28% 13% 3%

Surface or NearSurface 62% 58% 76%

Possible Machining

Damage

0% 0% 0%

Selected specimens were examined in an SEM since,

in general, machining damage is very difficult to discern with optical microscopy [10,11,12].

A summary of the fractography results is shown

in Table 3. Some of the flexure specimens fractured

from flaws located in the volume and near the edge

and chamfer. Examination of the fracture surfaces in the SEM did not reveal any evidence for fracture

initiation from machining damage. Associated with

the majority of failures were near-surface or surface

flaws. The fracture origins were associated with

pores, agglomerates, and inclusions in RBSN and

with inclusions, pores, and regions of incomplete

densification in SRBSN. For example, the fracture surface of an SRBSN specimen is shown in Figure

12. Although the fracture origin is very close to the

surface, there is little evidence of machining damage.

Because a region of porosity can be readily

identified at the fracture origin, it can be assumed

that machining damage did not play a significant role

in the fracture of this specimen.

Discussion

As stated in the introduction, the goal of our

research program is to obtain data and information

288

for the optimization of the grinding process for

ceramics. To achieve this objective, grinding should

be performed at the highest possible removal rates consistent with maintaining the desired strength of

the part. A wide range of removal rates have been

explored in this investigation (in both Parts I and II). The results clearly indicate that the fracture strengths

of the two materials are relatively unaffected by the

grinding conditions, although material removal rates were varied by a factor of 60. Since a broad range of grinding conditions were

used, it was expected that a wide range of machining

damage would be introduced. However, characteristic strength was not significantly affected by any of

the grinding conditions used. This indicates that the

pre-existing flaws in the two materials had a larger

influence on characteristic strength than the flaws

produced in grinding. The implications of these

findings is that the optimum grinding conditions

have not yet been obtained. Therefore, further

testing is warranted at even higher removal rates.

Figure 12. Fracture surface of an SRBSN specimen at a) low magnification and b) high magnifi

cation.
A key requirement for a production viable

grinding process is a high material removal rate [13].

Creep feed grinding has this characteristic and may

be the technique of choice for grinding advanced

ceramics. Creep feed grinding techniques can produce parts with higher strength and better surface finish

than conventional grinding [14]. Normal forces,

however, are much higher for creep feed grinding

and may require machine tools with higher stiffness. Surface roughness measurement and evaluation

of ground surfaces in the SEM indicated that surface

integrity is influenced by the grinding condition. In

particular, creep feed grinding with coarse grit

wheels produced a rougher surface in both materials.

In SRBSN, there did not appear to be any grinding

damage. However, the brittle fracture process, which

dominated the removal process in RBSN, could have

resulted in the formation of microcracks. Nevertheless, a detrimental effect on strength was not observed.

It must be emphasized that the data presented in

this paper are only applicable to the two types of

silicon nitride materials tested in the present investigation. Application of these results to other types of

silicon nitride materials could be erroneous. For

example, silicon nitride materials with a reduced

flaw population, compared with the RBSN and

SRBSN, are expected to be more sensitive to machining damage. Recent studies [15,16] have shown

that grinding related surface damage can be detri- mental to the strength of hot isostatically pressed

silicon nitride, even on longitudinally ground tensile

specimens.

A question of the sensitivity of the flexure

experiments to machining damage may also be

brought up by these results. Previous studies have

shown that strength is reduced by 30-60% when the

specimens were ground in a direction perpendicular

to the long axis of the flexure specimens [17]. This

reduction in strength is attributed to the formation of

median cracks by grinding, such that the crack plane

is perpendicular to the tensile axis during flexure

testing [18,19]. It is then possible that specimens

ground in the transverse direction are more sensitive

to grinding damage, particularly, at high removal

rates.Although the results in this paper showed that the strength was not affected by the grinding conditions used, the variation in strength was increased by

the grinding condition in which spark -out was used.

During spark -out passes, the diamond particles slide on the workpiece without removing much material.

This added abuse may have amplified some of the

damage from previous grinding, therefore, randomly

introducing machining flaws that reduce strength in some specimens. Another concern in using spark -out

is that the stiffness of the machine tool must be

sufficiently high to allow very low (and therefore

precise) down feeds [20]. Otherwise, dynamic instability can result and produce undesirable surface

integrity'. A possible explanation for the observed

variation in strength with the spark -out passes is an

increased level of machining damage and interaction

of this damage with the pre-existing flaws. Such

interactions would obscure identification of the cause of fracture when using fractography. Nevertheless,

since the surface roughness was not improved, the

spark-out passes did not provide a beneficial effect.

Conciusions

1. Varying the volumetric removal rate in surface

grinding from 310 to 516 mmVmin for RBSN and

from 310 to 1737 mmVmin for SRBSN by employing conventional and creep-feed grinding conditions

had no significant effect on the characteristic

strength of four-point flexure specimens made from

these materials.

2. The grinding condition that used spark-out

reduced the Weibuli modulus of the flexure specimens, i.e., increased the variation in flexure strength.

3. Surface roughness and topography were affected

by the grinding conditions. A coarse grit wheel used

in creep-feed grinding generated a rougher surface

than surfaces produced with finer grit wheels.

Acknowledgments

This project was conducted under the auspices of

the Ceramic Machining Consortium, which is a joint

research program between NIST and several companies, universities, and government agencies. The

authors are grateful to James Edler of Eaton Corporation and John Mangels of Ceradyne, Inc., for

providing the silicon nitride materials used in this study. The financial support received from the

Department of Energy — Office of Transportation

Materials, is gratefully acknowledged. We thank

Peter Blau, Ceramic Machining Project Coordinator,

Oak Ridge National Laboratory, for his encouragement and his assistance.

Information on product names, manufacturers, or

suppliers are included in this report for clarification. This does not imply endorsement of the products or

services by NIST.
References

[1] K. Subramanian and S. Ramanath, "Mechanism

of Material Removal in the Precision Grinding of

Ceramics," Precision Machining: Technology and

Machine Development and Improvement. PED-Vol.

58, ASME Winter Annual Meeting, November 8-13,

pp. 1-19, 1992.

[2] S. Jahanmir, L. K. Ives, A. W. Ruff, and M. B. Peterson, Ceramic Machining: Assessment of Current

Practice and Research Needs in the United States. NIST Special Publication 834, U.S. Government

Printing Office, Washington, DC, 1992.

[3] S. Jahanmir, T. J. Strakna, G. D. Quinn, H.

Liang, R. L. Allor, and R. D. West, "Effect of

Grinding on Strength and Surface Integrity of Silicon

Nitride: Part I," Proceedings of the International

Conference on Machining of Advanced Materials. Gaithersburg, MD, July 20-22 1993.

[4] S. Malkin, Grinding Technology. New York,

NY, pp. 188, 1989.

[5] ASTM Standard C 1161-90, "Standard Test Method for Flexural Strength of Advanced Ceramics

at Ambient Temperature," American Society for Testing and Materials, Philadelphia, PA, 1990.

[6] E. P. Whitenton, M. B, Peterson, and L. K. Ives,

"Method for Quantitative Measurement of Galling

Damage," Metal Transfer and Galling in Metallic

Systems. H. D. Merchant and K. J. Bhansali (Eds.),

The Metallurgical Society, Warrendale, PA, pp. 155-

170, 1987.

[7] G. D. Quinn, "Twisting and Friction Errors in Flexure Testing," Ceramic Engineering and Science

Proceedings ." Vol. 13, No. 7-8, pp. 319-330, 1992.

[8] J. E. Ritter, Jr., N. Bandyopadhyay, and K. Jakus,

"Statistical Reproducability of the Dynamic and

Static Fatigue Experiments," Ceramic Bulletin . Vol.

60, No. 8, pp. 798-806, 1981.

[9] G. D. Quinn, "Flexure Strength of Advanced

Structural Ceramics: A Round Robin," Journal of the

American Ceramic Society . Vol. 73, No. 8, pp. 2374-

2384, 1990.

[10] MIL-HDBK-790, "Fractography and Characterization of Fracture Origins in Advanced Structural

Ceramics," U. S. Department of Defense, July, 1992.

[11] G. D. Quinn, J. Swab, and M. J. Slavin, "A

Proposed Standard Procedure for Fractographic

Analysis of Monolithic Advanced Ceramics," in Fractography of Ceramics and Glasses. V. Frechette

and J. Vamer, (Eds.), Vol. 22, 1976.

[12] R. W. Rice, S. Freiman, J. Mecholsky, Jr., R. Ruh, and Y. Harada, "Fractography of Silicon Nitride

and Silicon Carbide, " Ceramics for High Perfor- mance Applications II ." J. Burke, E. Lenoe, and R. Katz, (Eds.), Chestnut Hill, MA, pp. 669-687, 1978.

[13] K. Subramanian, S. Ramanath, and Y. O.

Matsuda, "Precision Production Grinding of Fine

Ceramics," First International Conference on Manufacturing Technology, Chiba, Japan, 1990.

[14] W. Konig, L. Cronjager, G. Spur, H. K. Tonshoff, M. Vigneau, and W. J. Zdeblick, "Machining of New Materials," Processing of Advanced

Materials . No. l,pp. 11-26, 1991.

[15] M. R. Foley and V. K. Pujari, "Tensile Testing

in the Development of Processing Methods for High

Strength/High Reliability Silicon Nitride," Ceramic

Engineering and Science Proceedings. 16th Annual

Conference on Composites and Ceramics, Fall 1992.

[16] R. H. Licht, M. R. Foley, L. C. Sales, and D. M. Tracey, "Surface Integrity in Advanced Structural

Ceramics," Presented at the Workshop on Superabrasives and Grinding Wheel Technology for

Machining Ceramics, Oak Ridge National Laboratory, Oak Ridge, TN, May 1992.

[17] H. Miyasato, H. Okamoto, S. Usui, A.

Miyamoto, and Y. Ueno, "The Effect of Grinding on

Strength of Hot-Pressed Silicon Nitride," ISIJ In- ternational. Vol. 29, No. 9, pp. 726-733 1989.

[18] R. W. Rice and J. J. Mecholsky, "The Nature

of Strength Controlling Machining Flaws in Ceramics," The Science of Ceramic Machining and Surface Finishing II . National Bureau of Standards

SP562, U. S. Government Printing Office, Washington, DC, pp. 351-378, 1976.

290
[19] B. R. Lawn, A. A. Evans, and D. B. Marshall, [20] K. Subramanian, Personal Communication,

"Elastic/Plastic Indentation Damage in Ceramics: The 1993.

Median/Radial Crack System," Journal of the American Ceramic Society. Vol. 63, No. 9-10, pp. 574-

581, 1980.

291

PRECISION MACHINING

293

TOOL FLANK WEAR IN SINGLE POINT

DIAMOND TURNING

J. S. STRENKOWSKI and W. C. LARSON

North Carolina State University, Raleigh, NC

In this paper, wear mechanisms on the flank of diamond tools are investigated. In particular, wear due to solid state diffusion of carbon atoms from the tool into the workpiece and abrasion wear are studied for their relative contribution to overall flank wear. Flank wear in single point diamond turning is of particular importance because degradation of the flank can lead to surface roughness and damage in the machined workpiece

surface.

The wear of diamond tools during machining is highly influenced by the temperature

and velocity of the chip/tool interface. Using a 2-D Eulerian model of orthogonal cutting, the chip velocity and the temperature distribution in the chip and the tool is determined. By coupling this information with a model of wear, the relative contributions of

diffusion-controlled wear and abrasion effects can be assessed.

It is found that flank wear is due to solid-state diffusion. Computed tool temperatures

are consistent with temperatures required for producing a prescribed level of observed

wear from experimentally-measured worn tools.

1. INTRODUCTION

Diamond tools can wear in a variety of ways in- cluding mechanical attrition and thennal degradation. Mechanical attrition is the abrasive removal

of material from the tool by the workpiece. Wear

due to abrasion is a function of the relative hard- ness of the workpiece and the tool, which is a function of the crystal orientation of the tool. It is negligible when machining soft materials, but it is the dominant mechanism for tool material re- moval for hard materials (1) such as during diamond lapping of a diamond tool. Thermal

degradation occurs at elevated temperatures when

the diamond crystal oxidizes and/or tums into

graphite (graphitization). These reactions occur at high temperatures and are dependent on the crystal structure of the diamond, the presence of oxygen, and pressure. Diamond tools begin to oxidize and graphitize at 1500 °C in an inert environment (2). In the presence of oxygen this reac- tion can start to occur at 500 °C (3). However,

experimental observations and theoretical cutting

models indicate that temperatures of this magnitude are not reached in single point diamond

turning (SPDT) of copper (4), so that thermal

degradation is an unlikely wear mechanism.

Diamond tools exhibit wear on both flank and

rake faces during single point turning. Wear on

the rake face is characterized by crater pattems

along the rake and almost negligible wear near the

cutting edge. In contrast, flank wear is characterized by the appearance of a wear land on the tool

flank parallel to the surface of the workpiece,

which is relatively smooth. Flank wear is of particular interest because it can lead to surface damage and roughness in the machined workpiece.

The two most common forms of diamond tool wear include abrasion and chemical wear. In chemical wear, solid state diffusion of carbon

atoms from the tool occurs directly into the

workpiece. This can be seen in the machining of

materials such as steel in which carbon is highly

soluble. For these materials, diamond tools wear

295
very quickly (5) and it is suspected that solid state difftision is the dominant wear process. Because

solid state diffusion depends exponentially on the

temperature, small changes in temperature can

have large effects on the wear rate. This has been

observed experimentally in wear patterns of diamond tools, in which deep wear pits occur in re- gions of high temperature, which are adjacent to unworn areas of low temperature (6). Therefore,

there is some basis for diffusion as the dominant

chemical wear mechanism for diamond tool wear.

The temperature must be accurately predicted in

order to predict the wear due to diffusion. There

have been numerous attempts to evaluate the temperature in diamond turning. The average temperature in SPDT was direcUy measured by Iwata

(4), but a discrete temperature distribution along

the rake and flank faces has yet to be determined

by experiment. The temperature distribution can be evaluated using advanced numerical techniques, such as the finite element method. A

model of orthogonal cutting has been developed

by Strenkowski and Carroll (7), (8) and by

Strenkowski and Moon (9) for chip geometry and

temperature prediction. These models can be

applied to micromachining in single point diamond turning with additional refinement.

Specifically, modifications must be made to the

boundaries because they are close to the cutting zone so that the temperatures within the workpiece

and tool are extremely sensitive to the thermal

boundary conditions. This sensitivity can be

eliminated by incorporating a reduced thermal

conductivity along the tool boundary that ac- counts for the remaining portion not included in

the tool model.

This paper focuses on the understanding of diamond tool wear mechanisms for cutting of copper. In particular, wear due to solid state diffusion

and abrasion wear are investigated for their

relative contribution to the overall flank wear.

Predicted flank wear is compared with measured wear rates for cutting of copper, and excellent

agreement was found. Temperature contours in

the tool confirm that flank wear is largely a diffusion-controlled mechanism.

2. FORMULATION OF A VISCOPLASTIC

CUTTING MODEL

An Eulerian finite element technique can be used

to simulate the complex interaction between the

workpiece and the tool in terms of basic engineering quantities such as strain, strain-rate, flow stress, pressure, temperature, and tool forces. In this technique, the grid is spatially fixed and the

workpiece behavior is treated as viscoplastic. The

governing equations can be derived as follows.

For a body undergoing plane strain steady-state

deformation, the equations of equihbrium can be

written in the form:

[L]{o} + {f} = 0 in V (1)

where ( is a vector of stress components, [L]^ is a matrix of first-order differential operators in x and y, and (f} is a vector of applied body forces.

Because the workpiece undergoes large deformations, it is more appropriate to describe its behavior in terms of strain-rate and particle velocity,

rather than strain and displacement as in conventional solid mechanics problems.

For extremely large strain-rates, the workpiece deforms in a viscous manner. In general, the deviatoric stress {cj } is linearly related to the strain-rate

as:

(a) = [D]je) and P] =

2 0 0

0 2 0

L 0 0 1 J (2)

where jx is the viscosity of the workpiece and is given by

1^ = - +

(3)

where £ is the second invariant of the strain-rate tensor or the effective strain-rate, o is the effective

flow stress, and jJ. accounts for work hardening of

the workpiece. For the special case of ideal plasticity, U becomes zero.

The stress can be written in terms of the deviatoric

stress {o } as:

{o} =|a') -(m)p and {m)T = { I i o) (4)

The pressure p can be eliminated by recognizing

that it is related to the rate of volumetric strain Ey

by

p =Gt€v =oc{m}^[L]{u} where a=lC^^

(5)

where E is Young's modulus, and {u} represents

the particle velocity components. This is the so- called penalty function formulation (10) which is widely used for dealing with incompressible or

nearly incompressible problems. The advantage

of using a penalty function method is that the

pressure can be expressed in terms of the strain296
rates and eliminated from the governing equations. This reduces the number of degrees of

freedom to only nodal velocities. However, a drawback of this approach is that a suitable value

for the penalty number a must be chosen. The

value of k will vary depending on the application

and the size of the computer word. Note that a very large value for k will lead to iU-conditioning

or over-constraint of the governing equation. In

contrast, a small value for k does not satisfy the in- compressibihty constraint. For the cutting model,

it was found that a k equal to 5 was suitable.

The equations of equilibrium may now be written

in terms of the particle velocities by substituting

equations (2), (4), and (5) into equation (1), which

gives

-[LF(m}a{m)T[L](u}-h[L]T[D][Hu) + {f} = 0(6)

which is applicable for steady-state cutting. The

finite element equations may be derived by using

a Galerkin approach (11). In this method, the

governing differential equation for the flow field

is satisfied over the domain, where the weighting

functions are chosen to be the same shape functions used for discretizing the velocity,

[u) = [N]{U} (7)

where [N] is the shape function matrix and {U} is the vector of nodal velocities. Substituting equation (7) into (6), multiplying the resulting equation by [N]"^, and integrating over the volume results in the discretized equations:

[[K]-h[KP]]{U}={F]

(8)

where [KP] contains the penalty term a. In solving equation (8), a penalty function formulation

must be used in which a is identified as the

penalty num.ber. Because a approaches infinity

for an incompressible material, reduced integration techniques must be used to evaluate [KP] as compared to [K] to eliminate element locking.

Note that the viscosity can become very large and

exceed the allowable real number range of the

computer. Therefore, a suitable cut-off value

must be supplied to avoid numerical error. For

the cutting model described above, a cut-off in the

C If) range of 10 - 10 was found to be suitable.

Thermal Model

During metal cutting, a significant amount of heat

is generated due to the plastic work in the primary

shear zone and frictional sliding along the

tool/chip interface along the rake and flank faces. The high temperature generated during cutting

can directly influence the resulting quality of the

machined parts and the life of the cutting tool. For example, crater wear and flank wear are

greatly effected by the temperature.

The highest temperature and temperature gradients occur in a small region close to the cutting

edge. A two-dimensional finite element therm^

model can be coupled to the cutting model described in the previous section. The resulting velocity, strain-rate, stress, and chip/tool contact

length as calculated by the cutting model are used

as input data for the thermal model. The thermal

model for a two-dimensional, steady-state, heat

conduction-convection problem is governed by

the following partial differential equation:

dx dy (9)

where k is the thermal conductivity, p is the mass

density, Cp is the specific heat, and Q is the volumetric heat generation rate defined as:

0 = Ty + TfVs (10)

where t is the effective flow stress, 7 is the effective

strain-rate, Xf is the interfacial shear stress, and Vg

is the sliding velocity of the chip along the rake

face. The first term in this equadon is the heat

generated due to the plastic work in the entire

workpiece, and the second term is the heat generated due to the frictional sliding along the

tool/chip interface. Because only a smaU portion

of the workpiece and the tool is modeled, the

thermal boundary conditions must be chosen very

carefully. In particular, the boundary of the tool

in the model is actually an internal boundary.

Therefore, prescribing a room temperature condition on this boundary is not realistic. However,

this problem can be overcome by reducing the

thermal conductivity of the elements along the

boundary of the tool to account for the portion of

the tool that is not included in the model.

The thermal finite element equations can be

derived from the temperature equation (9) by us- ing a standard Galerkin approach. Expressing the

temperature in terms of the same shape functions

as in equation (7) leads to

[Kp{T}e = [FY (11)

where {T} is the nodal temperature vector and

[K]e = ( ,(3[NlTa[N]^9[NlT3[Nl^

J dx dx dy dy

J x,y
J pCp.N,V_^.v^,dxdy

and

{F}e =

1 [NJ'TQ dxdy

J-y (13)

Note that the workpiece, chip, and tool are included in the thermal model, so that a complete

mapping of the temperature field in the workpiece

and tool can be determined.

3. DIAMOND TOOL WEAR MECHANISMS

In the next two sections, the relative contribution

of abrasion and diffusion-controlled wear to flank wear will be investigated.

3.1 Abrasion Wear

The high speed flow of workpiece material across

the flank of the tool provides an ideal condition

for abrasive wear. Abrasive wear is governed by

Archard's wear law:

W = KiPS (14)

where W is the wear volume, P is the normal load,

S is the sliding distance, and Kj is the wear volume per unit load per unit sliding distance, or wear constant. This equation can be used to find

the wear rates by replacing the sliding distance S

by the relative sliding velocity.

To investigate the contribution of abrasive wear to overall flank wear, previously reported diamond

turning experiments with worn tools were carried

out in which changes along the cutting edge and

clearance face were measured (12). A straight

edge tool with a 0° rake angle and a 6° clearance

angle was used to plunge cut a plated copper

workpiece to a depth of 1 mm. The width of the cut was normalized to 1 mm. The wear was determined by directly measuring the wear land and

indirectly determining the tool edge radius from

the measured tool forces. The worn tool geometry can be determined from the wear land length,

tool edge radius, rake angle, and clearance angle,

which can then be used to compute the wear vol- ume. Table 1 summarizes the changes in the tool

geometry as a function of the cutting distance for

diamond turning of copper. The change in tool

geometry is shown schematically in Figure 1.

The volume worn away can be converted to wear

rate W by:

Table 1: Worn tool geometry variation.

(Km)

Fiit^p Radius;

(mm)

Wpar T Jinrl (mm)

16 0.10 0.9

20 O.lS l.l

40 0.59 1.^

70 0.84 2.7

100 0.95 3.3

where A is the measured wear area, w is the width

of cut, V is cutting velocity, and d is the distance cut between measurements. Using this equation

for a 10 m/s cutting speed and cutting width of 1 mm, the wear rates for various elapsed distances can be calculated, which are summarized in Table

2. The last column of Table 2 shows the wear rate divided by the average wear land length. Note

that the relatively constant normalized wear rate implies a wear process that is dependent on the area of contact.

Values of Ki of 1.4 to 4.1 * lO'^^ mm^/Nm have

been determined using slow speed rubbing tests of

diamond on copper (1). Using these values and

the measured thrust forces, abrasive wear rates of

6.02 to 17.6 X 10' ^2 mm^/s are predicted by

equation (14). The range of values of Ki are due

to the uncertainty in the crystal orientation of the

diamond tool. Note that these computed wear

rates are much lower than the measured wear rates. Therefore, abrasion is not the dominant wear mechanism and the bulk of the wear must be ac- counted for by another mechanism.

o

1 -

0 2 4

micrometers

w = AwV/d (15)

298

Figure 1: Worn tool profile.
Table 2: Tool wear rates.

Case

(Km)

Total

Area

(mm^)

Wear Rate

1 ^ (10-12

mm^/s)

Norm'lzed.

Wear Rate

(10-9

mm2/s)

10-

20

0.0556 54 45

10-

40

0.315 158 U

40

- 70

m 79

70-

100

0.^64 5ii ^5

where Qool is the concentration of carbon atoms

in the tool and Idia is the lattice parameter of diamond. Note that the carbon concentration in the

chip remains essentially zero due to the short time

of contact with the tool. The flux required to produce a given wear depth is :

tmach (19)

where Jreq is the required flux, dwear is the depth

of wear damage, and tmach is the machining time.

Equations (16) through (19) can then be solved

for the temperature (Treq) necessary to account

for the observed wear: 32 Diffusion-Controlled Wear

Solid state diffusion of carbon atoms from the di- amond tool to the workpiece is governed by Pick's

First Law, which is given by

J =

(16)

where J is the diffusion flux, C is the concentration

of carbon atoms, and D is the diffusion coefficient. In this equation, the diffusion coefficient

depends on the temperature T as.

D = DoexprAEp

RT (17)

where Dq is the frequency factor, AEd is the acti- vation energy, and R is the universal gas constant. These equations are coupled to the thermal analysis through the temperature dependence of the

diffusion coefficient. Once the temperature distri- bution has been determined in the tool, the temperature dependent diffusion coefficient along the

tool/chip interface can be found. Solving Pick's

Pirst law gives the diffusion flux, which is direcfly

related to diffusion wear of the tool.

The previous equations can also be used to calculate the temperature required to produce an experimentally observed wear pattern due to diffusion.

The carbon concentration gradient can be esti- mated by setting the separation between the tool

and the workpiece to be equal to the atomic lattice parameter. The concentration of carbon atoms in

the tool is fixed and the concentration of carbon

atoms in the chip is zero, so that the concentration

gradient can be approximated as: ^ _ Ctool

5x Idia

-AED

R

log dwear ^dia ^mach

log Do

(20)

where l^ia is the lattice parameter of diamond,

dwear is the depth of wear damage, and t^ach is the machining time. Por the interstitial diffusion

of carbon atoms into copper. Do = 74. x 10"^

m2/s and AEd = 37.9 kcal/mole (13). This

equation can be applied to the previous wear ex- periments by setting R = 1.98 cal/mole K, tmach

~ 9000 sec, and l^jia = 0.35 nm. Using equation

(20), a temperature of 254 *^C along the flank of

the tool is required to account for all of the observed wear for a distance of 10 to 100 Km.

4. PREDICTION OF TOOL TEMPERATURES

Diffusion-controlled wear depends exponentially on the temperature in the tool. Therefore, an ac- curate assessment of the role that diffusion plays

in flank wear is highly dependent on accurate determination of the tool flank temperatures. Tool

temperatures can be determined with the finite element model described in a previous section.

Pigure 2 shows the finite element model of the

workpiece, chip, and tool. An enlarged view of

the region near the cutting edge is shown in

Pigure 3 in which the predicted temperature contours are also shown. Note that the maximum

temperature of the tool is 175 °C and it occurs near the cutting edge.

(18)

299
CfflP

WORKPIECE

Figure 2: Finite element model of tool and

workpiece (with flank).

Figure 3: Predicted temperature contours (°C)

near tool edge (without flank).

Figure 4 shows the temperature contours with a

tool flank included. Note that this figure corre- sponds to an edge radius of 0.59 mm and a flank

land length of 1.6 mm. The inclusion of the flank

in the model results in a 75 increase in temperature along the flank to 250 ^C. This is in

good agreement with the required temperature of

254 °C for the observed wear due to diffusion.

300

Figure 4: Predicted temperature contours ( C)

near tool edge (with flank).

The temperatures predicted by the finite element

model can be verified by calculating the temperature rise along the flank of the tool due to friction as follows. The energy due to frictional heafing on the flank Eflank is given by:

Eflank = B2FffVc (21)

where Fff is the measured frictional force on the | flank and Vc is the cutting speed. Assuming that the average temperature rise in the workpiece and

tool are equal, the fraction of frictional heat along

the flank face that flows into the workpiece is given by Jaeger (14) as:

B2= : / i! kwp + kt ^22)

where k^p and kt are the thermal conductivities of

the workpiece and tool, respectively. For a ther- i mal conductivity of 7.03 W/cm °C for type 1 diamond, and 3.95 W/cm for the copper workpiece, this fraction equals 36%. Note that because

of the large thermal conductivity of diamond as compared to copper, the majority of the heat

(64%) flows into the tool, which is the opposite of

what occurs in conventional cutung where the

bulk of the heat flows into the workpiece.
The bulk temperature rise due to frictional heating

along the flank can now be calculated using the

energy balance:

^J _ Eflank

mCpp

(23)

where AT is the temperature rise, Cp is the specific

heat, and p is the density. The mass flow rate ih is found by assuming that the workpiece can be

modeled as two strips of material. One strip represents the chip with a depth equal to the depth of

cut, and the other is imbedded in the workpiece

subsurface. This strip carries heat into the flank

region of the woikpiece. Because of the short duration that the workpiece is in contact with the tool

flank, it is reasonable to assimie that this portion

of the heat is confined to only the subsurface re- gion of the workpiece.

This equation can now be used to determine the

temperature rise due to frictional heating along

the flank. Using a measured frictional force on

the flank of 0.71 N/mm for a 10 m/s cutting speed

gives a temperature rise of 74 °C for a mass flow

rate based on a strip of material equal to the depth

of cut. Note that the incoming temperature of the

workpiece in the flank area is predicted from the

finite element model with no flank to be 175 °C,

as shown in Figure 3. Therefore, the total temperature of the tool flank is 250 °C, which agrees

with the required temperature calculated in equation (20) for the observed wear and the temperature predicted by the finite element model with a

flank, as shown in Figure 4. Therefore, diffusioncontrolled wear is a likely mechanism of flank wear for a diamond tool.

5. CONCLUSIONS

Wear due to soHd state diffusion and abrasion wear have been investigated for their relative contribution to overall flank wear. A finite element

model of the cutting process was used to determine the temperatures in the tool. In addition, a thermal analysis was conducted to determine the

temperature rise due to friction heating along the

flank of the tool. By coupling this temperature

information with a model of wear, the relative

contributions of diffusion-controlled wear and

abrasion effects were assessed.

It was found that computed tool temperatures were

consistent with temperatures required for producing a prescribed level of observed wear from ex- perimentally-measured worn tools. Because the

required temperatures were based on a diffusioncontrolled model of wear, it can be concluded that solid-state diffusion of carbon from the diamond

tool is a plausible mechanism for flank wear.

REFERENCES

1 . D. Crompton, W. Hirst. M.G.W. Howse,

"The Wear of Diamond", Proc. R. Soc.

Lond. A. 333, pp. 435-454, 1973.

2. G. Davies, T. Evans, "Graphitization of

Diamond at Zero Pressure and at a High

Pressure", Proc. R. Soc. Lond. A. 328,pp.

413-427, 1972.

3. J. Wilks, E. Wilks, Properties and

Applications of Diamond. (ButterworthHeinemann Ltd, Oxford), 1991.

4. K. Iwata, T. Moriwaki, K. Okuda, "A

Study of Cutting Temperature in UltraHigh Precision Diamond Cutting of

Copper", 15TH North American

Manufacturing Research Conference

Proceedings. 1987 SME Manufacturing

Technology Review. 2, pp. 510-515,

1987.

5. C.J. Evans, R.L. Lundin, E.S. Paul, R.S.

Polvani, D.D. Stewart, E. Whitendon,

Proceedings of the Seventh Annual

Meeting of the American Society for Precision Engineering, pp. 57-60, 1992.

6. J. Drescher, T.A. Dow, in Precision

Engineering Annual Report CNorth

Carolina State University, Raleigh,

NC),Vol. VIII, 1990.

7. J.T. Carroll III, J.S. Strenkowski, "Finite

Element Models of Orthogonal Cutting

With Application to Single Point Diamond

Turning,", International Journal of

Mechanical Sciences^ Vol. 30, pp. 899-

920, 1988.

8. J.S. Strenkowski, J.T. CarroU III, "An

Orthogonal Metal Cutting Model Based on

an Eulerian Finite Element Method,"

Manufacturing Processes. Machines and

Systems. Proceedings of the 13th NSF

Conference on Production Research and

Technologv . pp. 261-264, 1986.

9. J.S. Strenkowski, K.J. Moon, "Finite

Element Prediction of Chip Geometry and

Tool/Workpiece Temperature Distributions

in Orthogonal Cutting", ASME Journal of

Engineering for Industrv . Vol. 112, pp. 313-318, 1990.

10. O.C. Zeinkiewicz, P.N. Godbole, "A

Penalty Function Approach to Problems

of Plastic Flow of Metal with Large

Surface Deformations", J. of Strain

Analvsis . Vol. 10, p. 180, 1975.

301
1 1 . R.C. Cook, D.S. Malkus, M.E. Plesha,

Concepts and Applications of Finite

Element Analysis. (John Wiley & Sons,

New York, New York), 1989.

12. J.D. Drescher, PhD Dissertation, North

Carolina State University, 1992.

13. D. Bergner, "Diffusion of Heavy

Interstitials in Metals". Metallurgia I Odlewnictwo . Vol. 12, No. 4, pp. 379-400,

1986.

14. J.C. Jaeger, "Moving Sources of Heat and

the Temperature of Sliding Contacts", Proc

of the Roval Society. NSW. Vol. LXXVI,

pp. 203-224, 1942.

302
DEPTH PROFILING OF RESIDUAL STRESS ALONG

INTERRUPTED TEST CUTS IN MACHINED

GERMANIUM CRYSTALS

R. G. SPARKS

Alcoa Technical Center, Alcoa Center, PA

M. A. PAESLER

North Carolina State University, Raleigh, NC

Residual stress profiles in machined germanium have been determined by

measuring the longitudinal optical phonon spectra using micro-Raman

spectroscopy. Lateral spatial resolution of 1 \xm results from this direct nondestructive technique. By measuring the Raman spectra with several probe

wavelengths, axial resolution of 10 nm is obtained through differential

absorption profiling, or DAP. The stress field in these machined brittle

semiconductors is characterized by a surface layer under compression, an

abrupt compressive-to-tensile transition, and a deeper-lying region of tension

which ultimately relaxes in the unstressed bulk. The abrupt transition is indicative of a change from plastic to elastic deformation. The DAP technique

has been applied to interrupted test cuts in diamond-turned germanium to reveal

that fractured regions of the machined surface possess higher tensile stress that occurs at shallower depths than unfractured regions. A qualitative mechanism

of the machining process consistent with these residual stress data is presented.

Introduction

Previous Raman investigations of

residual stresses in thin films have assumed

that the stress field is contained only in the

film with very little penetration of the stress into the sample substrate. Such studies

also generally assume that the magnitude of

the stress is constant throughout the film

with no variations occurring axially. (1-4)

Stress in these films will result in Raman

spectra that are shifted by an amount which

is indicative of the changes in the

frequencies of the normal modes of

vibration in the crystal. Since the stress is constant throughout the depth of the film

there is little or no evidence of asymmetry

in the shifted Raman peaks and the analysis

of the stress magnitude is straightforward.

In samples that display a continual variation

in the stress field such as that occurring in a polished or single point diamond turned

material, the above assumptions can no

longer be made.(5, 6) The stress in this

case is again contained in a thin surface

region analogous to the case of a grown or deposited film but there is no distinct

boundary for the termination of the stress

field. In such cases the stress field may

also be a continuously varying quantity

with respect to depth into the sample. In such cases the Raman peaks are not only

observed to be shifted, but are also seen to be asymmetrically broadened by the

superposition of a continuum of stress

states occurring axially in the material. The

observed lineshape of the resulting Raman

signal can be quantitatively analyzed by

mechanisms based on the convolution of

the penetration depth of the light and the

depth of the stress field.(5 - 9)

The occurrence of asymmetrically

broadened Raman peaks can be understood

if one assumes that the residual stress field

varies continuously with depth. By

probing greater and greater depths in the

material through the use of successively

303
longer wavelengths of light one obtains

stress information from a broad overlying

region of the sample. This leads to a superposition of a continuum of Raman

peaks, with each infinitesimal peak shifted

by the stress field. A Raman spectrum

obtained in such a material measures this

superposition of Raman signals. The

resulting spectrum displays a broadening of

the features. Superposition broadening can

be exploited to obtain high axial resolution

depth profiles of residual stress by

computer profiling of the stress based on

the broadened and shifted Raman

features. (5 - 7)

1200 1

290 300 310 FREQUENCY SHIFT (WAVENUMBERS)

Figure 1. Asymmetrically Broadened

Raman Spectra. Micro-Raman spectra

obtained using 514.5 and 488.0 nm probe

wavelengths are displayed for a single

point diamond turned Ge crystal. Note the

respective shifts of the resulting peaks from

the unstressed position of 300 cm-^ as well

as the asymmetric broadening induced in

each case.

Figure 1 shows micro-Raman spectra

obtained in the ductile region of a feed cut

in germanium for both 514.5 and 488.0 nm

probe wavelengths.(5, 6) The asymmetric

broadenings of the Raman lines are clearly seen. An important detail of the spectra is the evidence of a shift of the peak position

towards compressive stress in the 514.5

nm (green) data and a shift of the peak in

the opposite direction (tensile stress) for the

488.0 nm (blue) data. Both green and blue

spectra show clear evidence of asymmetry.

The asymmetry is also clearly seen to be

wavelength dependent since the green peak

broadens toward lower absolute

wavenumbers (tensile stresses) while the

blue data broadens toward higher absolute

wavenumbers (compressive stresses). If the broadening were due to spectrometer

error one would expect both green and blue

data to show asymmetries that were similar

to one another toward either higher or lower wavenumber. Figure 1 clearly

indicates that this is not the case. (5, 6, 10)

Differential Absorption Profiling of

Residual Stresses

When Hght of intensity Iq penetrates a

solid, the intensity, I, at any depth z will be

determined by an exponential decay of the

incident field. In the case of Raman

scattering, the light must penetrate to a depth z before being scattered and it must

return to the surface in order to be

collected; therefore, the intensity of the

Raman signal from any depth, z, will be

proportional to e'^^^. The absorption

coefficient for the scattered light is actually

very slightly different from that of the

incident light but these slight variations do

not affect the analysis and will not be

considered. The value of the absorption

coefficient, a, depends on both the material

of interest and the energy of the incident

light.(ll) The exponential decay of the

intensity of the Raman signal with depth

combined with the variation of the

absorption coefficient with photon energy

makes determination of stress profiles from

Raman spectra in opaque semiconductors

possible. Other investigations of polish

induced surface strain in semiconducting

materials such as InP and GaAs have

assumed that the stress field is an exponentially decaying function of

depth. (8, 9) While such an assumption

may be valid for stress induced damage

occurring at the surface of a polished

sample it is not valid, in general, for the

process of single point diamond turning,

which typically displays little or no surface

damage in regions of ductile material

removal. (5, 6, 10) The stress fields

generated are thus qualitatively different

304
from those produced in samples with

polish induced surface damage only.

In our analysis of Raman spectra, we

found that in test cases, the stress profiling

was not a function of the choice of the

analytical lineshape chosen for the Raman

spectrum at a given depth. That is, both a Lorentzian and a Gaussian assumption

resulted in essentially the same stress

profile for a given set of input data. The

lack of response of the differential

absorption profiling technique to the

changes in lineshape may be attributed to the similarity of both Gaussian and

Lorentzian lineshapes near the peak

position. Significant variations between

these lineshapes occur at positions more

widely separated from the peak location

where the effects of instrument noise may

be expected to dominate any slight effects

from the lineshape differences. The

convergence of the DAP analysis is much

slower when using the Lorentzian, thus in

order to expedite convergence a Gaussian

lineshape was used in this analysis. (10)

The analytical expression for the intensity

of the Raman signal, R(co) is then given by;

R((0) = A exp(-2az) exp

-(o>-0)z)^

2o2

dz

where (Oz = cOq + S(z) This expression

consists of an integral over all depths and

contains an exponential decay which

modifies the assumed Gaussian lineshape.

The Gaussian is located at angular

frequency coz and has an associated natural

line width la^. The amplitude factor. A, is included to scale the integral to the

measured Raman data. The depth

dependence, z, enters expression in the

exponentially decaying intensity factor and

in cOz- The term cOz consists of the location

of the unshifted Raman signal, cOq. plus the

shift induced due to the stress at each

incremental depth dz. This shift function,

S(z), will hereafter be referred to as the

stress profile and is the quantity of interest

in this investigation.

A depth, Zo, is chosen such that

beyond this depth the stress may

realistically be presumed to vanish. The

integral in the Raman function is naturally

divided into two regions, greater and less

than zq. Where the stress is zero (z > zq),

the Gaussian lineshape can be taken outside

the integral and the remaining integral can

easily be solved analytically. For 0 < z <

zq, the stress is changing and the integral can only be solved numerically. Numerical

integration involves dividing the interval

into dz elements and evaluating the function

at each incremental depth dz. Therefore the

stress function, S(z), must be evaluated at each dz regardless of its functional form.

The stress at each depth can be made a parameter that is free to vary. The function

to be fitted becomes

R((0)=^exp(-2azj)exp{-^"-^^|^^'^^>'}dz .

jA 1 J:(OMOO)^1 ^xp[-2azo]exp|—

—

|+bk

The first term is the integral over a depth of zero to zq of the assumed

Gaussian lineshape. The second term is the integral of the Gaussian lineshape from

a depth Zq to infinity and is constant for a given value of zq. An additional term, bk,

allows for a background correction. This

correction amounts to removal of a constant

number of dark counts from the Raman

spectra. The only parameters free to vary

are the S(zj)'s, the residual stresses at each

depth; thus the stress profile can assume

any form, even if it is discontinuous.

Least squares data analysis involves

minimization of the sum of the squares of

the differences between the experimental,

Yexp, and the calculated data, Yfit. This

sum of squares is referred to as j}, and

11

(yexp-^fit)^

where r|2 is the experimental uncertainty

associated with each data point. Since the

Raman data are collected in a photon

305
counting mode, the data follow Poisson

statistics and the uncertainty, r\ is equal to the reciprocal of the square root of the

number of counts (typically T] < 0.01 near

the Raman peak).(12) The interval from

zero to Zq is divided into an appropriate

number of intervals, each having width dz.

Hie Raman function is approximated by the

above sum, and the stress at each depth is varied until the calculated Raman signal

corresponds to the one obtained

experimentally by minimization of

is a function of the experimental

uncertainties and the parameters S(zj).

These parameters span a j dimensional

parameter space. There are a number of

standard methods for minimizing the

function x^, which involve searching

parameter space for the smallest value of

x2-(12) A grid search least squares

technique is employed which involves the

independent variation of each parameter,

S(zj), until is minimized. The next

parameter is adjusted, following a grid in

parameter space. The above procedure is repeated until the incremental change in x^,

that is AyJ^, is less than some specified

value (typically 1%). The disadvantage of

this method is that convergence is very

slow since each parameter is adjusted

independently while holding all the others

constant. It is also implicidy assumed that

the various x^'s are optimal, or near

optimal. In the initial stages of the fit (i.e. far from the minimum x^), this is not the case and this effect contributes to the slow

convergence of this method. The criterion

for a good fit is determined by minimizing

X^. If data are taken with green and blue

laser lines, it, is possible to define a X^blue

and X^green to measure the deviations

from the blue and green data respectively.

The simplest solution to this problem is to minimize the total x^ defined as the sum of

X^blue and X^green, or

X^tot = X^blue + X^green

•3.00e+8 1

I Iu

I IB —r—r""t—t—

0.0 0.1 U OJ 0.4 Oi 0.6

DEPTH(Hfli)

Figure 2. A depth profile of residual stress

in machined Ge. The depth profile was

obtained using 514.5 and 488.0 nm probe

wavelengths from the spectra shown in

Figure 1. The surface displays a zone of

compression with a counterbalancing region

of tension beneath. The abrupt change from

compressive to tensile stress is indicadve of

the plastic/elastic boundary.

A typical stress profile obtained ft^om

machined Ge using the technique described

above is shown in Figure 2. The single

crystal of Ge was prepared by single point

diamond turning the surface of the crystal

via a facing cut in the form of a flat region.

The facing cut was performed using a -10°

rake angle, 6° clearance angle diamond tool

with a maximum depth of cut of 240 |im.

The infeed rate utilized was 2.5 |im/rev,

with a spindle speed of 6.28 X 10^ rad/sec.

The depth resolution of this technique is determined by the distance between the

S(zj)'s. A value of 10 nm is typically used

for this spacing and changes in the residual

stress have been observed over this scale,

indicadng that this value is a good upperbound for the depth resolution in this

system. The maximum depth resolution

for this technique is determined by the

value of the absorption coefficient for the

light used, with higher absorption

coefficients giving greater depth resolution.

Since the stress is changing over the

volume sampled, it is inappropriate to

306
determine the magnitude of the stress from

the peak value of the Raman signal alone.

Initial depth profiles were generated using

only 514.5 and 488.0 nm wavelengths.

Though this procedure resulted in

reasonable depth profiles the convergence

time of this method was extremely slow.

In an effort to decrease the convergence

time as well as provide more reproducible

data the number of light wavelengths used

to probe the stress field was increased.

The use of 514.5, 496.5, 488.0 and 457.9

nm excitation wavelengths decreased the

convergence time of the deconvolution

routine tenfold. Stress profiles generated

using a greater number of wavelengths

usually display fewer signs of numerical

instability than do those obtained using the

514.5 and 488.0 nm wavelengths alone.

Residual Stresses Along Tool

Shoulders in Interrupted Test Cuts

The fabrication of optical quality or

precision surfaces in materials is of prime

interest in both the electronics and optics

industries. Classical techniques of

grinding and poHshing have allowed many

ductile materials such as metals to be

formed into figures with complex profiles.

Recent advances in materials fabrication

have used diamond crystals in either a grinding or turning mode for the

production of high quality surfaces in

materials such as ceramics and glasses as well as metals. In the process of single

point diamond turning a diamond crystal is used like a conventional lathe tool to form

surfaces of optical quality. The process of

single point diamond turning has also

allowed a number of otherwise brittle

materials such as semiconductors and

glasses to be machined in a ductile mode.

In such cases the diamond turning

technique can produce surfaces of optical

quality with complex figure profiles such

as aspheres (e.g off-axis parabolids) with

the introduction of little surface damage

into the workpiece.(lO)

The process of material removal via

the technique of single point diamond

turning in brittle materials, however,

introduces stresses into the workpiece that can in some instances lead to undesirable

effects such as figure deformation,

subsurface damage, or fracture. The

fabrication of an optical quality surface

with significant residual stress fields and/or

subsurface damage is, therefore,

counterproductive to the manufacture of a useful device. A non-destructive probe of

residual stress states and subsurface

damage in machinable brittle materials is thus indicated as an important tool for the

evaluation of the quality of such machined

surfaces. There exist a limited number of

suitable non-destructive methods for

characterizing the residual stress states in

such machined materials. The technique of

laser Raman spectroscopy has been proven

to be useful for the non-destructive

determination of stresses in thin films

where the region of stress is typically

contained in the film only. Brittle

semiconductors such as silicon and

germanium are easily measured using

Raman spectroscopy and the Raman

spectra are sensitive to the residual stress in

such samples. In addition a nondestructive probe of residual stresses may

be used to confirm the validity of various

theoretical mechanisms of the machining

action. (5, 6, 10)

The relation of the residual stress

field to the transition between ductile and

brittle fracture can play a significant role in understanding the physics of the material

removal process in the machining of brittle materials such as germanium. It is hypothesized that a region of high

compressive stress indicative of plastic

deformation should exist at the leading

edge of the diamond tool used to machine

brittle materials such as germanium. The

extent of this plastic deformadon is expected to be much greater in magnitude

than the stress field present in the finished

surface due to the excessive deformation

that must accompany material removal in

the process of single point diamond

turning. A further factor which can play an important role in the fabrication of smooth

surfaces is also the depth and magnitude of

the compensating zone of residual tension

occurring in the machined sample.

Although one might expect the stress field to display increasing surface tensile stress in regions of fracture, the results of this

investigation have proven that this is not

necessarily the case. Variations in the

307
magnitude and depth of the residual stress field might also expected for different

machine parameters. High negative rake

angle tools, for instance, have been found

to result in stress fields of greater

magnitude than tools with lower negative

rake angle.(13)

An important method for the

evaluation of the residual stress fields as well as the surface induced damage

occurring during the process of single point

diamond turning has been the formation of

interrupted test cuts. Interrupted test cuts

are made by rapidly removing the diamond

tool during the process of machining. The

rapid removal of the tool leaves behind a

profile of the tool in the surface of the

machined sample. The behavior of the

remaining machined surface may then be

studied as a function of position along the

resulting profile. Such profiles typically

display a range of behavior from purely

ductile removal at the base of the profile at the level of the finished surface to a region

of pure brittle fracture at the upper regions

of the profile where no material removal

has taken place. The profile remaining in

the sample surface is often referred to as a

tool shoulder and should not be confused

with the actual diamond tool used in the

machining process.

Depth profiles of the residual stress states as the probe spot is moved up the

tool shoulder (toward the uncut portion)

remaining in the machined sample after

rapid removal of the diamond tool in

interrupted test cuts were obtained using

four wavelengths (476.5, 488.0, 496.7,

514.5 nm) at each probe location. The

depth profiles of the residual stress

remaining in the sample were investigated

using the differential absorption profiling

technique that was previously discussed.

The depth of penetration of the

compressive stress field as well as the

depth of the plastic/elastic boundary has

also been investigated as a function of

position on the tool shoulder. The

penetration of this stress field can then be

related to the transformation from ductile to britde material removal in single point

diamond turning.

NO

TM^ 1 LASER PROBES (1-4)

IN UNCUT

REGION 2 I

UNFRACTURED

REGION

Figure 3. Probe positions for invesdgadon

of residual stress fields along interrupted

test cuts remaining in the Ge sample after

removal of the diamond tool. The laser

probe positions are indicated by the arrows. The reladve dimensions of the

depth of the tool shoulder to horizontal

position are not drawn to scale. In this case the diamond tool was moving from

right to left (cross-feed direction) before

being removed from the Ge workpiece.

The workpiece motion due to spindle

rotation is in a direction normal to the plane

of the page.

The investigation of residual stresses

in the interrupted test cuts remaining in the

machined germanium samples were

undertaken by locating the microscopic

probe spot at various positions along the

shoulder as shown in Figure 3. The probe

spot was located at an indicated position

and a micro-Raman spectrum was obtained

using the longest wavelength (514.5 nm)

probe. The laser was then tuned to the next

wavelength without moving the probe

posidon to e;isure that each wavelength

probed the same microscopic spot. The

laser probe position was then moved to the

next probe locadon where the process was

repeated. All spectroscopic data were

analyzed using the differennal absorption

profiling algorithm that has been previously

discussed. The depth profiles which were

308
0 0.2 0.4 0.6 0 0.2 0.4 0,(

DEPTH (|im)

Figure 4. Depth profiles generated for -45° rake angle tool with low cutting speed.

The profiles are shown in order of increasing distance from the uncut region. The

bottommost profile was thus obtained in the ductile cut region with no surface damage

evident. The displayed stresses can be considered to be accurate to within ± 10.0 MPa.

309
DEPTH ( ixm

)

Figure 5. Depth profiles generated for the -30° rake angle tool. The profiles again

are shown in order of increasing distance from the uncut region. Though similar to the case of the -45° rake angle tool several differences are evident in both the

magnitude and depth of the resulting stress field. The displayed stresses can again be

considered to be accurate to within ± 10.0 MPa.

310
generated from this data are reproducible to a high degree, i.e subsequent runs of the

deconvolution program with small changes

in background and amplitude parameters

resulted in insignificant changes in the

resulting depth profiles.

Several samples which were

machined using several different rake angle

tools (-45° and -30°, 2.5 |im/rev infeed

rate, 240 - 500 |im depth of cut, 6°

clearance angle) were investigated using

this technique. The germanium wafer

machined using the -45° rake angle tool

was probed in two test cuts of different

cutting speed (6.28 X 10^ rad/sec, 15.7 X

10^ rad/sec). The depth profiles resulting

from these cutting speeds are approximately equivalent. The depth

profiles generated for interrupted test cuts

remaining after cutting by a -45° rake angle

tool shown in Figure 4 show similarities to those generated for the -30° case shown in

Figure 5. In both cases a general region of

compressive stress is found at the sample

surface. In the ductile cutting regime

where there is little evidence of surface

fracture the magnitudes of the compressive

stress field for the -30° rake angle tool are » 4.0 X 10^ Pa. Significant differences in

the maximum magnitude of the surface

compressive field are evident as one moves

up the tool shoulder into the visibly

fractured region as shown in Figure 6. The

compressive surface stress increases in

magnitude until the average compressive

component has roughly doubled the value

found in the absence of fracture.

The maximum value of compressive

stress is seen to occur at approximately the

middle of the tool shoulder where the

frequency of pitdng and fracture is highest.

The compressive field then diminishes as the probe position is moved toward the

transition region where no cutting has taken

place. A significant difference between the

-30° and -45° rake angle tools is seen to be

the relative magnitudes of the compressive

surface stresses. In the fractured regions

the higher negative rake angle tool is seen

to produce a stress field of greater

magnitude than the lower negative rake

angle tool. In the absence of an applied

force the total integrated stress must vanish

in the surface regions. The existence of a surface compressive stress field thus

implies the existence of a counterbalancing

tensile field lying beneath the compressive

zone.(14, 15) The magnitude of the tensile

field is directly proportional to the

magnitude of the compressive surface

zone. The maximum magnitude of the

tensile field typically occurs at the

approximate depth of the boundary

between plastic and elasdc deformation as

illustrated in Figure 4. The magnitudes of

the tensile stress field are plotted as a funcdon of probe position in Figure 7 for

the three cases studied. The same general

trend as evidenced in Figure 6 for the

surface compressive stress is again present

in the deeper tensile field. Regions of high

fracture possess a high tensile stress field

which accompanies the greater surface

compressive stress.

FRACTVRED UNFTIACTURED

, ^ . REGION I REGION

4.00e-f i

|

I

I ,

I

I

I , 0 1 00 200 30 0 PROBE POSITION FROM UNCUT REGION ( Mm ) " -30 CBg, Rake "A- - 4 5 CBg Rake —« -45 CEG *KEHiaiSPffi)

Figure 6. Surface compressive stress as a funcdon of posidon from the uncut region.

The approximate transidon from ductile to britde material removal is indicated by the

vertical line. The compressive surface

stress is seen to reach a maximum in the

center of the fractured region. The

displayed stresses are accurate to within ±

0.1 X 108 Pa.

Fractured regions are further

characterized by the depth at which the

transition from compressive to tensile

stress occurs. In the regions which display

311
no fracture the transition from compressive

to tensile stress occurs at a depth of 0.2 |j,m

beneath the sample surface as indicated in

Figure 8. Fractured regions typically

display transitions that occur at shallower

depths of 0.1 \im. The depth of transition

is approximately constant for all probe

positions measured in fractured regions.

The transition from ductile cutting to fracture is thus characterized by two

distinct features. The first is the magnitude

of the stress fields (both compressive and

tensile), the second being the depth at which the transition from plastic

(compressive stress) to elastic (tensile

stress) deformation has occurred.

Figure 7. Maximum tensile stress as a function of probe position. The maximum

magnitude of the tensile stress filed

occurring beneath the compressive zone is plotted as a function of position from the

uncut region. The stress is seen to follow

the same general pattern of the compressive

stress shown Figure 4. The vertical line

indicates the transition form brittle to ductile material removal. The displayed

stresses are accurate to within ±0.1 X 10^

Pa.

Surface Failure from Crack

Nucleation at the Plastic/ESastic

Boundary

The behavior of the residual stress

field in the region of transition from ductile

to britde material removal suggests that the

formation of a smooth surface with litde or no evidence of surface fracture depends

directly upon two important considerations.

The first factor to be considered in the

evaluation is the relative magnitudes of the

compressive and tensile stress fields

encountered in the sample. It has been well

established by other researchers that

machining, grinding and polishing

processes give rise to a radially

compressive residual field with a region of

tangendal tension outside the plastic zone which surrounds the contact site.(15) Our

depth profiling investigations have

confirmed these results, and it is suggested

from the study presented. in this section that

if the magnitudes of the stress field reach a

certain cridcal value then fracture ensues. Such a conclusion is consistent with the

Griffith criterion for crack growth i.e when

the critical value of stress is reached a crack

will begin to propagate in an unstable manner. (16)

The depth at which the plasdc/elastic

boundary occurs is also demonstrated as an

important factor in the formation of surface

damage in a machined pait. The process of

plastic deformation, as discussed

previously, relies on the motion of

dislocations and defects through the crystal

structure. At the transition from plastic to elastic deformation it has been noted by

many researchers that the maximum values

of the tensile stress in the elastic field occur

at this transition. It has, in addition, long

been known that dislocation density

increases with increasing stress. The

presence of high stress gradients in the

vicinity of the plastic/elastic boundary thus

give rise to a high defect and dislocation

concentration in this region. (17, 18) As

previously discussed the fracture

thresholds for median and lateral cracks

(both subsurface varieties) are found to be

nucleation limited because of the lack of

preexisting flaws to serve as nucleation

sites. The presence of dislocations and

defects at the plastic/elastic boundary can then provide nucleation sites for

microcracking to occur in either the median

or lateral system. (19) The existence of

microcracks alone, though adding to the

degree of subsurface damage, does not in

itself constitute a sufficient condition for

312
the propagation of cracks into the finished

surface. The additional presence of a tensile stress field which lies at a

sufficiently shallow depth, however, may

serve to further open these microcracks and

provide a driving mechanism for their

unstable growth.(15)

0 100 200 PROBE POSITION FROM UNCUT REGION ( nm

P -MDegRake -A. -45D«gRak« 4SD«gH^3F>««d

Figure 8. Transition depth versus probe

position. The depth of transition from

compressive to tensile stress is plotted as a function of position from the uncut tool

shoulder edge. The vertical line again

divides regions of fracture and nonfracture. The depth resolution in this case

is ± 5 nm.

The investigations of residual stress fields occurring in interrupted test cuts

points to the following mechanism for the

transition from ductile to brittle material

removal in the process of single point

diamond turning. Ductile material removal

takes place due to the large plastic

deformations occurring in the material

which accommodates plastic flow. The

portion of material that is removed in this

fashion is concentrated in a small region in

the vicinity of the tool tip. The majority of

material is removed high on the tool

shoulder in a brittle fashion due to radial

cracking which is activated from

preexisting surface flaws in the material.

Since the region of this brittle material

removal is subsequently subjected to the

high plastic deformation of the tool tip the

regions of brittle fracture are removed

leaving a surface formed by plastic flow.

If the compressive field arising from the

plastic zone at the tool penetrates deeply

enough, the effect of the residual tensile

field as a crack opening force is counterbalanced. Microcracks may

nucleate from dislocations and defects at the plastic/elastic boundary but will not

derive sufficient driving force from the

tensile field for their subsequent growth

and propagation into the finished surface.

If however the zone of plastic deformation

penetrates to a more shallow depth the

underlying tensile field will serve to drive

the microcracks nucleated from the high

defect and dislocation concentration at the

plastic elastic boundary into the finished

surface.

This mechanism provides a

qualitative explanation for the improvement

of surface finish and allowed depths of cut

with increasing negative rake angle tools.

Tools of large negative rake angle result in

plastic deformation zones which penetrate

more deeply than those of tools of lower

rake angle. Since the resulting

plastic/elastic transition lies more deeply in

the sample than the transitions associated

with lower negative rake angle tools the

formation of microcracks lies more deeply

beneath the surface which increases the

ability of the surface to resist fracture

formation. This trend may only be

expected to apply for a range of negative

rake angle tools since the accompanying

stress magnitudes also increase with

increasing negative rake angle. If a

sufficiendy large tensile stress magnitude is produced in the surface a point will

eventually be reached at which the stress will result in unstable crack growth even

though the transition boundary is relatively

deep.

Summary

Micro-Raman spectroscopy has been

used to measure the residual stresses in

machined surfaces with a lateral resolution near the diffraction limit. We have

extended this technique to incorporate axial

resolution (i.e. depth into the sample) by

using several excitation wavelengths,

which, by virtue of their differing

penetration depths, allow determination of

313
stress profiles as a function of depth into

the sample. Greater depths in the material

may be probed by using longer

wavelengths of light thus leading to a superposition of a continuum of Raman

peaks. Depth profiles of residuals stress can then be obtained using computer

analysis of the resulting Raman spectra.

Surface fracture is characterized by

high residual stresses as well as a shallow

plastic/elastic transition depths. The high

residual stress fields typically occur with a layer of compression at the surface with an underlying tensile region beneath. The

transition depth from compressive to tensile stress is indicative of the transition from

plastic to elastic deformation. The presence

of high defect and dislocation

concentrations in the vicinity of the

plastic/elastic transition serve as nucleation

sites for microcracks of the lateral and/or

median system. When the plastic/elastic

transition lies at a sufficiently shallow

depth the residual tensile field will act to further open these microcracks resulting in

their propagation and growth into the

finished surface. If the stress fields are of

low magnitude the constraints on the

necessary penetration depth may be relaxed

somewhat, i.e low stress fields may allow

the plastic/elastic boundary to occur at shallower depths without surface failure

than would higher stress fields.

REFERENCES

1. S. Lyon, R. Nemanich, N. Johnson,

and D. Biegelsen, "Microstrain in laser

recrystallized silicon islands on fused

silica," Appl. Phys. Lett., 40, 316 (1982)

2. Th. Englert, G. Arbstreiter, and J. Pontcharra, "Determination of Existing

Stress in Silicon Films on Sapphire

Substrate using Raman Spectroscopy,"

Solid State Elect., 23, 31 (1980)

3. Y. Cheong, H. Marcus, and F. Adar,

"Raman microprobe measurements of

residual strains at the interfaces of Si on

quartz," J. Mater. Res., 2, 902 (1987)

4. S. Brueck, B. Tsaur, J. Fan, D.

Murphey, T. Deutsch, and D. Silversmith,

"Raman measurements of stress in siliconon-sapphire device structures," Appl. Phys.

Lett., 40, 895 (1982)

5. R. G. Sparks and M. A. Paesler,

"Micro-Raman analysis of stress in

machined silicon and germanium," Prec.

Eng. 10, 191 (1988)

6. R. G. Sparks, W. S. Enloe, and M.

A. Paesler, "Development of a high axial

resolution Micro-Raman Technique for

studying the effects of machine parameters

in machined semiconductors," Microbeam

Analysis 1989, 149 (1989)

7. W. Enloe, Masters Thesis, North

Carolina State University, unpublished

8. H. Shen and F. Pollack, "Raman

study of polish-induced surface strain in

<100> GaAs and InP," Appl. Phys. Lett.,

45, 692 (1984)

9. Z. Hang, H. Shen, and F. Pollack,

"Comprehensive investigation of polishinduced surface strain in <100> and <111>

GaAs and InP," J. Appl. Phys, 64, 3233

(1988)

10. R. G. Sparks, Ph. D Thesis, North

Carolina State University, unpublished

11. S. Sze, Physics of Semiconductor

Devices, pp. 40-44, (John Wiley & Sons,

New York, 1981)

12. P. R. Bevington, Data Reduction and

Error Analysis for the Physical Sciences , pp. 189-194, (McGraw-Hill, New York,

1969)

13. H. R. Philipp and E. A. Tafi,

"Optical Constants of Germanium in the

Region 1 to 10 eV," Phys. Rev., 113 (4),

1003 (1959)

14. J. I. Pankove, Optical Processes in

Semiconductors , pp. 34-87, (Dover, New

York, 1971)

15. D. Johnson-Walls, A. Evans, D.

Marshall and M. James, "Residual Stresses

in Machined Ceramic Surfaces," J. Am.

Ceram. Soc, 69 (1), 44 (1986)

314
16. S. Chaing, D. Marshall and A.

Evans, "The response of solids to elastic/plastic indentation, II. Fracture

Initiation,"!. Appl. Phys., 53 (1), 298

(1982)

17. C. Barrett, W. Nix and A. Tetelman,

The Principles of Engineering Materials, pp.

226-236, (Prentice-Hall Inc., Englewood

Cliffs, 1973)

18. F. Mclintock, A. Argon, Mechanical

Behavior of Materials, (Addison-Wesley,

Reading, 1966)

19. S. Chaing, D. Marshall and A.

Evans, "The response of solids to

elastic/plastic indentation, I. Stresses and

residual stresses," J. Appl. Phys., 53 (1),

312 (1982)

315

DESIGN OF ULTRAPRECISION GRINDING MACHINES FOR

DUCTILE MODE GRINDING OF BRITTLE MATERIALS

H. SHIMURA

Tokyo Metropolitan College of Aeronautical Engineering, Tokyo, Japan

F. HASHIMOTO

The Timken Co., Canton, OH

J. YOSHIOKA

NUTR K. K., Chiba, Japan

The field of ductile mode grinding of brittle materials using grinding wheels

with a grain size larger than the ductile-brittle transition value dc is defined as nano-gr inding technology. In the present paper, new design concepts

are introduced for a vertical spindle surface grinding machine, aiming at developing an ul traprecision and high stiffness machine tool for finishing of

silicon wafer with 8 inch diameter. This concept is also applied to develop an ultraprecision centerless grinding machine. Both ultraprecision grinding

machines have been built and evaluated through grinding tests. Effects of

application of these concepts on motion accuracies and stiffness are shown by

experimental results.

1. Introduction

(1) Technical problems in ductile mode

grinding

In the case of mass production,

ultraprecision grinding technology for

brittle materials is principles of motion

copying, ductile mode grinding with high

copying accuracy and crack-free machining

(grain size > dc), predictability of wheel wear in ductile mode, and programming of

grinding processes in brittle mode and

ductile mode for improving cost

performance.

(2) Development of ultraprecision grinding

machines and trueing cut grinding machines

The motion accuracies and system stiffness

required for the machine tools and the

trueing accuracies of grinding wheels for

achieving the criteria of the ductile mode

grinding are specified, and the following

trial targets can be derived using the measure : feed resolution : dc/10,

straightness of carriage : dc /travel range,

work and wheel support stiffness : no more deformation than the feed resolution

under grinding load,

vibration level of work and wheel support

system : dc,

trueing accuracies : < dc,

height distribution of cutting points : < dc. these items by M. Miyashi ta< i

2. The applied design concepts

2.1 Force compensated hydraulic linear

actuator

Two types of force compensated hydraulic

linear actuator units are applied.

(A) Positioning system of wheel head with

dual hydraulic linear actuators

Figure 1 shows the positioning servo system with dual force operated hydraulic

linear actuators'^) . in Figure 2, a pressure difference between two cavities of

the positioning actuator (Pi-Pa)

corresponding to difference between command

input and output of feedback displacement

transducer actuates the piston of the

positioning actuator to the place at the

command position. Grinding wheel head slide

is fed directly by a couple of actuators

without any lead screws or gears in order

to eliminate backlash. This system is one

of multi-points drive. xhe actuators

face each other and they have their shares

7
Input

couunand Servo

ampl i f ier

Servo

valve

liydraul ic actuator

for

load compensation

Hydraul ic actuator

for positioning fore

Wheel head

Output

posi tiofi ~f >

Displacement transducer ^

Figure 1 Block diagram of positioning servo system with dual hydaraulic linear actuators

Displ acement ]-'-'-

t ransducer

N

Hydrostatic bearing

Figure 2 Hydraul ic

1 inear actuator

Pully Cringing whael Hydrostatic Hydrostatic

1 Thrust Bearing Bearing Racesa Hydrosta tic Dearlng Recaas Wheel Unit

,

1 I 1 I

• Fixed

Spl nd l6

Whe«X Haad

Wheel Unit

rully

MydroataCic

Thruoc Bearing

Grinding Wheel

Fixed

Spwidli

I mi

m

3—i c

Rotary Table

(a) Construction of both

end fixed spindle

(b) Construction of one

end fixed spindle

Figure 4 Construction of wheel unit on stationary spindle

Figure 3 Hyraulic actuator with

L. C. unit — Nanoact —

( Designed by F.Hashimoto ) 318
in load of the wheel head. One of them is

for positioning and the other is for load

compensation in order to reduce load of

the positioning actuator. A pressure

difference between two cavities of the

positioning actuator (pi-pa) is measured

with pressure sensors, and its output

signal is input to servo amplifier for the

load compensation. A pressure difference

(Pi'-Pa') is supplied to the actuator for

load compensation. Load share rate Kg is

ratio of load shared by the actuator for

load compensation to load of the

positioning actuator. The actuator for load

compensation shares load of Ks/(1+Ks) and

load of the positioning actuator is reduced

to l/d+Ks). It is expected high

resolution and high repeatability in

positioning under heavy friction load. This

system is applied to the surface grinding

machine.

(B) Nanometer positioning actuator with

load compensation unit

Figure 3 shows the nanometer

positioning actuator with load compensation

unit referred to as Nanoact. It has two

pistons. The inner piston is for

positioning and the outer piston is for

load compensation. They are coaxial, and

the piston for positioning is built in

the piston for load compensation.

Therefore, Nanoact can be be very compact.

Both pistons are supported by hydrostatic

bearing. They are connected to the

connecting plate at each front end. A pressure difference between two cavities of

the inner positioning actuator (pi-ps) is

picked up with pressure sensors and its

output signal is input to the servo

amplifier for load compensation. A pressure

difference (pr-p2') equal to a pressure

difference of the positioning actuator (pi

P2) is supplied to the outer actuator for

load compensation. The effective sectional

area of outer piston is K5 times as wide as one of inner piston. Ks is 2.4 in Figure 3. The piston for load compensation shares

load of Ks/(1+Rs)- Load of positioning

piston is reduced to 1/(1+Ks). High

resolution of positioning is expected.

Nanoact is used for the ul traprecision

centerless grinding machine.

i 2.2 Stationary spindle structure

Figure 4 shows a construction of

stationary spindle structure. <3) One or

both ends of the spindle is fixed on the

housing. Stationary spindle structure is a closed force flow system and gives high

structural stiffness.

2.3 Composite bearings guideway

In Figure 5, a composite bearings

guideway is composed of a couple of plain

bearing guideway and hydrostaic bearing

guideway.''" It enables to build a uniform

distribution of oil film thickness and to eliminate any play between the guideways.

It is expected to improve dynamic stiffness

and motion accuracies.

2.4 Hydrostatic support of wheel unit

A wheel unit is hydros tat ical ly

supported on a stationary spindle, and a high accuracy and repeatability of runout are expected.

3. Design of ul traprecision vertical

spindle surface grinding machine

The ul traprecision vertical spindle type

surface grinding machine for ductile mode

grinding of brittle materials has the

following characteristics:

1) positioning system of grinding wheel

head with dual force operated linear

actuators as shown Figure 1, 2) stationary spindle structure, which is

fixed at one end of spindle on wheel head,

3) composite bearings guideway is applied to the guideway between column and

wheel head slide. Also it is applied to the

guideway between X slide and bed,

4) grinding wheel unit hydrostat ical ly

supported on the stationary spindle.

Figure 6 shows the schematic of

ul traprecision vertical spindle rotary

table type surface grinding machine. In Figure 6, the hydraulic actuator for

positioning is fixed upon the bottom of

the column, and the actuator for load

compensation is fixed on the top of the

column in order to reduce load of the

positioning actuator. Each piston rod is connected directly with wheel head without

any lead screws or gears. Displacement

transducer for position feedback is fixed

on the column. Wheel head is fed by a tandem drive system. Wheel unit which is

supported hydrostat ical ly on the stationary

spindle is driven with a belt. The rotary

table is supported on the X slide with

hydrostatic bearings.
Plane bearing guideway

One end fixed spindle

Grinding

wheel unit

Grinding wheel Rotary table

X slider

Hydraulic actuator

for load compensation

, Columny Safety device

Hydrostatic bearing

gTjideway

Grinding wheel head

Belt

Hydraulic actuator

for positioning

' ] I I I I I

^1 I I I I

D. Q

llfnd - Rolnry lablc • Sntnrj tnlilc A I nil I tiR licnil

5U0 750 iOOO

Load (N)

Figure 7 Stiffness record of

the surface grinding machine

Figure 6 Scheme of the vertical spindle

^ Displacement:

surface grinding machine Upwards /

r 10 s ^ M

Figure 8 Record of wheel head

displacement on the surface grinding

machine for stepwise feed of 10 nm

Figure 9 Kinematic drive of grinding head slide

320
Figure 7 shows the stiffness record

between the wheel head and the rotary table

of the surface grinding machine for loads

up to 1000 N. The stiffness at a radius of

100 mm from the center of spindle is higher

than 160 N/^m. Figure 8 is the record of

grinding wheel head displacement for

stepwise feed of 10 nm. It shows the

positioning resolution of 10 nm and little

generation of lost motion at reversing

point of feed direction.

4. Design of ultraprecision and high

stiffness centerless grinding machine

The ultraprecision and high stiffness

centerless grinding machine for ductile

mode grinding of brittle materials was designed based on the fol lowing*^)

(1) Guidelines of design

1) High stiffness and high precision

2) Stability of thermal deformation

characterist ics

3) Minimum outside diameter of ground

workpiece 0. 1 mm

4) Full automatic centerless grinding

machine with auto-setup function

(2) Applied new technology

1) Use of low thermal expansion materials

and a symmetric structure

2) Kinematic drive of grinding head slide

(shown in Figure 9)

3) Hydraulic actuator with load

compensation unit — Nanoact —

4) Composite bearings guideway

5) Fixed spindle structure

6) Controlled 2-axes work support blade

with auto- lock mechanism.

Table 1 shows the specification of the

centerless grinding machine.

5. Motion accuracy and stiffness of the

centerless grinding machine

The experimental test results on the motion

accuracies and stiffness are summarized as

fol lows

(1) Performances of the positioning servo

cntrol system of the force operated

hydraulic linear actuator (Nanoact)

Nanoact is applied to feed the

grinding wheel head slide. The total

stiffness is about 180 N/^m as shown in

Figure 10. The resolution of positioning

servo control system feeding the grinding

wheel head slide is 8 nm/step under

Coulomb's friction of 784 N, as shown with

experimental result of Figure 11. (2) Trueing accuracies and remountabi 1 i ty

The diamond wheel was trued on a specialized grinding machine equipped with

the hydrostat ical iy supported wheel on the

stationary spindle for trueing cut to an accuracy of 100 nm runout and was remounted

on the centerless grinding machine to the same accuracy.

6. Grinding test results

Through-feed centerless grinding tests were carried with diamond wheels on cylindrical

test pieces of ZrOo, BK7, SIALON and

HP-Si3N4 shown in Table 2. Typical records

of the roundness and the surface roughness

are shown in Figure 12. Size of the test

pieces, grain size of diamond wheels, and

the resulted grinding accuracies of

roundness and surface finish are listed on

the Table 3. Figure 13 shows that ductile

mode grinding is superior to lapping in

shape generation. Surface grinding tests

for silicon wafers were carried out by the

newly designed surface grinding machine and

conventional machines. Damaged layer on

the test pieces ground by the designed

machine was thinner than one by

conventional machines. <6) The flatness

of ground 6 inch silicon wafer was less

than TTV 0.6 jim.

7. Summary

(1) The surface grinding machine

The designed machine has feed

resolution of the order of 10 nm. The loop

stiffness more than 160 N/jum were obtained.

Grinding tests of silicon wafers revealed

that

1) the subsurface damage was in the order of

a half micron.

2) The flatness of a 6 inch silicon wafer

was less than TTV 0.6 iim. (2) The centerless grinding machine

The designed machine has feed

resolution of 8 nm. The total stiffness is

about 180 N/fim. Through-feed centerless

grinding tests were carried with diamond

wheels for cylindrical test pieces of IrOo,

BK7, SIALON and HPSN. The roughness for

ground test pieces of ZrOs was 10 nmRa,

and the roundness was 10 nm.

321
ioo 1000 ijoo

I- u a d f (J r c n I" ( N )

Base

HcgulaLiiig whcol licacj

Oil hydraulic aclualor Grinding »liccl licad

(a) Repeatal)ilily of positioning for

Stroke of 8 /im

-J J 1 On m

1 Os ec

(b) Positioning for stepwise

feed of 8 n ni

Figure 10 Total stiffness of the

centerless grinding machine

Figure 11 Positioning characteristics

of the slide table driven by Nanoact

Table 1 Design specification of the centerless

grinding machine

Table 2 Vickers hardness

of the test pieces

Name NANOTRON

Dimensions 1 085^ X

460^"' X 450 "

Ma ter ia 1

Ultra-Low Thermal Expansion Castiron

Grinding Wheel

D W Diamond Wheel or CBN Wheel 150 x50

Regulating V)heel

D W Ceramics or Steel Wheel 150 x50

Wheel Support Hydrostatic on Fixed Spindle

Work Support

Blade

2-Axes Controled Work Support Blade

with Auto-Lock Mechanism

Workpiece Brittle Materials, 40. 1"^ (J)20mm Machine Stiffness Total Stiffness of Grinding System

about 160 N/ um Slide Feed Resolution of Positioning 0.001 urn provided with NANOACT

Material Hardness

Zr02

BK7 (glass)

SIALON

HP-SN

HV 1200

HV 700

HV 1900

HV 1700

Table 3 Grinding accuracies

Material Size Roundness Rougliness

Ir02 5* 2.5 X loL 10 nm 10 nm Ka 0K7 ^ 0 X llL 34 nm 15 nm Ra SIALON s* 4 . 1 X 29l- 15 nm 28 nm Rs HPSH 9i 10 X lOL 2\ nm 55 nm Ra

Zr02,BK7 : 3000 resined bonded wheel

SIALON: D4/8 vitrifieded bonded wheel

HPSN : 800 vitrifieded bonded wheel

322
s. ly m FILTER IIIIIDOll 5. 0' FILTER

15 LIPR

nnx 0,003 ,>.^»

«vi: o.CiOl O 0.027 VIZ 0.090

/. 24 0

»

/

v

V.

0. u...

\

\

O.OOl AVC 0 . 000 - O 0.010 \

li 0. 0?0

/

£— 2-t O" / Z' 0 . I £:«.

J

0. !>... /

Roughness of llnmRa

:a) No filter

Roughness of lOnmRa

(b) Filter 15 UPR

Figure 12 Typical records of the roundness for test

piece of ZrOo ( size: 02.5*10.51- )

UIHOOU

s. e- ^\

SL'JPE (MX o.om

ftVO 0.001 O

E

• 0.0*1

_ Z

0. i;.c

O.lv

FILTER

13 UPR

SB

(a) Grinding

SB

(b) Lapping

Figure 13 Roundness records for grinding

and lapping ( material: ZrOo )

Reference

F. Hashimoto, and M. Miyashita,

"Ul traprecision Grinding — Construction of Positioning System of Wheel

Head — " , Autumn Meeting of JSPE, pp. 55-56, 1986. ( in Japanese)

J.Yoshioka, F. Hashimoto, M. Miyashita,

A. Ranai, T. Abo, and M. Dai to, "Ul traprecision Grinding Technology for

Brittle Materials: Application to Surface and Centerless Grinding

Processes", M. C. Show Grinding Symposium

ASME, PED-Vol. 16, pp. 209-227, 1985.

J. Yoshioka, K. Koizumi, M. Shimizu,

H.Yoshikawa, M. Miyashita, andA.Kanai,

"Surface Grinding with a Newly

Developed Ultra Precision Grinding

Machine", SME Mfg. Eng. Trans. , pp. 18-

28, 1983.

K. Abe, N. Yasunaga, M. Miyashita,

J.Yoshioka, and M. Dai to, "Study on

Ul traprecision Grinding Technology for

Brittle Materials (1)", Autumn Meeting

of JSPE , pp. 551-552, 1992. (in Japanese)

J. Yoshioka, M. Miyashita, A. Ranai,

M.Dai to, and F. Hashimoto, "Duct i le Mode

Grinding of Brittle Materials with

Newly Designed Centerless Grinding

Mach i ne"Proc. of ASPE Annual Conf

, pp. 143-144, 1990.

1 M. Miyashita, "The Way to Nanogrinding

Technology", SP IE Proc. , Vol. 1333, p. 13,

1990.

2 H. Shimura, J. Yoshioka, M. Takahashi,

323
Il

I
PRECISION FORM TRUING AND DRESSING FOR

ASPHERIC CERAMICS MIRROR GRINDING

T. KURIYAGAWA, K. SYOJI, and L. ZHOU

Tohoku University, Sendai, Japan

The goal of this paper is the development of a precision truing and dressing device for an extremely fine-grit round-profile diamond wheel. A new and compact truer is designed

and built to achieve the desired profile without detachment of the diamond wheel from the

grinding machine. This "Arc Truer" has a rotation unit consisting of a silicon carbide cup

wheel, and a swing/infeed mechanism for the cup wheel. The truer has the advantage that dressing after truing is not necessary, and that sharp cutting edges protrude from the wheel

surface. The profiled wheel is for ductile-regime grinding of aspheric ceramic mirrors. As

the wheel is wider than conventional ones for form generation grinding, it is expected that the wheel wear and grinding characteristics will be uniform widthwise. In this paper, a resin bonded diamond wheel is used, and truing tests are carried out with a CNC ultra-precision

surface grinding machine. A new grinding method for form generation is proposed and truing

characteristics and accuracy of the Arc Truer are investigated.

1. Introduction

Hot isostatic processed silicon carbide (HIPSiC), which has high heat conductivity and a small

coefficient of linear expansion, is the material

normally used for constructing mirrors for reflecting

short wavelength radiation (e.g. X-ray, etc.). Advanced technologies for both mirror surface

finishing and aspheric machining are demanded in the field of mirror grinding.

Many trials of mirror grinding or ductileregime grinding have been made. Analyses of

ductile behavior of brittle materials, and development of CNC ultra-precision grinding machines and

diamond wheels have been useful in this regard

^^°\ Recently, several attempts of grinding of

aspheric ceramics mirrors have been made

In mirror grinding, wheel life is usually short, and

it is difficult to keep a grinding characteristic of the wheel constant. Furthermore, truing and dressing

are the most important processes which affect the wheel performance. However, there are few studies

of precision truing and dressing for mirror grinding.

This paper proposes a new grinding method

for aspheric mirrors that uses an extremely fine grit

diamond wheel with a round profile. A precision

truing and dressing device, called the Arc Truer, is developed to obtain a wheel having round crosssectional profile on a surface grinding machine.

The truer demonstrated excellent truing characteristics.

2. Principle of Form Generation with a Round Profile Wheel

In the field of three-dimensional ultra- precision mirror grinding, form generation of a workpiece is carried out with a 4- or 5- axis

controlled grinding machine as shown in Fig. 1. With this grinding method, form accuracy of a ground workpiece depends on the total accuracy of

the grinding machine. If an ideal precision grinding

machine can be developed, three-dimensional ultra- precision machining can be realized in which part

accuracy is proportional to the machine accuracy. When grinding a mirror of a difficult-to-grind

material such as silicon carbide, loading and dulling

of a wheel can easily occur, decreasing both the machining accuracy and the wheel life, since the

325
Swi ng

Whee

Fig. 1

Workp i ece

Conventional method of form generation

same part of the diamond wheel is used in grinding.

In this paper, a new method for form

grinding is proposed by means of a diamond wheel

with a curved profile. Since the workpiece profile

generated by this method is an envelope of a family

of arcs as shown in Fig. 2, any convex or concave

profile can be produced. As the grinding point is changed along the wheel profile with every grinding

pass, the entire width of the wheel can be effectively used and wheel life is expected to increase.

In this method, form accuracy of a ground workpiece depends on the accuracy of not only the grinding machine but also the wheel profile.

Therefore, a device capable of precisely truing an

arc profile on a wheel is required.

3. Wheel Design and Accuracy of

Ground Surface

As shown in Fig. 3, the xyz coordinate

system axis is perpendicular to the plane of

paper and in the direction of the table feed) is fixed on the table of the grinding machine. The workpiece is symmetric with respect to the xz -plane (y = 0), and h is the width. The cross section of the workpiece is expressed by the equation z = f (y), where (y, z) = (0, 0). B is the wheel width and R is the radius of

the arc cross section. When grinding a convex

workpiece with the wheel, the following restrictions apply between R and B: 2R ^ B

0 < < 00

,(1)

(2)

(3)

Wheel

Workp i ece

Fig. 2 New method of form generation

z=f(y)

Wo rkp 1 ece

Fig. 3

where

Arc profiled wheel and workpiece

- arcsin|^j, i/r^ - arctan f(b

When grinding a concave workpiece, one

further restriction must be added:

R ^ pI r. (4)

where p^^^,, is the minimum value for the radius of

curvature f(y)

The accuracy with which a workpiece can be

ground by a wheel that satisfies the above conditions was then considered. This accuracy can be

expressed in terms of the profile error, surface

roughness and waviness of the workpiece. The

profile error and surface roughness depend on the grinding machine used and the grit size of the

326
E

ino

0.5 mm

Fig. 4 Waviness generated on workpiece

(/?=40mm, 5=20mm, r=100mm, i)=50mm)

wheel. For example, in our surface plunge grinding

tests, we achieved a surface roughness of == 0.25 jxm {R^ = 25 nm) in the case of SD1500

diamond wheel, and R^^ = 0.05 ^im {R^ = 5 nm)

in the case of SD3000 diamond wheel. Therefore, we will discuss the waviness on the yz plane only.

The following assumptions were used: The wheel

profile was ideally trued into an arc and the profile was transcribed to the workpiece during grinding.

The cross feed. Ay, is given in such a manner that the density of grinding points along the wheel

profile is uniform, namely the increment of At|) in the center angle i|) is constant. For simplicity, the shape of the workpiece is limited to a cylinder with

radius r and width b as shown in Fig. 3. In this case, condition (3) is R <. rBlb. Profile generated on the yz cross section of the workpiece (r = 100

mm, b = 5Q mm) can be calculated as shown in Fig. 4. Peak-to-valley height of the profile is defined as the waviness, Figure 5 shows the influence of Ai|) and R on the value of E^.. It is clear that the value of E^^ decreases as R increases

or as Ai|j becomes smaller.

In designing a wheel, R is taken to be as large as possible in the range of conditions (1)

through (4) keeping in mind the general consideration of grinding machine table size and wheel

manufacture.

4. Arc Truer

An Arc Truer for a cylindrical grinding

machine was developed and its truing accuracy

investigated ^^'^\ The stiffness of the truer, the infeed accuracy of the green silicon carbide (GC)

cup wheel, and the control method all exert a great

influence on truing accuracy. The Arc Truer in this experiment is an improved one, and trued an extremely fine grit diamond wheel precisely.

The principle of arc truing is shown in Fig.

200 Number of divisions300

0.125 0.1 0.075 Alp rati

0.05

Fig. 5 Relationship between waviness E^ and

increment Aip

Diamond wheel

GC cup wheel

(a) Truing principle

@ Swing mechanism

® Infeed mechanism

® Rotation unit for GC cup wheel

(b) Schematic drawing of Arc Truer

Fig. 6 Arc Truer

6(a). The rotating GC cup wheel reciprocates in a direction vertical to the plane of the paper and

327
(a) Arc Truer and its controller

(b) Swing during truing

Fig. 7 Photographs of the Arc Truer

swings around the center O. Fig. 6(b) shows a schematic drawing of the Arc Truer. It is composed of a rotation unit for a GC cup wheel ®, an

infeed mechanism @ and a swing mechanism ®.

The arc profile of the wheel is trued in combination

with the above three motions and the table reciprocation of the grinding machine ®. The radius, R,

of the trued profile can be changed from 10 mm to 50 mm. Fig. 7(a) is a photograph of the Arc Truer

and its control unit. Table 1 shows its specifications. Fig. 7(b) is a photograph of the Arc Truer

truing a wheel on the grinding machine. Fine

adjustment of the relative position of the Arc Truer

and the diamond wheel is needed during set up on 328

Table 1 Specifications of the Arc Truer

® Rotation unit for GC cup wheel

• GC cup wheel

a Rotfltionfll speed

• Peripheral speed

• Camming of wheel surface

» Perpendicularity of wheel axis and swing axis

• Intersection of wheel axis and swing axis

6A2, 90° X 10"^ 60~600 rpm 17~170 m/min ((J)90) 0.005 mm / <^90

0.002 mm / 90 mm

0.01 mm

® infeed mechanism

» Drive

• Stroke

• Distance between wheel surface and swinp nxi<:

• Minimum infeed

» Feed speed

• Straightness of infeed

• Perpendicularity of infeed and swing axis

• Thrust stif&iess

Stepping motor (direct) 40 mm

10-50 mm

0.001 mm

150 mm/min

0.003 mm / 40 mm

0.005 mm / 40 mm

10 N/|xm

(D Swing mechanism

• Drive

• Maximum swing angle

• Minimum swing angle

• Swing speed

• Parallelism of swing axis and

grinding machine table

Stepping motor + Worm

±60 deg

0.0072 deg/step 0.2~1.2 rpm

0.01 mm / 40 mm

50 [im

Fig. 8 Typical stereo SEM photograph

(Truing conditions are shown in Table 2)

the table. In this experiment, the adjustment is made using a laser beam with an axis that coincides

with the swing axis. Since truing and dressing occur simultaneously during truing with a GC cup

wheel, the process offers a significant advantage

that dressing is not necessary after truing.

5. Truing Tests

5,1 Grain Density

Truing tests for a straight wheel were carried out to determine optimal conditions. An extremely

fine grit diamond wheel (SD1500R125BRC lAl,

diameter: 205 mm, width: 15 mm, diamond layer

depth: 5 mm) and an ultra-precision surface
2000

1500

1000 s

500 Wheel: SD1500,F„=1800 m/min

Truer: GC1000,K,=150 m/min

Depth of cut: 2 pim Table feed rate: 0.1 m/min

1 10 100 Flow rate of truing fluid

1000 10000 mL/min

Fig. 9 Relationship between flow rate of

truing fluid and grain density

2000

B 1500

1000 c

-ac

'Su

500

Wheel: SD 1 500, K„= 1800 m/min

Truer speed: 150 m/min

Depth of cut: 2 [im Table feed rate: 0.1 m/min

Truing fluid: 5 mL/min

# 2000 #4000 # 6000 Mesh number of GC cup wheel#800)

Fig. 11 Relationship between mesh number of

GC cup wheel and grain density

2000

s

S 1500

^ 1000 ea

c 500

'«ua

Wheel: SD1500,K„=1800 m/min

Truer: GC1000,K,=150 m/min

Depth of cut: 2 pim Truing fluid: 5 mL/min

0.1 0.5 1 5 10 Table feed rate m/min

Fig. 10 Relationship between table feed rate and grain density

grinding machine were used. Synthetic fluid was

sprayed on a GC cup wheel as truing fluid. The

wheel was prepared so that a piece could be

partially removed and its working surface was

observed by stereo SEM. Figure 8 shows a typical stereo SEM photograph of the wheel surface just

after truing. In this photograph, a number of

diamond grains can be seen protruding from the wheel surface. At the same time, holes from which

diamond grains were released can be observed.

Grain density on the wheel surface was measured

and used as an estimate of truing performance.

Fig. 9 shows the influence on grain density of

various truing fluid flow rates. In the case of an extremely small flow rate, loading of loose GC

abrasive occurred not only into the GC cup wheel,

but also into the diamond wheel and many diamond

grains were pulled off, thus lowering grain density.

At a flow rate of 5 mL/min or more, however,

2000

ES 1500

1000 c

C3uo

500

Wheel: SD1500,F,^=1800 m/min

Truer: GC6000,K,= 150 m/min

Table feed rate: 0.1 m/min

Truing fluid: 5 mL/min

Depth of cut ^m

Fig. 12 Relationship between depth of cut and

grain density

loading did not occur and grain density was

increased. Grain density was almost constant throughout this range.

Figs. 10 and 11 show the influence of table feed rate and grain size of the GC cup wheel on

grain density. It is clear from the figures that slower table feed rate and finer grain size result in higher grain density. Fig. 12 shows the influence

of depth of cut on truing, A,. Grain density was

almost constant in the range of A, < 2 [im. In the case of A, = 0.25 (im, however, the volume of

loose GC abrasive was too small to true the wheel

efficiently. In contrast, at A, > 2 [xm, chatter marks

occurred on the wheel periphery. In light of the above results, optimal truing conditions were as shown in Table 2.

5.2 Truing Accuracy

Arc truing tests were carried out by an Arc

329
Table 2 Optimal truing conditions

Diamond wheel peripheral speed 1800 m/min

GC cup wheel speed 600 rpm

Flow rate of truing fluid 5 mL/min

Table feed rate 0.1 mm/min

Grain size of GC cup wheel #6000

Depth of cut 1 [xm

().5mm

5mm

Fig. 13 Cross-sectional profile of trued wheel

Truer and a CNC ultra-precision surface grinding

machine. In this truing test, an arc profile {R =

30.700 mm) was generated from a straight profile.

The tested wheel was a SD1500R125B (diameter:

305 mm, width: 15 mm, diamond layer depth: 5

mm) and the truing conditions in Table 2 were

applied. First, rough truing was performed using

GC cup wheels whose mesh numbers are #400,

#800, and #1000, in that order. Finally, finish truing was performed using a #1500, #3000 or #6000 GC cup wheel.

Fig. 13 shows the arc cross sectional profile

of the trued wheel. Profile measurement was

carried out through a transcription method. Fig.

14(a) shows the parameters used to evaluate truing accuracy. The difference between the value of the expected radius and the value of the approximated

radius of the measured cross section is defined as the radius error, E^. The deviation from the approximated arc is defined as the profile error, E^,

and the difference between the maximum and

minimum values of E^ is defines as the maximum

profile error, E^^.

Fig. 14(b) shows a typical trace of the profile

error, E^, along the wheel profile. It is clear that the truing accuracy widthwise is within ±0.4 \im

(£^„ = 0.8 fxm). Fig. 15 shows the influence of

depth of cut A, and mesh number of GC cup wheel

on the maximum profile error, E^^, in finish truing.

In the case of A, = 2 \im, finer grain size (larger

mesh number) resulted in smaller value of E^^. In the case of A, = 1 fxm, however, truing with

GC6000 cup wheel was a little unstable, because it

(a) Profile error and radius error Er

2 mm

(SD1500, V„=450 m/min, GC1500, V, =150 m/min,

A( = l fim, D,=0.1 m/min)

(b) Profile error

Fig. 14 Truing accuracy

!<J 3

E 1 s

1 1 Wheel: SD1500,K^=450 m/min

Truer speed: 150 m/min

Table feed rate: 0.1 m/min

Swing speed: 0.6 rad/min

At=2 \im

' ' '

' #2000 #4000 # 6000 Mesfi number of GC cup wheel

Fig. 15 Influence of depth of cut and mesh

number of GC cup wheel on maximum

profile error

was difficult to supply truing fluid uniformly on the swinging GC cup wheel. Therefore, value of

GC6000 cup wheel became larger than that of

GC1500.

Measured values of R were in the range of

30.627 mm to 30.711 mm. As the set up value of

R was 30.700 mm, an of less than ±0.08 mm

could be achieved. In the actual grinding of an

aspheric mirror surface, it is necessary to compensate for the wheel depth of cut by inputting data

regarding the measured wheel profile to the CNC

controller.
6. Conclusion

A new grinding method was proposed for aspheric mirrors that uses an extremely fine-grit

round-profile diamond wheel. A precision truing

and dressing device called an "Arc Truer" was also developed to obtain a curved profile wheel cross

section. An examination of truing accuracy of the Arc Truer revealed that profile and radius errors were within ±0.4 |im and ±0.08 mm, respectively.

References

11. H.Suzuki, "Development of CNC Ultra Precision Aspheric Grinder and 3-Dimensional High Precise Machining of Brittle Materials," J. of the JSPE, Vol. 56, No. 5, pp.815-819, 1990.

12. T.Osaka, K.Unno, A.Tsuboi, Y.Maeda, and

K.Takeuchi, "Development of High-Precision

Aspheric Grinding/Turning Machine," Proc. of the 6th IPES, Braunschweig, pp.95-102, 1991.

13. Z.Zhong, and T.Nakagawa, "New Grinding

Methods for Aspheric Mirrors whh Large Curvature Radii," Annals of the CIRP, Vol. 41, No. 1, pp.335-338, 1992.

14. L.Zhou, K.Syoji, and T.Kuriyagawa, "Truing of Super-abrasive Wheel for Foun-grinding," Proc. of 7th IPES., in press, 1993.

L T.G.Bifano, T.A.Dow, and R.O.Scattergood,

"Ductile-Regime Grinding: A New Technology for Machining Brittle Materials," Trans, of the ASME,

Vol. 113, May, pp.184-189, 1991.

2. T.ADow, and R.O.Scattergood, "Ductile/Brittle Transition and Development of Ductile Mode

Grinding Technology," J. of the JSPE, Vol. 56, No. 5, pp.794-798, 1990.

3. T.G.Bifano, and S.C.Fawcett, "Specific Grinding

Energy as an In-Process Control Variable for Ductile-Regime Grinding," Precision Engineering,

Vol. 13, No. 4, pp.256-262, 1991.

4. M.Miyashita, "Ductile mode Grinding Technology

for Brittle Materials: The Way to Nanogrinding

Technology," J. of the JSPE, Vol. 56, No. 5, pp.782-787, 1990.

5. W.S.Blackley, and R.O.Scattergood, "DuctileRegime Machining Model for Diamond Turning of

Brittle Materials," Precision Engineering, Vol. 13, No. 2, pp.95-103, 1991.

6. K.E.Puttick, and A.Franks, "The Physics of

Ductile-Brittle Machining Transitions: Single-Point

Theory and Experiment," J. of the JSPE, Vol. 56, No. 5, pp.788-792, 1990.

7. N.J.Brown, "Lapping: Polishing and Shear Mode

Grinding," J. of the JSPE, Vol. 56, No. 5, pp.800-804, 1990.

8. P.AMcKeown, K.Carlisle, P.Shore, and R.F.

J. Read, "Ultra-Precision, High Stiffness CNC

Grinding Machine for Ductile Mode Grinding of

Brittle Materials," J. of the JSPE, Vol. 56, No. 5, pp.806-813, 1990.

9. K.Carlisle, and P.Shore, "Experience in the Development of Ultra Stiff CNC Aspheric

Generating Machine Tools for Ductile Regime

Grinding of Brittle Materials," Proc. of the 6th IPES, Braunschweig, pp.85-94, 1991.

10. H.Ohmori, I.Takahashi, and T.Nakagawa, "Minor

Surface Grinding by Metal Bonded SuperAbrasive WTieel with Electrolytic In-Process

Dressing," Proc. of the 6th IPES, Braunschweig,

pp.153-165, 1991.

331

ELECTROCHEMICAL DRESSING OF BRONZE BONDED

DIAMOND GRINDING WHEELS

E. WELCH, Y. YI, and T. BIFANO

Boston University, Boston, MA

An experimental apparatus was constructed to study electrolytic dressing on bronze-bonded diamond grinding wheels. An EDM power supply was used to drive the electrolysis. Experiments were performed to explore the qualitative

microstructural features of the dressing, and to evaluate the quantitative rates of oxide film formation and wheel erosion. For one choice of the relevant pro- cess parameters (open circuit electrode voltage, electrode gap size, electrolyte,

and pulsed voltage duty cycle) it was found that an oxide film extended both out

of and into the wheel surface. The rate of growth into the wheel of the electrolytically affected layer was found to be approximately constant (1.5 nm/s).

Other combinations of the relevant electrolysis parameters are currently un- der investigation.

INTRODUCTION

It has been demonstrated that ceramics

can be fabricated using fixed-abrasive

grinding without generating sub-surface

damage (1-3). The absence of damage, a requirement in many structural and optical applications, demands sufficient control of the machine tool that a critical load

(or depth of cut) is not exceeded by the di- amond grains in the grinding wheel.

A problem encountered in such fine

grinding is that the wheel is not generally

found to be self-dressing. That is, the

grinding process itself does not erode the

grinding wheel bond at a rate sufficient to expose new diamonds before old diamonds

become worn out. As an example, Figure 1 shows the progression of specific grinding energy as a function of material re- moved for a particular ultraprecision

grinding process. In the experiment from

which this graph was obtained, a 4-8 \m\

grit wheel was first dressed to expose new

diamond edges, then used to blanchardgrind CVD silicon carbide under conditions that would ensure a damage-free

surface on the ceramic work piece(4).

After a total removal volume of 3 cubic

millimeters, the specific grinding energy

began to rise rapidly, indicating that the

diamond grains were no longer effectively removing material. Scanning

Electron Microscope (SEM) photographs of

the wheel surface suggested three possible

contributors to increased grinding en- ergy: dulling of the diamond grains,

loading of the wheel with grinding swarf,

and loss of diamonds from the wheel sur- face. Under such conditions, the wheel

must be re-dressed to restore its original

cutting capability.

There are various techniques available

for dressing, which can be divided into

two categories: intermittent and in-process. Intermittent dressing occurs while

the wheel is not grinding, typically at the

end of a grinding pass across the work

piece. A common method of intermittent

dressing involves pressing a porous ce- ramic "dressing stick" against the wheel

surface. The porous ceramic erodes the

wheel bond without significantly affecting the cutting edges of the newly exposed

diamonds. In-process techniques have

witnessed recent progress involving

333
glass-fiber dressing media bonded to the

work piece surface (5), and electrolytic

in-process dressing (ELID) (6-8). ELID has

been found to be an effective tool for production-level machining of ceramics to optical quality finishes using cast iron

fiber bonded wheel (6).

Currently, such wheels are available

from only one source. In the US, the most

common bond material used in precision

grinding of ceramics is bronze. Our re- search is aimed at exploring the applicability of the ELID process to bronzebonded diamond grinding wheels.

MVA.

X =

coo

o

a c

V) m

T

'

1

' r

0 12 3 4

Material Removed (mm^ )

Figure I: Specific grinding energy as

a function of total material removed in a blanchard-style micro-grinding operation on CVD silicon carbide. The sharp in- crease in specific energy is a result of

diamond wear.(6)

BACKGROUND

In the ELID process, an electrolytic cell

is created with the rotating grinding

wheel as the anode, immersed in a flowing

electrolyte and opposed by a fixed,

graphite cathode. The basic theoretical

equation governing wheel recession rate as a function of electrolysis parameters

can be derived from Faraday's Laws of

electrolysis and Ohm's Law (9):

zFlp^pA

where M is the molecular weight of the

anode material, V is the applied voltage

across the gap, Ac is the active surface

area of the cathode, z is the valence of the

anode material, F is Faraday's constant, / is the inter-electrode gap distance, pa is the

density of the anode, p is the resistivity of

the electrolyte, and A-^ is the surface area of the wheel. Often, instead of simple stock

removal according to this equation, the

electrolysis generates an oxide film on the

anode. If the conductivity of the film is low, it may act to retard the electrolytic

material removal rate. ELID techniques have been under development for ceramic microgrinding for

about five years (10). Most of the early

work in this field was performed on cast

iron fiber bond grinding wheels. It has

been proposed that the mechanism of ELID

consists of the formation of a passivated

oxide film on the surface of the cast iron

wheel, followed by a gradual removal of

the oxide layer due to abrasive wear between the oxide and the work piece (6). As

the non-conductive oxide layer is re- moved, the electrical resistance between

the fixed cathode and the rotating anode

(i.e. the grinding wheel) is reduced,

causing increased electrolysis and a re- building of the oxide layer. The oxide layer

acts as a feedback controller in the dressing process. Without it, electrolytic ero- sion would take proceed at a nearly constant rate for a particular set of ELID

conditions. The formation of a passivated

oxide layer inhibits further electrolysis,

linking the wheel erosion rate to the

abrasive wear rate of the oxide.

In practice, ELID is accomplished using

a DC, AC, or pulsed DC power supply. While

all three have been used, the relative

merits or deficiencies of the selected type

of power source are not well understood

(10).

334
ELECTROLYSIS FOR BRONZE-BONDED

GRINDING WHEELS

In 1992, we initiated an investigation

of electrolytic in-process dressing of

bronze-bonded diamond grinding wheels

for ceramic micro-grinding. We began

with a phenomenological study designed

to provide insight into what occurs on the

surface of the wheel during electrolysis.

Using scanning electron micrography, we

examined the surfaces of bronze-bonded

grinding wheels subjected to electrolysis.

A second series of experiments were di- rected at quantifying electrolysis pro- cesses on bronze wheels in terms of the

rate of growth of the electrolytically af- fected layer. These experiments were con- ducted on a rotating grinding wheel,

without any grinding taking place (i.e., no work piece contact). The results will serve

as input for a planned set of experiments

involving electrolysis analysis during

actual grinding.

Surface Features in Bronze-Bonded

Wheel Dressing

In our initial qualitative study, a bronze grinding wheel was cut into eight

segments that could each act as an anode

(these small anodes were sized for viewing

under the SEM). A block of graphite was

used as the cathode. The two electrodes were positioned parallel to one another,

spaced apart by non-conducting shims, in an electrolyte bath (a 1 Molar NaCl solution). A 20 watt power supply was used to provide the current across the electrode

gap. A digital storage oscilloscope was used

to measure both voltage and current. The

anode was subjected to electrolysis for a brief period, and was then inspected un- der the SEM. This cycle was repeated several times. A series of fiducial markings

engraved into the anode surface prior to the beginning of the experiment permitted re-viewing of a particular location at various stages of the experiment.

Figure 2 shows the surface of one of

the small anodes at 500X magnification,

prior to any eliectrolysis. Figure 3 shows

the same location after 2.5 minutes of

electrolysis at 4 volts, with a 380 \im gap. It was observed that etching did not occur

uniformly. Rather, the electrolytic process appeared to cause localized pits on the

bronze surface.

Dynamics of Film Formation and

Wheel Recession

Figure 4 is a schematic of the experimental set-up used to evaluate electrolysis

dynamics. The bronze bonded grinding

wheel is rotated by a servo-controlled DC

motor. The bottom third of the wheel is immersed in the electrolyte bath. A steady

flow of new electrolyte is continually

added to the bath, and the overflow is discarded. Temperature and conductivity

probes in the electrolyte bath are continuously monitored. The anodic wheel is maintained at ground potential using a contacting graphite brush. The graphite

cathode is positioned with its surface

parallel to the face of the cup-shaped

grinding wheel. The electrode gap is adjusted using a micrometer-driven linear

ball-bearing stage. Electrode power is provided by a pulsed EDM power supply,

with specifications outlined in Table 1.

Table 1: Specifications of the EDM power

supply.

Voltage 150 Volts (max)

Current 35 Amps (max)

Pulse On Time 0.1 to 1000 us Pulse Off Time 0.1 to 1000 \is Duty Cycle 0.1 to 99.9%

Before mounting the wheel on the test apparatus, it was trued and two 6 mm wide,

20 f.im deep grooves were cut into its cup

face, using a separate ultraprecision

grinding machine. These grooves serve as reference surfaces for measuring oxide

film thickness and the electrolytically af- fected depth. Before each electrolysis test,

the wheel face was returned to its original

condition by chemically etching it with a three-part solution of distilled water, hydrogen peroxide (3%), and ammonium hydroxide (30%), mixed in equal proportions.

Subsequently, a capacitive displacement

335
Capacitive

Displacement Gauge

Wheel (anode)

Figure 2: SEM photo of bronze surface

prior to electrolysis.

Stage

Graphite (cathode) EDM Power

Supply

Brush

Contact

Figure 4. Experimental Set-up

gauge, with 10 nm resolution, 40 kHz

bandwidth, and 50 |jm range, is used to measure the face of the cup wheel. The

gauge records a profile of the rotating

wheel face; this profile, which includes

the grooves, serves as a reference against

which future profiles can be compared.

To protect the grooves during a test, a hard, removable film was applied to the

grooves. Cellulose nitrate (clear finger

nail polish) was found to be non-conducting, rugged, transparent, and soluble in

acetone, making it ideal for this masking

operation.

After masking the grooves, the trough

built around the lower third of the

grinding wheel was flooded with electrolyte from a large reservoir using a constant flow-rate, one-pass coolant delivery system. The wheel was then rotated at 200 rpm, and the entire apparatus was al- lowed to run for 10 minutes without applied voltage before the beginnmg of

each test.

The electrolysis tests were conducted

using the parameters shown in Table 2. A

total of thirteen tests were made, varying

only the duration of the test, up to 180

minutes. Conductivity and temperature,

which generally remained constant m all tests, were recorded at one-minute intervals.

Figure 3: SEM photo of bronze surface

showing pitting brought about by electrolysis.

Maximum and minimum values of the

voltage and current waveforms were mea336
Table 2. Settings for electrolysis tests

Parameter Setting

Voltage

Electrode Gap

Duty Cycle

Cathode Area

Electrolyte

Flow Rate

Wheel Speed

Wheel Type

Wheel Geometry

130 Volts

250 [im

50 %

55 mm^

0.0002 M NaCl

140 ml/min.

200 rpm

Bronze, 10-20

Concentration 100

101 mm outer diam.

89 mm inner diam.

Test Duration 0-180 minutes

sured and stored at three second intervals.

After the test was corrpleted, the electrolyte was drained from the system. At this

point, a hard oxide film had built up on the

wheel surface (A film forms even in the

absence of electrolysis, due to surface

oxidation in the 10 minute warm-up

period). The wheel was then rinsed with

methanol and allowed to air dry, after

which the protected grooves were

uncovered using acetone. The displacement gauge was then used to measure

a profile of the wheel. When the original

profile was subtracted from this new

profile, a measure of the oxide film thick- ness could be made. A problem with this

technique is that, in principle, the oxide

film should not be conductive, and would

not permit accurate profile measurement

using the capacitance gauge. In fact, measurements of resistance across the film

consistently showed that the film continued to be conductive, even when it had

grown to considerable thickness. After

storing a profile of the oxidized wheel, the

reference grooves were again protected

with cellulose nitrate. The film on the sur- face of the wheel was removed by a twostep application of acids. The acids were

applied to the rotating wheel with a sponge. The first application used 7% HCl

for two minutes. The second application

used 2% H2SO4, also for two minutes. After

each acid application, methanol was applied to the wheel to rinse residual acid

and dry the wheel. HCl and H2SO4 were

chosen because they are used in industrial

pickling solutions that remove oxide films

from bronze without attacking the metal

itself. Resistance of the bronze bond to chemical attack by the acid solutions was

confirmed experimentally.

After removal of the oxide film, the

protective coating on the reference

grooves was again removed, and another

profile of the wheel face was measured

and stored. From this new profile, the

original profile was subtracted, and the

difference was displayed on an oscilloscope. Figure 5 shows one of these

"difference profiles" in the vicinity of a reference groove. The peak of the differ- ence profile corresponds to the reference

groove on the grinding wheel. Since this area was unaffected by the electrolysis

process, the difference profile is set to zero at this peak. From this zero level, the

distance to the to the base of the differ- ence profile is a measure of the electrolytically affected depth. With the oxide film

removed, it is also a measure of the recession of the bronze. Our estimate of the

probable measurement error in calculating the depth of the electrolytically af- fected depth is +/- 1 micrometer. This in- cludes errors due to: non uniformity of the

oxidation and acid etching processes,

gauging/fixturing errors, and setting the

electrode gap.

TOc

TO

'CO ^

o —

_ !C S P

«} >

Q) Q

-3

12

Electrolytically

- Dressed p'^'

Surface 1 j ^ Protected

Surface

i

1

1 1 1 • 1 1 30 40 50 60 70 80

Circumferential Position on

Grinding Wheel (mm)

Figure 5: Difference profile, resulting

from subtraction of final wheel profile

(with the oxide film removed) from refer- ence profile. Data is for a 70 minute electrolysis test. 337
RESULTS AND DISCUSSION

Figures 6 and 7 display data for the series of timed electrolysis tests. Figure 6

shows the electrolytically affected layer as

a function of time, while Figure 7 shows

the height of the oxide film above the

original surface of the wheel, again as a function of time.

Figure 8 shows a representative plot of

voltage and current from one our tests. Unlike the results reported by others for

cast iron grinding wheels, the voltage and

current remain relatively constant

throughout the electrolysis test. This is due to the fact that our oxide film is conductive, and does not inhibit the electrolysis process. Also, at <50 micrometers in

thickness, the oxide film is much thinner

than the 3-4 mm oxide layer reported for

cast iron wheels ( 10).

At the same time, the film advances into

the bronze binder in the wheel, convert-

•

-•-

•

•<

0 50 100 150 200

Electrolysis Time (min.)

Figure 7. Oxide film height (above original

surface) as a function of time.

a

.Si g100 150 200

Electrolysis Time (min.)

Figure 6: Electrolytically affected depth

versus time for electrolysis tests. The electrolytically affected depth is a measure of

the recession of the bronze bond. In this

graph, the solid line is a linear regression

of the data, indicating a slope of 1.5 nm/second.

From Figure 7 one can see that the oxidized surface of the wheel does not recede

with electrolysis. Because of the growth of

an oxide film both into and out of the

original surface, the film on the wheel

face actually advances toward the cathode.

1 50

O 100 - >

a>O)

<S 50 - o>

---Q^D-Di>D-CHD-n

0.15

-0.12 (Aa

E

-0.09 <

Voltage (Volts)

Current

—

r- 10 20 —

T" 30

0.06 c

0.03 5

0.00

Electrolysis Time (min.)

Figure 8: Maximum voltage and maximum

current as a function of time for a 30

minute electrolysis test.

ing bronze into oxide. The depth to which

this oxide film extends is a parameter that

we label the "electrolytically affected

depth." In the absence of work piece

abrasion of the oxidized film, the film

might eventually grow thick enough to obscure the diamonds in the wheel. Thus,

338
the rate at which the oxide film wears in an actual grinding process is a critical

parameter.

For the conditions chosen in our electrolysis experiments, the rate of oxide film

growth into the wheel is approximately

constant at 1.5 nm/s. If the grinding process is such that the erosion rate of the

oxide film due to abrasion is greater than

the rate of oxide film formation, one would

expect the film generation rate (1.5 nm/s)

to be equal to the dressing rate (or rate of

wheel recession). As the abrasion wears

through the oxide layer, the work piece

will contact the bronze sub-layer, which

is difficult to further abrade. Gradual re- building of the oxide layer through electrolysis would allow continued dressing at a rate equal to the film growth rate. Currently, it is unclear to us whether the

oxide film will be more or less abrasion

resistant than the underlying bronze

bond.

If, however, the grinding process is such that the removal rate of the oxide

film due to abrasion is less than the rate of

oxide film formation, one would expect the

rate of growth of the film to outpace the

rate of wheel recession. In this case, the

oxide abrasive wear rate will determine

the rate of wheel recession, independent

of the growth rate of the electrolytically

affected layer.

The film observed on the surface of the

wheel does not prevent, nor even slow

down the electrolytic process; rather the

boundary between the bronze and the

electrolytically affected layer on a bronze-bonded wheel continues to ad- vance into the wheel at a constant rate with dressing time. The rate for the particular combination of factors investigated

here 1.5 nm/s. It is expected that this rate will be decrease linearly with decreasing

cathode voltage magnitude and will de- crease inversely with electrode gap. Our

plan is to alter the test parameters systematically to determine if these relationships will hold, and to install this system

on an ultraprecision grinding apparatus.

CONCLUSIONS

A relatively simple test procedure for

analyzing electrolysis dynamics has been

described. In a series of experiments, we

have found that electrolysis results on bronze-bonded grinding wheels are quite

different than those reported for cast iron

grinding wheels. We have found that a thin, conductive oxide film forms on the

surface of the wheel for the conditions

tested. The rate of growth of this film into

the grinding wheel is remarkably constant for up to three hours of electrolysis.

The dressing process may be controllable

using this constant film growth rate, depending on the abrasive wear characteristics of the film in actual grinding conditions. Further work varying the electrolysis parameters and including the effects

of actual grinding are planned in the near

future.

REFERENCES

1. T. G. Bifano, T. A. Dow, and R. O.

Scattergood, "Ductile-Regime Grinding: A

New Technology for Machining Brittle

Materials," ASME J. Eng. for Industry,

[113], pp. 184-189, 1991.

2. T. G. Bifano, Y. Yi, and W. K. Kahl,

"Ductile-Regime Grinding of Ceramic

Mirrors," in review, J. Precision

Engineering, 1993.

3. T. G. Bifano and S. C. Fawcett, "Specific

Grinding Energy as an In-Process Control

Variable for Ductile-Regime Grinding,"

Precision Engineering, [131:4, pp.

256-262, 1991.

4. T. G. Bifano, Y. Yi, and W. K. Kahl, "Fixed

Abrasive Grinding of CVD SiC Mirrors",

Accepted for publication, J. Precision

Engineering, 1993.

5. K. Abe and N. Yasunaga, "Cleaning Technique for Super Abrasive Grinding Wheel

with FRP Cleaner", Proceedings of ASPE

1992 Annual Meeting, pp. 83-86, 1992.

6. H. Ohmori and T. Nakagawa, "Mirror

Surface Grinding of Silicon Wafers with

339
Electrolytic In-Process Dressing", Annals

of CIRP. 1990.

7. R. Polvani and C. J. Evans, "Preparing

Ultraprecision Grinding Wheels", ASPE

1992 Annual Meeting, October, 1992.

8. R. Polvani and C. J. Evans, "Electrical

Assisted Grinding of Precision Glass and

Ceramic Parts", Optical Fabrication and

Testing Workshop, OSA Technical

Digest Series Volume 24, pp. 120/WB2-

1, 1992.

9. J. A. McGeough, Principles of Electrochemical Machining, Chapman and Hall,

1974.

10. K. L. Blaedel and P. J. Davis, Foreign

Trip Report, (Visit to Japan), Lawrence

Livermore National Laboratory, June 1990.

340
EFFECT OF DUCTILE-REGIME GRINDING ON THE

STRENGTH OF HOT-ISOSTATICALLY-PRESSED

SILICON NITRIDE

L. K. IVES, C. J. EVANS, S. JAHANMIR, R. S. POLVANI,

T. J. STRAKNA*, and M. L. MCGLAUFLIN

National Institute of Standards and Technology, Gaithersburg, MD

Hot isostatically pressed silicon nitride was surface ground in the ductile regime. The

specimens were in the form of standard flexure test bars thereby permitting the evaluation of both

the surface finish and the influence of grinding on fracture strength. A high precision grinding

machine equipped with air-bearing spindles was employed for ductile-regime grinding. In-process

electrochemically assisted dressing of the 2000 grit diamond metal-bond wheel permitted the

simultaneous grinding of multiple specimens. For comparison, similar specimens were also

ground by conventional means employing 400 grit and 900 grit wheels. Specimens ground in the

ductile regime exhibited improved characteristic strength and superior surface finish.

Introduction

Silicon nitride ceramics, along with a number of

other advanced ceramics, are increasingly important

materials for applications which demand high

hardness, high strength, excellent wear resistance,

and extreme chemical inertness. These materials

become even more attractive when exposure is at elevated temperatures. Unfortunately, the properties

that make advanced ceramics so useful, coupled

with very limited ductility, render them difficult to machine. In general, abrasive machining methods

are employed with diamond grinding being the most

flexible for achieving complex parts geometries,

high precision, and satisfactory removal rates [1,2].

The finish and integrity of conventionally diamondground surfaces are suitable for many applications.

For rolling-element bearing components and other

highly critical applications the as-ground surface

may not be adequate. When surfaces must be

I extremely smooth and the presence of machining

damage minimized, an additional finishing step,

I abrasive polishing, is required. Conventional

polishing methods, however, usually involve

different fixturing, give relatively low material

* Graduate Research Assistant, University of

Maryland, College Park, MD 20742

removal rates and have considerable difficulty in maintaining specified tolerances for other than the

simplest geometries.

Ductile-regime grinding is an attractive alternative finishing method which can minimize or perhaps entirely eliminate the need for final polishing [3, 4]. With this grinding method it is possible to create components having extremely

precise geometric form and at the same time produce

surface finishes on a nanometer scale of roughness.

In addition, near surface and subsurface damage is expected to be comparable with that achieved by

fine abrasive polishing. The success of ductileregime grinding has been demonstrated in the

fabrication of aspheric glass optical components [5]. Ductile-regime grinding has also been demonstrated

in the case of semiconductors [6] and several

ceramic materials [3,4,7]. In practice, ductile-regime grinding is accomplished by limiting grit penetration to a depth

less than that at which fracture is initiated. While

grit penetration in polishing is controlled largely by

using abrasive media of a sufficiently small grit size, in grinding, control is a result of achieving and

maintaining extremely precise displacement of the

tool with respect to the workpiece.

To operate within the ductile regime, grit

penetration must not exceed a threshold depth, about

341
0.1 |im or less, determined primarily by the hardness

and toughness of the material being ground [4]. Material removal is confined to the extent of grit

penetration and is described by the cross sectional

profile of the grit in the direction of motion. Plastic

flow is considered to play a critical role and the chip

formation process is likened to that which occurs with ductile materials. When the ductile threshold

depth is exceeded, microcracks are introduced into

the surface and material removal occurs by brittle fracture processes that may extend far beyond the

grit penetration profile [8]. The purpose of the present investigation was

threefold: 1) to show that ductile-regime grinding of

a high strength, hot-isostatically-pressed silicon

nitride material could be achieved; 2) to assess the

influence of ductile grinding on the strength of this material; and 3) to compare ductile-regime grinding

to two more-or-less conventional grinding methods

with respect to influence on strength and surface

finish. The experimental approach was similar to that taken in two related investigations [9,10] concerned

with the effect that grinding might have on the

strength of different types of silicon nitride and on

the influence of grinding parameters on strength and

surface finish. In those investigations and similarly

in the present study, the workpieces or specimens were four-point flexure test bars. By applying the

grinding condition of interest to the critical tensile

surface of the flexure bar, the influence on strength

of that grinding condition was determined by

conducting a flexure test. Subsequently,

fractography, carried out by means of light

microscopy and scanning electron microscopy

(SEM), was used to locate and examine the fracture

origin, the intent being to determine if the

originating flaw could be associated with grinding

induced damage. The surface topography was

evaluated by means of SEM, stylus profilometry,

and light optical methods.

Experimental Procedure

a) Specimen Material

The material employed in this investigation was

a bearing grade, hot-isostatically-pressed silicon

nitride (Noralide NBD-200, a product of Cerbec,

Inc.**). Specifications provided by the manufacturer

are given in Table I. To minimize the possible

influence of processing related variations, the

specimens were cut from the same production lot of

Table 1. Properties of NBD-200 silicon nitride*

\_\_'iVir Wol 1 1\_/1^ 1 \trl 0/- \Ani^ 1 1 Wl. /O Mgi^, 1 bal. SijN^

j DENSITY (g/cm^) 3.16

POISSON'S RATIO 0.26

HARDNESS, VICKERS 10 kg, (GPa) —^- -\ 16.6 STRENGTH PROPERTIES (o! 20°C -FLEXaJRAL, 4-point bend, (MPa) 800 -WEIBULL MODULUS 9.7 -TFNSIT F CMPa'> 400 -COMPRESSIVE (GPa) 3.0 -HERTZ COMPRESSIVE (GPa) 28

(ball on flat) FRACTURE TOUGHNESS, K,. 4.1 (MPa.m'")

THERMAL PROPERTIES

-THERMAL EXPANSION COEF.

(10-*/°C) 170°C - 20°C 0.43 20°C - 1000°C 2.9 -THERMAL CONDUCTIVITY

(W/m.K)

100°C 29.3 500°C 21.3 1000°C 15.5 -MAXIMUM USE TEMPERATURE 1000

1 (°C)

*Manufacturers data

** Certain trade names and company products are identified in order to adequately specify the experimental

procedure. In no case does such identification imply recommendation or endorsement by the National Institute of Standards and Technology, nor does it imply that these products are necessarily the best available for the

purpose.

342
material. The starting material consisted of 5 tiles, each having dimensions of 9 x 58 x 59 mm.

Flexure bar specimens for each grinding condition were distributed equally among all tiles to minimize

variations that might be associated with differences

among the tiles.

b) Grinding Methods

Grinding was carried out to obtain flexure test bars with final dimensions of 4.0 x 3.0 x 50 mm in compliance with configuration B of ASTM CI 161-

90 Standard Test Method for Flexural Strength of

Advanced Ceramics at Ambient Temperature [11].

The tensile surface of each flexure bar was ground:

1) under conditions resulting in a ductile removal

mechanism, 2) using a conventional grinding method

employing a 400 grit wheel that satisfied the

requirements of ASTM CI 161-90 designed to minimize or eliminate the effect of surface

preparation on flexure strength, and 3) under

carefully controlled conditions with a 900 grit wheel

to produce a smooth surface, and also thought to minimize grinding damage. The conditions for each

of these grinding methods (designated ductile, 400

grit, and 900 grit) are summarized in Table 2. In all cases grinding was parallel to the long axis of the

bar.

Initial preparation was the same for all bars. Preparatory grinding was done on the five tiles to obtain flat surfaces and appropriate starting

dimensions. A relatively coarse 150-180 grit resinbond diamond wheel was used for this purpose.

Eighteen bars were then cut from each tile with a resin-bond diamond saw to a size of approximately

4 mm thick, 6 mm wide, and 51 mm long. This

oversize (0.5 mm on each face) was sufficient that the effects of process-related variations at the

original tile surface and damage introduced by

preliminary grinding and cutting could be eliminated

by subsequent surface grinding.

One set of 30 bars was then surface ground in compliance with ASTM Standard CI 161-90. The

standard requires that grinding on each face be

conducted in two stages. In the first stage material

is to be removed at a rate of no more than 0.03

mm/pass and that at least 0.06 mm remain for

removal in the final finishing stage at a rate not to exceed 0.002 mm/pass. The same requirements also

apply to the chamfers which must be placed on the

long edges of each bar. Grinding is to be carried out with a wheel having diamond grit in the 320-500

size range. In this investigation, a nominally-400

grit resin-bond diamond wheel was employed.

Additional grinding parameters are listed in Table 2. Finished in this way, the 30 specimens were then

Table 2. Grinding parameters used to finish flexure bar tensile surface

PARAMETER METHOD

400 GRIT 900 GRIT DUCTILE-REGIME

WHEEL

-TYPE -BOND

-GRIT -CONCENTRATION

Periphery Resin 400

100

Periphery

resin 900

Cup

Bronze 2000

50

DRESSING METHOD

Powered SiC grinding wheel

In process

electrochemical

COOLANT

Water-base

synthetic 20:1

Water-base

synthetic 20:1

Water-base emulsion

80:1 WHEEL SURFACE SPEED (m/s) 29 16

FEED RATE (cm/min.) 0.2

DEPTH PER PASS (mm) 0.0025 0.0001 DEPTH REMOVED (mm) 0.05 0.005

343
ready for flexure strength measurement and surface

finish characterization.

Except for what would be the tensile face, the

above grinding procedure was also used for the two

other sets of 30 bars. For the thirty bars finished by means of the 900 grit method, final grinding of the

tensile face employed wheels of progressively finer

grit size, with the last 0.005 mm removed by means

of a 900 grit resin-bond diamond wheel.

For the bars to be finished by ductile-regime

grinding, the thickness was reduced only to 3.076

mm with the final 0.076 mm to be removed from

the tensile face according to the various steps

associated with the ductile-regime grinding method

described below.

Ductile-regime grinding was performed on a Moore M-18 Aspheric Generator retrofitted with an

integral, AC motorized, Professional Instruments Co.

BlockHead 4R air bearing grinding spindle. The

machine configuration is shown schematically in Figure 1 The axes of rotation of the grinding and

main work spindles lie in a plane parallel to the

machine's xz-plane. The grinding spindle axis was

tilted at an angle of about 10° from the work spindle

axis. Bronze-bond cup wheels, 100 mm in diameter, were employed for grinding.

Main spindle center line

Z-axis

Rotary table

(B-axis)

Cup wheel

I

1 Grinding spindle

'. center line

Figure 1. Machine configuration for ductile-regime

grinding (top view).

To minimize the possible effects of process

variability between flexure bars and to emulate, to some extent, practical grinding approaches, all 30

bars were machined simultaneously. This was

accomplished by mounting the bars to an interface

plate attached to a face plate on the M-18 spindle.

Prior to mounting the bars, the interface plate was

single point diamond turned to eliminate run-out and

obtain a flat, parallel surface. A low-temperature

melting-point wax was used for mounting the bars. Fifteen bars were arranged in three groups of

five, 120° apart, such that the groups effectively

formed the sides of an equilateral triangle. Thus," the

central portion of each bar, where the stress is highest in the four-point bend test, had its machining-lay parallel to the axis of the bar. The

remaining fifteen bars were also arranged in three

groups of five. However, they were oriented parallel

to the medians of the equilateral triangle formed by

the first three groups of bars and were effectively

ground perpendicular to the long axis of the bar. At

the completion of grinding, these bars were set aside

for a separate study, not part of the present

investigation.

As an initial step, preliminary experiments were

performed to identify parameters which would

ensure grinding within the ductile regime. Based on

these experiments, the following conditions were

found to be suitable: a 2000 grit (6-10 nm) 50

concentration diamond wheel, grinding spindle speed

of 3000 rpm (5 m/s wheel surface speed), work

spindle speed of 10 rpm, 0.1 ^im nominal depth of

cut, and 2 mm/min feed rate. Satisfactory results were obtained using both a mineral oil coolant and

a water-base emulsion coolant. Initial machining

experience suggested that wheel wear would be

relatively rapid, requiring frequent dressing, so it was decided to grind exclusively in water-base

coolant. This allowed the use of electrochemical in- process dressing [12].

Initial truing of the wheel was done with a 400

grit (40-50 ^Lm) metal-bond diamond wheel. In the

grinding configuration used, i.e. the face of the cup

wheel at an angle of ~10° to the face plate, there was

nominal line contact between the workpiece and the outer edge of the freshly trued wheel. Wear,

however, produced a conical shaped wheel and an accompanying increase in contact area. This

resulted in a higher contact force and increased

machine deflection, adversely affecting grinding

precision and finish. Thus, retruing was required at intervals throughout the grinding process.

In grinding such hard ceramics as silicon nitride,

especially with metal-bond wheels, maintaining the

wheel in a well-dressed condition is of critical importance. Rapid grit wear requires that the dressing procedure be applied frequently, or better. 344
continuously by means of an in-process method.

Electrochemical dressing, employed in this investigation, is not only an effective means of

dressing metal-bond wheels but also has the

advantage that no mechanical contact is required

with the wheel. Moreover, the rate of dressing can

be controlled by adjusting the applied current. Compared to abrasive dressing methods,

electrochemical dressing also avoids the introduction

j of abrasive particles and the damage that these

j particles might cause to the workpiece.

j The dressing apparatus itself consisted of a flat metal electrode positioned parallel to the face of the

cup wheel. The electrode mounting fixture was

attached to the housing of the grinding spindle.

As-mounted, the flexure bar blanks were nonplanar by over 10 ^lm. Preliminary grinding was

performed using a 400 grit wheel to level the bars. A semi-finishing cycle using a 1200 grit (15-20 |im)

wheel produced a surface with some residual fine

fractures which were removed in the final ductile

grinding mode using the parameters previously

described.

The 400 grit and 900 grit grinding methods

employed conventional horizontal-spindle,

reciprocating-table grinding machines. The grinding was done by a commercial shop experienced in machining ceramic materials. Operational

parameters are summarized in Table 2. Only limited

information was available for the 900 grit method

which was considered to be proprietary.

d) Flexure Test Procedure

Flexure tests were conducted using a four-point

bending configuration. The steel flexure fixture was

fully articulating to facilitate accurate alignment.

Load was applied by means of a screw driven

universal tension / compression test machine. The

applied force was sensed with a load cell and the

I amplified signal monitored by means of a strip-chart

J recorder. Load was applied at a constant cross-head

j speed of 0.5 mm/min. An initial preload of 5 to 10% of the expected failure load was first applied

and displacement halted while inspection was made

for proper specimen alignment. If the alignment was not satisfactory (this rarely occurred) the load was released and the needed adjustments made.

With proper alignment in the preload condition, the

displacement was resumed and continued until

I failure. The load at failure, specimen width and

I thickness, and flexure span were used to calculate

I the failure stress, i.e., the fracture strength of the

' specimen. A Weibull statistical analysis was carried

out on the data. The resulting Weibull modulus and

characteristic strength (defined as the stress at which

63.2% of the specimens failed) were used to assess

the influence of the grinding method. A more

detailed description of the test procedure and data

analysis approach can be found in the paper by

Jahanmir et al. [9]. The fracture surfaces of each specimen were examined by means of light optical microscopy and

in selected cases by means of SEM to identify the location and possible nature of the fracture origin.

Results and Discussion

a) Surface Finish

Scanning electron micrographs displaying the

typical finish of surfaces ground by the ductile, 400

grit, and 900 grit methods are shown in Figure 2-4. Apart from a few pits and an occasional scratch, the

ductile ground surface, at least in the central region

of the bars, was nearly indistinguishable from that produced by careful polishing. A representative

SEM micrograph of the central region is shown in Figure 2a. At the edges of the bars, however, there was a band of fine pits. This indicated that grinding was to some extent in the brittle regime at that location. Figure 2b shows the pitted region adjacent

to the chamfered edge of the bar. The surface prepared with the 400 grit method.

Figure 3, is seen to consist of a relatively high

concentration of pits distributed over a pattern of

parallel grooves lying along the direction of

grinding. The pits are a consequence of

microfracture that occurred during material removal.

The smooth areas between pits suggest that plastic

flow and perhaps ductile removal also had occurred.

The finish of the surface ground with the 900

grit method, Figure 4, differs from that prepared by

the 400 grit method in that grinding grooves are very shallow and are barely detected in this SEM

micrograph (the grooves are more apparent in the

topographic map. Fig. 6b). Pits similar to those

associated with the 400 grit method are present,

however. Complete details of the 900 grit grinding

method were unavailable precluding the development

of additional conclusions regarding the appearance

of the surface. The shape of the ductile ground surfaces was

evaluated by means of a phase measuring

interferometer. Along their central long axes, the

bars were flat to approximately 1 |lm. Lateral traces also showed that the bars were flat near their ends.
Figure 2. SEM micrographs of surfaces ground by ductile-regime method, a) near center of bar, b) adjacent to chamfer edge showing evidence of micro-fracture.

Figure 3. SEM micrograph of surface ground by Figure 4. SEM micrograph of surface ground by means of the 400 grit method. means of the 900 grit method.

346
The centers of the bars, however, were crowned

laterally by about 3 \Lm. This surface geometry can

' be explained by the fact that the M-18 was not designed to be a stiff production grinding machine.

I Thus, because of machine defection, stock removal

will not equal nominal infeed. At the center of the

bar, this causes no problem, nor at the ends of the

bar where the grinding direction is at an angle

relative to the bar axis so that the wheel contact

bridges several bars. As the wheel approaches an

; edge - particularly at the mid span of the bar at the

i gap between adjacent bars caused by the

chamfers - the situation changes. The wheel contact

area decreases resulting in a higher contact pressure, and more material is removed causing the crowning,

j The surface finish of the ductile ground bars was measured with a phase measuring

j microinterferometer using a 20X objective. This

j gave a 0.5 mm square aperture and a spatial

I resolution of approximately 2 [im per pixel. A

I representative example of a ductile ground surface

• is shown in Figure 5. The indicated roughness of

; the ductile ground surfaces was found to be

approximately 3 nm R3. Stylus profilometry was employed to generate

i three-dimensional maps of surfaces prepared with

] the 400 grit and 900 grit methods and to determine

i| roughness values. Representative examples of

j surface maps are shown in Figure 6. Surface

j roughness values are given in Table 3. These values

j were determined by averaging over a 1 mm by 1

'! mm area of the surface.

Dis'tanca (Microns)

I

' Figure 5. Three-dimensional interferon! etry

representation of ductile ground surface.

I

Table 3. Surface roughness values

GRINDING METHOD SURFACE

ROUGHNESS, R, (pm)

400 GRIT 0.30

900 GRIT 0.018

DUCTILE REGIME 0.003

The grooved topography of the surface ground

with the 400 grit method is quite apparent in Figure

6a. As mentioned previously, only shallow grooves

can be distinguished on the surface prepared with

the 900 grit method shown in Figure 6b. The pits

observed in the SEM micrographs result only in small local variations in the surface maps due to the

relatively low lateral resolution employed (10 \Lm

between traces), and the limited capability of the

stylus (10 \Lm tip radius) to trace such features. The

surface ground with the 900 grit method is significantly smoother (0.018 ^im R^) than the

surface ground with the 400 grit method (0.30 |im

RJ.

b) Fracture strength

Weibull plots displaying probability of failure

vs. fracture stress are shown in Figure 7 for the three

grinding conditions. Straight lines representing a

linear least square fit for each data set are shown.

The slope of each line gives the corresponding

Weibull modulus, and the point on the line at which

63.2% of the specimens fail defines the

characteristic strength. Figure 8 compares the

characteristic strength and Weibull modulus values

for each condition. Within the 95% confidence

limits indicated in Figure 8, it is clear that the

ductile-regime grinding method resulted in a higher

characteristic strength compared to the 400 grit and

900 grit methods. Statistically, the 400 grit and 900

grit methods resulted in about the same characteristic

strengths. It is interesting to note that the strength

values in the range 900-1000 MPa obtained here are significantly higher than the value of 800 MPa given

by the manufacturer in Table 1. Although detailed

information was not available, the lower strength is apparently the result of greater damage introduced

during specimen preparation.

As seen in Figure 8, there is no statistical difference in Weibull modulus values determined for the different grinding conditions. That is, the spread

0.0025600 700 800 900 1000

Fracture Stress (MPa)

200

Figure 7. Weibull plots of flexure strength

comparing 400 grit, 900 grit and ductile-regime

methods.

in the distribution of fracture stress values was

approximately the same for the three grinding

methods.

There are four factors which could contribute to the greater strength of the ductile ground specimens.

These are 1) the lack of grinding induced flaws, 2)

smaller flaws at the ductile ground surface compared

to the other grinding methods, 3) a smoother surface

from ductile grinding and consequently the

elimination of notch effects, and 4) a more favorable

residual stress state at the ductile ground surface. Examination of the fracture surfaces in the

immediate vicinity of the fracture origin was

employed to assess the influence of factors 1), 2),

and 3) above. Analyses were not carried out to determine the possible influence of the fourth factor,

the surface residual stress state. The identification

of a larger compressive residual stress at the ductile

ground surface, for example, might be regarded as providing greater resistance to failure. Scanning electron micrographs of a mating pair

of fracture mirrors obtained on one of the fractured,

ductile ground specimens are shown in Figure 9. The mirrors intersect the ground surface in a nearly

semicircular shape. This is a good indication that the fracture origin was at the surface of the

specimen. The centers of the mirrors at their

intersection with the ground surface are shown at higher magnification in Figure 10. The specimens were tilted so that both the fracture surface and

ground surface could be imaged simultaneously.

The presence of an inclusion, large grain, or

evidence of a pre-existing pore could not be

identified. By a process of elimination, this suggests

a microcrack or pit in the surface, perhaps

introduced by machining, might have acted as the

fracture origin. The geometry of the fracture

surface, shown in Figure 10, could indicate that the

microcrack or pit was oriented at an angle to the

tensile axis. Of the 15 ductile ground specimens tested, all had fracture mirrors that were similar in appearance,

apart from size and small geometric variations, to that shown in Figure 9. All were semicircular in shape

D

Qc

u

(A

u

o

D

SI

X)o

400 GRIT 900 GRIT DUCTILE Grinding Condition

400 GRIT 900 GRIT DUCTILE Grinding Condition

Figure 8. Influence of grinding method on a)

characteristic strength and b) Weibull modulus.

349
Figure 10. SEM micrographs of fracture origins in Figure 9 at higher magnification. Specimens are tilted at 45° to show ground surface.

350
indicating an origin in the immediate vicinity of the

ground surface. In two instances the fracture mirror was at the specimen chamfer. In no case was there

conclusive evidence that the initiating flaw was an

inclusion, large grain, or pore, nor was it possible to associate a machining related surface scratch or pit

with the origin. As was noted previously, a few

scratches and pits were present in the ductile ground

surfaces. Moreover, the pitted zone near the edge of

the bar was not a preferred location for fracture

initiation. Therefore, the obvious pits and scratches were not necessarily the sources of fracture

initiation. It is possible that large cracks that are closed or lie slightly below the surface, and

therefore not detected in the SEM, are the

predominant source of failure. Examination of fracture surfaces of specimens

prepared with the 400 grit and 900 grit methods was

similarly indeterminate with respect to associating

the fracture origin with a specific flaw. The fracture mirrors were similar in appearance to those on the

ductile ground specimens. However, the fact that the average strength of specimens prepared with the

400 grit and 900 grit methods was less than that of

the ductile ground specimens suggests that machining related flaws did have an adverse effect on failure. The only alternative, as mentioned

previously, is that a more favorable surface residual

stress state of the ductile ground specimens resulted

in the higher strength.

Conclusions

(1) By utilizing a high precision grinding machine,

ductile-regime grinding of a high-strength hotisostatically-pressed silicon nitride was

demonstrated. Moreover, simultaneous grinding of

multiple specimens was achieved by employing in- process electrochemical dressing of the metal-bond

grinding wheel. Insufficient machine axial stiffness,

however, did lead to a brittle zone near the edge of

the specimens.

(2) Ductile-regime grinding led to a higher flexure

strength than conventional grinding with a 400 grit

resin bond wheel, where the conventional method

met the requirements of ASTM Standard CI 161-90.

A grinding method employing a 900 grit wheel

produced a smoother surface than the 400 grit

method; however, the flexure strength was not improved. Approximately the same Weibull

modulus was obtained for the three grinding

methods.

(3) Fractographic analysis of the failed flexure test specimens was not able to identify the nature of the

originating flaws. All of the examined fractures originated in the immediate vicinity of the surface. Even though a few pits and scratches were present on the ductile ground surface, these visible defects

did not act as preferential sites for failure.

Ackn owled 2emen ts

This project was conducted under the auspices of the

Ceramic Machining Consortium, which is a joint

research program between NIST and several

companies, universities, and government agencies.

Financial support received from the Defense

Advanced Research Projects Agency is gratefully

acknowledged. We thank Dr. Bill Coblenz of that agency for his support and encouragement.

References

[1] S. Jahanmir, L. K. Ives, A. W. Ruff, and M. B. Peterson, Ceramic Machining: Assessment of

Current Practice and Research Needs in the United

States . NIST Special Publication 834, U.S.

Government Printing Office, Washington, DC, 1992.

[2] R. W. McEachron and S. C. Lorence,

"Superabrasives and Structural Ceramics in CreepFeed Grinding," Ceramic Bulletin Vol. 67, pp. 103 1- 1036, 1988.

[3] K. Mitzutani, T. Kawano, and Y. Tanaka, "A

Piezoelectnc-Drive Table and Its Application to Micro-Grinding of Ceramic Materials," Precision

Engineering. Vol. 12, pp. 219-226, 1990.

[4] T. G. Bifano, T. A. Dow, and R. O. Scattergood,

"Ductile-Regime Grinding: A New Technology for Machining Brittle Materials," Trans. ASME Vol.

113, pp. 184-189, 1991.

[5] I. F. Stowers, R. Komanduri, and E. D. Baird,

"Review of Precision Surface Generating Processes and Their Potential Applications to the Fabrication

of Large Optical Components," Proc. SPIE Vol. 966

in Advances in Optical Fabrication and Metrology

Including Large Optics. J. B. Arnold and R. E. Parks, eds., pp. 62-73, 1988.
[6] P. N. Blake and R. O. Scattergood, "DuctileRegime Machining of Gennanium and Silicon," J. American Ceramics Soc. Vol. 73, pp. 949-957,

1990.

[7] P. Blake, T. Bifano, T. Dow, and R. 0.

Scattergood, "Precision Machining of Ceramic

Materials," Ceramic Bulletin , Vol. 67, pp. 1038-

1044, 1988.

[8] A. G. Evans and D. B. Marshall, "Wear

Mechanisms in Ceramics," in Fundamentals of

Friction and Wear of Materials, D. A. Rigney, ed.,

ASME Press, New York, NY, pp. 439-452, 1981.

[9] S. Jahanmir, T. J. Strakna, G. D. Quinn, H.

Liang, R. L. Allor, and R. D. West, "Effect of

Grinding on Strength and Surface Integrity of

Silicon Nitride, Part I," in Proceedings of the

International Conference on Machining of Advanced

Materials , Gaithersburg, MD, July 20-22, 1993.

[10] S. Jahanmir, T. J. Strakna, G. D. Quinn, R. N.

Kopp, C. S. Yoon, and K. Kumar, "Effect of

Grinding on Strength and Surface Integrity of

Silicon Nitride, Part II," in Proceedings of the

International Conference on Machining of Advanced

Materials . Gaithersburg, MD, July 20-22, 1993.

fill 1992 Annual Book of ASTM Standards. Vol.

15.01, Refractories; Carbon and Graphite Products:

Activated Carbon, ASTM, Philadelphia, PA, pp. ; 333-339, 1992. |

[12] H. Ohmori and T. Nakagawa, "Mirror Surface

Grinding of Silicon Wafers with Electrolytic In- Process Dressing," Annals of the CIRP. Vol. 39, pp. l 329-332, 1990. I

352
INFLUENCE OF DUCTILE MODE GRINDING ON THE

STRENGTH OF SILICON BASED CERAMICS

J. M. BOETTGER

United Technologies Research Center, East Hartford, CT

M. K. KER

Cranfield Institute of Technology, Bedford, England

P. SHORE

Cranfield Precision Engineering, Ltd., Bedford, England

D. J. STEPHENSON

Cranfield Institute of Technology, Bedford, England

Test bars of silicon carbide and silicon nitride were diamond ground with various machining

conditions. Ultra fine "nanometric" surface textures were obtained using particular

machining methods.

Three and four point bend tests were carried out to reveal any improvement of fracture

strength through the precision grinding trials. A method of predicting any potential

increase in fracture strength of a ceramic through precision grinding was verified.

Work described in this paper further illustrates that good quality surface texture is no

indication of high strength or reliability of a ceramic component.

INTRODUCTION

Manufacture of precise ceramic components

i often involves some finish machining. This finish

machining is commonly carried out by diamond

grinding. Previous research in this field (1,2) has

shown that the strength properties of ceramics components can be greatly affected by any finish grinding. It has also shown that the direction of the grinding is influential when using conventional grinding

practices. Higher strengths were found when the

direction of "lay" was parallel to the tensile stress

axis.

i

! More recent research (3) has proposed that a

significant increase in fracture toughness and Weibull modulus can be obtained by using "ductile

regime" or "ductile mode" grinding methods. Using

this mode of grinding, the dependence on the direction of machining is considered to be less influential.

I Ductile regime grinding of ceramics is described as

; a machining process which does not introduce or

j propagate microcracking into the ceramic surface,

j: This ductile mode of material removal is possible

only if the stressed volume of material and the

associated forces arising from individual grains

interacting with the ceramic are controlled and

remain below critical levels.

In order to ensure that the mode of material

removal is maintained in a ductile regime it is necessary to accurately control the machining forces. This can be achieved by use of controlled force grinding

or, by using highly stiff and precise machine tools.

Generally, research into ductile mode grinding

of glasses and ceramics has involved the use of fine

grain size grinding wheels (4). It is well understood

that difficulties in achieving a stable process are encountered with such wheels. Much reliance is placed on the dressing technique and recent developments in dressing technology have ensured that

long term stability can be assured when using fine

grain abrasive wheels (5). The electrolytic

in-process dressing (ELID) method allows the use

of fine grain hard metal bond wheels for ductile

mode grinding.

In addition to ductile grinding, there are other

possible methods to modify the surface of ceramics

and therefore influence mechanical properties. The
effect of a flash oxidation treatment on the strength

of engineering ceramics has been well documented

(6). The fracture strength and Weibull modulus can

both be improved considerably by the appropriate

thermal exposure and this is generally believed to

result from a crack blunting mechanism. If this is the

case, then the level of strength enhancement could

provide an indication of the possible improvements

resulting from ductile grinding assuming that the

surface defects can be minimized.

EXPERIMENT

Ceramic test bars (7) were initially produced

using a conventional grinding machine and grinding

technology. Various types of silicon nitride and sili- con carbide were evaluated, as listed in Table 1.

Table 1. Silicon Based Ceramics

Material

Name Supplier Type

Grain

Size Density

NT 230 Norton/

TRW

SiSiC 5[xm >99%

UF15 T&N

Technology

SSiC 5fxm >99%

Issue 4

Si3N4

T«&N

Technology

HIP'ed

Si3N4

l\im >99%

SiSiC - Siliconized SiC

SSiC - Sintered SiC

A number of the test bars were assessed for frac- ture strength and Weibull modulus in the Mil specification condition.

The remaining test bars were either;

1) Precision ground using a high precision and

highly stiff grinding machine.

Or

2) Heat treated using flash oxidation conditions

to oxidize the surface.

These test bars were subsequently tested for

fracture strength.

Precision grinding methods.

The test bars were initially prepared to the Mil

specification using a standard Jones & Shipman 540

peripheral wheel surface grinding machine. The

direction of the "lay" for this initial machining was

parallel to the length of the bars, see Figure 1.

Longitudinal lay Transverse lay

(Parallel to length) (Perpendicular to length)

Figure 1. Lay of grinding on test bars.

Subsequent precision grinding carried out in the

ductile regime was undertaken using a CPE CNC

cup wheel grinding machine, shown in Figure 2. The

specification of the grinding wheels used are listed

in Table 2.

Figure 2. CPE precision grinding machine.

35A
Table 2. Grinding Wheel Specification.

Type

of

Grinding

Grit

Size

Spec.

im)

Wheel

Spec.

Bond

Type Supplier

Conventional

46 D46/C40

/X53

Metal

(brass)

Diamant

Boart

Precision

3-6 D3/6/C75

/FX3

Metal

(cast

iron)

Sintobrator

Precision

3-6 D3/6/C75

/B912

Resin Diamant

Boart

Fine grain grinding wheels, 3-6 [xm, were necessary to produce nanometric quality surfaces.

These wheels are often difficult to use and require

frequent dressing to maintain adequate "openness"

and prevent "loading". Two methods of dressing were used. An aluminium oxide abrasive stick was

used for the resin bonded wheel, and ELID (5) was employed with the fine metal bond wheel.

show similar trends as summarized in Figure 3. This

figure illustrates the change in fracture strength as the surface finish improves for HIP'ed Si3N4 and

SSiC. Each batch of ceramic exhibits only a small or no increase of strength even though ultrafine sur- faces (0.02 [xm Rmax measured on a Wyko TOPO

3D, 40LX head) were produced by ductile mode

grinding on all three materials.

Mode of testing - a, b-3 point bending

c-4 point bending 1000

Q.S 800

xf

Ductile grinding

""„

.. , ^. Resin bond wheei nAnwf>niinnnl DucUle gnnding _ L^nvcmiorai

Cast Iron bond votteel T grinding

(a) I

600

!0

3 400

li. 200

OudHc grinding Casi iron bond whe«l

_L

0.01 0.1 1 Surface finish, [im Rmax

10

Figure 3. Strength as a function of surface finish.

Due to the configuration of the CPE grinding

machine the precision "ductile" mode grinding was

carried out with the "lay" perpendicular to the

length of the test bars.

Oxidation tests.

To investigate the influence of a flash oxidation

treatment on fracture strength, batches of silicon

nitride and silicon carbide test bars were oxidized in

air at 1000° C for 25 hours and 100 hours respectively. These conditions were established to provide

the optimum improvement in properties for each

batch of material.

RESULTS & DISCUSSION

Previous studies reported in the literature concerning the grinding of engineering ceramics have

suggested that there is a good correlation between

surface finish and fracture strength (8-10). It has

also been proposed that although strength increases

as the surface finish improves, there is little benefit

in producing surfaces to better than 1 ^im Rmax

(8-11). The preliminary results in this study also

Examination of the fracture surfaces revealed

that almost all test pieces had failed from surface

defects consistent with machining damage even though the ultrafine surfaces had been produced

under ductile conditions which were known not to

initiate surface or subsurface damage (12,13). This

clearly indicated that critical surface defects were

still present at the surface after ductile grinding.

This was further demonstrated by the large

improvement in strength obtained by the flash

oxidation treatment. Table 3 shows that strength

increases between, 16% and 53% are possible for

the conventionally ground test bars, and that the

increase in strength is highest for the Si3N4. This

increase in strength due to oxidation has been

associated with a crack tip blunting mechanism

which minimizes the effect of surface defects. As

shown in Figures 4 and 5, the majority of failures for

this batch were from internal defects within the bulk

material rather than from machining damage at the

surface. The Weibull modulus also increased typically from 6 in the ground condition to 14 after

oxidation.

355
1

Table 3. Influence of Oxidation on the Strength of Silicon Based Ceramics

(four point bend testing)

Material

Oxidation

TiiTie at 1000°C

(hours)

oirengm in iviu Spec. Condition

(MPa)

Strength in Mil

Spec. Condition

(after oxidation)

% Increase in

Fracture

Strength

(MPa)

NT 230 100 254 295 16

UF15 100 245 338 38

Si3N4 25 360 550 53

Figure 4. Failure from internal defect in

SiSiC.

Figure 5. Higher magnification of defect in Fig. 4.

The modest improvement in strength resulting

from the ductile grinding and the significant benefit

of a flash oxidation treatment both point to the pres- ence of surface damage which was identified as microcracks and a heavily deformed layer produced

during the conventional grinding process (12). The

poor improvement in strength following ductile

grinding clearly results from the presence of surface

damage introduced during previous grinding operations. As the depth of cut required for effective ductile grinding is small, typically 1 ^m or less ,the total stock removed to produce an ultrafine surface finish

is relatively low. Hence, although a high quality sur- face finish may be obtained, this does not guarantee

the removal of surface and subsurface damage

which will control mechanical behavior.

To illustrate this effect the influence of stock

removal on strength enhancement has been studied.

It was considered that the Si3N4 was the most

appropriate for this work since the flash oxidation

treatment had resulted in the highest increase in strength suggesting that this material would be most

sensitive to stock removal. The depth of damage fol- lowing conventional grinding can be estimated from

detailed structural and residual stress measurements. (12). The depth of damage in the form of

microcracks and the level of residual stress varies

from surface to surface depending on the prior

grinding history. For the 0.5 ^m Rmax and the 2.5 jim

Rmax surface finishes considered here, the total depth of subsurface damage was determined to be

40-50 \xm and 80-90 \xm respectively.

The crack length at failure can be estimated

from the fracture strengths. For a beam subjected to bending, (14),

Kic 2.40r

356
The fracture toughness of the Si3N4 was estimated

to be 4.7MPa /m from indentation tests (12). Using

the average strength values given in Figure 6 for

2.5 [xm and 0.5 ^m Rmax surfaces, critical crack

lengths of approximately 90 |im and 50 ^im are

obtained.

This suggests that the magnitude of subsurface

damage produced by conventional grinding is of the same order as the critical defect size and therefore

determines the strength of the ceramic. Thus, to ensure optimum strength, the layer of damaged

material must be removed by ductile grinding.

Figure 6 illustrates the strength enhancement

obtained for both surface finishes as more material

is removed by use of "ductile" mode grinding. The

fracture strengths for these tests were measured

under 3 point bending to maximize the number of

data points obtainable from the batch of test bars.

Two tests per bar were possible for the three point

testing, whereas only one test per bar was afforded

with four point testing. The optimum strength

increase was obtained following the removal of

about 90 [xm and 50 |im for the 2.5 \im and 0.5 ^m

Rfnax surface finishes respectively, this is the same

as predicted from the fracture mechanics approach.

Material -5% Y2O3HIP SI3N 900

Mode of testing - 3 point bending

500

A 0.5 p.vn Rmax P'''^'' to ductile grinding

2.5 fj.m Rmax prior to ductile grinding

0 20 40 60 80

Total depth of material removed,

Figure 6. Strength as a function of material

removed in a ductile mode (using

resin bond wheel).

This improvement in strength as a function of

the depth removed by ductile mode grinding highlights two important conclusions.

Firstly, surface roughness gives no indication of

the strength of the ceramic, this is determined more

by the level of subsurface damage introduced or remaining from the conventional grinding.

Secondly, material removal by ductile grinding can be undertaken without initiating or propagating

microcracks. The fact that the fracture stresses for

the two surface finishes shown in Figure 6 approach

a similar value is to be expected since the optimum

strength is now determined by the intrinsic defect

size in the material. Following the complete

removal of the original layer of surface/subsurface

damage by ductile grinding, new surface defects will be produced which were originally present in the

bulk of the material. Thus, failures should be

expected to result from either surface or internal

locations depending on the probability of finding a

critical defect. Indeed, examination of the fracture

surfaces indicated that the majority of failure were

from internal sources.

These observations demonstrate the potential

benefits of ductile grinding but also show that there

is an upper strength, limited by the intrinsic properties of the ceramic rather than the processing of the

surface. Clearly, the greatest benefits from ductile

grinding will only be realized for high quality ceramics which exhibit high intrinsic strength.

A measure of the potential benefits attainable

by ductile grinding can be obtained from a flash

oxidation treatment which minimizes the influence

of surface defects and should provide a simple means of estimating the upper strength level of the

material.

CONCLUSIONS

1 . Ductile mode grinding can enhance the strength

of engineering ceramics by reducing the size of

surface defects. The maximum level of strength

enhancement is determined by the incidence of

internal defects which will ultimately arrive at the surface as material is removed.

2. Although ultrafine surface finishes are possible

by ductile mode grinding, surface finish has little effect on strength. Rather, it is the level of sur- face/subsurface damage introduced by conventional grinding that determines mechanical

behavior.

357
3. A simple fracture mechanics approach can be

used to estimate the level of stock removal

required in order to remove the original grinding damage.

4. A flash oxidation treatment, used to minimize

the effect of surface defects, gives a good indication of the possible improvements in strength

attainable by ductile grinding.

REFERENCES

1. I. Inasaki. "Grinding hard and brittle materials," Annal of CIRP, Vol. 36, No. 2, 1987.

2. T. J. Whalen, L. L. Terner, "Surface condition

and strength of silicon carbide -silicon materials," The Science of Ceramic Machining and

Surface Finishing II, NBS Special Publication

562, 1978.

3. P. Shore. "State of the art in "damage free"

grinding of advanced engineering ceramics,"

British Ceramics Proc , Vol. 46, 1990.

4. T. G. Bifano. "Ductile regime" grinding of

brittle materials, Ph.d Dissertation, North Carolina State University, 1988.

5. H. Ohmori, T. Nakagawa, "Mirror surface

grinding by metal bonded super abrasive wheel

with electrolytic in -process dressing," Progress

in Precision Engineering, IPES6, 1991.

6. D. W. Richerson, Modern Ceramic Engineering, Marcel Dekker, 1982.

7. U. S. Department of Defense, "Flexural

strength of high performance ceramics at ambient temperature," MIL-STD-1942A, 1990.

8. T. Kubo. Conf. Proc, Japan Ceram. Soc, 1978,

p. 114.

9. S. Ito, M. Nakamura, and W. Kanematsu,

"Machining of high performance ceramics,"

Bulletin Japan Soc. of Prec. Engrg., 1987, p. 167.

10. M. Ota, and K. Miyahara, "Influence of grinding on flexural strength of ceramics," SME 4th Intl. Grinding Conf., Dearborn, Michigan, 1990, p. 538^

11. P Hancock, T.N. Ho, M-K Kerr, D.J. Stephenson. Processing Adv. Mat., 1991, 1, 145.

12. M-K. Ker, Phd Thesis, School of Industrial and

Manufacturing Science, Cranfield Institute of

Technology, 1992.

13. M. J. Ball, N. A. Murphy, P Shore, "Electrolytically assisted "ductile" mode diamond grinding

of BK7 and SFIO optical glasses," SPIE,

Vol. 1573, No. 30 ,1991.

14. F W. Smith et al., ASME J. Appl. Mech, Vol. 34,

1967, p. 953.

358
EFFICIENT AND PRECISION GRINDING TECHNIQUE FOR

CERAMICS WITH ELECTROLYTIC

IN-PROCESS DRESSING (ELID)

H. OHMORI

The Institute of Physical and Chemical Research, Tokyo, Japan

Research on efficient and precision grinding technique for

ceramics is stimulated by the necessity for mass - production of

ceramic parts. The establishment of new grinding technique for

ceramics realizing high efficiency and quality has been expected by

the employment of tough metal bonded superabrasive wheels and

rigid grinding machines . However , the maintenance of the

grindability of ceramics could not be achieved because of its

hardness and toughness. The author together with coresearchers

has proposed the new grinding technique utilizing electrolytic in —

process dressing (ELID), and applied this technique to ceramics

grinding successfully. Mirror surface quality was obtained by

using a # 4000 cast iron fiber bonded diamond (CIFB — D) wheel

for representative ceramics. Stabilization and reduction of grinding

forces have been achieved even in rough grinding operations of

ceramics with efficiency of 6000 mm^/min by using a # 140 CIFB — D wheel and an ordinary surface grinder. Applications of ELID

technique on ceramics grinding such as cylindrical grinding,

aspheric generation, grooving and cutting were also realized.

1. Introduction finished by lapping or polishing

Application of advanced materials

such as fine ceramics is spreading

widely for replacement of con —

ventional materials. Ceramic parts

are machined by grinding and are

processes. However, recent develop —

ments of advanced ceramics are

making its effective machining

processes more and more difficult.

For rapid production of high pre —

cision ceramic parts, new machining

359
techniques, realizing efficiency and

quality simultaneously, is needed.

High reduction of grinding costs

must be done.

The author has proposed the mirror

surface grinding technique with

electrolytic in — process dressing for

silicon wafer (i), and later applied

the technique to ceramics successfully

(2). This technique was proved to

be effective for both finish grinding

and rough grinding of ceramics (3).

This paper describes an advanced

grinding technique for ceramics

utilizing electrolytic in — process

dressing (ELID).

2. Problems in Grinding of

Ceramics

Grinding of ceramics is becoming

more difficult, because of increasing

toughness of the new advanced

ceramics. In ordinary grinding op —

erations, utilizing resinoid or vitrified

bonded super abrasive wheels, lower

grinding ratios and unstable

grindabilities have been inevitable

for ceramics machining (4). Efficient

grinding, utilizing a rigid machining

center and metal bonded diamond

wheels , has been proposed for

ceramics machining ; but its reliability

and adaptability in practical appli —

cations for ceramic parts manufac -

turings have not been established (5).

This is attributed to the lack of

adequate dressing technique. A

possible solution for the above —

mentioned problems is utilization of a new grinding technique known as "ELID — Grinding" which can control

abrasive protrusions before and

during grinding of ceramics.

3. Electrolytic In - Process

Dressing (ELID) Technique

1) Principle

The ELID - grinding method is not

a complex machining process, but an

efficient grinding technique where

electrolytic in — process dressing is

used to maintain the abrasive

protrusions during grinding. This

phenomenon is not continuous ; it

is rather "adaptive" to wear of the

abrasives . Figure 1 shows the

principle of the proposed ELID —

grinding process (1). The grinding

wheel is connected to the positive

terminal of a power supply with a smooth brush contact ; and a fixed

electrode is made negative. In the

small clearance of approximately

0.1mm between the negative and the

positive poles, electrolysis occurs

upon supply of a suitable grinding

fluid and an electrical current.

2) Behavior

Figure 2 shows the electrical

360
behavior during dressing and ELID

grinding. At the time pre — dressing

starts, the surface of the trued wheel

has a good electrical conductivity.

Therefore, the current is as high as has been set on the power source,

and the voltage between the positive

pole and the negative pole is low

(denoted by ®). For several minutes,

the bond material (mainly iron in

cast iron) is removed by electrolysis.

It is mostly ionized into Fe^^, forming

hydroxides, Fe (0H)2 or Fe (0H)3 . Next, these substances change into

oxide substances such as Fe203.

After these reactions occur, the

electro — conductivity of the wheel

surface is reduced with the growth

of the insulating substances . Thereby, the current decreases, and

the working voltage becomes as high

as that which has been set as the

open voltage (denoted by ®).

3) Mechanism

Figure 3 shows a schematic

illustration of the mechanism of the

ELID — grinding process . After

grinding with the predressed wheel

begins, the protruding grains grind

the workpiece. Accordingly, as the

grains are worn, the oxide layer on

the bond is removed. The wear of

the oxide layer causes an increase

in the electro — conductivity of the

wheel surface. Thus electrolysis in -

creases (denoted by (3) in Fig.2) and

the oxide layer can be recovered. The

protrusion of the grains, therefore,

remains constant in a general sense. The cycle described can be changed

according to the grain size of the

employed wheel. ELID works most

effectively if the thickness of the

oxide layer and the depth of the

etched layer on the wheel surface

are balanced.

4) Advanced Construction of ELID

The ELID — system is composed of

such elements as a wheel bond

material, a power supply, and a

grinding fluid. The selection of

elements determines the properties

that can be realized by ELID

Electrical dressing itself is not a new

technique ; it is said that the history

started in the 1960's in the USA (e).

Among many electrical dressing

techniques, ELID is the only one that

utilized a non — linear electrolyzing

phenomenon with an insulating layer

on the wheel surface. Therefore, the

standard ELID - system is composed

of a cast iron bonded wheel, a direct

current pulse generator as a power

supply, and an ordinary chemical

solution type grinding fluid diluted

by water as an electrolyte, which

upon combination generates an in —

sulating layer easily. For rough

grinding, the composition of the
elements for ELID which makes the

oxide layer thinner is better from

the viewpoint of higher dressing rate.

4.ELID - Grinding for Mirror

Surface Machining of

Ceramics

1) Experimental Equipments

Mirror surface grinding experi —

ments were conducted by utilization

of the equipments shown in Table

1. A rotary surface grinder with an

ELID — system and cast iron fiber

bonded diamond (CIFB - D) wheels

were employed. Grinding chips of

cast iron by # 80 electroplated wheel

are used as the "fiber" for the bond

material of the employed wheel (7).

The machine has oil hydrostatic

bearings for the wheel spindle and

the rotary table. Figure 4 shows a schematic of the ELID — grinding type

and a close — up view of the machine.

For ELID — power supply, a pulse

generator was employed. Represent —

ative ceramics such as SiC.SisN^, Zr02,

AI2O3 — TiC were selected. Table 2

shows the properties of the employed

ceramics.

2) Characteristics

All wheels were trued by a car — borundum wheel. Figure 5 shows

changes of ground surface roughness

with and without ELID using a

# 4000 (the distribution of the grit

size is 2 — 6 /i m, the average grit

size is 4 m) CIFB - D wheel.

When ELID is applied , surface

roughness was improved during the

grinding operation, and becomes a

mirror quality at approximately

Rmax40 — 50nm, where Rmax means

maximum peak to valley of the

surface roughness. On the other

hand, surface roughness could not

be improved without ELID. This is

because ELID enables the fine dia — mond particles to protrude and

engage in fine removal of the ground

surface. Without ELID, the bond

material of the CIFB - D wheel made

the ground surface rougher by

rubbing caused from the strong

friction between the wheel bond

surface and the work surface. Figure

6 shows that a stable grinding forces

was maintained during the ELID —

grinding operation using a # 4000

CIFB - D wheel.

3) Examples

Several types of ceramics have

been finished by using the ELID

technique with fine grit CIFB — D

wheels. Figure 7 shows examples for

SiC,Si3N4,Zr02 and AI2O3 - TiC ceramics

finished by ELID with a #4000 grit

wheel. Difference of surface rough —

ness among the examples is very

small. High surface quality of Rmax
30 — 50nm, Ra5 — 6nm (Ra means

average value of peak to valley in

surface roughness) has been ob —

tained. Roughness patterns of the

ground surfaces as in Figure 8

suggest fine grinding marks on the

finished surfaces.

4) Ductile Mode Grinding

Figure 9 shows the ground surfaces

observed at high magnifications. In

mirror surface grinding of S'hN^ and

Zr02 ceramics, mostly ductile mode

ground surfaces were obtained with

the # 4000 wheel. In SiC ceramics,

a mixed ductile and brittle fracture

modes was observed. In AI2O3 — TiC

ceramics , a fine brittle fracture

ground surface was obtained. It is

important to recognize the differences

in removal modes in ground surfaces

finished at similar surface roughness

values. By the employment of ELID

in fine grinding of ceramics, ground

surface characterization is easier and

more reliable than by conventional

grinding, because a fine surface finish

can be maintained.

5.ELID - Grinding for Rough

Machining of Ceramics

1) Experimental Equipments

Rough and efficient grinding ex -

periments were conducted by util —

ization of the equipments shown in

Table 3. A reciprocal surface grinder

with an ELID — system and cast iron

fiber bonded diamond (CIFB — D)

wheels with a coarse grit were

employed. Figure 10 shows the ELID — grinding principle of the appli —

cation and the close — up view of the

grinding machine. For ELID — power

supply, a specialized pulse generator

was employed . Representative

ceramics such as SiC,Si3N4„ Sialon, WC

were selected for the grinding tests.

Table 4 shows the properties of the

employed ceramics.

2) Characteristics

Figure 1 1 shows the normal

grinding force in grinding of Si3N4

ceramics with and without ELID.

When ELID was not applied, the

normal grinding force increased

gradually to approximately 450N.

On the other hand, ELID reduced the

normal grinding force to about 200N.

Table 5 is a summary of several

typical ELID - grinding conditions.

In traverse grindings as in Fig. 11,

the material removal rate of 1800mm^

/min is not high compared with

ordinary grinding operations. But,

in plunge grinding conditions such

as in Table 5, a removal rate as high

as 6000mm'''/'min could be achieved

with ELID - technique. Figure 12

shows an effect of ELID in reduction

of the normal grinding force. Even
when ELID started to be supplied

during grinding operation, grinding

force could be reduced in the course

of the operation. Figure 13 shows

the effect of grinding conditions and

the grinding direction on the

grinding force for SiC.

Figure 14 shows an influence on

grinding force by working current

of ELID. Higher working currents

reduces the grinding force. But Fig.

14 suggests that the differene of the

grinding forces will be smaller after

large stock removal. This is because

CIFB wheel has a non — linear

electrolyzing feature, and so dressing

rate was decreased by a small

amount. Modification of the fluid

for ELID will be effective in

increasing the dressing rate.

3) Difference on Materials

Table 6 shows the normal grinding

force obtained in ELID grinding of

different work materials. The hardest

material is SiC, and the toughest

material is WC. WC tends to load

the wheel surface with chips. Even

when ELID is applied, the grinding

force is higher than that of other

ceramics. SiC ceramic is the most

brittle material among the employed

ceramics, and its grinding force is

the lowest.

S.Application of ELID -

Grinding for Ceramics

1) Cylindrical Mirror Surface Grinding

with ELID

Figure 15 shows application of

ELID for cylindrical external surface

grinding. Figure 16 is an example

of cylindrical surface grinder

equipped with an ELID - system.

Cup wheels can be used for

cylindrical grinding if the wheel

spindle rotates vertical to the work

spindle. But the application shown

in Fig. 16 is popular in production

shops. Figure 17 shows examples

finished by a CIB - D (Cast Iron

Bonded) and a CIB - CBN straight

wheel of # 4000 with ELID on a

cylindrical grinder. The surface

roughness of Si^N, ceramic cylinders

was approximately Rmax40nm, and

the roundness of the finished cylinder

of 0 50mm was 0.1—0.2 fi m.

2) Curved Mirror Surface Grinding

with ELID

Figure 18 shows a schematic

diagram for spherical and aspherical

ELID — grinding. A straight wheel

is used in order to obtain desired

geometry with a CNC grinding

machine. Figure 19 shows an

example of surface roughness of

finished aspherical lens [BK — 7 glass].

Good mirror surface quality has been

achieved by using ELID and a
# 8000 CIFB - D wheel. Figure 20

shows machined and finished

spherical and aspherical mirrors of

ceramics with ELID — grinding

technique. Even on a conventional

machining center, good mirror surface

quality has been obtained. Figure

21 shows stability of the grinding

force on spherical mirror machining.

An ultraprecision aspheric generator

which had airspindles and high

positioning resolution of lOnm was

recently employed, and an excellent

surface accuracy was achieved.

3) Efficient Shape Grinding with ELID

Figure 22 shows an effectively

machined ceramic sample with ELID — plunge grinding using a #140

CIFB - D wheel. Blades of 0.3mm

width could be made without

chipping at a high removal rate of

eOOOmmVmin.

With respect to another application

of rough grinding with ELID, high

speed cutting using ELID — grinding

has been tried. Figure 23 shows

the schematic diagram. Figure 24

is an example cut by a #120 CIFB — D blade wheel of 1mm width and

ELID. Figure 25 shows the stability

and the reduction of normal

grinding force in cutting operation

using ELID.

7. Conclusions

This paper described ELID —

grinding technique and its effects

and characteristics in efficient and

precision machining of ceramics.

Mirror surface quality was esta —

blished with fine grit wheels and

ELID. Moreover, grindability and

stability in rough grinding of

ceramics were greatly improved with

coarse grit wheels and ELID — system.

Finally, the author applied ELID —

technique to curved surface mach —

ining such as cylindrical, spherical

or aspherical, and grooved surface

successively. Practical manufacturing

of ceramics parts is expected to

become effective widely from rougher

to finer or ultrafine processes.

Acknowledgements

The author wishes to thank Prof.

Takeo Nakagawa ( University of

Tokyo), Mr. Ichiro Takahashi (The

Institute of Physical and Chemical

Research), NACHI - FUJIKOSHI Corp.,

SINTOBRATOR Ltd .

, YAMAZAKI

MAZAK Corp., MAKING Milling Co.

Ltd., Stanley Electric Co. Ltd., Sodick

Co. Ltd., Kuroda Precision Machinery

Co. Ltd.. Mitutoyo Co. Ltd., SHIN - ETSU HANDOTAI Co. Ltd., NIKON

Co. Ltd., MITSUI SEIKI Kogyo Co.

Ltd., Yushiro Chemical Industry Co.

Ltd., Olympus Co. Ltd., OKUMA

Machinery Works Ltd.. NORITAKE

Co. Ltd., Fuji Die Co. Ltd.. and
TOPCON Corp. for their support and

cooperation.

References

[1] H. Ohmori and T. Nakagawa : Grinding of Silicon using Cast

IronFiber Bonded Wheel (3rd report),

Preprint for Autumn Conference of

JSPE. (1987) 687. (In Japanese)

[2] H. Ohmori and T. Nakagawa : Mirror Surface Grinding of Hard and

Brittle Materials Using Cast Iron

Fiber Bonded Wheel, Preprint for

Autumn Conference of JSPE, (1988)

355. (In Japanese)

[3] I. Takahashi, H. Ohmori and T.

Nakagawa : Highly Efficient Grinding

of Hard and Brittle Materials (2nd

report ) , Preprint for Spring

Conference of JSPE, (1989) 355. (In

Japanese)

[4] K. Takahashi and Y. Sekiya : Characteristics in Matching of

Diamond Wheel — Influence of Filler

Type in Resinoid Bond — , Preprint

for Spring Conference of JSPE, (1988)

657. (In Japanese)

[5] T. Uematsu, K. Suzuki, T. Narita

and T. Nakagawa : Grinding of Hard

and Brittle Materials Utilizing

Machining Center (9th report).

Preprint for Spring Conference of

JSPE, (1987) 589. (In Japanese)

[6] K. Okano : Electro — chemical and

Electro — discharge Dressing of Super — abrasive Wheel, Journal of JSPE,

Vol.55, No.6, (1989) 998. (In Japanese)

[7] T. Nakagawa, M. Kimura and K.

Suzuki : Manufacturing of Diamond

Wheel with Matrix Using Fiber

Obtained by Grinding , Preprint for

Spring Conference of JSPE, (1985) 633.

(In Japanese)

[8] H. Ohmori and I. Takahashi : Rough and Efficient Grinding of

Electrolytic In — Process Dressing

(ELID), Preprint for Autumn Conference of JSPE, ( 1992 ) 543. (In

Japanese)

[9] H. Ohmori and T. Nakagawa : Mirror Surface Grinding of Silicon

Wafers with Electrolytic In — Process

Dressing, Annals of the CIRP Vol.

39.1,(1990) 329.

[10] H. Ohmori. I. Takahashi and T.

Nakagawa : Mirror Surface Grinding

by Metal Bonded Super — abrasive

Wheel with Electrolytic In — Process

Dressing , Progress in Precision

Engineering (IPES6/UME2) , ( 1 99 1)

152.

[11] H. Ohmori and T. Nakagawa:

Electrolytic In — Process Dressing

(ELID) for Mirror Surface Grinding,

Proceedings of the 10th International

Symposium for Electro — machining,

(1992) 553.

[12] H. Ohmori : Electrolytic In —

Process Dressing (ELID) Grinding

Technique for Ultraprecision Mirror

Surface Machining, Int. J. JSPE, Vol.

26, No.4. 273.
CIFB-D/CBN wheel

Chemical-solution-type Nozzle ^ ^ grinding fluid

A. System construction B. Electrode detail

Fig. 1 Principle of ELID-grinding

<

5

Wheel ; 5^150 X WlOmm straight, if 4000CIFB-CBN

Condition ; £"0=60"^, /p = 10A, ron = roff = 2/iS,

Rotation ; 900rpm, Electrode ; 1/6 copper, Gap ; 0 . 1mm

(electrode area is 1/6 of the wheel surface) :o I

-

10 15

Electrical dressing

80

>

60

40

TO "c>

20 c

0

25 30 mm

ELID-grinding

Fig. 2 Behavior of electrolytic dressing and ELID

367
©Dressing started (3A~) ©Dressing completed (~-lA)

Protruding

grain

n Elid-cycle

® ELID stabilized {~2A)

(Hydroxide, oxide)

Insulating layer

I

©ELID started (~1 .5A)

7

/ Oxide layer

Worn removed Contacted Scraped oxide

Fig. 3 Mechanism of ELID-grinding

Tablel Specification of equipments

Grind ing

mach ine

Vertical rotary surface grinder

*Both the wheel spindle and the rotary table

are oil nyurosLaLic oearings.

Grinding

wheel

tl600CIFB-D wheeKfor rough grinding) , Cone. 100,

tl4000ClFB-D wheeKfor fine grinding) . Cone. 50

Dimentions: 02OOXW5 Cup

Power

supply

Direct current pulse generator

*Open voltage: GOV (square wave)

*Peak current: 30A

*Pulse width: 1-5// s(on-time, off- time)

*Duty factor:maximum 50%

Fluid Chemical solution type grinding fluid

(X50 di lutin by water)

Truer ttlOO Carborundam wheel ( 0 80 XW20)

and brake truer

MeasurementSurface roughness measuring instrument using

5/imR diamond stylus

368
Table 2 Properties of ceramics workpieces

^^^^

oil 0 ; M 0I3N4 zrU2 A 1 D -'\ ;r A I2U3 1 IL

Vicl<ers hardness Hv Z/UU 1 /UO lyuu

rrdCLure Louf^rinebb i\ic

MPam^''^

3.0 7.0 6.8 5.0

Remarl^ PSZ TiC 30%

Sintering • normal

C1FI5 wheel Cp |l'.,wer|

C()olaiU<2,

\\'(jrk Rotary table

a B

o

JtHiilary iiidin>;) C

Carbon brnt>h

Fixture

Isdlatiir

tlectrode

A) ELID-Grinding application

B)Close-up view of grinding machine

Fig. 4 Rotary ELlD-grinding principle and the employed machine

369
400

E 300

c

CO^ 200

O

3JO

CO

100

0

•o

.O" •o, o \ o

\Qrdinary Grinding

WA Dress

[ V lOOOm/min, '\

f 44mm/min, d2 /zm ] *

Work: SiC

Mirror Surface

'Achieved Eiid Grinding

0 1 2 3 4 5 6

Number of Pass

7 8 9 10

(v, f and d mean wheel velocity, feed rate, and depth of cut)

Fig. 5 Changes of ground surface roughness for SiC workpiece

using tt4000CIFB-D grinding wheel

120

# 4000CIFB-D. ^'=-1000m/min, /=50mm/min. d=2;xm

206

Grinding time s

Fig. 6 Change of grinding force with ELID for WC

412

370
Fig. 7 Examples of finished ceramics

371
Ra-0.005/2m

A)SiC

RTMa.=0.030//m .

B)Si3N.

5 T "

^"t Ra=0.005Azm

I R.a.-0.n34/zm

Ra-0.006/im

RTMa.=0.048/2m

C) ZrO^ D)AU03-TiC

Fig. 8 Surface roushness obtained by ELID-grinding

A)SiC B)Si:

C) Zr02 D)Al203-TiC

(Optical microscopic photograplis) 50//m

Fig. 9 Surface quality by ELID-grinding

372
Table 3 Specification of equipments

Grinding

machine

I Reciprocal surface grinder

*Maxumin table speed is 25m/min, and maximum

spindle rotation is 2550rpm.

*The spindle motor capacity is 2. 2kW.

Wheel ttl40,ttl70CIFB-D straight wheel, Cone. 100

Dimentions: 0 150XW1Q

Power

supply

Direct current pulse generator

*Open voltage :60,90V

*Peak current: 10, 30A

*Pulse width:l. 5-5. 5/zs (on-time, off-time)

^Maximum duty factor: 50%

Fluid Chemical solution type grinding fluid

(X50 diluted by water)

Truer mo Carborundam wheel ( 08OXW2O)

and brake truer

Measuring

instrument

2-axes dynamometer for measuring grinding

forces and analizing system of grinding forces

on a personal computer

Table 4 Properties of ceramics workpieces

SiC Si3N4 Sialon WC

Vickers hardness Hv 2700 1700 1740 1200

Fracture toughness Kic

MPam^''^

3.0 7.0 7.7 12

Remark Co 12 %

Sintering normal

373
A) EL ID-grinding application B)Close-up view of grinding macliine

Fig. 10 Reciprocal ELID-gr inding principle and the employed machine

- £, = 60V, /p = 30A. r„„=r„ff = 1.5/.s

^0 500 1000 1500 2000 2500 3000

Stock removal mm^

Fig. 11 Changes of normal grinding force;

Whee 1

: ttlTOCI FB-D, Work : S i sN^ (50 x 50)

374
Table 5 Plunge grinding condition

~"^--.-„l\'ork material

Conditions

SiC Si3N4 Sialon WC

Grinding

conditions

Wheel velocityiv m/min 1200

Feed rate :f m/min 20

Grinding widthiB mm 10

Depth of cut:d ij. m 30 10

Removal rate mmVmin 6000 2000

ELID

conditions

Open voltageiEo V 90

Peak current: Ip A 10

On-time : r on p.s

Off-time : r off us

2

2

4 80*

3 2 01

16 04

-I I I [ V

. <vl200m/min, f20in/min, d30/zm, WlOmm, Work;SiC>

r iPI MI i:' i' i'''' if' ''if|f

472

Grinding Time S

944

Fig. 12 Reduction of grinding force due to ELID for SiC

375
y 0 ^

300

(vnOOm/min, [20in/inin, WlOrnm, WorkiSiC)

, .iJ.ii^..I...Dlli!".Cit.

(1_5 //JB^

000 2000 3000 4000

Stock Removal V mm'

5 0 0 0

Fig. 13 Influence of grinding condition and grinding direction for SiC

C

LL

2000 3000 4000

Stock Removal V mm^

5 0 0 0

Fig. 14 Influence of current on grinding force for Si3N4

376
Table 6 Normal grinding force and working current for ELID

grinding of different work materials

Removal Rate

mm^/min

SiC SiaN^ Sialon WC

Fn N Iw A Fn N Iw A Fn N Iw A Fn N Iw A

2000 170 2.2 280 2.2 300 2.2 450 6.0

4000 250 3.0 500 6.0 500 4.2

6000 350 4.0 650 7.5 650 5.0

Electrode I 1

r

Coolant

(-)

(+)

Center

Cylindrical work

Fig. 15 Application of ELID for cylindrical external surface grinding

377
.p^curve.

4.

Xc =B. 25mm ! c' i X3.. : : A^:

Fig. 16 Example of cylindrical surface grinder with ELID

Ra4nm, Rinax40nm -

-X"

><288800i:,:-il»l: m . 1

.1. . PS P

XI00 -'A'r' m.,M„:j.;.., :

SiaN. (tt4000ClB-D.ELID) Steel (tt4000CIB-CBN. ELID)

A) Surface roughness

.!! :!. .'!! jl. Jil ' ..„\\i,i„,,uuh, „*.. awvu. ,\\\\\ B)Finished workpices

Fig. 17 Examples finished by cylindrical grinder with ELID

378
Fig. 18 Schematic diagram for spherical and aspherical ELID-grinding

A) Surface roughness " B) Ground surface

Fig. 19 Surface quality of finished aspherical lens (BK-7)

379
A)WC mirror (speherical) B)SiC mirror (asphericai)

Fig. 20 Finislied spherical and aspiierical ceramics mirrors

,,^^^"5^"^ : z;, 1500 m /mm, /J 00 mm /mm, d 2 /im

Conditions ' n.Ar^ i • ^r^n Qrrjn f'vv 300 rpm, 7? 100mm . work . ^loOmmbi

Conditions- ^oHOV, Ip5A. ronS^s. roffl. 7

Time

Fig. 21 Stability of grinding force on spherical mirror machining

for si 1 icon

380
A)Sialon B)Si3N„

Fig. 22 Ceramics specimens machined with ELID-plunge grinding

A)Cutting with ELID B)Near view of machine

Fig. 23 Schematic of high speed cutting with ELID — and the employed machine

I
A)Al203

(Cylindrical cutting sample)

B)Si3N. OAUO,

(Groove shaping samples)

Fig. 24 Examples cut by ttlOO/120 CIFB-D blade wheel with ELID

o

o

c

.1—1 -a

c:

1—1

SLh C3

o

60

50

40

30

20

10

0,

[ GrindingCond.: ^^igScT^Bio' 65?^f R.fsm/ minHidth 1 ] - [ ELID Cond.: E06OV. UlOA^ ronl. r o-rrLS/zs]

^ t Ordinary Grinding

t ELID Grinding

<Work:Si3N4. 50X50X20

-J L_l \ L

I

10 13 2 0 2 5 Grinding Length m

J 1 1 I I I 1 1 0 250 500 750 1000 1250

Stock Removal mm'

Fig. 25 Stability and reduction of normal grinding force

in cutting operation with ELID for SisN^

382
FREE ABRASIVE MACHINING

383

ULTRASONIC SHAPING: AN ECONOMICAL PROCEDURE

TO MACHINE COMPLEX CERAMIC COMPONENTS

R. HAHN and C. LOST

University of Tubingen, Tubingen, Germany

P. SCHULZE

Erosonic AG, Wattwil, Switzerland

Anew technique is presented to fabricate complex ceramic components using a multipart tool

system for ultrasonic machining. Mechanical oscillation combined with an abrasive slurry causes the vibrating sonotrode to reproduce its shape into the hard and brittle material. The

result of this machining operation is an accurate component of almost any geometry.

In this study, the influence of the ultrasonic shaping procedure on the surface quality, bending

strength and fatigue behaviour of a feldspatic porcelain and an experimental monolithic

alumina was investigated. It was found that diamond grinding of porcelain yielded a rough

surface and a decrease in bending strength and fatigue. In contrast, ultrasonic shaping resulted

in a smoother surface. However the smoothest surfaces were achieved by lapping, although,

bending strength and fatigue behaviour of the conventional feldspathic porcelain were

significantly increased by ultrasonic machining. The influence of the work gap to the process

is discussed as an effective instrument to control the material removal mode, either for

shaping or for finishing in the same operation.

Introduction

Due to the dimensional distorsions during firing,

complex ceramic components have to be (re-)

machined to control shapes, tolerances and surface

finish. Material removal rates, surface quality and

subsurface integrity depend on the microstructure

and toughness of the machined ceramics (1).

Conventional diamond coated grinding instruments

are suitable for grinding porcelains efficiently (2).

However, for the most high toughness monolithic

ceramics, ceramic matrix composites and single

crystals, efficient removal rates can hardly be

achieved by diamond grinding (3). Additionally,

accurate geometrically complex shapes, such as prototypes, unique medical implants or tooth restorations, and close dimensional tolerances can not be

realized (4, 5).

Alternatively, honing and lapping are efficient

refinishing procedures for polishing flat surfaces

with diamond slurries (6). Chemical etching is a

slow process, which is hard to control for vertical

removal, although future development of laser

induced etching may increase the applications.

Laser or elektron beam machining of brittle

ceramics yield thermal gradients and explosively

induce internal stress, wich may cause spontaneous

damage (7). In addition, shaping three-dimensional

surfaces is almost impossible (8). Electron discharge machining requires an electrical conductive

material or the addition of conductive elements (9),

which yields a black color and decreases the original strength behaviour. Both ultrasonic machining

and ion milling, using ionized argon in a high

vacuum, are suitable for shaping ceramics gently

(10, 11). However, because ion milling requires

expensive machinery, it is not economical for

shaping unique components or small volumes, such

as medical implants or tooth restorations. Thus,

ultrasonic shaping remains as the only technology
particularly suited for shaping high toughness cera- mics i.e. for tooth restorations (12).

The presented processing technique is based on an

accurate wax or resin model ofthe target component

(10). Using an adjusting device the positive model

is favourably orientated to the z-axis of the machine, unilaterally fixed and contralaterally coated

with a special polymer up to the widest circumference (equator). The resin die is then locked at a

tool holder, which is aligned parallel to the z-axis of

the machine. After hardening and insulating of the

polymer, the remaining model surfaces are completely resin embedded and also fixed to a contralateral

aligned tool holder. The obtained polymer presonotrodes can be separated and the original model

removed completely. A mould form is made enclosing the presonotrodes with a refractory investment

material and conventionally prepared to casting.

The actual sonotrode patterns are directly cast onto

the tool holders using a tough metal (13).

For ultrasonic machining, as detailed below, the

sonotrodes are axially aligned, spatially locked and

unilaterally coupled to the oscillating unit of the

machine, in succesion. In polyphase operations the

tools are sunk into both sides ofthe preform, thereby

the surface properties of the component can be

characterized by using defined anisotropic blanks

and modelling the sinking depth of the first operation (Figs. 1, 2).

The aim of this study was to compare the surface

quality, bending strength and fatigue behaviour of

ultrasonically shaped ceramics to that after conventional grinding and lapping.

Test Procedure

In this study the machining induced influences on a feldspatic porcelain and an experimental monolithic aluminium oxide ceramic (ISO 6474 standard) were compared. Seventytwo rectangular porcelain

test bars of size 3.0 mm x 5.0 mm x 22.0 mm were

fabricated following ISO 6872 standard. Twentyfour randomly selected specimens each were alter- natively surface treated by lapping, diamond grinding or ultrasonic machining.

A vertical standardized device served to cut the

surface plan-parallel to 2.00 mm in thickness. The

lapped samples were preground to a thickness of

2.30 mm, followed by diamond lapping to 2.00 mm,

using water based diamond slurry with decreasing

abrasive grain sizes to 3 |j,m.

For grinding, an experimental set up was used,

mounting a conventional D 91 diamond wheel to a high-speed electromotor. The cutting speed was

constant at 35 m/s. A standardized normal force of

20 N (linear spring) was applied acting the sample

against the rotating tool. For the aluminium oxide

ceramic the wear of the diamond wheel was extremly high, which made grinding of the particular test bars almost impossible. Consequently, diamond

ground or lapped alumina specimens could not be

evaluated.

For ultrasonic machining a prototype machine was

used as detailed below. Another twentyfour ultrasonically shaped alumina test bars (ISO 6474) of

equivalent dimensions served as controls.

Evaluating the critical bending strength and fatigue

behaviour, two subgroups of 12 specimens each were randomly divided from the porcelain and alumina test bars. Using a compression testing machine, a three-point bending test was applied in air at room temparature (cross head speed = 0.05 mm/

min, span between bearers = 15.00 mm, roller

bearing surfaces = 1.6 mm in diameter) measuring

the applied load over the flexion of the sample (ISO

6872 standard). Following the fracture, the machined porcelain surfaces were investigated using profilometry. Rmax (maximum peak-to-valley height),

Rz (mean peak-to-valley height) and Ra (arithmetric mean roughness value) were measured following DIN 4768 (mesured length = 4.8 mm,

diamond cone 2 mm, 60°). Additionally, Rk (core

height of the profile excluding the highest peaks and

deep valleys) and the affiliated bearing ration curves (Abbott-Plots) were quantified in accor- dance with DIN 4776 to show the ratio between the

bearing surfaces and the profile depths. Six spots of

all specimens were analyzed to calculate median

and percentiles.

Each surface of the test samples was measured

twice, arranging a 90° angle between the measured

length which was located parallel to the edges.

For fatigue, cyclic load was applied to the remaining 12 samples (50 000 cycles, 60% of the critical

bending stress, stress ratio = 0.1), followed by

386

i
critical fracturing as detailed above. The data were

j analyzed with a statistical software. The median and

corresponding quartiles of the roughness values are summerized in Table 1, and bending strength data

are displayed as a box plot (Fig. 7). An ANOVAwas

used to prove for significance (95%).

Ultrasonic shaping

: The machining process (as specified in DIN 8589

I part 15) uses ultrasonically induced vibrations (f =

20 + 0.5 kHz) amplified by a booster and delivered

I to a designed tool (sonotrode). This, combined with

an abrasive slurry, causes the tooling to reproduce

its shape into any hard and brittle material, in

negative (12). The material removal process is based on an indirectly applied mechanical energy

made by defined longitudinal strokes of the form

I sonotrode coupled with a slurry, such as boron

I carbide (F 280) suspended in water (1:10). The fast

I motion of the tool works the abrasive into the face

! of the workpiece, thus microscopically chipping

! and grinding away the material (Fig. 3a). Moving

i down the tool, the abrasive filled fluid causes an

'! elastic-plastic distortion of the sonotrode face and

i creates flat microcracks in the surface of the workpiece (Fig. 3b). Swinging the sonotrode up, the

'! work gap is rinsed with slurry, supplying the abra-

; sive and removing the chips (Fig. 3c). Because the

carrying fluid is successively compressed and

'.' expanded between sonotrode and workpiece, a

I cavitational erosion also occurs.

! The process can be controlled by varying the gap

I between the tool and the workpiece. During the

1 actual shaping period, the distance between the

sonotrode and the ceramic is kept constant by tuning

the vertical pressure (p = 0.5 - 1.5 bar) and the

I amplitude (A^ = 5-20 |xm) of the master sonotroI de to an optimum, which is specific for each indivij duell sonotrode and material. Thereby the material

I removal rate can be raised to a maximum, decreasing the process time and the wear of the tool

surface. Exceeding a critical vertical pressure, a

sufficient acceleration of the abrasive is impossible,

; limiting an economic removal performance,

i Increasing the gap between the tool and the machined surface by tuning the vertical pressure to 10 %

of the individuell applied maximum, immediately

before achieving the target shape, the energy adaption becomes even more indirect by visco-elastic

muffling of the abrasive movement in the slurry

suspension. Thereby refinishing can be applied to the machined surfaces (Fig. 3d).

Results

The diamond ground surfaces revealed superficial

chipping, such as acute indents, spallings and sharp

incisions, which yielded a rough porcelain surface

microstructure (Fig. 4). On the other hand the ultra- sonic machining provided a leveled surface microrelief (Fig. 5), although diamond lapping decreased

the porcelain roughness values the most (Table 1, Fig. 6).

When compared to diamond grinding and lapping,

ultrasonic machining resulted in a significantly

increased bending strength of the porcelain test bars

(Fig. 7). However, this particular effect of the ultra- sonic treatment on porcelain significantly decreased after applied fatigue. There were no significant

differences in the critical bending strength of the

lapped and the ground porcelains, although after

applying fatigue the ground specimens were weakened (Fig. 7). The experimental alumina revealed

significantly the highest bending strength (Fig. 7).

Discussion

The functional performance of ceramic components is based on the surface structure and their

specific chemical or physical properties. The

influence of applied surface treatment or machining

procedure on the surface microrelief was recently

investigated. Thus, SEM studies, as well as profilometry, exposed glazing and conventional polishing,

i. e. using fine-grained finishing diamonds suitable

for achieving smooth surfaces, although overglazing revealed the smoothest microrelief (14).

However, the surface quality was dependent on

both, the microstructure of the ceramic and the

applied surface treatment. Fracture mechanic

studies and stress analysis experiments exposed

surface and subsurface defects of ceramic components, such as initial microcracks, acute indents,

sharp incisions or delamination as the most crucial

factors limiting strength and fatigue behaviour (15,

16, 17).
As a characteristic feature of ceramic surfaces, the

small indents, pinholes and microcracks are not

accessible for conventional profilometry using the

diamond stylus technique (5 \xm in tip, 90°) (18).

The translucency of feldspathic porcelains complicates optical profilometry such as laser scanning

systems (19). Therefore, a special diamond cone

(2|Am in tip; 60°) and a filter system following DIN

4768 were found as being compatible for ceramic

surface analysis.

Applied diamond grinding induces a rough surface

and subsurface damage resembling delamination

(20). While refinishing using conventional

diamond lapping, the delaminated layer can be

removed and the surface relief leveled, although

deep radial microcracks which were induced during

diamond grinding are retained.

On the other hand, ultrasonically machined cera- mics show a rougher surface than lapped specimens, however the rounded microrelief is favourable for carrying applied load (21). Additionally, the

oscillating hammering of the boron carbide against

the ceramic cause surface inherent residual stresses. Thus, the fracture resistance can be increased by

offsetting the applied stress (22), even though

fatigue partly decreases this effect during cyclic

loading.

Due to their low tenacity, this effect is favourably

featured when ultrasonically machining fragile glass

matrix composites, such as feldspathic dental porcelains. In contrast, high tough ceramics, including

monolithic alumina are primarily more resistant to crack growth, fatigue and fracture (4).

However, the principle to reinforce ceramic materials by inducing inherent residual stresses is well

known from tetragonal zirconia polycrystals which

are used for reinforcing alumina (23, 24). In conclusion, applied load can be statistically

concentrated on surface inherent flaw tips, such as

superficial sharp indents or subsurface damage, on machined ceramic components. Depending on the

crack growing energy, a blind fissure or a critical

progressing crack can be initiated leading to fatigue or fracture (4, 25). Networking surface and structurally inherent flaws by statistically growing microcracks, the fracture resistance may decrease,

leading to fatigue failure (26).

Conclusions

1. The described ultrasonic shaping procedure facilitates the machining of almost any geometry

and provides an economical removal rate, while

treating the surface layer gently.

2. Bending strength and fatigue behaviour of ultra- sonically treated porcelain test bars are significantly better than the ground or lapped controls.

References

1. I. Inasaki, "Grinding of Hard and Brittle

Materials," CIRP Annals. Vol. 36, pp. 463-

471, 1987.

2. W. Mormann and M. Brandestini, "Verfahren

zur Herstellung medizinischer und zahntechnischer alloplastischer, endo- und exprotheti- - scher PaBkorper," Patent, EPA 0054785,

1981.

3. O. W. Richardson, "Modern Ceramic Engineering," Marcel Decker Inc., New York, 1982.

4. R. Hahn and C. Lost, "Konventionelle Dentalporzellane versus bruchzahe Hochleistungskeramiken," Dtsch. Zahnarztl. Z., Vol. 47, pp. 659-664, 1992.

5. M. Week, R. Hartel and U. Bottger, "Anforderungsprofile an Fertigungseinrichtungen zur Ver- und Bearbeitung von keramischen

Werkstoffen," VDW Forschungsbericht 0150. Aachen, 1987.

6. G. Spur, "Keramikbearbeitung," Hanser,

Miinchen, 1989.

7. C. Emmelmann and H. Semraw, "Keramikbearbeitung mit dem Laser," Laser Magazin. Vol. 3, pp. 8-14, 1987.

8. R. F. Firestone, "Lasers and other Nonabrasive

Machining Methods for Ceramics," Mat.

Science Comp.. Vol. 32, pp. 133-138, 1987.

9. U. Panten, "Funkenerosive Bearbeitung von

elektrisch leitfahigen Keramiken," Tech Diss. TH Aachen, 1990.

10. R. Hahn, "Verfahren zur Herstellung von

Zahnersatzteilen mit Hilfe zweier Sonotroden,

Vorrichtung zum formgebenden Bearbeiten von Werkstiicken unter Verwendung einer

Sonotrode und Verfahren zur Herstellung

einer in der Zahntechnik verwendbaren

Sonotrode," Patent. DE 39 28 684, 1989.
11. J. Orloff, "Focused Ion Beams," Sci. Am. . Vol.

10, pp. 96-101, 1991.

12. R. Hahn and C. Lost, "Sonoerosive Fertigung

keramischer Zahnrestaurationen," Dtsch.

Zahnarztl. Z.. Vol. 47, pp. 734-739, 1992.

13. R. Hahn and C. Lost, "Sonoerosive Fertigung

keramischer Einlagefullungen," Quintessenz, accepted for publication, 1992.

14. C. J. Patterson, A C. McLundie, D. R. Stirrups

and W. G. Taylor, "Refinishing of Porcelain

by Using a Refinishing Kit," J. Prosth. Dent..

Vol. 65, pp. 383-388, 1991.

15. S. E. Brackett, J. M. Leary, K. A. Turner and R.

D. Jordan, "An Evaluation of Porcelain

Strength and the Effect of Surface Treatment," J. Prosthet. Dent.. Vol. 61, pp. 446-

451, 1989.

16. W. Pfeiffer, W. Berweiler and T. Hollstein,

"Optimierung ausgewahlter keramischer

Hochtemperaturwerkstoffe bins ichtlich

Festigkeit, Ermiidung und Oberflachenzustand," IWM-Bericht. W 10/90, Fraunhofer

Inst, fiir Werkstoffinechanik, Freiburg, 1990.

17. Y. Sano, M. Fujisawa, K. Shinohara, T. Uetani,

H. Tadano, H. Horiguchi, M. Adachi,

N. Wakamatu, T. Goto and I. Sekine,

"Influence of Surface Properties on the

Bending Strength of Optec Porcelain," Gifu

Shika Gakkai Zasshi. Vol. 16, pp. 577-581,

1989.

18. G, Henzold, "Rauheitsmessung mit elektrischen Tastschnittgeraten," DIN Normenheft.

Vol. 12, Beuth, Berlin, 1971.

19. R.-J. Ahlers, "Optische Verfahren zur

Oberflachenpriifung," Z. Ind. Fertig.. Vol. 74,

pp. 465-468, 1984.

20. H. Siemer and H.-G. Wobker, "Randzonenanalyse an keramischen Werkstoffen,"

Sprechsaal. Vol. 123, pp. 282-293, 1990.

21. T. Hollstein, D. Meyer and W. Pfeiffer,

"FertigungseinfluB auf Festigkeit und Bruchverhalten von Hochleistungskeramik,"

Materialpriif. . Vol. 33, pp. 332-337, 1991.

22. D. Johnson-Walls, A G. Evans, D. B. Marshall

and M R. James, "Residual Stresses in

Machined Ceramic Surfaces," J. Am. Ceram.

Soc. Vol. 69, pp. 44-47, 1986.

23. N. Claussen, "Umwandlungsverstarkte keramische Werkstoffe," Z. Werkstofftech. . Vol.

13, pp. 138-147, 1982.

24. F. Wakkai, T. Iga and T. Nagano, "Effect of

Dispersion of ZrO^ Particles on Creep of FineGrained Al^O^," J. Ceram. Soc. Jpn. Inter.. Vol. 96, pp. 1176-1178, 1988.

25. K- Jakus, D. C. Coyne and J. E. Ritter, "Analysis of Fatigue Data for Lifetime Predictions

for Ceramic Materials," J. Mater. Sci.. Vol.

13, pp. 2071-2080, 1978.

26. H. Richter, "Langsame RiBausbreitung und

Lebensdauerbestimmung - Vergleich

zwischen Rechnung und Experiment," Ceram.

Forum Int. . Vol. 57, pp. 10-12, 1980.
Table 1: Surface quality of a variably treated feldspatic porcelain. Median ( x ), minimum (x^J

and maximum (x

of the measured roughness values (Rmax, Rz, Ra and Rk) in relation to the applied surface treatment.

Xmm Xmdx X

Ground Surfaces

Rmax 14.27 36.26 22.80

Rz 13.09 21.79 16.03

Ra 2.10 2.65 2.33

Rk 7.12 8.72 7.60

Lapped Surfaces

Rmax 0.20 0.47 0.32

Rz 0.17 0.28 0.22

Ra 0.018 0.022 0.019

Rk 0.062 0.067 0.064

Ultrasonically

Machined

Surfaces

Rmax 8.19 15.78 11.94

Rz 7.13 10.36 8.64

Ra 1.08 1.23 1.14

Rk 3.40 3.87 3.54

390
Controlled

Work Gap

Ceramic

Preform

Mounting

Fig. 1: Operation 1

Diagrammatic view of the first operation ultrasonically shaping an artificial ceramic tooth crown, as an example of a very involved geometrical component.

Fig. 2: Operation 2

By sinking the second sonotrode into the contralateral surface of the unilaterally machined preform, the actual

component is shaped (diagrammatic). In this operation the first sonotrode acts as a shaped seat for the ceramic

blank. The spatial relation of the sonotrodes is clearly defined by the exact alignement of the length axis and

the square position. aoi
t.=o

Sonotrode

iilMilir-i

Workpiece

2 A Amplitude max r

Gap: "Shaping"

Boron Carbide Slurry

Fig. 3a:

Schematic description of the ultrasonically machining operation. The transmission of the linear oscillation

to the master tool results in a physical stroke of the sonotrode of a few microns at a rate of 20 kHz. During

the actual shaping period, the tooling never touches the ceramic surface. The abrasive grains are propelled

indirectly by the mechanical motion of the oscillating tool and pummel the surface of the workpiece.

Elastic-Plastic Distorsion

of the Tool Surface

Boron Carbide Slurry

Microcracks

Fig. 3 b:

Because of the contrasting elastic properties of the tool surface and the ceramic, the abrasive is indirectly

hammered onto the machined surfaces.

3/4 X

Rounded Grains

Tool Wear

Fragmented Grains

Cavitation

Microchips

Fig. 3c:

Material is removed by erosion and grinding as they are created by the abrasive grains, the carrying fluid and

additional cavitation.

392
I

Implosion

Gap: "Finishing"

Rounded

Partly Fragmented Grains

Microchips

Fig. 3d:

Expanding the work gap between the sonotrode and the ceramic results in a increasingly indirect energy

adaption, refmishing the machined surfaces.

Bearing Surface (%)

Fig. 4: Surface density of a representative diamond ground porcelain specimen. The Abbott-Plot relates specifically

to the applied surface treatment. This ground sample showed a "total" surface of 100 % in a depth of -23.12

microns.

393
Profile

Depth

(l^m)

-9.35

80 100

Bearing Surface (%)

Fig. 5: ^ ^ ^

Surface density of a representative ultrasonically machined porcelain specimen. The almost symmetrical

coarse of the Abbott-Plot indicates a leveled microrelief with a maximum roughness of -9.35 |^m in this

specimen.

0

Profile

Bearing Surface (%)

Fig. 6: Surface density of a representative diamond lapped porcelain specimen (Abbott-Plot). The used

lapping procedure reproducibly maintained flat microsurfaces (-0.243 |i,m).

394
450

400

Gf Lf Uf Hf

Fig. 7: Box-Plot of bending strength of porcelain (Biodent Inlay Ceramic) and experimental alumina test bars in

relation to the applied surface treatment, before and after cyclic loading. The box displays the median as well

as the 25% and the 75% quartile as the belonging scattering value. The whiskers show the 10% and the

90 % quantile; values far outside are plotted as empty circles. For both critical bending strength and fatigue

behaviour, the ultrasonically machined specimens show the highest values.

G = ground porcelain

Gf = ground porcelain after cyclic loading

L = lapped porcelain

Lf = lapped porcelain after cyclic loading

U = ultrasonically machined porcelain

Uf = ultrasonically machined porcelain after cyclic loading

H = ultrasonically machined experimental alumina

Hf = ultrasonically machined experimental alumina after cyclic loading

395
n
A MODEL FOR HYDRO AND HYDRO-ABRASIVE JET

MACHINING OF CARBON FIBER REINFORCED

PLASTIC COMPOSITES

R. GROPPETTI and A. CATTANEO

Universita degli Studi di Perugia, Perugia, Italy

Hydro Jet Machining (HJM) and Hydro-Abrasive Jet Machining (HAJM) have

demonstrated high effectiveness and potential for multilayer Carbon Fibre Reinforced

Plastic (CFRP) composites. A deep understanding of the material removal mechanism

for machining is needed for prediction of productivity and quality, and for prediction and

avoidance of machining damage. The problem of the quaUty of the cut in the kerfing area of a CFRP workpiece is of great interest, because functional behaviour highly depends on the level of damage and defects along the edge of the workpiece, due to delamination, spalling and surface roughness.

This paper presents and discusses a theoretical global energetic model and its experimental vahdation, that is a contribution to the study of the mechanism of

delamination occurring during multilayer CFRP composite machining by means of HJM

and HAJM. A predictive model of edge damage caused by HJM and HAJM can be

useful in order to evaluate machining conditions, to predict the machining results in terms of kerfing conditions without time-consuming and expensive experiments.

1. INTRODUCTION

Among the machining processes for Carbon

Fibre Reinforced Plastic (CFRP) multilayer

composites, Hydro Jet Machining (HJM), or

briefly Water Jet, and Hydro-Abrasive Jet Machining (HAJM), or Abrasive Water Jet, have

demonstrated high effectiveness and potential.

Many efforts are devoted both in industry and in research institutions to study these processes in order to get a deep understanding of the material

removal mechanism, to predict machining results

in terms of productivity and quality, and to

predict and to avoid machining damage.

HJM and HAJM processes are based on a high

speed jet of water or water and abrasive particles,

focused, mixed and accelerated, to impinge on

the workpiece in a restricted area. The material

removal mechanism is complex and involves

erosion, abrasion, cavitation, deformation, and

fracture. The problem of the quaUty of the cut in the kerfing area of a CFRP workpiece is of great

interest, because functional behaviour highly

depends on the level of damage and defects along

the edge of the workpiece, due to delamination,

spalling and surface roughness. Both HJM and

HAJM offer virtually net shape and finished

workpieces without heat affected zone. But,

when machining conditions are not set properly,

delamination, spalling and low surface finish can

take place. While many results have been

published regarding structural analysis and

synthesis methods based on theoretical and

experimental mechanics of composites, rather

few studies analyse CFRP multilayer composite

machining problems.

Multilayer Carbon Fibre Reinforced Plastic

(CFRP) composites are widely used as structural

components because they offer high performance

and cost effectiveness. Many efforts have been

done in order to define design rules and

methodologies for an optimal design, in terms of

number and orientation of plies, their relative

397
positions, or the thickness of the boundary

between two plies, etc. [1, 3, 10, 17]. It has been

demonstrated also that the optimal design of the

CFRP structural components depends on many

factors and decisions, that have to be taken into

consideration concurrently, because they are highly coupled, particulary functional

requirements and factors relevant to their

manufacturing process. Therefore new

approaches, methodologies and techniques have

to be followed to execute this decision process,

as concurrent engineering, sensitivity analysis,

design of the experiments etc. [3, 9, 17, 21, 23].

Many results have been published regarding

structural analysis and synthesis methods based on theoretical and experimental mechanics of

composites. Rather few studies analyse

composite manufacturing problems, particulary

machining, piercing and cutting of composite

materials [10, 11, 15]. Because multilayer CFRP

composite products have to be finished and

trimmed by piercing and cutting, some

undesirable effects may occur as a low surface

finish or high roughness of the edge of the part, i

e. of the side of the kerf, delamination and

spalling [1,20,21,23]. The quality of the cut is particulary important for the behaviour of CFRP

structural components because delamination may

also take place due to edge effects, as stress concentration near the edge [3] when the part is in operation.

HJM and HAJM are processes, based on a high speed jet, for HJM a water jet only, for

HAJM a hydro abrasive jet formed by a stream of

abrasive particles focused by a secondary nozzle or focuser, mixed and accelerated by the high

speed hypersonic jet of water, obtained by the

output from a primary nozzle of water highly

pressurized usually up to 400 MPa. The hydro or

hydro-abrasive jet impinges on the workpiece in a

restricted area, giving origin to complex micromachining interactions with the material, as erosion, abrasion, cavitation and brittle fracture

of the material. The paper presents a theoretical

analysis and experiments of the mechanism of

delamination, occurring when CFRP multilayer

composite cutting and piercing by means of HJM

and HAJM. A delamination predictive model of

the dependence of delamination on the cutting

process variables, based on a global energetic

approach [4,10], is discussed and the main results

of this vahdation are analysed as a contribution to the study of the mechanism of delamination due

to the machining process. A suitable model of the

edge damage due to machining can be useful in order to predict and optimize kerfing conditions,

without time-consuming and experiments, when

cutting and piercing CFRP laminated composites.

2. HJM / HAJM DELAMINATION

MODELLING

When cutting and piercing CFRP by means of

HJM and HAJM the principal structural damage

is delamination [5,10,11,15]. Hydro and hydroabrasive jet stability perpendicular to its own axis

is limited and it can be disturbed within the

cutting kerf easily. The cutting surface shows the

characteristic bundles of parallel visible grooves,

which are bent opposed to the fibre direction.

The fibres embedded in the matrix act as obstacles, therefore HJM jet is deflected and

erodes the soft component of the material first. That leads to the typical appearance of washedout matrix protuding fibres [15]. This occurs mainly at the upper end of the cut. The jet erodes

the matrix material only, whereas the fibres

remain as solid links across the kerf. The main

action of the jet is in the direction of the fluid

flow and generally perpendicular to the surface.

The more the jet penetrates into the material, the

less is the force, due to the absorption of jet

energy during penetration, and delamination

appears at the exit side. Water pressure between

the layers is a cause of material damage of CFRP

with HJM. An increase of the pressure gives

origin to a higher feed rate and an improvement

of the surface quality. CFRP surfaces machined

by HAJM show a much higher uniformity than

workpieces machined at lower feed rate by HJM

and no washed-out single fibres had been

obtained [10,14].

Cut surface is usually smoother in the upper part

of the kerf than in lower part, that presents

waviness.

For the piercing process it is demonstrated

that at high pressures, depending on material

thickness, the jet loading close to the bottom of

the hole causes delamination in the bottom layers.

A failure analysis of HJM piercing of multilayer
composite has been performed to study the

11 process of delamination [5]. With the help of

energy balance, a simple analytic equation is

i

I proposed in order to choose the water pressure in

I! such a way that no delamination occurs. In the

following section this model will be summarized

and considered critically. The jet impinges on the

laminae and each layer is penetrated by the jet.

The model assumes that the jet thrust force acts

: in the centre of a circular plate and the single

laminae responds to the acting force with flexible

bending. During the kerfing process the uncut

material thickness becomes thinner, and the force

, distributes on a decreased number of laminae

' connections. A first delamination stage occurs when the laminae bending forces are exceeded by

the jet thrust force and the laminae starts to come

' off and a first crack arises. During a second

delamination stage, water penetrates into this

J J crack and due to water pressure crack

I propagates. According to this model,

delamination can be avoided if crack initiation is suppressed, and delamination occurs when the

strain exceeds the critical value of the laminae

bonding force at the crack initiation site or at the

' tip of the crack.

A critical water jet pressure can be

I determined, that should not be exceeded to avoid

r delamination, following an approach of linear

elastic fracture mechanics [19], based an energy

' balance equation. This equation relates the

; energy release rate with the thrust force of the

I

i jet, the displacement and the stored strain energy,

l! using the classic plate bending theory and the

assumption of a concentrated load on a circular

ij plate with clamped ends. The phenomenon can

'f be led to an erosion, loading and impact of the

impinging jet on the workpiece surface, with

matrix cracking and erosion into the laminae,

jl, delamination, fiber breakage and erosion.

I' However from the analysis of the experimental

results, considering the nature of this

|; phenomenon, the plastic energy dissipated in the

:| contact zone is a good parameter to measure the

degree of damage received by the laminae [20].

j Moreover fi"om this analysis it was noticed that

i delamination occurs mainly cutting by pure water

I jet. Therefore a semi-empirical predictive model,

J based on an energetic approach has been defined

! [10], in order to correlate the extent of the

delamination extent to the energy of the jet at a defined depth inside the kerf, corresponding to the bottom layer, where delamination takes place.

The istantaneous energy can be expressed as a function decreasing at the different depth inside

the multilayer composite and a proportional

relation can be assumed between the

instantaneous energy applied to the part and the

material volume removal rate, as demonstrated in [2] and [10]. Using the Levenberg-Marquardt

Method for the estimation of the parameters of

nonlinear models [18], the specific energy E, or

residual energy on the bottom ply, for unit length

and width equal to dj at the kerfing depth, can be

expressed by the following semi-empirical

equation:

The average extent Del of the delamination

damage, expressed into length unit, can be

evaluated along a direction orthogonal to the

kerf, i.e. to the feed of the cutting operation. The

specific energy, expressed by the previous

equation (2.1), can be correlated to this average

delamination damage extent Del A single

negative exponential relation between Del and E

can be assumed as follows, where p and y are

suitable parameters:

Del = yc (2.2).

This simple equation, that represent a delamination avoidance model suitable for

practical applications, has been justified with a good correlation factor by a regression analysis on experimental results, as it will be discussed in

the following section.

Equations (2.1) and (2.2) allow the definition

of the domain of HJM process variables that do

not arise delamination damage. While this energy

approach seems to have a general validity,

because experimental validation has been

executed on different materials, the validity of

equations (2.1) and (2.2), as semi-empirical

models, is restricted to the domain of the

experiments. The analysis of the results of the

I
experiments has demonstrated a good correlation

among the average delamination damage extent

and the main process variables, as water pressure,

feed rate, standoff distance of the nozzle from the

workpiece surface and the residual energy.

3. MODEL VALIDATION AND

DISCUSSION

Experiments were designed to understand the

mechanism of structural damage in HJM and

HAJM piercing and cutting of CFRP and vaMdate

the proposed delamination avoidance model. The

results of these experiments were correlated with

the process variables and analysed in order to

validate the proposed models.

3.1 Experimental Procedure

The experimental setup was based on an advanced HJM/HAJM cell installed at Dipartimento di Meccanica, Politecnico di Milano

[4, 8].

The test parts were made by a composite that

consists of cloth prepreg carbon fibre laminae,

which are stacked one upon the other and

embedded in a conventional epoxy matrix. The

carbon fibres are united in bundles and

crosswisely woven with each other within the

layer. The structure of the laminae is similar to

texile tissues at which the fibre boundles are interwoven rectangularly (90°). The fabrication

was performed while curing laminae and matrix

by using the autoclave method with standard

process conditions. The main specifications of the

test parts are summarized in Table 1. The

specimens were fabricated in rectangular plates

with a length of 200 mm and a width of 100 mm

For the various experiments three thicknesses were used, respectively 1.6, 2.2, 4.4 mm. The

surfaces are different on both sides of the

material: one side is rough and without

recognizable pattern, whereas the other one is smooth and the cloth weave design can be

detected. The rough side has been used as the jet

entrance side. When HAJM was investigated,

natural garnet (mesh 80) was chosen as abrasive

material.

The experiments have been designed for both

HJM and HAJM, cutting and piercing. To

warrant a time optimized execution of the

experiments, a meaningful cut and hole partition on the specimens as well as a time planning of the [ single experiments were adopted. Every edge of \ the specimen possesses 12 cuts. In Table 2 the ' process parameters adopted during the HJM

experiments have been reported. These values

; have been chosen to study the behaviour in limit

; conditions. During the experiments water i pressure, feed rate, standoff distance and primary

nozzle diameter were considered as parameters.

The specimens for the piercing experiments

has been prepared in a particular way. A

protection layer on the surface has been j necessary, in order to avoid the impact of the jet ; on the workpiece surface when the high pressure

j valve is actuated, which effects like a shock. The

arrangement for the piercing experiment was '

different from the cutting experiment ! arrangement: three different nozzle diameters i were used. j These experiments were executed with HJM, , because HAJM did not produce delamination

under standard experimental conditions.

3.2 Results and Discussion I The investigation of the delamination

behaviour for HJM is meaningful, in general, for

small material thicknesses, because for higher

material thicknesses the cutting ability of jets

without added abrasives is not sufficient. The

processed material is not cut through perfectly

and resulting surface is often very rough.

However, the knowledge about the possibility to avoid delamination damage, HJM offers

alternatives, compared to the relatively more

expensive HAJM

Beside delamination other damage arises

during the machining process which influence the

quaUty of the cut, e. g. spaUing formation or kerf

shape and conicity. Although all these damages

have been fully analysed, here only delamination

is reported. The area of delamination along the

cuts, :. e., where adjacent laminae separate from

each other, are recognizable due to a dent

formation on the bottom layer. As a rule the

visible delamination is greater at fibres bundles

crosswisely orientated to the cutting direction

than at bundles in parallel arrangement.

Therefore, the extent of the separated layers is not constant and fluctuates along the entire kerf
length. However, these observations exclusively

refer to the outer layer at the jet exit side. The

test-cuts demonstrate that delamination is always

connected with the appearance of spalhng. The

experiments have made clear that every cut

crosses three bundles of the bottom laminate

layer which are oriented perpendicularly to the

feed rate direction. At these places the extent of

the delamination is always wider than at bundles

that run parallel to the cutting direction.

Therefore, for each cut there are six maximum

delamination zones. Since for the usage of a machined workpiece only the largest extent of the

damage is crucial, these six areas of the largest

lamina separation have been measured. In order

to record a width of delamination, the

recognizable border of lamina separation has

been marked and, using a microscope with

Unescale, the width of delamination has been

measured. Due to the observed fluctuations of six

measured values for the extent of delamination, a medium width of delamination Del as the

arithmetic average values of these six

measurements has been defined. The

consideration of the complete cuts shows that

delamination never occurs between laminates at the jet entrance side. The cutting surface at the

upper zone of the kerf is rather smoother without

any protrusion of fibre or cavities. With growing

penetration depth of the jet, the cutting surface

becomes increasingly rougher

The measurement of the damage of piercing

takes place similar to the one of cutting. For the

maximum extent of delamination the maximum

radius of delamination was registered and for

spalling the largest distance between hole rim and

damage border. The technical execution and

devices used for the value recording were

generally the same as the ones used for the

cutting experiments.

The produced holes were generally of very

poor machining quality. Nearly every hole can be

insufficient concerning the chosen criteria for

delamination, spalling, and hole shape. Properly no piercing experiments exist that gives a

satisfactory and usable hole. The poor quality of

the pierced holes appears in the delamination

formation most clearly. For all pierced holes

damage in form of delamination occurs at a great

extent. The boundary between separated and

non-separated layers runs not circulary.

Delamination takes place at a rim of the hole,

whereas at the other no damage occurs. The maximum extent of damage was very large

for each hole and did not start for less than 2 mm.

All the delamination layers of the piercing tests were connected with the workpiece body. This

behaviour was also observed during the cutting

tests. Hence no tipped up separated material

layers appear, and it looks as if the width of crack

is rather narrow between the loose laminae.

A dependence of delamination on the

arrangement of the fibre bundles was not

observed. Therefore, it is not necessary for the

direction of maximum delamination radius to be

identical with the orientation of the fibre bundles,

which surround a hole.

The delamination problem is exclusively

confined to the specimen surface at the jet exit

side. Due to the small diameter, it is not possible

to cut open the hole. Therefore, an investigation

of the occurrence of delamination between the

laminae within the material cannot be executed.

Despite the great dispersion of the measured

values for the maximum delamination radius, two

fundamental tendencies can be observed: on one

side, for all three nozzle diameters an increase of

the delamination extent with growing working

pressure is to be observed. On the other side, the

maximum radius of delamination arises with a larger nozzle diameter.

An influence of the standoff distance on the

delamination behaviour was not observed in the

beginning for the executed piercing tests. The

great dispersion of the maximum delamination

radius does not give the possibility to observe a principle tendency.

Also, spalling and hole shape were analysed in order to validate the proposed model [5],

demonstrating a dispersion of the results. The

analysis demonstrated that the experimental

results concerning the influence of the nozzle

diameter on the delamination correspond to the

model in principle. Therefore, the critical pressure

decreases with increasing nozzle diameter so that

the delamination starts at lower pressures. In this way, the prerequisite is created that the laminate

separation increases with greater nozzle diameter

and constant pressures which are above the

critical value. This behaviour was observed in the
experimental results, but a discrepancy between

the experimental results and the theory was

observed, probably due to the fact that jet

diameter was assumed to be equal to nozzle

diameter. In fact considering the measured hole

diameter at the jet entrance side of the specimen

instead of the nozzle diameter, the predicted

values approach the experimental results in a

better way.

Systematical experiments were executed also for

HAJM, and demonstrated to remove the two

limitations of HJM, i. e., the maximum workpiece

thickness that can be cut and the need of a careful

selection of cutting conditions to avoid

delamination.

Suitable parameters were used for HAJM tests and the analysis was carried out in the same manner as for HJM. The results showed that aU

specimens were completely cut through for all parameter combinations. But in contrast to the

former HJM experiments, there is no

delamination damage for any of the cuts. A

limited spalling was observed. The cuts were very

clean and showed an equally smooth surface and

the width of the kerf presented a small variation

along the whole cut length. No delamination was

recognizable within the domain covered by the

experiments [10].

From the energetic point of view the HAJM

results are an accordance with the HJM results,

when at the higher energy level of the hydro jet,

delamination does not take place. However due

to the intrinsicly higer complexity and cost of

HAJM, when low thickness CFRP components

have to be cut, HJM is generally adopted if suitable cutting conditions can be used to avoid

delamination damage.

The results of the experiments, executed by

means of HJM, were analysed in order to validate

the proposed energetic model (eqs. (2.1) and

(2.2)) and to correlate the average delamination

extent damage with the main process variables, as water pressure, feed rate, standoff distance of the

nozzle from the workpiece surface, and with the

residual energy. Four main vsiriables were

adopted in the HJM cutting experiments on

specimens at three different thickness (see Table

2). An analysis of variance was executed using

the Yates Method [12]. Factorial experiments were designed with five factors at two levels

(A=water pressure, B=feed rate, C=specimen

thickness, D=standoff distance, E=nozzle

diameter). Table 3 shows the HJM delamination

experimental data in factorial form. The

experimental variance was estimated, assuming

that third order interactions among factors are neghgible or unimportant. A suitable significance

level was set equal to 2.5% . Table 4 shows the

Yates Method on HJM delamination

experimental data. A pressure increase from 270

MJPa (AO) to 350 MPa (Al) causes a decrease of

delamination in all the experimental conditions

(average value= 1.32 mm). A feed rate increase

from 500 mm/min (BO) to 1000 mm/min (Bl)

always causes a delamination increase (average

value= 0.52 mm). A specimen thickness increase

from 1.6 mm (CO) to 2.2 mm (CI) causes the

increase of the delamination (average value= 0.69

mm).

From this analysis all the effects were significant

for the single factors, except E, and the CD, AE,

CE, DE double interactions. An increase of

pressure reduces delamination, while an increase

of feed rate and thickness increases also

delamination. Double interactions suggest that

the effect of C is influenced and reduced by the

effect of D. It is interesting to observe that, even

if E alone is not significant, it has an influence on

A, C, D. The AE and ED interactions mean that

the influence of A and D factors is limited by a superior level of E, that enhances C factor. With

the remark that these results are restricted to the

domain of these experiments, it can be noticed

that pressure and feed rate have a great influence on the average delamination extent. This

conclusion can be seen as a justification of the

dependence of delamination on the absorbed

energy, that is mainly increases with the velocity

of the jet, i. e. with its kinetic energy, and with

the supply pressure, while this energy decreases

when feed rate increases.

A regression analysis of the experimental results

for specimen thickness equal to 1.6 mm

demonstrates, with a good correlation, that

delamination increases when pressure decreases

and feed rate increases: e.g. Fig. 1 shows the

dependence on feed rate of the average extent

Del of the delamination damage, expressed into

length unit, along a direction orthogonal to the

kerf, ai constant supply pressure. Also a multiple
regression analysis on the two variables pressure

and feed rate has confirmed this behaviour. The

analysis of the results on specimens of larger

thickness demonstrates that delamination

increases for the same process conditions.

Finally a systematic analysis was executed of

the experimental results in order to validate the

proposed energetic model. By means of the

Levenberg-Marquardt Method [18] the

parameters of the proposed model were

estimated. Therefore the residual energy E at the

bottom ply was evaluated (eq.(2.1)) and

correlated with the average extent Del of the

delamination damage, expressed into length unit,

along a direction orthogonal to the kerf, i.e. to the feed of the cutting operation. Assuming for

the kerf depth of the bottom ply the specimen

thickness reduced by the ply thickness, the

parameters of eq. (2.2) were estimated with good

correlation.

The regression analysis function curve of the Del

average delamination versus the E residual

energy is shown in Fig. 2. From the observation

of this behaviour it can be inferred that, at a low

energy level, bending stresses occur related to

flexural deformation causing delamination, and,

in order to avoid delamination, cutting conditions

have to assure a high residual energy level on the

bottom ply, in order to allow localized net fiber

breakage and matrix cracking due to high

transverse shear stresses without delamination

[10,20].

4. CONCLUSIONS

The residual energy of the impinging jet on the

bottom plies presents a fundamental role on the

delamination mechanism of multilayered Carbon

Fibre Reinforced Plastic (CFRP) composite

during a cutting process by means of Hydro Jet Machining (HJM). A semi-empirical delamination

avoidance model has been defined that identifies

suitable machining conditions that assure a

certain minimum energy level during operation.

While Hydro Abrasive Jet Machining (HAJM)

allows generally results without structural

damage to CFRP laminated workpieces with both

low and high thickness, the less expensive and

simpler Hydro Jet Machining (HJM) can be

applied advantageously to low thickness.

provided that these suitable machining

conditions are selected.

The energetic model, that has been proposed

and validated, can be applied in order to predict

and avoid delamination when HJM is adopted for

CFRP cutting in order to obtain composite parts

of high quality. This model not only provides a better understanding of this phenomenon but also

allows a concurrent engineering of CFRP

laminated composite structural components, and

it is suitable to be implemented in a CAD/CAE/CAM tool for multilayer CFRP

composite components.

ACKNOWLEDGEMENTS

This work was carried out with the funding of

the Italian MURST (Ministry of University and

Scientific and Technological Research ) and CNR

(National Research Council) and the support of

Istituto di Energetica, Universita^ degli Studi di Perugia and Dipartimento di Meccanica,

Politecnico di Milano. The authors are grateful to UHDE GmbH - Werk Hagen, Hagen (Germany),

SOITAAB s.a.s, Ronco Briantino, Milano (Italy),

NUOVA OSAI S.p.a., Ivrea (Italy), FTS

RACING s.r.l, Passignano sul Trasimeno,

Perugia (Italy) for their support. The

experimental facilities were partially supplied

within commodatum contracs between

Dipartimento di Meccanica, Politecnico di Milano, UHDE GmbH and SOITAAB s.a.s.

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No. 4, pp. 207-214, 1989.

NOMENCLATURE

m Water jet mass flow rate P Water pressure

V Impinging jet velocity

Vq Threshold velocity

dj Primary nozzle diameter

^max Maximum kerf depth for quasi-static

feed rate h Workpiece thickness

E Residual energy on the bottom ply u Feed rate Del Average delamination damage

(3, Y Delamination avoidance model parameters

Prepreg Weight fg/m^i 345

Elastic Modulus [GPal 63

Ultimate Strength fMPal 643

Elongation [%] 1.05

Fibre Content [%V1 52.4

Cured Ply Thickness fmml 0.23

Specimen Thicknesses [mm] 1.6 , 2.2 , 4.4

Table 1

- Test part CFRP specifications.

405
Water Pressure [MPal 150 , 190 ,230 ,270 ,310 ,350

Feed Rate fmm/minl 500 , 1000 , 2000 , 4000 , 6000 , 8000

Standoff Distance fmml 1,2,4,8

Nozzle Diameter fmm] 0.08,0.12,0.16,0.20

Specimen Thickness fmml 1.6,2.2,4.4

Table 2 - HJM experimental conditions.

EG El

CO Ct CO CI

AO A1 AO A1 AO A1 AO A1

DO

BO 1.32 0.1 2.76 0.5 0.78 0 2.72 1.82

B1 2 0.53 3.13 1.82 1.38 0.11 3.68 2.85

D1

80 1.8 0 1.88 0.2 0.7 0 1.25 0

81 2.06 0.58 2.15 0.52 1.18 0 1.76 0.35

Table 3 - HJM delamination damage experimental data in factorial form [mm]

(A=water pressure, B=feed rate, C=specimen thickness,

D=standoff distance, E=nozzle diameter).

406
Comb. Effetti Sigma Fsper 5% 2.5%

(1) 1.32 1.42 3.95 12.16 21,35 39.93 2.495625 49.8252 1158,304 1 1 a 0.1 2.53 8.21 9.19 18.58 -21.17 -1 ,323125 14.0053 325,5858 1 1 b 2 3,26 4,44 13.34 -12,85 8.27 0.516875 2.13728 49,68609 1 1 ab 0.53 4,95 4,75 5.24 -8.32 0.01 0.000625 3.1E-06 7.26E-05 0 0

c 2,76 1.8 2.27 -6.26 4.23 14.85 0.928125 6.89133 160,2052 1 1 ac 0.5 2.64 11.07 -6.59 4.04 -1,37 -0,085625 0.05865 1.363531 0 0

be 3.13 2.08 1.88 -3.78 1.07 1.99 0,124375 0.1Z375 2,876934 0 0

abc 1.82 2.67 3,36 -4,54 -1.06 1.81 0.113125 0.10238 2.380022 0 0

d 1.8 0.78 -2.69 2.8 4.57 -11.07 -0.691875 3,82953 89,02644 1 1 ad 0 1,49 -3.57 1.43 10.28 -1.09 -0.068125 0,03713 0,863131 0 0

bd 2,06 4,54 -3.28 2.7 -0.91 -2.73 -0.170625 0.2329 5.414384 0 0

abd 0,58 6,53 -3.31 1.34 -0.46 -0.55 -0.034375 0,00945 0.21976 0 0

cd 1,88 0,7 -2.05 0.7 0.33 -11.27 -0.704375 3,96915 92.27236 1 1 acd 0,2 1.18 -1.73 0.37 1,66 -0.25 -0.015625 0,00195 0.045405 0 0

bed 2.15 1.25 -1.88 -0.42 0.93 -1.73 -0.108125 0,09353 2.174283 0 0

acbd 0,52 2.11 -2.66 -0,64 0.88 -1.71 -0.106875 0,09138 2.124301 0 0 e 0,78 -1,22 1.11 4,26 -2,97 -2.77 -0.173125 0.23978 5.57421 0 0 ae 0 -1.47 1,69 0,31 -8,1 4.53 0.283125 0,64128 14.90803 1 1 be 1,38 -2.26 0.84 8.8 -0.33 -0.19 -0.011875 0,00113 0.026226 0 0

abe 0.11 -1.31 0.59 1.48 -0.76 -2.13 -0.133125 0.14178 3.295968 0 0 ce 2.72 -1.8 0.71 -0,88 -1.37 5.71 0.356875 1.01888 23.68623 1 1 ace 1.82 -1.48 1.99 -0.03 -1,36 0.45 0.028125 0.00633 0,147112 0 0

bee 3.68 -1.68 0.48 0.32 -0.33 1.33 0.083125 0.05528 1.285071 0 0 abce 2,85 -1.63 0.86 -0.78 -0.22 -0.05 -0.003125 7.8E-05 0.001816 0 0

de 0.7 -0.78 -0,25 0,58 -3,95 -5,13 -0.320625 0.8224 19,11871 1 1 ade 0 -1.27 0.95 -0,25 -7.32 -0,43 -0.026875 0.00578 0.134326 0 0

bde 1.18 -0.9 0.32 1,28 0.85 0.01 0.000625 3.1 E-06 7.26E-05 0 0

abde 0 -0.83 0.05 0.38 -1.1 0,11 0.006875 0.00038 0.00879 0 0

cde 1.25 -0,7 -0.49 1.2 -0.83 -3.37 -0.210625 0.3549 8.250563 1 0

acde 0 -1,18 0.07 -0,27 -0,9 -1.95 -0.121875 0.11883 2.762441 0 0

bcde 1.76 -1,25 -0.48 0.56 -1.47 -0.07 -0,004375 0.00015 0.00356 0 0

aebde 0.35 -1.41 -0.16 0.32 -0.24 1.23 0.076875 0.04728 1.099092 0 0

Table 4 - Yates Method on HJM delamination damage data.

407
0

P=310 MPa

u [m/min]

P=350 MPa

8

Fig. 1 - Dependence of delamination damage on feed rate and water pressure

(HJM, primary nozzle diameter=0.15 mm, CFRP thickness=1 .6 mm).

100 200

E [J/m2 io6]

300 400 500

Fig. 2 - Dependence of delamination damage on jet residual energy at kerf bottom

ply (HJM, primary nozzle diameter=.15 mm, CFRP thicknesses=1 .6,2.2 mm,

standoff distance=2 mm, water pressure=270/31 0/350 MPa.feed rate=500,

1 000,2000,3000,4000,6000,8000 mm/min).

408
EFFICIENT MECHANOCHEMICAL POLISHING FOR

SILICON NITRIDE CERAMICS

T. UEMATSU

Toyama Prefectural University, Toyama, Japan

K. SUZUKI and M. H. WU

Nippon Institute of Technology, Saitama, Japan

K. SUZUKI

Toyama Prefectural University, Toyoma, Japan

O. IMANAKA

Gunma Polytechnic College, Gunma, Japan

Instead of conventional loose abrasive mechanochemical polishing, a fixed

abrasive tool or a polishing disc for mechanochemical polishing was used here

for obtaining a high polishing efficiency and a better working environment.

In experiments, SUN^ disc workpieces were polished with a polishing disc

made of CrjO, on a vertical spindle surface grinding machine. By choosing

an appropriate polishing pressure and rather high rotating speeds of the tool

and the workpiece, initial surface roughness of Rmax=0.17!im(maximum peak

to valley height) was improved to Rmax=40nm only within 30 seconds. Even

for initial surface of Rmax^O.Tjim, surface finish of better than Rmax=50nm

was obtained in about 30 seconds.

1. INTRODUCTION

Adoption of ceramic parts into actual

products has been restricted because of poor

grindability of these materials. But the

application of ceramic parts has already started

because of the fact that mechanical properties,

especially the toughness, reliability, and

homogeneity of ceramics have improved, and

new grinding methods(l) have been developed.

Requirements of obtaining better surface

quality of ceramic materials for mechanical use

are increasing. But in many cases conventional grinding methods can not cope

with these requirements. In order to improve

surface finish, lapping or polishing with very

fine diamond abrasives is usually applied. But

it is not possible to obtain satisfactory surfaces

free from scratches and/or damaged layer.

Mechanochemical polishing method(MCP in

short, Fig.l)(2) utilizes mechanochemical

reaction between soft abrasive grains and

harder work material for providing smooth

surface finish. Accordingly, the method

extends the benefit of solving various problems

Direction of Movement

Hard Workpiece

Fig.1 Model of mechanochemical reactlon(2).

such as damaged work surface layer and

appearance of scratches found in lapping or polishing with hard diamond abrasives.

This method has been used as an important one in the field of precision finishing of

electro-magnetic and optical materials. But

there still remain problems to be solved such

as poor productivity, insufficient working

environment, and restriction in geometry of

engineering ceramic parts to be finished,

especially with complex shapes and of rather
large size. This paper describes a new

mechanochemical polishing method proposed

by the authors to solve these existing

problems.

2. PROBLEMS IN EXISTING MCP AND

THEIR SOLUTIONS

2.1 MCP with a Fixed Abrasive Tool

Problems in the existing MCP method are rather poor working environment, low finishing

efficiency and restriction in geometry to be

applied. These problems are caused mainly by

using loose abrasives in MCP. Some of these

problems can be solved by the development of

a fixed abrasive tool for MCP by T.Suga(3).

The polishing tool consists of chromic oxide

abrasives with a resinoid bond.

In order to apply the MCP method to ceramic components of rather large size, it is essential to develop a highly efficient MCP

method by investigating polishing conditions,

polishing tools and polishing machines(4).

2.2 Conditions for High Efficiency MCP

Figure 2 shows a schematic illustration for

high efficiency MCP of a silicon nitride

ceramic disc with a resinoid Cr^Oj polishing

disc on a specially designed machine. It is essential to find suitable conditions for

promoting mechanochemical reaction between

the tool and the workpiece, leading to high

removal rates. For example increasing the

working pressure and the relative peripheral

speed of the tool and the workpiece may

increase the temperature at micro-reaction zones and increase the polishing distance,

resulting in a high removal efficiency. But

an upper limit of working pressure is restricted

by the strength of the polishing disc. If the

rotational speed is too high, accumulation of

heat in the workpiece and the polishing disc

may damage both the tool and the workpiece.

3. EXPERIMENTAL APPARATUS AND

CONDITIONS

A vertical spindle surface finishing machine

was used instead of an existing lapping

machine to rotate a polishing disc and a disc

type workpiece with a rather high speed(Fig.3).

Workpiece

Fig. 2 Schematic MCP method for disc type

workpiece.

Fig. 3 Set-up for high efficiency MCP.

Table 1 Experimental devices and conditions.

Machine :Vertical surface finishing machine

(Air controlled constant pressure

polishing machine)

Polishing tool iCrjOj polishing disc(D=80mm,b=5mm)

Workpiece illjSijN^ disc(D=34mm,t=1 .5mm,Rmax=0.17-0.8A/m)

(2)SiC disc (D=34mm,t=2mm,Rmax- 0.6pm)

:Pre-grinding: Cup wheel(SD2500), Plunge grinding Measurement :Surface roughness profilometer(stylus radius=5pm)

Conditions ;Tool peripheral speed : Vs=0. 42-8.3 m/s

;Work peripheral speed: Vw=0.83-5.0 m/s

:Normal pressure : P =0.70-2.5 MPa

:Fluid : Chemical soluble type
A polishing disc with 80 mm diameter was

rotated at a peripheral speed of Vs=0.42 - 8.3

m/s. A sintered silicon nitride workpiece with

a diameter of 34 mm and a thickness of 2 mm

was rotated by a sub-spindle installed in the

machine bed at a peripheral speed of Vw=0.83

- 5.0 m/s. The maximum values of these

speeds are several times higher compared to existing lapping machines. The normal

pressure was varied from P=0.7 - 2.5 MPa,

this being rather high compared to existing

methods on conventional lapping machines.

Chemical solution type coolant was used to cool the polishing disc and the workpiece.

Table 1 lists the experimental devices and

conditions.

E

03 200

DC

E

cc

CO

ift Q)c

SI

C31 DO

tr

CDO

to t13 CO

160

!20

40

Disc(Cr2O3),Vs=8.3m/s,Vw=0.83m/s,t==30s. Work(Si3N4,D=34mm,t=1.5mm,Rmax-==170nm) • . ...

Rmax \

_ \

\

• • •

- Ra

1 'Q (i) 0.7 1.3 1.9 Normal Pressure P (MPa)

2.5

Fig. 4 Effect of normal pressure on surface finish.

4. EXPERIMENTAL RESULTS

4.1 MCP Conditions for Silicon Nitride Disc

Workpiece

Effects of normal pressure and rotational

speeds of the polishing disc and the workpiece on the improvement of surface finish were

investigated.

(1)Effect of Normal Pressure

Effect of normal pressure was investigated

under constant tool peripheral speed Vs and

work peripheral speed Vw at a polishing time

of t=30s (Fig.4). When low normal pressure,

such as P=0.7 MPa was adopted, surface finish

was not improved in a short polishing time of

t=30s. But rather high normal pressure, larger

than P=1.3 MPa was used, surface finish

improved from the initial roughness

Rmax=170nm to Rmax=55nm.

(2)Effect of Tool and Work Peripheral Speeds

Figure 5 shows a relation between tool

peripheral speed Vs and the surface roughness

attained at t=30s. Figure 6 shows a relation

between work peripheral speed Vw and the

surface roughness attain at t=30s. It is clear

from these diagrams that both Vs and Vw do

not influence the surface finish so much. But

the smoothest surface, having Rmax=25nm,

was obtained when a high work speed of

Vw=5.0 m/s and a low tool speed of

Vs=2.1m/s were adopted.

E

n3 cr 100

Disc(Cr2O3),t=30s, - Work(Si3N4,D=34mm,t=1 .5mm,Rmax - 1 70nm)

Vw=0.83m/s & P=2.5MPa

0 0.42 2.1 4.2 8.3 Tool Peripheral Speed Vs (m/s)

Fig. 5 Effect of tool peripheral speed.

Ea

cri

E

cc

if> tn0)c

x:O)o

QC

03 t

80

60

40

Disc(Cr203),P=

Work(Sl3N4,D=

1.3MPa,t=30s,

34mm,t=1 .5mm, Rmax " 170nm)

•^..^^^^Rmax Vs=8.3m/s ^ — —

A-

. o —

Vs=2.1m/s ^

Ra

A3 ZD

^

m

0 0.83 1.7 3.4 5.0 Work Peripheral Speed Vw (m/s)

Fig. 6 Effect of work peripheral speed.

411
(3)Effect of Polishing Time

The change of surface roughness as a function of polishing time was investigated for

three combinations of tool and work peripheral

speeds(Fig.7). Surface finish was improved

rapidly during initial 20 seconds, and the

initial surface of Rmax=170nm was improved

to Rmax=35 - 60 nm at t=60s.

(4) Effect of Pregrinding Conditions

Two workpieces having almost the same

surface roughness value(Rmax=0.7 - O.Sjim), one of which was pre-ground with a straight

wheel on a surface grinding machine and the

other was pre-ground by plunge grinding with

a cup wheel on a vertical spindle grinding

machine, were polished. In both cases surface

roughness was improved rapidly at the initial stage of MCP as shown in Figure 8. The

roughness of the surface pre-ground with a

straight wheel improved from Rmax=0.8!im to Rmax=60nm within t=90s, in spite of a rough

initial surface; but it required a long time to

n: 1, wci- vuu

:i> wc:l wcb wcl

WCf>

w

wc

\

[Initial, Rmax=0.17/im] [MCPed,Rmax=40nm]

Fig. 10 Mirror finished sample MCPed under

optimum conditions.

Disc(Cr203), Vs=2.1 m/s,Vw=5.0m/s,P=1 .3MPa,t=30s,

Work(Si3N^,D=34mm,t=1 .5mm,lnitial Rmax=0.17)unn)

200-

^§E

cc 150

-&100 eno

cr

50

Disc{Cr203),P=1.3MPa

Work(Si3N4,D=34mm,t= 1 .5mm,Rmax - 1 70nnn)

Vs/Vw(m/s) Rmax Ra

8.3 / 0.83 n

2.1 / 5.0 A

0.42/ 5.0 • O

1 10 20 30 40 50 60

Polishing Time t (s)

Fig. 7 Effect of polishing time.

I

0.8

cc cri 0-6 E

CC

in§ 0.4

.c

:3o

cr

0)o

to tr

CO 0

,^Disc(Cr203). Vs=2.1 m/s,Vw=5.0m/s,P=1 .SMPa

Work(Si3N4,D=34mni,t=1 .5mm)

ii

^\ [Pre-ground] Rmax 1 Ra

\\ Straight wheel

Hi (Reciprocationj_

\\ Cup wheel

-\\ (Plunge grind'g) - 1 \

• 1 O

1 A

-K

— • — n

0 30 60 90

Polishing Time t (s)

Fig. 8 Effect of pre-grinding conditions.

Polishing Time

>- 50pm

[t=0s]

(a)Pre-ground by reciprocation

grinding with a straight wheel

(b) Pre-ground by plunge grinding

with a cup wlieel

Fig. 9 Change of surface profiles for

different pre-ground workpieces.

412
remove scratches completely. On the other

hand the roughness of the surface pre-ground

by plunge grinding with a cup wheel became

less than Rmax=50nm within only t=30s.

Figure 9 shows the change of the surface

profile in both cases. It is found from the

Figures 8 and 9 that the rate of improvement

of surface roughness depends not on the initial surface roughness but on the initial profile of

the pre-ground surface.

4.2 Effect of Optimum Conditions

By adopting appropriate polishing conditions

such as Vs=2.1 m/s, Vw=5.0 m/s, and P=1.3

MPa, surface finish was improved within a very short time: for example, from

Rmax^O.lTjum to Rmax=40nm, only in 30

seconds as already shown in Figure 7. Figure

10 shows a mirror finished sample and surface

profiles of the initial and the finished surfaces

after 30 seconds of polishing. This result

indicates that a very high finishing efficiency can be realized by MCP, which has been

thought to be poor in polishing efficiency, if optimum finishing conditions and an

appropriate method are adopted.

4.3 Application of MCP Method to Silicon

Carbide Workpiece

The above-mentioned MCP method was

also applied to sintered silicon carbide

workpieces under almost the same conditions

as silicon nitride discs. Figure 11 shows that

surface finish was improved to Rmax=0.1|xm

from the initial surface roughness of

Rmax=0.6!im. It was found in the MCP of

silicon carbide discs that it takes longer,

attainable surface roughness is rougher, and

scratch marks in pre-grinding are left on the

finished surface. This indicates that finishing

of silicon carbide material by MCP is more

difficult than silicon nitride workpieces, and

more investigations concerning polishing

conditions in MCP, polishing disc materials

and so on should be conducted.

5. CONCLUSIONS

In order to solve problems in existing

mechanochemical polishing, a highly efficient

^ 0.8

0.6'

0.4

E

CC

COQ)c

D)

Z) IT 0.2

0)o

CO t

CO

Disc(Cr203), Vs=2.1 m/s,Vw=5.0m/s,P=2.5MPa

"Work(SiC,D=36mm,t=2mm, Initial Rmax=0.6/[;nn)

(T0 3 5

Polishing Time t (min)

Fig. 11 MCP for SiC workpiece.

MCP method has been developed. By

applying this method to flat portion of a

silicon nitride ceramic product, satisfactory

results were obtained. This method has

potential for use in finishing ceramic

components.

ACKNOWLEDGEMENTS

The authors would like to express their

heartfelt thanks to Messrs. Sodick Co.Ltd.

Nissei Ind. Co.Ltd. and Mitsutoyo Co.Ltd.

They also thank to Mr. K. Sakashita, J. Takagi

and K. Nomura for their assistances in the

experiment.

REFERENCES

(1)T. Nakagawa, K. Suzuki and T. Uematsu,

"Highly Efficient Grinding of Ceramics and

Hard Metals on Grinding Center,

" Annals of the

CIRP 35/1, pp.205-210,1986.

(2)N. Yasunaga, A. Ohara and O. Imanaka, J. of JSPE, Vol.44, p.77, 1978(in Japanese). ~

(3)T. Suga, S. Suzuki and K. Miyazawa,

"Mechanochemical Polishing of Sintered Silicon

Nitride," J. of JSPE, Vol.55, p.2247, 1989(in

Japanese)

(4)K. Suzuki, T. Uematsu, et al., "Development

of a New Mechanochemical Polishing Method

with a Polishing Film Method for Ceramic

Round Bars," Annals of the CIRP 42/1,

pp.339-342, 1992.

413

MAGNETIC ABRASIVE FINISHING OF WC-CO CURVED

SURFACES

M. ANZAI and E. KAWASHIMA

University of Tokyo, Tokyo, Japan

H. OTAKI

Nihon University, Narashino, Japan

T. NAKAGAWA

University of Tokyo, Tokyo, Japan

Summary

Most of mold manufacturing procedures have been automated by the

introduction of NC machine tools and CAD/CAM systems. But the surface

finishing of the mold must be still done manually by well-skilled

workers. Though many attempts were made to eliminate this hand work,

automatic finishing of mold surface with 3-dimensional curvature has

not been achieved yet. A new magnetic abrasive finishing process is

proposed as a potential method for the automation of surface

finishing. In order to improve the grindability of magnetic abrasive

process for WC-Co alloys, a new magnetic abrasive and finishing

procedure was developed. This magnetic abrasive and finishing procedure

was experimentally confirmed to have a sufficient grindability to

finish EDMed surface of WC-Co alloys into mirror surfaces. This

magnetic abrasive finishing may contribute to the automatic finishing

of WC-Co alloy.

1. Introduction

Generally, EDMed mold surfaces are

finished by skilled manual labor. In

this study, magnetic abrasive finishing

process was used to polish the surface

of WC-Co alloy. New types of magnetic

abrasives were developed by sintering,

and were confirmed to have remarkably

high grindability in comparison with

existing abrasives. [ 1 ] This study aims

at investigating the grindability of

these magnetic abrasives for a WC-Co

alloy and confirming the possibility of

automatic surface finishing using a NC

machine tool.

2. Experimental

2.1 Magnetic Abrasive Finishing Machine

The outline of the magnetic

abrasive finishing machine is shown in

Fig.l.[2] This set up includes an NC

415

milling machine,a steel rod(0.45%C),a

coil and a D.C. electric source. The

diameter of copper wire covered with

phenol resin is 1 mm and the number of

turns in the coil is 3500. The gap

between the rod and the surface of

workpiece can be easily changed. The

abrasives are attracted to the

magnetized rod. The workpiece is finished by the force of attraction

between the rod and the workpiece. As is evident from Fig. 1, the finishing set up

is so simple that this might be employed

on any kind of machine tools. Magnetic

abrasive finishing experiments were

performed with this machine using newly

developed magnetic abrasives. The

stock removal was checked by measuring

the difference in weights of unfinished

and finished workpieces. The surface
roughness and the line profiles were

checked by a surface profilometer, while

the surface properties were observed

with scanning electron microscope(SEM).

2.2 Magnetic abrasives

2.2.1 Finishing Method using Diamond

Paste and Magnetized Iron Powder

The method for the preparation of

magnetized iron powder + diamond paste

I KC Milling

I machine

Fig.1 Outline of aagnetic abrasive

finishing aaciiine and it's

f i n i sft i ng pricipie

Rod(magnet)

luagnetic force

Fig. 2 Principle of new finishing

Method using diamond paste and

aagnet i zed iron powder brush

which uses a brush and slurry abrasives

is shown in Fig. 2. [3] The iron powder

particles are attracted to the magnetic

rod and also the workpiece which is magnetized by this rod. Slurry abrasives

are held by the pores between iron

powder and other particles. The slurry

abrasives are attracted towards the

workpiece due the magnetic force which

results in polishing of the workpiece.

As the diamond paste has high

grindability in comparison with

conventional slurry abrasives, this

finishing method is suitable for

polishing of hard materials like WC-Co

alloy. But the abrasive life is short

and this is not suitable for polishing

of 3-dimensional shapes. Therefore iron

bonded diamond magnetic abrasives were

developed by sintering process.

2.2.2 Iron Bonded Diamond Magnetic

Abrasives Prepared by Sintering

Figure 3 shows the preparation

process for iron bonded diamond

magnetic abrasive by sintering method.

[4] At first, iron and graphite powders

are mixed in a stainless steel pot and

balls for 10 min.Then these powders and

diamond particles and 1 Wt% of zinc
stearate are mixed in the stainless

steel pot and balls for 10 min. These

raw materials for magnetic abrasives are formed into the deep cup with a wall

thickness of 2.5 mm under a compacting

pressure of 490 MPa. Then the green

compact is sintered at 1423K for 30 min

in NH3 cracking gases. Finally, sintered

product is pulverized by hammer and

whose powders are passed through

sieves.

3. Results and discussions

3.1 Fundamental Finishing Characteristics of Diamond paste + Magnetized

Iron Powder Brush

The effect of increasing revolution

of magnetic tool on the stock removal

using diamond paste and iron powder is shown in Fig. 4. Workpiece used for

polishing is WC-Co alloy (8-16%Co) which

could not be finished by conventional

slurry abrasive. The stock removal

increases lineally with the rotational

speed of the tool. This value is higher

than the value of stock removal using

conventional slurry abrasive on carbon

steel workpiece under similar finishing

conditions. [5]

The effect of finishing gap and

magnetic flux density on the stock

removal using diamond paste and

magnetized iron powder is shown in

Fig. 5. The stock removal increases with

200

J- 150

^100

t 50

0

Magnetic flux dens ity -0.657

^

Gi^p: 1 . Omm

y Finishing t i me : 15]n i d

I L I I I 1 0 1000 2000 3000

Rotational speed of rod { rpm)

Fig. 4 The effect of rotational speed of

rod on the stock reaoval / workpiece'-

tfC-Co al loy / Abras ive - i ron powder (2g)

+diaaond paste(0.7g)+stear ic acid(0.3g)

it I TOD powder (2 or 6 w ) +grapb i te\

(1 .0.1 .2wtl) J

I (^Milling for 20mTn^

it-dai aMODd(35 , 40. 45vol%)

Zinc stearate(lwt%

)

{^Milling for lOmin^

{Moulding (490MPa 1)

S i nter i ng(1150 'C , 30 win

in NH3 craking gas ^

I (Sinterd product^ L-== f (Pu 1 ver i z i Dg)

(^la ssification^

^Iron bonded diamond

magnetic abrasives

25

Fig. 3 Preparat ion of iron bonded

diaaond magnetic abras i ves by

sintering

the finishing gap, but the stock removal

under finishing gap of 1.2 mm is low. In

the finishing conditions of smaller gap,

the pores between iron powder and

particles which hold the diamond paste

decreases according to increasing of

magnetic flux density. On the other

hand, the stock removal decreases in the

'=^200

150

* 100

_ 2 750rpm

50

0

Finishing time: 15 win

0.6 0.8 1.0 1.2

(0. 75T)(0. 7TX0.65T) (0. 6T)

FiusbiD gap (be) S

Magnetic flux density

Fig. 5 The effect of gap and Magnetic

flux density on the stock removal

/workp iece- HC-Co al loy /Abrasive '- ironpowder(2g)+diaKond paste (0. 7g)+

stearic ac i d (0 . 3g)

417
finishing conditions of larger gap. This

is caused by weak force of attraction

between the iron powder and workpiece

with decrease in magnetic flux density.

The effect of diamond paste content

on the stock removal is shown in Fig. 6. The stock removal increases lineally

with the diamond paste content. However,

the stock removal may be saturated with

70 vol% of diamond paste content.

Figure 7 shows the relationship

between surface roughness(Rmax:maximum

surface roughness, Rz: 10 point average

surface roughness, Raicenter line average

surface roughness) and diamond paste

content. The surface roughness after

finishing is under Rmax=0.1|nm on the

workpiece of WC-Co alloy using diamond

paste including sub-micron diamond

particles. The best surface roughness

was obtained at about 50 to 65 vol% of

diamond paste. Original roughness

remained on the workpiece under 50 vol%

of diamond paste content. In the volume

content of over 65%, surface roughness

was not good which is caused by the

over polishing of workpiece surface.

The effect of feed speed of

workpiece on the surface roughness is shown in Fig. 8. The feed speed of 3 mm/min is suitable for giving an

acceptable surface roughness.

soo

Abrasive: iron powder 2g

stearic acid 0.3g+ ^-'^

^ 200

§ 150

100

50

Gap'- 1

Re volution '2750rpa

Finisiiing ti we -'1510 in

Magnetic flux dens i tj '• 0 . 8T

20 30 40 50 60 70

Diamond paste content (vol^)

Fig. 6 The effect of diaaond paste

content on the stock removal

/ workpi ece: ffC-Co alloy

20 30 40 50 60 70

Diamond paste content ( voll)

Fig. 7 The effect of diamond paste

content on the surface roughness/

Iforkp iece : WC-Co alloy/Abras i ve iron

powder (2g) -i-stear ic ac ld(0 . 3g) -t-d laaond

paste/Magnet Ic flux dens i tyO.8T/2750

rpa/Gap : 1 . Oaa/Fin ish ing t iae ' 15a in/

Or ig i na I roughness •' 4 /i aRaax. 3/i aRa

0.5/iaRa

^0. /

'^O

to.

"^0.2

rv

^

0.

0

Mriginal roughnes s ( ,a

Rmax : 4 . Rz: 3 , Ra : 0 . 5 • Rmax

MRz ^a

Vy/

70

60

50

40%

302,

20

10

0

418

0 13 5 10 15

Feed speed of worisp i ece

(mm/min)

Fig. 8 The effect of feed speed of

workpiece on surface roughness/

Iforkp iece : IfC-Co alloy/Abras i ve : Iron

powder (2g)+stear Ic acld(0. 3g)+diaaond

paste(0.4g)/Magnet ic flux dens ity:0.8T

/2750 T pa/Gap: 1 .Oaa

Figure 9 shows the SEM images of

surface of WC-Co alloy and it's surface

roughness profiles. (A) is as fabrication

and (B) is after finishing using diamond

paste and magnetized iron powder brush.

The finishing conditions are : *composition of abrasives = iron

powder(KIP300AS:2.0g) +diamond paste

(0.4g) + stearic acid(0.3g)

*finishing gap between rod and

workpiece = 1.0 mm
*rotational speed of rod = 2750 rpm

^magnetic flux density = 0.8 T

*finishing time = 15 min

The surface with the original roughness

of Rmax=4jum was finished in 15 min

operation to Rmax=under O.l/Ltm as is

evident from Fig. 9. Figure 10 shows the external

appearance of WC-Co alloy plate and

finished one. The finishing conditions are: *composition of abrasives = iron

powder(KIP300AS:2.0g )+diamond paste

(0.4g) + stearic acid(0.3g)

*finishing gap between rod and

workpiece = 1.0 mm

^rotational speed of rod = 2750 rpm

*feed speed of workpiece = 3 mm/min

*magnetic flux density = 0.8 T

^finishing time = 15 min

As is evident from Fig. 10, the surface of

finished sample shows the mirror

surface. Unfortunately the finishing

method has several weak points that are responsible for a short abrasive life

and it is too expensive. It is well

known that diamond has the highest

hardness among all the abrasives, but

when used in paste form the life is short. Therefore, iron bonded diamond

magnetic abrasives by sintering was

developed.

1 QHm X2 , 000 1 0 m

0 . 2 mm

F/g.9 SEM iaages of surface of IfC-Co

al loy and surface roughness profiles

/(A): as fabricated (B) -after finishing

Fig. 10 External appearance of tIC-Co

plate /(A) -as fabricated (B): after

finishing

3.2 Finishing Characteristics of Iron

Bonded Diamond Magnetic Abrasive

for WC-Co Curved Surfaces

The effect of finishing gap and

magnetic flux density on the stock

removal using iron boned diamond

magnetic abrasive, which consisted of 35

vol% diamond particles is shown in Fig.

11. The average diameter of diamond

particles is 11.5 jim. The workpiece of

WC-Co alloy machined by wire-cut EDM,

and whose radius curvature of concavoconvex was 40 or 80 mm.

The finishing conditions were : ^magnetic abrasives = iron bonded

diamond / grain size = 214-425 ium

*abrasive weight = 1.5 g

^rotational speed of rod = 1000 rpm

*feed speed of workpiece = 5 mm/min

^moving of finishing tool= stroke for

3 times with 1 mm pitch

*shape of rod top = radius curvature

8 mm with slit

The stock removal decreases lineally

with the finishing gap for any radius of

curvature. This is caused by the

decreasing magnetic flux density.

419
100

^80

^60

l40

20

0

% 40S con vex

O 40S concave

'M80S convex

'\J,80S concave

I

(0.86)

(0. 98)

(0.88)

(0. 90)

0.8 1.0

Finishing gap (. ^(1.04) (0.92)

0(1.28) (1.10)

M(1.08) (0.94)

0(1.12) (0.98)

Magnetic flux density (T)

Fig. 11 The effect of finishing gap and

Magnetic flux density on the stock

reaoval Workp iece - w i re-cut EDUed 9CCo al loy with curvature

The effect of increasing rotational

speed of rod on stock removal is shown

in Fig. 12. The finishing conditions of

this experiment were based on the one

of Fig. 11 except a finishing gap of 0.8

mm was used. The stock removal increases

linearly with the rotational speed at

any radius of curvature. The difference

in stock removal for different radius of

curvature may be due to the addition of

mechanical force which was generated by

the mass of abrasives closely attached

at the point of the smallest gap. This

problem can be solved by the moving

finishing tool(rod) to a normal line.

The effect of feed speed of

workpiece on the stock removal is shown

in Fig. 13.

The finishing conditions are : ^magnetic abrasives = iron bonded

diamond / grain size = 214-425 /im

*abrasive weight = 1.5 g » ^finishing gap = 0.8 mm

^rotational speed of rod = 1750 rpm

*moving of finishing tool = stroke for

3 times with 1 mm pitch

*shape of rod top = radius curvature

8 mm with slit

The stock removal decreases linearly

^140 40R convex

408 concave

120\-m80S convex

v^80R concave 2100

80

60

40

20

0

500 1000 1750

Rotational speed of rod (rpm)

Fig. 12 The effect of rotation I speed

on stock reaoval /workp iece- wi re-cut

EDMed ffC-Co a I loy with curvature

200

tie ^160

^120

:^ 40

0

%40S convex

C40R concave

M80R convex

[J80R concave

±

3 5 10

Feed speed of worlcpi ece (mB/'min)

Fig. 13 The effect of feed speed of workp

iece on the stock reaoval/workpiece: wire

-cut EDMed WC-Co a I loy with curvature

with the feed speed of workpiece for

any radius of curvature. This result is reasonable because of the finishing time

per unit area become short according to

increase of feed speed of workpiece.

The effect of rotational speed of

rod on the surface roughness is shown

in Fig. 14. - ^

The finishing conditions are : *magnetic abrasives = iron bonded

diamond / grain size = 214-425 am

*moving of finishing tool = stroke for

3 times with 1 mm pitch

*shape of rod top = radius curvature

8 mm with slit

*abrasive weight = 1.5 g

^finishing gap = 0.8 mm

*radius curvature of workpiece =40 mm

/ convex • -

420
'^0. 04V MSa

0 JL

500 1000 1750

Rotational speed of rod (rpm)

Fig. 14 The effect of rotation I speed of

rod on the surface roughness

workpiece: wire-cut EDUed WC-Co al loy

with curvature

The roughness value of Rmax = 0.5 (im

was obtained under finishing condition

of 1000 rpm. But this roughness value is worse than the value for using diamond

paste and magnetized iron brush. It is possible to solve this problem using

finer diamond particles in the magnetic

abrasives like a sub-micron order in the

diamond paste.

Figure 15 shows the external

appearance of s-shape WC-Co alloy which

was machined by wire-cut EDM and it's

roughness profiles. (A) is as machined

and (B) is after finishing with iron

bonded diamond magnetic abrasive using

NC program.

The finishing conditions are : *moving of finishing tool = stroke for

3 times with 1 mm feed pitch

^magnetic abrasives = iron bonded

diamond / grain size = 300-425 am

*radius curvature of workpiece = 100

mm / concavo-convex

*shape of rod top = radius curvature

8 mm with slit (Figure 16)

*finishing gap between rod and

workpiece = 1.4 mm

^average diameter of diamond particle

= 60-70 lum *diamond content in abrasive = 35 vol%

*feed speed of workpiece = 10 mm/min

*rotational speed of rod = 2750 rpm

^magnetic flux density = 0.8 T

*abrasive weight = 2.0 g

5um

0 . S"""

Fig. 15 External appearance of wi re-cut

EDUed s-shape HQ-Co al loy /(A) • as

Machined (B): after finishing

Fig. 16 External appearance of rod top

3.3 Polished Example of WC-Co Alloy

Punch

Figure 17 shows an outline of the

magnetic finishing apparatus for WC-Co

alloy punch using Sm-Co magnet and also

the external appearance of finished

sample. The diamond paste and iron

powder brush were used as abrasives

in the finishing of hard metal. It was

possible to finish the WC-Co alloy punch

to a surface finish under Rmax = O.lym.

However, magnetic abrasive finishing

is not suitable for polishing of nonmagnetic ceramics like a SiC in

comparison with WC-Co ally. Moreover,

this finishing method has geometric

limitations for corner of cavity or groove. But, this magnetic abrasive

finishing has the advantage of cost

reduction using conventional NC machine

tools and this method is suitable for

magnetic substances like a steel.

421
4. Conclusions

Based on the facts described above, the

authors can conclude that magnetic

abrasive finishing process may be

considered as one of the potential

methods for the flat or curved surface

finishing of wire-cut EDMed WC-Co alloy.

In this finishing, the most important

factors are development of magnetic

abrasives according to materials and

design of optimum magnetic circuit. The

magnetic abrasive finishing will be

used in practice in the field of

automation of WC-Co alloy mold or punch

surface finishing with the development

of a suitable magnetic abrasives and new

magnetic finishing process by the

authors.

Acknowledgment

The authors would like to express their

thanlvs to the considerable assistance

of DAIDO STEEL CO., LTD., MAKING MILLING

MACHINE CG.,LTD , BRGTHER INDUSTRIES

LTD. , KAWASAKI STEEL CG., LTD. and

STANLEY ELECTRIC CO., LTD. for supplying

experimental equipments and materials.

The authors would also like to express

their thanks to, Mr. Torn Sudo and

Koichi Masaki (STANLEY ELECTRIC CO.,

LTD., Ex-trusteed researcher of IIS,

University of Tokyo), Mr. Yoshinori

Ayabe (student of Nihon University), Mr.

Hiroyuki Noguchi (staff of Nakagawa lab,

IIS, University of Tokyo), Seizo Takahashi

(lecturer of Nihon University).

References

[1] T.Shinmura:Tool Engineer, 40, 10

(1988) 48

[2] M.Anzai,T.Sudo and T. Nakagawa:

Proceedings of the 2nd International

Conference on Die and Mould

Technology, (1992) 145

Chuck of fUC Sm-Co

milling machine magnet

Fig. 17 External appearance of WC-Co

al loy punch after finishing with Sm-Co

aagnet and coapund of diaaond paste and

iron powders

[3] M.Anzai, T.Sudo , H.Otaki , Takuya

Obana and T.NakagawaiJournal of the

Japan Society of Powder and Powder

Metallurgy, 39(1992) 510

[4] M.Anzai, Hisanori Otaki,Toru Sudo

and T. Nakagawa: Paper being for

publication in International Journal

of the Japan Society Precision

Engineering, (1992)

[5] M.Anzai,K. Masaki and T. Nakagawa:

Journal of the Japan Society of

Grinding Engineers, 33, 4( 1989) 34

422
TURNING AND MILLING

423

MACHINING CERAMIC MATERIALS FOR DENTAL

RESTORATIONS

E. D. REKOW

University of Maryland at Baltimore, Baltimore, MD

G. M. ZHANG

University of Maryland at College Park, College Park, MD

V. P. THOMPSON

University of Maryland at Baltimore, Baltimore, MD

Ceramic dental restorations must perform in a demanding environment including

dynamic loading with high-stress in a wet and corrosive environment. Patient

expectations for a restoration include both esthetics and long life. In this regard,

ceramics are appealing but their brittle nature is a significant limitation. All of the

fabrication techniques for creating dental restorations damage the ceramic. However,

the smoother the surface produced, the less likely it is damaged sufficiently to compromise its clinical performance. The objective of this investigation is to determine

the interaction between machining parameters and machining environments on surface

finish. Additionally, an overview of the oral environment and design specifications for

dental restorations as well as a brief review of clinical failures of current ceramic

restorations are presented.

Dentistry provides a demanding

environment in which to investigate ceramic

materials, their properties, and their

performance. Patient are no longer satisfied

with gold restorations, particularly for their

front teeth. Increasing numbers of patients are demanding esthetic restorations for both the

front and back teeth. In this regard, ceramics

are appealing for dental restorations because of

their esthetics as well as their strength, wear

resistance, chemical durability, and

biocompatibility. Unfortunately, the brittle

nature of current esthetic materials is a

significant limitation for widespread use in

dentistry. All fabrication techniques damage

the ceramic surfaces. In the wet and corrosive

environment of the mouth, microscopic damage

from fabrication can propagate, resulting in

clinical failure. For any given ceramic, the

smoother the surface produced, the less likely

it contains damage and the more probable it

will survive in the clinical application. This

paper presents results of investigations to

elucidate the influence of machining parameters

and machining environment on surface finish.

Additionally, it provides an outline of the oral

environment and design specifications for

dental restorations as well as a brief review of

clinical failures of current ceramic restorations.

Design Requirements for Dental Ceramics.

The oral environment to which these materials are subjected is demanding. The teeth (and any

restorations associated with them) are continually bathed in saliva - a composite of

secretions of a number of salivary and mucous

glands, the gingival crevice, and tluid

transudate from the oral mucosa. It contains

epithelial and blood cells, bacteria, and food

debris in water (1). The average fiow rate is approximately 0.3 ml/minute but this value

rises considerably during eating or salivation.

In a normal person, between 700 and 800 ml of

saliva is secreted daily.

Instantaneous temperature ranges of

65°C (150°F) degrees are common during the

425
course of an average meal (2). Thermal

conductivity and the coetTicient of thermal

expansion for restorative materials are important properties to be considered in

preserving the pulp of the tooth and minimizing

thermal changes in the restorative material

(which can cause tooth sensitivity and/or open

gaps between the restorative material and the

tooth which can induce decay). The pH

ranges typically experienced during a normal

day may range from 0.5 (when drinking soda,

for instance) to 8 (with some toothpastes) (3).

Forces generated during normal chewing

functions are typically 77 kgm (170 pounds)

but these vary markedly from one area of the

mouth to another. In the molars, the range is from 40-90 kgm (90-200 pounds) but only 9-25

kgm (20-55 pounds) in the incisors, the most

anterior teeth. Biting forces are generally

higher in males than in females and are usually

greater in young adults than in cliildren. A

force of 75-80 kgm (160-190 pounds) applied

to a tooth cusp with area of approximately

0.039 cnr (0.006 in-) generates a compressive

stress of 193 MPa (28,000 psi) (2).

Instantaneous forces may be substantially

higher than the values reported from these

studies.

Loading occurs on a cyclic basis with an average of 1400 cycles per day (4). For people

who grind their teeth, however, that value is 5-

10 times this amount per day.

Wear of restorative materials in the

mouth is an important and complex

phenomenon (5). Wear of a restorative

material must be measured against enamel, as well as against other restorative materials in

other teeth. Normal enamel to enamel wear is approximately 30-40 microns/year depending

upon the region of the mouth; it is highest for

posterior teeth (6,7). Excessive wear of a

restorative material may cause a tooth to supererupt causing interferences during

function. Supereruption also accelerates

breakdown of the restorative material. An

often overlooked aspect of wear is the wear

that occurs between adjacent teeth. This occurs because teeth are free to move independently

Parameter/Characteristic

Oral Environment: Temperature Range^

Load^

Cycles/day"" pH Range^

Tooth Properties:

Modulus of Elasticity

Values

65°C (150T)

70 kgm (170 lbs) 1400 0.5-8

\/fn]?ir ffipntin)^ 11 . y/\ ID'' L\j MPa ivir

i / X lu psi

\/f r\I ^1 r ( f^rf^ ma] ^

^.0 A lU Mra

0/ \ iu psi

L l\JU\Jl UUIIoi 1.U111L iVlUIol l^UCllUIl^ 140 Mra v..l,UUU psi)

iVl UJ di ^ CI ] d. 1 1 1 CX / "^^d MPn ^77 ^>ci^ ^-.-4 Mra ^jZ,jUU pSlJ StrengthMolar (dentin)- 305 MPa (44,200 psi) Molar (enamel)^ 79 MPa (37,800 psi) Coeff. of Linear Expansio'n^ 11.4 KxlO-*

Density-^ 2.9 gm.cm"-*

Specific Heat" 0.18 (cal/gm . C)

Conductivity-^ 0.0022 (cal.cm/ cm-. sec. C)

Diffusivity-' 0.0042 (cm-. sec' Color (Munsell scale) Hue- 7.5 YR - 2.7Y

Brightness^ 5.8/ to 8.5/ Chroma- 1.5-5.6

Restoration Requirements:

Surface Roughness'^

Fit^ Expected Life'"

Table 1.

<2 micrometers' < 60 microns

10-1- years

Oral Environmental Conditions

and Tooth Properties

during function. If excessive wear occurs between two teeth, the posterior teeth move

forward causing disruptive changes in the bite.

Because of their relatively high hardness

when compared to enamel, roughness in

ceramics can cause very high wear rates on any

opposing tooth structure. To minimize wear,

surface roughness must be less than 2 microns

(8). Wear effects are moderated, to some

degree, by saliva and adsorbed proteins.

However, this effect is not yet well understood.

Fit of any restoration is critical for its success. If a restoration projects over the edge

of the tooth, especially near the gumline, it will

entrap food debris. The by-products of the

breakdown of this debris by the microflora in

the mouth can result in both potential damage

to tooth supporting tissues through

inflammation and in decay of the tooth itself.

426
If the restoration does not adapt well creating

a gap between the tooth and the restoration,

leakage beneath the restoration is likely to occur. Again, food debris is likely to be

trapped or imbedded in the area of improper

fit. The breakdown of that debris will result in decay. For long term clinical success, the

adaptation at the edge between the tooth and

the restorative material must be to within 60

micrometers (0.002 inches) (9). If this

criterion is met, most ceramic restorations are expected to last at least 10 years (10).

Color is paramount for reestablishing

esthetics of a tooth that is to be restored. In general, the Munsell data for hue range is from

7. SYR to 2.7Y, brightness from 5.8/ to 8.5/,

and chroma from 1.5 to 5.6 (2)

Biocompatibility and toxicity are also

critical to the long term success of any

restorative materials to be used in the mouth.

Few if any dental materials are totally inert

from a physiological standpoint. Ideally the

material should be harmless to the pulp of the

tooth and to the soft tissues surrounding the

tooth. It should contain no toxic diffusible

substances that can be absorbed into the

circulatory system. It must be free of potential

sensitizing agents that could lead to allergic

reactions. And it must have no carcinogenic

potential. The American Dental Association

has developed a specification for evaluating the

toxicity of any materials to be used for dental

restorations (ANSI/ADA Document 41,

Recommended Standard Practices for Biologic

Evaluation of Dental Materials and Devices).

Table 1 summarizes the oral

environmental conditions and provides

information regarding the physical properties of

human teeth and restorative materials.

Clinical Performance of Dental Ceramics.

Dental ceramics have been around for decades,

provided excellent esthetics in the anterior

regions of the mouth. Restorations for

posterior regions of the mouth, fabricated from

the same materials, failed almost immediately.

Higher strength ceramics with aluminous cores

or cast ceramics were introduced to eliminate

the problem. Unfortunately, the clinical failure

rates remain high. This is especially true in the high stress posterior regions of the mouth.

Malamut et al. reported a failure rate of

2.9% in 616 restorations with up to 4 years of

service life, 13.5% in 369 restorations with up

to 7 years of service life (11). Christensen et

al. found failure rates as high as 80% for

multi-unit all ceramic bridges (12). Little data

is yet available on clinical failures of machined

ceramic restorations. However intracoronal

machine-ground glass ceramic and feldspathic

porcelain restorations are vulnerable to fracture

at the edges in 0.5% of the restorations as well

as bulk fracture in 1-2.5% of the restorations

within three years (13).

The primary clinical failure mode is fracture. The initiation site of the fracture is generally on the internal surface of the

restoration. The internal surface is fabricated

to fit a prepared (ground) tooth. When cast ceramics are used, the surfaces are sandblasted or ground to remove refractory materials. (The

external surface is then polished but the internal

surface is left unpolished to improve the

mechanical advantage of the adhesives which

hold the restoration to the tooth.) The

sandblasting and/or grinding processes

introduce defects and/or expose porosities

which function as nucleation sites for fracture.

Machining, while less widely used until now,

is likely to induce similar defects. The

microcracks initiated during fabrication, grow

to a critical size when exposed to the oral

environmental effects of repeating loading,

wear, and saliva. When this critical size is reached, fractures occur (14).

Fractography of failed cast all-ceramic

(Dicor) restorations suggests that failures are caused by surface defects located, primarily, on

their internal surface (15-18). Most of the

effects of loads delivered during damaging

surface contact (fabrication-related processes,

chewing, and wear) take place below the

ceramic surface. Micro-cracks extending

normal to the contact surface have been

427
identified as being the failure origin for both

structural ceramics (19,20) and for dental

ceramics (21). Finite element analyses have

demonstrated similar failure mechanisms (17).

Machining Ceramic Dental Restorations.

Recently, CAD/CAM systems have been

introduced to speed the production of dental

restorations (22). The primary fabrication

technology used by these systems is machining

(either grinding or milling). Dental

restorations have an extremely complex threedimensional shape. Milling rather than

grinding offers advantages for rapid shaping.

To be economically sound, machining time for a crown should be under 20 minutes. For

comparison, a typical crown for a posterior

tooth can be approximated as a cylinder with

diameter of 10 mm and height of approximately

5-7 mm. Both the internal and external

surfaces must be shaped. As discussed above,

not only is the speed of material removal

important, the resulting surface roughness and

degree of damage are critical to the success of

the restoration.

Ceramic materials currently used for

dental restorations are predominantly either

feldspathic or reinforced porcelains and glass

ceramics (23). Two new materials are the

primary focus of investigations for CAD/CAM

produced dental restorations: (1) Dicor-MGC

(Dentsply) which is a two-phase micaceous

glass and (2) Vita Porcelain which is a

feldspathic porcelain (24).

Unfortunately, machining of ceramics is generally application dependent (25). Almost

nothing has been reported addressing the effects

of milling parameters on surface finish for any

dental CAD/CAM ceramic materials. What

has been published has generally related to grinding (21,25).

The objective of this investigation was

to establish a preliminary empirical relationship

between surface finish and milling parameters

(spindle speed, depth of cut, and feed rate) and

the machining environment. The investigation

reported here utilizes the factorial design (26).

The results of this effort provide an important I basis for more in-depth investigation of a very

complex problem.

Method and Materials. To investigate the

effects of machining parameters on surface

finish, a 2^ factorial design was selected. This

approach uses orthogonal arrays to provide

insight into the relationship between surface

finish and machining parameters. The three

parameters under investigation were feed, depth

of cut, and cutting speed. Additionally,

different machining . environments were

investigated in a second 2^ factorial design.

Cylinders of Dicor MGC (Dentsply)

with diameter of 16 mm and length of 12 mm

were mounted on a metal base plate using an epoxy resin (figure 1). A Matsuura MC-510

Figure 1. Dicor MGC in Machining Fixture

machining center was used to perform the

milling operations. High-speed steel end mills were used for the machining because of their ease of availability and cost. Machining was

done with the cutting fluid introduced at the

cutting tool-Dicor interface. Two different

fluids were used: distilled water and distilledwater with a chemical additive. Machining

parameters (summarized in Table 2) were

selected based on previous machining

experience with Dicor MGC as follows: feed

rate: 10-20 mm/minute, depth of cut: 0.05-0.10

mm, spindle speed: 200-400 rpm. Surface

roughness was measured using a Taylor428
Test Number Feed Rate

(mm/min)

Depth of Cut

(nun) Spindle Speed(rpm)

Cutting Fluid

Type

1 10 . 0 0 . 05 200 water

2 20.0 0.05 200 water

3 10 . 0 0. 10 200 water

4 20.0 0. 10 200 water

5 10 . 0 0 . 05 400 water

6 20 . 0 0.05 400 water

7 10 . 0 0 . 10 400 water

8 20.0 0 . 10 400 water

9 10 . 0 0 . 05 200 chemical

10 20.0 0 . 05 200 chemical

11 10 . 0 0. 10 200 chemical

12 20.0 0. 10 200 chemical

13 10. 0 0 . 05 400 chemical

14 20.0 0. 05 400 chemical

15 10. 0 0. 10 400 chemical

16 20.0 0 . 10 400 chemical

Table 2. Machining Parameters

Horbson profiiomter (resolution 0.01 micron).

Surfaces were also qualitatively evaluated under

scanning electron microscoDv.

Tiie factorial design permits generation

of an empirical formula quantifying the relative

effects of the machining parameters - both

alone and in combination. For this

investigation, the parametric equation is given

by the following equation which considers the

linear effects:

R, = ME+ 1 /2 (FMEx, +DMEx2 + SSMExj +

IFDX1X2 + IFSSx 1X3 + IDSSX2X3 +IFDSX1X2X3)

where:

R, = Surface roughness

ME = Mean effect = average of all surface roughnesses

FME = Main effect of feed rate

Xi = Feed rate DME = Main effect of depth of cut

X2 = Depth of cut

SSME - Main effect of spindle speed

X3 = Spindle speed

IFD = Interaction between feed and

depth of cut

IFSS = Interaction between feed and

spindle speed

IDSS = Interaction between depth of

cut and spindle speed

IFDS = Interaction between feed, depth

of cut, and spindle speed

It should be noted that such equations are conditional on the choice of test setting and

cannot be extrapolated beyond the range of

tested parameters. Also, the parameters Xi, Xj,

and X3 are normalized with respect to the mean

and the range of parameter settings.

Results and Discussion. Results are summarized in table 3 and figures 2 and 3.

429
Test Number Measured Value Ra Avg

1 2 .71 2 80 2 . 63 2.7

2 5 . 64 5 42 5 .80 5.6

3 8 .36 8 41 7 . 86 8 . 1

4 4 .42 4 58 4 . 27 4 . 4

5 3 .92 3 70 3 81 3 . 8

6 6 . 14 6 28 5 98 6 . 1

7 1 .90 1 85 1 74 1.8

8 7 .93 7 30 7 61 7 . 6

9 0 . 97 102 1 08 1 . 0

10 4 . 35 4 . 21 4 06 4 . 2

11 3 . 16 2 . 90 3 04 3 . 0

12 2 .81 2 . 70 2 93 2 . 8

13 0 . 76 0. 79 0. 85 0.8

14 4 .40 4 . 05 4 . 22 4 . 2

15 3 .71 3 . 90 4 . 08 3 . 9

16 5 40 5. 82 5. 61 5.6

Table 3. Average Surface Roughness for

Dicor MGC

Figure 4 presents representative sets of profiles

taken from tests numbers 1 and 9 (feed rate 10.0 mm/min, depth of cut 0.05 mm, spindle

speed 200 rpm, and cutting fluids water and

water with a chemical additive respectively).

Figure 5 is a three-dimensional reconstruction

of a 30 micron by 30 micron section of

surfaces machined under identical conditions in

water and in water with a chemical additive.

0.1 -(+)

0.05 t-(-)

Figure 2.

0.1 -(+)

0.05 i-(-)

10.0 Feed (mm/nun) 20.0

Surface roughness for machining

in water only

10.0 Feeddnin/min) 20.0

Figure 3. Surface roughness for machining

in water with a chemical additive

ME FME DME SSME IFD IFSS IDSS IFDS

Water only 4.96 4.52 -1.07 -0.37 0.93 -1.78 0.28 1.43

Water plus 3.39 1.23 1.13 2.08 -0.73 -0.23 2.08 1.43

Table 4. Main and Interactive Effects

Main effects and interaction values for

the test conditions were determined using a

table of contrast coefficients (26) (table 4). For

samples machined in water spray only, surface

roughness (RJ ranged from 1.8 microns (test 7: feed = 10 mm/min, depth of cut=0.10 mm,

spindle speed =400 rpm) to 8. 1 microns (test 3: feed = 10 mm/min, depth of cut=0.10 mm,

spindle speed =200 rpm). At first glance, it would appear that spindle speed is the greatest

determinant in surface roughness. However,

when other comparisons where only the spindle

430
/V -v- s^.^^ H

Ra=1.2|im Ra=0.9|im

Ra=0.8nm

r

0

Ra=0.8|im — J

Ra=l.l|im Ra=0.7^im

Ra=0.9|im

(a) Distilled Water Only

Ra=0.8|im

(b) Chemical Additive Used

Figure 4. Representative Surface Roughness Profiles

431
|i.m

(a) Surface Topography (Distilled Water Only)

(b) Surface Topography (Chemical Additive Used)

Figure 5. 3-D Reconstruction of Surface

Topographies (feed rate 10

mm/niin, depth of cut 0.05 mm,

spindle speed 200 rpm)

speed is changed are investigated, the

differences in roughness are quite small: 1.1

micron between test 1 and test 5, 0.5 microns

between tests 2 and 6. For test 4 and 8,

however, surface roughness difference is 3.2

microns. Obviously, the relationship between

surface fmish and machining parameters is not

denendent upon only a single variable.

Calculating main and interactive effects

gives the values summarized in Table 3. Using

these variables, the equation for machining

Dicor MGC in a water-only environment

becomes:

= 4.96-f 2. 26x1-0.54x2-0. 19x3 -f-0.47x,X2-

0.89x,X3-l-0. 14x2X3 -f 0.72x1X2X3.

This empirical model provides an

estimate for surface roughness as a function of

machining parameters. The negative values of

the main effects associated with depth of cut

and spindle speed indicate that increasing these

parameters will reduce the surface roughness.

On the other hand, increasing the feed rate will

make the surface rougher. Small changes in

feed rate will impact the surface finish more

quickly than changes in either the depth of cut or feed rate. The numerical values associated

with two-factor interactions elucidate the

interplay between the three machining

parameters on surface tlnish. The interplay

between feed rate and spindle speed is far more

important than the interplay between the other

combinations of parameters. However, even

this interaction has less of an influence on

surface finish than feed rate. For samples machined in water with a chemical additive, surface roughness (RJ on

average was less than for water-only samples.

The average surface roughness for all tests with

water only was 4.96 microns vs 3.39 microns

with water plus a chemical additive.

With the additive, roughness ranged

from 0.8 microns (test 13: feed = 10 mm/min,

depth of cut=0.05 mm, spindle speed =4(X)

rpm) to 5.6 microns (test 16: feed =20

mm/min, depth of cut=0.10 mm, spindle

speed =400 ipm). Note that the parameter

yielding maximum and minimum surface

roughness are not the same as for maximum

and minimum roughness with the water-only

spray. This implies that changing the cutting

environment alters the material removal

mechanism. Other investigations indicate that,

with chemical additives, the most prominent

change is the ratio of the tangential cutting

force component to the feed force component

of the cutting force; the ratio changes

significantly, without any change in the

resultant cutting force^"*. Calculating main and interactive effects
gives the values summarized m table 3. Using

these variables, the equation for machining

Dicor MGC in water with a chemical additive

, becomes:

R, = 3.39+0.62x,+0.57x2+ 1. 04x3-0.37x1X2-

0.12x1X34-1.04x2X3+0.72x1X2X3.

It is interesting to note that with the

I chemical additive, the relative contributions of

the various parameters change dramatically.

For instance, with water only, feed rate has the

greatest impact of single factors on the surface

finish (FME=2.26). For water plus chemical

I additive, spindle speed makes the greatest

impact on finish (SSME= 1.04). Additionally,

for water only, the main effects of depth of cut

and spindle speed are significantly less than for

feed. With water plus additive, however, feed

and depth of cut make nearly equal

contributions, and they are much more nearly

equal to the effect of changes in spindle speed

\ than with water only.

j With water plus chemical, the major

interactions is between depth of cut and spindle

I speed (vs. feed rate and spindle speed for water

!i only). Another difference is that not only do

I the values of the main effects and interaction

I': coefficients change, the signs of some of them

|l reverse. For instance, in water only,

I increasing the depth of cut improves the surface

I (DME<0) whereas with water plus the

I chemical additive, increasing the depth of cut

j roughens the surface (DME>0). Similarly,

ii increasing the spindle speed will improve the

I surface in water; whereas, the same change

will make the surface rougher in water plus a

I chemical.

;i Contour maps were created from the

surfaces presented in figures 5a and 5b. The

jl average roughness was calculated. Then a map

was created for all areas of the surface with

surface roughness of +1.0 micron, +0.5

I micron, the average roughness (0.0), -0.5

micron, and -1.0 micron. Only two areas (both

! for the water-only environment) were found to have a profile roughness of +1.0 micron. In

figure 6, all surfaces with roughness of at least

+0.5 microns is shown as black areas. For a dental restoration, these represents the first points of contact with the opposing tooth or

restoration. The more peaks at a given height,

the less likely there will be wear; the forces of

occlusion during jaw funcdon will be more

evenly distributed. Obviously the environment

in which a dental restoration is machined can have a significant impact in the likelihood for

clinical success.

10 20 30 40 50 60 70 80 90 100

(a) Contour Map (Distilled Water Only)

10 20 30 40 50 60 70 30 90 100

(b) Contour Map (Chemical Additive Used)

Figure 6. Contour Maps of Surfaces ir Figure 5.

433
Summary. The factorial design technique

provides insight into the interplay between

machining parameters and surface finish within

the range of variables investigated for one

machinable dental ceramic. In a water-only

cutting environment, feed rate is the major

factor influencing surface finish. For water

with a chemical additive, spindle speed

becomes the most important single factor.

However, within the range of the parameters

tested, both feed rate and depth of cut will

make significant differences on surface finish

and the interaction between the depth of cut

and spindle speed is critical.

The empirical model provides a method

to select machining parameters which will

produce a required surface finish. It also

defines subsequent investigations. For wateronly machining, there is little need to investigate large numbers of samples created

with different spindle speeds since little if any

change is likely to result from even substantial

changes in this speed. Alternatively, for

machining with water plus the chemical

additive, major emphasis must be placed on

fully characterizing the influence of spindle

speed on surface finish. Additionally, the

interplay between depth of cut and spindle

speed is important and warrants further

investigation.

Further elucidation of the most

important coefficients of the interplay between

machining parameters and surface finish must

continue. Other ceramic materials and cutting

tools need to be investigated. The empirical

equation needs to be made more robust by

extending the range of variables. The impact on clinical performance needs to be established.

The long range objective of our investigations

is to develop an empirical equation powerful

enough to establish a basis to efficiently

establish machining parameters as a function of

material properties. This will provide a method to achieve a specified surface finish for an array of ceramics which could be used for

dental restorations.

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Surface Finish II . MBS 526, pp. 351-

378, 1979.

21. J.R. Kelly, H. Luthy, A. Gougoulakis,

R.L. Pober, W.H. Mormann,

"Machining Effects on Feldspathic

Porcelain and Glass Ceramics:

Fractographic Analysis", Proceedings.

International Symposium on Computer

Restorations . WH Morman (ed)

Quintessence, pp. 253-274, 1991.

22. E.D. Rekow, "Dental CAD/CAM

Systems - What is the State of the

Art?", JADA, Vol. 122, No. 13, pp.

42-48, 1991.

23. R van Noort, "Dental Materials: 1990

Literature Review, J Dent, 19, 327-351,

1991.

24. D.G. Grossman, "Structure and

Physical Properties of Dicor-MGC

Class-Ceramics",

Proceedings.Internation al Symposium

on Computer Restorations . W.H.

Mormann (ed), Quintessence, pp. 103-

117, 1991.

25. S. Jahanmir, L.K. Ives, A.W. Ruff,

M.B. Peterson, Ceramic Machining:

Assessment of Current Practice and

Research Needs in the United States . NIST Report SP-834, 1992.

26. G.E.P. Box, W.G. Hunter, J.S. Hunter, Statistics for Experimenters: An

Introduction to Design. Data Analysis,

and Model Building . J Wiley & Sons,

1978.

27. M. Braden, "Heat Conduction in

Normal Human Teeth" , Arch Oral Biol . Vol. 9, pp. 279-486, 1964.

Acknowledgement: Tliis study was supported by NIDR

Grant R29 DE08455 to the University of Maryland.

435

TURNING OF FIBER AND PARTICLE REINFORCED

ALUMINUM

K. WEINERT and D. BIERMANN

University of Dortmund, Dortmund, Germany

The main problem for the conventional machining of metal matrix composites (MMC) is the tool wear which is caused by hard and abrasive reinforcements. To find suitable cutting

tool materials, several aluminium matrix composites have been machined with different

tools: cemented carbides, coated cemented carbides, polycrystalline diamond (PCD) and

a CVD-diamond coating on cemented carbide and on Si3N4-ceramic. The analysis of tool wear has shown that there are some basic rules concerning wear rate and the relationship

between hardness and grain size of the reinforcement and the hard phases of the cutting

tool material. Cutting tool materials which are suitable for different aluminium matrix

composites have been tested with a variation of cutting parameters.

1 Introduction

In comparison to conventional materials metal

matrix composites have many advantages, for ex- ample excellent strength-to-weight ratios, increased stiffness and wear resistance. With continuing improvements in process technology

MMCs, particularly short fibre and particle reinforced aluminium alloys, have become cost competitive with conventional materials. But to ex- pand the application of these materials cutting and

finishing operations also have to be investigated.

For conventional machining like turning, milling,

drilling etc., extensive tool wear is the main

problem.

2 Test Materials and Cutting Tools

During the investigations several composites with

a matrix of aluminium (AlMg3, AlMgSilMn,

AlMglSiCu, AlSil2CuMgNi) and a reinforcement of SiC- and B^C-particles or S-AljOj-short

fibres, were machined (Table 1).

An important point for the machinability of

MMCs is to find suitable cutting tool materials.

Therefore cemented carbide with different grain

sizes and coatings, polycrystalline diamonds

(PCD) and a CVD-diamond coating on cemented

carbide and on Si3N4-ceramic have been investigated. For all turning tests a neutral tool holder with

inserts of the geometry TPUN 1603 (clearance

angle a = 11°, rake angle y = 0° and 6°, entering

angle x = 90°, nose radius r = 0.4 and 0.8 mm)

have been used. The turning tests have been car- ried out on a CNC-lathe. Figure 1 shows the hard- ness of the reinforcement materials and the cutting

tool materials. It shows that SiC and B4C particles are harder than all cutting tool materials excluding

PCD.

437
Table 1 : Test materials

Matrix Reinforcement —^—— . .

Volume fraction

[vol. %1

Particle size/

fibre

^

size

AlSil2CuMgNi S-AI2O3 20 d=3|jm

1=0,05-0, 1mm

AlMgSilMn SiC 20 d(50)=13|am

AlMglSiCu SiC 25 d(50)=21|jm

AlMg3 B,C 10 d(50)=13|jm

d: diameter; 1: length; d(50): mean grain size

o

c

O w

.EE

«

Cemented

carbide

Ceramic

Coating

PCD

Fibre

Particle

medium grain

fine grain

AijOs

AiaOa+TiC

AlON

TiAIN

Saffil

SIC

I B^C

KXI Spread o( hardness

given by different producers

2000 4000

Hardness HV

6000

Fig. 1: Hardness of cutdng and reinforcement materials

for the tool wear of a machining process depends on many factors: work material, machining operation, cutting conditions, tool geometry etc.. From

a tribological point of view figure 2 shows a schematic representation of the elements and parameters of a cutting process. Solid bodies are the

different cutting materials which have been used.

Counter bodies are the metal matrix composites.

The surrounding medium is air. But in the cutting zone and in the contact area of chip and tool

vacuum can be assumed (2). Interfacial medium

can be particles from the solid body or counter

body or particles of a chemical reaction between

solid body, counter body and surrounding

medium. Because of the thermal and mechanical

activation during the cutting process chemical

reactions are possible.

The wear process at the rake face and clearance

face of the tool can be classified as sliding wear. It can be two-body weai-, where the particles or short

fibres are fixed in the metal matrix of the workpiece. Also possible is three-body wear, where the

particles are trapped between the contact surfaces

of solid body and counter body (figure 3).

3 Tool Wear

Selection of the cutting tool material is fundamental in determining the extent and type of wear to be expexted. The wear processes can be divided

into four main wear mechanisms: adhesion, abrasion, surface fatigue and tribochemical reaction

(1). The relative importance of these mechanisms

The reasons for tool wear when machining metal

matrix composites are the direct contact between

the particles or fibres and the cutting edge and

their relative motion to the rake and clearance

face. Therefore, the hardness of the reinforcements is a dominant factor for tool wear.

438
countarbody

( workpiece

surrounding medium

(air, vacuum

oounterbody

chip

interfacial medium

'abraded particles'^

reaction particles

lubricant

body (tool Fig. 2: Schematic representation of a

cutting process

600

500

400

300

200

100

0

^AIMg3/10vol.%B4C

Cutting material: Cemented carbide

Cutting speed: = 1 00 m/min Feed rate: f = 0,1 mm/rev

Depth of cut: ap= 1 mm

Tool geometry

'<

1 Y 1 /. 1 r 11" 1 0" 1 90' i 0.8mm AIMgSi1Mn/20vol %SiC

AISi12CuMgNi/20vol.% S-AljOg

200 400 600 800

Cutting time

1000 1200 s 1400

Fig. 4: Tool wear when machining

reinforced aluminium

2-Body Abrasion 3-3ody Abrasion

Fig. 3: Two-body and three-body wear

(1)

3.1 Tool Wear of Cemented Carbides

Figure 4 shows the effect of cutting time on flank wear of a cemented carbide with medium grain

when turning different reinforced aluminium al- loys. The wear rate increases with hardness of the

reinforcement material. Although the volume

fraction of B4C-particles is only 10 vol.% the wear

rate is extremely high. When machining S-AUOj

short fibre reinforced aluminium the wear is relativly low although the matrix consist of an

eutective aluminium alloy.

The topographies of the worn cutting edges lead

to the conclusion that the main wear mechanism is abrasion (figure 5). But there is a great difference

between the tool wear when machining S-AUOj

reinforced aluminium or SiC and B4C reinforced

aluminium. In the first case the hardness of the

reinforcement is lower than the hardness of the

WC-grains of cemented carbides. Therefore the

S-AljO, fibres or fractions of them cannot abrade

the cemented carbide by microcutting. But free

WC-particles which are seperated by microcracking and fatigue from the cemented carbide

insert can act like a polishing compound between

the chip, workpiece and the cutting tool. In the case of SiC and B4C the hardness is higher than

the WC-grains of cemented carbides. Therefore

the cemented carbides are abraded by microcutting which is indicated by the grooves at the

clearance face. For the machining of 6- AI2O3 reinforced aluminium the wear resistance of fine

grained cemented carbide is improved against the

standard medium grain. The reason for this result

is the higher toughness of the fine grained

cemented carbide by which the resistance against

microcracking and fatigue can be improved. In the case of SiC or B4C reinforced aluminium the

influence of the grain size of cemented carbide on

the flank wear rate is reversed. The main wear mechanism is abrasion by microcutting.

Therefore, wear rate can be decreased by using

coarse grained cemented carbide (3).
AISi12Cu MgNi/

20 vol.%6-Al203

AIMgSiIMn/

20 vol. % Si C

Cutting material: Cemented carbide (medium grain) Cutting speed: Vc= 100 nrVmin Depth of cut: ap= 1 mm

Feed rate: f = 0,1 mm/rev

AIMg3/10vol %B4C

10|im

Tool geometry a Y X

r 11° 0° 90° 0,8mm

Fig. 5:

3.2

SEM photographs of the flank

wear of cemented carbides

used for cutting reinforced

aluminium

Too! Wear of TiAlN-Coated Cemented

Carbides

The investigation for machining reinforced aluminium alloys with different cutting tool materials

has shown that for turning SiC or B4C particle

reinforced aluminium neither ceramics nor the

different coatings on cemented carbide (AlON,

TiN, TiAlN) offer advantages in relation to un- coated cemented carbides (4). But for turning

5-AI2O3 short fibre reinforced aluminium TiAlNcoated cemented carbides are suitable. Corresponding to the results when turning S-AUO,

reinforced aluminium with uncoated cemented

carbides, the wear rate of the TiAlN-coated fine

grained cemented carbide is lower than for the

medium grained material (4, 5). Because of the

good results of TiAlN-coating for machining

S-AUO, reinforced aluminium additional ex- periments with various cutting parameters have

been carried out. As to the wear rate, the results of

these investigations are summarized in figure 6. The increase of feed rate and use of a coolant

have almost no influence on the tool wear. But

the wear rate increases rapidly with an increase in

cutting speed. For turning with cutting speeds up

to v^, = 250 m/min, the wear rate is relative low

and a cutting tool life of 15 minutes for a Hank

wear limit of VB -- 0.2 mm is reachable. Although

the tool life of TiAlN-coated cemented carbide at these cutting conditions is lower than for PCD, the

application of this coating can be suitable because

of the cost of PCD tools, particularly for

geometrically complex tools like drilling tools.

Tool geometry:

a 7

r

If 0° 9Clf 0,8mm

Matrix material: AISI12CuMgNi

Reinforcement: 20 Vol.% O-AI^Og

Cutting material: TiAIN coated

fine grain cemented carbide

Cutting time: t • 450 s Cutting speed: v^, • 100 m/min

Feed rate: f » 0,1 mm/rev

Deptti of cut: ap • 1 mm

0,1 0,2 mm 0,3

Flank wear VB

Fig. 6: Tool wear when machining

5-AI2O3 reinforced aluminium

with TiAlN-coated fine grained

cemented carbide

440
3.3 Tool wear of PCD

With regard to the wear rate the best results for

machining different reinforced aluminium alloys

have been achieved with PCD (5). Some modifications of PCD inserts, different grain sizes (PCD

010: 10 |nm; PCD 025: 25 |jm), polished insert and

a positive rake angle 7 = 6° have been tested to optimize the application of PCD tools for turning

operations. Figure 7 shows the development of

wear rate of different PCD tools for the

machining of SiC and B4C reinforced aluminium.

When turning SiC or B4C reinforced aluminium

the wear rate decreases with the application of coarse grained PCD 025. Also the positive rake

angle and the polishing of the PCD insert reduce

the wear rate. The lowest wear rate can be

achieved with PCD inserts including all above

mentioned modifications. The influence of

polishing and positive rake angle on tool wear

can be explained by the reduction of friction that

leads to lower cutting temperatures and cutting

forces. Therefore, the thermal and mechanical

load on the PCD tool is reduced.

Figures 8 and 9 show SEM photographs of the

worn cutting edges. The PCD inserts used for

cutting SiC reinforced aluminium exhibit rough

surfaces at the worn rake and clearance face. The

wear of the PCD 025 insert shows no grooves and

at the rake face of the PCD 010 tool there are only

weak grooves visible. After the machining of B4C

reinforced aluminium the topography of the wear

land shows grooves especially at the rake face.

The fact that tool wear is less when using coarse

grained PCD corresponds to the results for turning an aluminium alloy with 18% silicon and quartz

reinforced epoxy (6). Also an investigation for

grinding PCD tools has shown that the abrasion of

PCD increases with the decrease of the PCD grain

(7). Because there is only a small difference in

hardness and toughness between PCD (HO and

PCD 025 the explanation for the results cannot be

given by the properties of the PCDs. In the fol- lowing two mechanisms are proposed:

400

(jm

300

ffls 200

C 100

Matnx malenal: AJMgSiIMn

Reinforcement: 20 vol,% SiC-parlicle Cutling speed: = 500 m/min Feed rate: f = 0.1 mm/rev

Depth of cut: a = 1 mm ^ PCD 010 ^

PCD 025

J PCD 025

polished; y = 6°

200 400 600

Cutting time

800 1000

400

|jm

300

\ 200

a>

100

Matrix material: AIMg3

Reinforcement: 10 vol °o B.,C-particle

Tool geometry: « 1 y 1 X 1 r 11" 1 0" 1 90' 1 0.8mm PCD 010 —* PCD 025 ^ PCD 025 Xr^-.— polibheu. Y = b

200 400 600

Cutting time

800 s 1000

Fig. 7: Tool v/ear when turning SiC

and B4C particle reinforced

alumminium with different

PCD inserts

a) The wear mechanism corresponds to the

observed mechanism for cutting SiC and

B4C reinforced aluminium with cemented

carbide. The wear resistance of PCD 025

is improved compared to PCD 010 be- cause the grains are larger than the

reinforcement particles of the composites.

Therefore, the wear resistance of PCD 025

is improved against an abrasion wear by

microcutting. This argument is supported

by the results of the investigations about

the wear mechanisms when machining

unreinforced aluminium (AlSil2) with

PCD (8). In this case the main wear

441
b)

mechanism is abrasion which produces

grooves on the PCD tools. This is refered

to the formation of alumina at the tool

which could be hard enough to produce

grooving wear.

Because of the lower hardness of SiC and

B4C compared to PCD, formation of the

wear is produced by microcracking and

fatigue. Free PCD particles lead to grooves at the inserts. This is underlined

by the different dimensions of the grooves

depending on the PCD grain size

(compare figure 9).

PCD 010

a 7

r 11" 0° 90° 0,8mm

Matrix matenal: AIMgSilMn

Reinforcement; 20 vol.% SiC-particle Cutting material: PCD

Cutting speed: 500 m/min

Deptti of cut: Zp= 1 mm

Feed rate: f = 0,1 mm/rev

Fig. 8: SEM photographs of PCD

inserts used for cutting SiC

reinforced aluminium

PCD 010

PCD 025

0,1 mm

a Y

r ir 0° 90° 0,8mm

Matrix material: AIMg3

Reinforcement: 10 vol.7o B4C-particle

Cutting material: PCD

Cutting speed: = 500 m/min

Deptfi of cut: ap= 1 mm

Feed rate: f = 0,1 mm/rev

Fig. 9: SEM photographs of PCD

inserts used for cutting

reinforced aluminium

B4C

3.4 Too! Wear of CVD-Diamond Coated

Tools '

The development of CVD-diamond coatings for

cutting tools leads to promising applications for

machining extremely abrasive materials of low

structural coherence, such as carbide and ceramic

green compacts, but also plastics and certain types

of Al-Si alloys (9). Also machining of MMCs

could be an field of application for diamond

coated tools. Especially for geometrically

complex tools this coating could be a viable

alternative to PCD tools. Substrate materials

should contain carbide forming elements like

tungsten, tantalum, titanium and zirconium to support the formation of diamond crystals. Si3N4-

442
ceramic is suitable as substrate material for

diamond coating (10).

First investigations for the inset of CVD-diamond

coated cemented carbides and Si3N4-ceramics

have been carried out in cooperation with a producer of diamond coatings. Figure 10 shows a SEM photograph of the CVD-diamond coating on

Si3N4-ceramic substrate. Because of the crystallization during the coating process the surface of

the diamond coating is rough. Single diamond

crystals are visible.

2520 20KU X2/,000 . 10Hm WD39

Fig. 10: SEM photograph of a CVDdiamond coating on Si3N4-

ceramic

Compared to the machining of 5-AI2O3 or SiC

reinforced aluminium with uncoated cemented

carbide, tool wear is reduced by the CVD-diamond coating (figure 11). Particularly, the diamond coating on a Si3N4-ceramic substrate leads

to a very low tool wear when machining SiC

reinforced aluminium. For turning B4C reinforced

aluminium the diamond coating is unsuitable,

because the coating has been abraded after a few

seconds cutting time.

Figure 12 shows SEM photographs of the worn

diamond coated Si3N4-ceramic inserts used for

turning 5-AI2O3 and SiC reinforced aluminium.

The flank wear is very low for both inserts.

Although 5-AI2O3 reinforced aluminium produces

a lower wear rate with uncoated cemented

carbides than SiC reinforced aluminium, the re- sults for turning with diamond coated Si3N4-

ceramic are reversed. The flank wear at the worn

cutting edge after machining 5-AI2O3 reinforced

aluminium is very smoth and shows only

polishing of the diamond coating. But at the

transition of rake to clearance face there are some

cracks and particularly bursts at the rake face. At

the insert used for cutting SiC reinforced

aluminium there is almost no wear after a cutting

time of t = 11.5 min. Only a slight abrasion of

diamond crystals at the cutting edge is visible.

200 Mm

150

03^ 100

,T 50

Mainx material: AISi12CuMgNi

Reinforcement: 20 vol,% 6-AljClj Cutting speed: = 100 m/min Feed rate: f = 0,1 mm/rev

Depth of cut: ap= 1 mm

Cemented carbide

CVD-diamond/

SI, N. -ceramic _ ^ ^ ^* - X—" ^

250

Cutting time

500 750

400

(jm

300

\ 200

0)

I 100

UdXm matenal: AlfVlgSil l\yin Tool geometry:

Reinforcement: 20 vol °o SiC-particle <'

1 Y 1 X 1 r 11" 1 0" 1 90" 1 0,8mm

• Cemented carbide

CVD-diamond/ Cemented carbide

CVD-diamond/

S13N4 -ceramic

250

Cutting time 500 750

Fig. 11: Tool wear when machining

reinforced aluminium with

CVD-diammond coated tools

The results for turning MMCs with diamond

coated tools are very promising and there are still possibilities to improve the wear resistance of the

diamond coating, for example by increasing of the

coating thickness or by polishing the diamond

coating.

4A3
„. „ . -^^---^ -^ww - ----v^. _ > AISi12CuMgNi;

AIMgSiIMn/

20 vol,% SiC

50 urn

Cutting material: CVD-diamond/

Si3N^-ceramlc

Cutting speed: \i^= 100 m/min

Depth of cut: ap= 1 mm

Feed rate: f = 0.1 mm/rev

0 7

r IV 0'^ 90" 0.8mm

Fig. 12: SEM photographs of CVDdiamond coated Si3N4-ceramic

used for cutting reinforced

aluminium

wear decreases with an increase of toughness of

the cemented carbides.

Using a TiAlN-coating on cemented carbide tool wear can be reduced for cutting S-AKO, reinforced aluminium. When turning at cutting speeds

up to = 250 m/min wear rate is relative low.

Although, the tool life of the TiAlN-coated

cemented carbide is lower than for PCD, the use of this coating can be suitable because of the cost

of PCD tools, particularly for geometrically

complex tools like drilling tools.

For the application of PCD to machining SiC or B4C reinforced aluminium this study has shown

that tool wear rate decreases with increasing the

grain size of PCD. Also a positive rake angle and

the polishing of the PCD insert reduce the wear

rate. The lowest wear rate can be achieved with

PCD inserts which include all three modifications.

The results for turning MMCs with diamond

coated tools are very promising. Compared to the

machining of S-AljOj or SiC reinforced aluminium with uncoated cemented carbide, the tool wear

is seriously reduced by using CVD-diamond

coating, particularly the diamond coating on a Si3N4-ceramic substrate led to a very low tool wear when machining SiC reinforced aluminium.

4 Conclusions

This investigation has shown that different cutting

tool materials should be used for turning MMCs

depending on the hardness of the reinforcement.

When machining SiC or B4C particle reinforced

aluminium with cemented carbides, the tool wear

is produced by microcutting, because the

reinforcements are harder than the hardphases of

cemented carbide. To reduce the tool wear, the

grain size of the hardphases should be larger than

the size of the reinforcement. The tool wear is caused by microcracking and fatigue for turning

5-AI2O3 reinforced aluminium. Therefore, tool

References

(1) K.-H. Zum Gahr, "Microstructure and

Wear of Materials", Tribology Series, Vol.

10, Elsevier Amsterdam, 1987.

(2) K.-H. Kloos, "Technische VerschleiBvorgange in tribologischen Systemen und

deren Simulation", Reibung und Verschleifi bei metallischen und nichtmetallischen Werkstoffen, DGM Informationsgesellschaft Verlag, 1990, publisher Prof.

K.-H. Zum Gahr.

444
(3) D. Biermann, A. Fischer and D. Meister,

"VerschleiBerscheinungsformen bei der

spanenden Bearbeitung von Verbundwerkstoffen auf Aluminiumbasis", Tagungsband zu dem Fachgesprach "Neue Werkstoffe -Anforderungen an die Schneide" , Universitat Dortmund, 09-12 June 1992.

(4) L. Cronjager and D. Biermann, "Turning

of Metal Matrix Composites", Proceedings

of Euromat 91 , European Conference on Advanced Materials and Processes, University of Cambridge, 22-24 July 1991.

(5) L. Cronjager, D. Meister and D. Biermann, "Spanende Bearbeitung von faserund partikelverstarkten Leichtmetall-Verbundwerkstoffen", Metall, Vol. 46, No.

10, pp. 1010-1017, 1992.

(6) P.J. Heath and M.E. Aytacoglu, "Erzeugung und Beurteilung von Schneidkanten

an Syndite PKD-Werkzeugen", Diamant

Information M38, De Beers, Diisseldorf,

1985.

(7) I.M. Kenter, "EinfluB der ProzeBparameter

auf den Materialabtrag beim Schleifen von PKD", IDR, No.4, pp. 246-252, 1991.

(8) W. Konig and A. Bomcke, "VerschleiBmechanismen beim Zerspanen mit PKD",

IDR, No. 2, pp. 78-87, 1990.

(9) T. Leyendecker, O. Lemmer, A. Jiirgens

and S. Esser, "Industrial Application of

Crystalline Diamond Coated Tools", Surface and Coatings Technology, Elsevier

Sequoia, Lausanne, No. 48, pp. 253-260,

1991.

(10) E. Lugscheider and U. Muller, "Kristalline

Diamantschichten auf Schneidwerkzeugen", Ingenieur-Werkstoffe, Vol. 4, No.

7/8, pp.42-44, 1992.

TURNING OF CARBON AND GLASS FIBER HYBRID

CLOTH COMPOSITE MATERIAL

N. L. CUPINI

UNICAMP/FEM/DEF, Campinas, Brazil

J. R. FERREIRA

EFEI/IEM/DPR, Itajuba, Brazil

Experiments were carried out to study the performance of different tool materials, like ceramics, SIALON, whisker-reinforced ceramic, cemented carbide, cubic boron nitrite (CBN),

and diamond (PCD). They were tested in turning of carbon and glass fiber reinforced fenolic resin

composite. Two different workpiece shapes were used: laminated plate and t"ilament wound tube.

During the experiments, tool wear, surface roughness, and feed force were measured, under

several cutting lengths, cutting speed, and fiber directions in composite. Experimental results showed that, due to the great influence of tool wear on workpiece surface roughness, only

diamond tools are suitable for use in finish turning. In rough turning, the cemented carbide tools

showed to be the best solution. Although, it showed larger wear than diamond, the cost/benefit

ratio is much lower and special precautions are not necessary.

INTRODUCTION

The use of composite materials based on fiber

reinforced plastics (FRP) is increasing, specially in aeronautic, aerospace, and automotive industries (1). As a consequence, new processes must be developed

to ensure high precision and good surface quality

of the components (2). The FRP are mold mainly in plates, laminated

and sheet shapes. Because of this, drilling and

cutting with abrasive discs and diamond saws are more frequent among the large number of

machining operations that can be used (3). Only a few number of papers and reports on turning,

milling and grinding of FRP can be found. The

main goal of this work is to contribute with

alternative applications of FRP, particularly for the case of turning (4). This work presents some turning experiments

carried out in carbon and glass fiber reinforced

fenolic resin (CGFRR). Laminated plates and filament

wound tubes were used as samples, as shown in Figure 1. The experiments were carried out to evaluate tool wear, surface roughness, and feed force

during several cutting lengths (cutting times) for

different tool materials. The influence of the relative

fiber direction and cutting speed were also evaluated.

EXPERIMENTAL PROCEDURE

TOOLS, MATERIALS AND EQUIPMENT

The properties of the CGFRR tested were: carbon fiber (70%) and glass fiber (30%) hybrid

cloth; fenolic resin (35 to 40% weight); 1,4 to 1,5

g/cm' density.

Cemented carbide (ISO KIO), ceramic

(AKOj+ZrO.), ceramic (AljOj + TiC),

whisker-reinforced ceramic (ALOj-l-SiC), SIALON

(ALOj-l-SijNJ, cubic boron nitrite (CBN), and

diamond (PCD) tools were used. The carbide, CBN,

and diamond tools had the geometry ISO SNUN

120412. The geometry of the ceramic tools was SNGN 120712.

Experiments were conducted in a CNC lathe with 30 HP power. The feed force were measured

by a dynamometer. Tool wear and surface

roughness were evaluated using conventional

equipment, and a vacuum cleaner was used to remove

the chip from the environment.

EXPERIMENTAL METHODS

Two sets of experiments were carried out:

447
CLOTH DIRECTION

SECTION A-A

o) FILAMENT WOUND TUBE

CLOTH DIRECTION

1

135 SECTION A-A

b) MACHINED TUBE SAMPLE OF LAMINATED PLATE

Figure 1 - Workpiece

wounded cloth with cutting conditions larger than

those usually used, and laminated plate with usual

cutting conditions (Figure 1). In the first case, all tool materials were used

to evaluate tool wear, surface roughness, and feed

force for the following cutting conditions: cutting

speed 320 m/min, feed 0.1 mm/rev, and depth cut 1.0 mm. This set of experiments was used only

to classify the tool materials based on their cutting

behavior.

In the second case, diamond and cemented

carbide tools chosen in the first set of experiments were tested. The cemented carbide tool wear growth was observed for the following cutting conditions:

cutting speeds 25, 100, 200, 320, and 450 m/min,

feed 0.1 mm/rev, and depth of cut 1.0 mm. The

obtained surface roughness after diamond tool tests was observed for the following cutting conditions:

cutting speed 200 m/min, depth of cut 1.0 mm, and

three different feeds 0.1, 0.2, and 0.3 mm/rev.

RESULTS AND DISCUSSIONS

The tool wear growth for several cutting lengths

are shown in Figure 2 for the first set of experiments.

For all the ceramic tools, a very high tool wear rate was observed. The SIALON presented the highest wear rate and the whisker-reinforced ceramic the

lowest one. The CBN presented a better behavior

compared with ceramics and cemented carbides. On

the other hand, diamond showed to be better then

CBN. If the cost/benefit ratio is considered only

for rough turning, the cemented carbide shows to be

the best among all the tool materials tested. In finish turning operation, the behavior of the

tested tool materials varies as shown in Figure 3. As

448
2 4 6 8 10 12 i4 i6 18 20 22

CUTTING LENGTH Ic (m x 10^)

-B-CARBIDE K10 H- PCD^

CBN --SIALON

XAl203+Zr03 -f-AlgOj + TiC -A-AlgOj+SiC

Figure 2 - Tool flank wear against cutting length.

s = 320 m/min

f = 0.1 mm/ rev

d = 1 mm

Filament wound tube

10 12 14 16 18

CUTTING LENGTH 1^. { m x 10^)

20

6- CARBIDE K10 4- P CD HK-CBN SIALON

KAIgOj + ZrOg --AlgOs+TiC -A-Al2 03+SiC

Figure 3 - Surface roughness (Ra) against cutting

length.

449
the composite material presents fiber delamination

and groove formation because of resin presence in FRP, the larger amount of wear larger incidence of

fiber delamination and groove formation on the

FRP resin and larger the influence on surface

roughness. Because of this, the more wear resistant

is the tool, the better is its behavior related to surface roughness. The results presented in Figure 3,

shows that the best tool material in finish turning

operation is the PCD.

It was also observed for all tool materials,

that the average surface roughness (Ra) varies rather

linearly with the cutting length. This could be | explained by the fact that the abrasion is the only

tool wear mechanism present. The polished flank

surface, the absence of crater wear and, mainly the

absence of groove wear, denotes that the active wear : mechanism is mainly based on the high abrasivity of

the composite material (Figure 4a). For ductile

steels, where other wear mechanisms are present ! during machining, the resulting surface roughness

is normally influenced by the groove wear variation '

(Figure 4b).

FLANK

a) Composlie iiiaieiiai luiiiiiij:

FACE

b) Ductile steels (SAE/AISI 1045) turning.

Figure 4 - Mechanism of tool wear

450
Figure 5 shows how the feed force increases

with cutting length. For some tool materials the rate of increase of feed force is remarkable (SIALON and

ceramics). For PCD and CBN, the feed force

remains almost constant. On the other hand, it had

been shown in Figure 2 how the flank wear varies

with the cutting length. So, the variation of feed

force could be explained by increasing of the flank wear with cutting length, and the absence of crater wear as shown in Figure 4a. The feed force intensity for fresh tools,

considered for PCD, was found to be close to 50 N.

This level of feed force is in the range obtained for

steels (5). It means that, despite of the low density of

CGFRR, it require large forces for machining,

probably because of its good mechanical properties.

For the second set of experiments, during the

turning of laminated plate, the cemented carbide

evaluated, and the influence of fiber direction on

surface roughness for the case of PCD tools was

tool wear growth during several cutting speeds was

investigate.

Figure 6 shows how the cemented carbide tool wear vary with cutting speed. It can be observed that up to a cutting speed of 325 m/min, the flank wear

increases slowly. For higher cutting speed, at least for 450 m/min, the flank wear increases morerapidly,

because of lower tool wear heat resistence (5). In Figure 7, it is possible to see that the surface

roughness is maximum when the fibers have 45 and

225 angles with the cutting direction. This fact occurs

because, due to the relative fiber position, the

possibility of delamination increases as cutting time

elapses. Also in Figure 7, it can be seen that there

isa strong influence of feed, similar to turning

of conventional materials.

UJo

tro

la.

Ul

UJ

1 s = 320 m/min

f = 0-1 mm /rev

d = 1 mm

Filament wound tube

) n

1 / i—-/—

. ^

i

: *

H h —^—

2 4 6 8 10 12 14 16 18 20

CUTTING LENGTH Ij. (m X 10^)

S-CARBIDE K10 H-PCD CBN -- SIALON

X AlgOj + Zr O2 -VAl203 + TiC -A-AlgOj + SiC

Figure 5 - Feed force against cutting length.

451
2i£

jJ 0.2

0.0 I \ \ \ \ \ \ 25 75 125 175 22 5 275 325 375 425

CUTTING SPEED S (m/minl

» |(. = 300m +|(.=600m

Figure 6 - Flank wear against cutting speed.

0 45 90 135 180 225 270 315

FIBER DIRECTION ANGLE CX (DEGREE)

f

= 0.1 mm /rev 0f = 0.2 mm/rev 03 f = 0.3 mm/ rev

Figure 7 - Surface roughness (Ra) against fiber

direction angle.

452
CONCLUSIONS ACKNOWLEDGEMENTS

Based on the results, it can be concluded that:

- The use of CBN and ceramic tools in the turning

of composite materials is not feasible based on a

technical and economical point of view.

- The workpiece surface roughness increases with

the tool flank wear until the fiber delamination.

- A large scatter in surface roughness measurements was found, because of a great number of defects

caused by delamination and groove formation. The

origin of these is related to the relative direction

of fibers widi respect to the cutting directions.

- In finish turning operations, only the diamond tool can ensure a surface roughness in the range of 2.5

to 5.0 /xm (Ra).

- For rough turning with cemented carbide tools, the

optimum cutting speed is between 320 m/min and

450 m/min. For cutting speeds less then 320

m/min, the wear evoluton is not so different. For

cutting speed higher then 450 m/min, large values

of wear occurs very rapidly.

- In rough turning operations the use of carbide tools

showed to be the best solution. In spite of

presenting bigger wear when compared to diamond

tool, it results in a better cost/benefit ratio; and

when it is used, there is no necessity of special

care.

- Despite of low density of CGFRR, it requires

higher levels of cutting forces in the range of

cutting forces used for steels.

- Because of flank wear and absence of crater wear,

feed force increases remarkably with the wear

evolution. Only for PCD and CBN tools, that are very wear resistant, no remarkably variation of feed

force was observed.

The authors would like to thank the

CTA-Aerospace Technical Center, DIFER Industrial

Diamond, SANDVIK, SKF Tools, BRASIMET,

WINTER, FINEP, and the student F. Ueda, whose

work has helped the progress of the researsh

described in this paper.

REFERENCES

1. T. J. Reinhart, "Engineering Materials

Handbook, Composites", ASM

International . Vol.1, Ohio, 1987.

2. H. Takeshita, K. Wehara, "Cutting

Mechanism of Some Composite Materials",

Second International Metal Cutting

Conference. Tokio University, Japan, 1985.

3. S. K. Malhotra, "High Steel Tool Wear

Studies in Machining of Glass Fiber

Reinforced Plastics", Wear, vol. 132, 1989.

4. U. E. Wunsch, G. Spur, "Turning of

Fiber-Reinforced Plastics", Manufacturing

Review, vol. 1, June of 1988.

5. D. Ferraresi, "Fundamentos da Usinagem

dos Metais", Editora Edgard Bliicher . Sao

Paulo-Brasil, 1977.

6. Y. Hasegawa, "Characteristics of Tool

Wear in Cutting GFRP", 5th International

Conference on Production Engineering . Tokio, 1984.

7. A. Koplev, A. Lystrup, "The Cutting

Process, Chips, and Cutting Forces in Machining CFRP", Composites . Vol. 14-4,

1983.

453

LASER-ASSISTED HOT MACHINING OF CERAMICS AND

COMPOSITE MATERIALS

W. KONIG and A. K. ZABOKLICKI

Fraunhofer Institute of Production Technology, Aachen, Germany

Laser-assisted turning of ceramic contributes to the expansion of tiie range of methods

which can be used to machine ceramic materials. It offers a greater degree of flexibility

than grinding in terms of the geometry of functional surfaces of the part. Parts and tools which are subjected to high levels of wear are clad with hard alloys such as

stellite 6 for protection. The only options currently available to machine these hard alloys are grinding and electrical discharge machining (EDM). Laser-assisted milling thus offers a very promising complement and alternative.

The basic principles of laser- assisted hot machining processes will be explained in this

contribution. Experimental results obtained during operations involving laser-assisted tur- ning of ceramic (silicon-nitride), and laser-assisted miUing of composites material (hotwork tool steel/stellite 6) will be compared and discussed and potential areas of application will be highlighted.

Introduction

Hot machining technology is enjoying a renais- sance, thanks to progress in laser techniques. The

flame-, friction- and resistance-heating media were

used in this field at the beginning of the 20th century (1). Recently, tests have been carried out

using plasma jet heating (1, 2). One disadvantage

which these methods of heating have in common

is that the power density of the heat source is generally insufficient to heat parts of the workpiece adequately. If the heat source is required to provide a pre-set power level, the heat transfer area must be large enough to heat the zone in front of the tool cutting edge to the appropriate

temperature. The requirements for hot machining

are fully met by the laser beam. The power density

of the laser beam can reach values in excess of 10^

W/cm^. Additional benefits to be gained are the

handling flexibility offered by the C02-laser and

the Nd:YAG-laser which are equipped with a focusing-optical system and an optical fiber, re- spectively, and the ability to form the laser beam

with the use of appropriate focusing optics

(mirrors or lenses) onto different spot shapes of

the point being machined.

Initial investigations into laser- assisted turning of

various materials were undertaken a few years ago

(3, 4). These studies indicate that laser-assisted

turning of materials such as titanium- and nickelbase alloys decreases the cutting forces and the

tool wear by 20% - 50%.

The results of preliminary investigations into la- ser-assisted turning of ceramic materials such as silicon-nitride are also available (5, 6, 7, 13).

In milling operations, in contrast to turning, the

cutting edge is subjected to mechanical and thermal alternating stress. The highly dynamic stress is caused by the interrupted cut and by the changing

chip thickness during the cutting operation. Thermally induced plastification of the material, using

laser power, appears in reducing mechanical stress

in machining to offer decisive advantages (8, 9).
Basic Principle of Laser-Assisted Hot

Machining

Laser-assisted hot machining is carried out using a

laser beam to heat the workpiece material selectively, i.e. heating only the chip cross-section, during the cutting operation, Figure 1 .

Figure 1: Principle and fields of application of

laser-assisted hot machining

The feed movement of the workpiece causes a certain area to be continuously heated and subsequently removed by a tool cutting edge. The laser

beam energy absorbed by the workpiece surface is converted into thermal energy and results in rapid

heating and often, in localized melting of the outer

layer, resulting in a steep temperature gradient.

The temperature balance throughout the workpiece is effected by thermal conductivity. Knowledge of the local and temporal temperature progression at the surface and in the depth of cut is essential for the determination of the optimal cutting and, above all, of the laser beam parameters

for the material removal rate required. The formation of isothermal areas in the workpiece can be

determined using mathematical simulation models

containing the solution of the thermal conductivity

equation (4). The temperature of the surface produced and of the laser interaction point are of

practical significance for hot machining. The

process parameters such as laser beam intensity

and interaction time must be selected so as to ensure that the temperature at a defined depth of cut456

does not result in damage to the outer layer of the

functional surface produced. j The two laser types now in common industrial i use, CO2- and increasingly Nd:YAG-laser can be | used as heat sources. Depending on the kinematics of the cutting operation, the workpiece moves with the cutting or feed I rate in the same way as in, for example turning or

milling operations in relation to the stationary or | oscillating laser beam focal spot which has a defi-

' ned geometry on the surface of the workpiece in front of the tool cutting edge (Figure 1). The temperature of the chip cross section directly in ad- vance of the cutting edge is dependent on the vo- lume removed and the energy introduced via the

laser beam. The total cutting energy is converted , largely into heat, due mainly to plastic strain in the ^ shearing zone and to friction between the cutting i edge and chip (10). The increase in total tempera- ; ture in the area around the cutting edge, causes | changes in the physical-mechanical properties of | the material. I

Metallic alloys and the silicon-nitride ceramics li- j sted in Figure 2 are examples of materials which

are particularly difficult to cut. |

g 800

0

ODS: Ox

^Nl-Cr-Mo-ste

Ide Dispersion Str

el

engthend

Inconel 718

TlAieVI r

SljNi, -ceramics (HPSN)

^---^^^

ODS-alloy ^

stelllte 6

1 1 1 1 1 1

4"%^

1 1 1

1 0 200 'OO 600 800 1000 'C UOO

temperature T

Figure 2: High-temperature strength of materials which are difficult to machine,

(source: DIN, WL, AFNOR, Knotek)

However, tensile strength decreases with in- creasing temperature, which result in the plastification of the material, and consequently making

the material considerably easier to cut by hot
machining operations.

Hot machining technology has the following advantages:

decrease of cutting forces,

reduction of tool wear,

avoidance of tool breakage, use of inexpensive tool materials,

decreased dynamic excitation (chatter vibration) and

higher material removal rates.

Laser-Assisted Turning of Silicon-Nitride

Ceramics

Silicon-nitride ceramics, by virtue of their combination of favourable characteristics such as low

density, high level of hardness, good high-temperature strength at temperatures above 1000°C, as well as thermal and chemical resistance, are in- creasingly being used in mechanical applications.

Finish machining is necessary in most cases due to

distortion, contraction and limited formability in hot-pressing or hot-isostatic-pressing operations

carried out on silicon-nitride ceramics. Machining

operations with geometrically undefined cutting

edge, such as grinding or lapping are currently

applied on an industrial scale in the manufacture

of ceramic parts. These machining operations,

however, are very time-consuming and are limited

with respect to the range of geometries that can

be produced. Attempts are, therefore, under way

to improve other technologies such as laser-assisted turning to a sufficient degree to permit their use in the machining of ceramics, in order to obtain higher removal rates and to increase flexibility

in terms of part geometry.

Machining silicon-nitride ceramics at room temperature with a defined cutting edge involves high

tool wear; this process variant is consequently of

little pracfical significance (11). When due ac- count is taken of the temperature-dependent

strength behaviour of silicon-nitride ceramics,

material-oriented machining technology is, however, conceivable even for tools with a defined cutting edge (c.f. Figure 2). Since amorphous areas

exist at the grain boundaries of the crystalline structure of silicon-nitride (12), partial plastification of the material is to be expected; thereby

providing the conditions necessary for chip for- mation (7, 13).

Test Procedure

The turning tests were carried out on machinery

consisting of a two-carriage lathe and integrated

2.5 kW-C02- and 1.5 kW-Nd:YAG-laser beam

source. One of the tool carriages on the lathe guides the beam, the other holds the tool, Figure 3 .

Figure 3: Experimental set-up for laser-assisted

turning.

The guidance system consists of a deflection and

focusing optic. In contrast, the Nd:YAG-laser beam is guided to the machining point by a flexible

optical fibre. The positioning of the focused laser

beam directly in front of the cutting edge is variable in both cases. Two types of silicon-nitride were used for the turning tests: HPSN-ceramics (hot-pressed silicon-nitride)

with a sample diameter of 30 mm and

HIPSN-ceramics (hot-isostatic-pressed sili- con-nitride) with a sample diameter of 10 mm.

The cutting tool materials in the turning tests were two polycrystalline tool materials: polycristalline

diamond (PCD) and cubic born nitride (CBN)

with negative tool cutting edge geometry.

Results and Discussion

The results of the turning experiments carried out
on silicon-nitride ceramics will be presented and

discussed in the following.

Cutting Forces

The value of the cutting force components is an important measure by which to evaluate the behaviour of a material during a cutting operation.

The cutting forces in laser-assisted turning operations are influenced to a great extent not only by

process parameters such as cutting speed, feed

rate and depth of cut but also by the temperature

of the workpiece at the cutting zone.

material: SijN^-ceranlcs (HPSN) tool naterial: PCB

feed rate: 6 m/Hln

laser power: 1100 M

« 100O

800

600

Kaarlns point

200

o no

Fc : cuttino forcel

F, : teed force

Fu : passive force

I

i

itO 45 50 m/inin 60 cutting speed

1)0 Its 50 n/nln 60 cutting sDced Figure 4: Machining temperature and cutting

force components in dependence on

cutting speed.

Figure 4 shows surface temperature and cutting

force components cutting force/direction of

primary motion, F^: feed force/direction of feed

motion and Fp: passive force/perpendicular to the

directions both of primary motion and of feed

motion) as a function of cutting speed at constant

C02-laser power. The workpiece surface

temperature on the chamfer, at an angle of 90° in front of the tool cutting edge was measured by the

two-color-pyrometer (temperature measuring

range: 800-3000°C). The results shown are mean

values of the temperature at the beginning and at the end of the turning process. The machining

operation begins with a pre-heating phase during

which the workpiece is heated to a surface temperature of approximately 1100°C in the area

around the first tool contact. The turning process

begins once the initial temperature has been

reached.

As shown in Figure 4, the surface temperature ri- ses with increasing cutting speed. Both the laser

interaction time and the time in which the same

point of the workpiece surface is reached by the

focal point of the laser drop as cutting speed is increased. This results in higher workpiece temperatures in the machining zone. The temperature

gradient is likely to be steeper in the direction of

chip thickness.

The cutting forces diminish as a result of the rise in temperature (Figure 4). This change in force

(e.g. in the case of cutting force by approximately 50 %) is attributed to increased temperature

in the chip cross section and to sufficient softening

of the workpiece.

The second important process parameter in tur- ning operations is the feed rate. The influence

exerted by this parameter on the surface temperature and the cutting force components is shown in Figure 5.

materiel: Sl,N,-cermlc8 (HPSN) tool nateriel: PCB cutting speed: 60 n/nln

laser power: 1100 w

1M00

5 6 mn/Bin 8 feed rate

(| 5 6 nm/niln 8 feed rate

Figure 5: Machining temperature and cutting

force components in dependence on

feed rate.

At constant C02-laser power, the surface temperature is observed to fall with rising feed rates. Increasing of feed rate results in higher cutting

forces. This is attributed to insufficient temperature (<1200°C) throughout the chip cross section.

458
The dependence of cutting force components on

feed rate and cutting speed at constant Nd:YAGlaser power is shown in Figure 6.

material

tool material depth of cut laser power

SI3N1, - ceranlcs (HPSN) CBM

1 nn

700 M (NdrYAG- laser)

7 im/mln

feed rate Vf

Figure 6: The cutting force components in dependence on feed rate and cutting

speed.

Here too, cutting forces are observed to rise with

increasing feed rate and fall with higher cutting

speed.

It emerges too, that the power requirement when

a Nd:YAG-laser (laser power: 700 W, Figer 6) is used as a heat source is approximately 35 % lower

than when a C02-laser (laser power: 1100 W,

Figer 5) is used.

Cutting Tool Wear

The forces acting on the cutting edge and the

temperatures in the workpiece and cutting tool

generated by the two heat sources - the self-induced heat as a result of friction and strain at the

machining point and the external heating by the la- ser beam - cause various wear mechanisms.

The progression of flank wear on a PCD-tool cutting edge as a function of the turning length is shown in Figure 7. After turning length of 100 mm, the maximum

flank wear is 0.08 mm.

0.10

0.06

0.04

0.02

material SljNi, - ceramics (HPSN) tool material PCD cutting speed 50 m/inln feed rate 6 mm/mln

depth of cut 1 mm

laser power 1 - 1,1 kW (C02-laser)

20 W 60

turning len*t If

100

Figure 7: The flank wear of tool cutting edge in dependence on turning length.

Figure 8 shows SEM photos (scanning electron

microscope) of the face and cutting edges of a PCD-tool cutting edge.

tool material: PCD machining oarameters

1,= 20 mm

Vc= 50 m/mln

Vf= 5 mm/mln

1 mm

1100 H

Figure 8: Types of tool wear (PCD-tool cutting

edge).

The PCD-tool cutting edge demonstrates flank

wear of 0.07 mm after turning length of 20 mm.

The two detailed photos (at the bottom of Figure

8) illustrate the tool cutting edge wear as results

of thermal-mechanical stress during turning operation. New cutting edges are formed by a deposit

of ceramics material on the flank and the face cut- ting edge.

The tool wear of PCD and CBN by laser-assisted

turning of silicon-nitride ceramics can be classified

as: flank wear. 459
crater wear on the tool face,

rounding of the cutting edge.

Chip Formation

Figure 9 provides information concerning the

shape and type of chips removed during turning

operations carried out on silicon-nitride ceramics.

3m

lower side of chip (detail 1)

; upper side of chip (detail 2)

material

SlsNj-ceramlcs

(HPSN)

tool material

PCB

cutting speed: 50 m/min

feed rate: 0.01 nm/rev. depth of cut: 1 - 1 .5 im

laser power: 1 .35 kW

Figure 9: Shape and type of chips produced in laser- assisted turning of HPSN-ceramics.

The chip morphology illustrate, that plastic deformation controls the chip formation process in hot machining. The rod-shaped Si3N4-grains

loosened from their original cohesion by intercrystalline sliding are visible on the surfaces (Figure

9, detail 1). In detail 2, a segmented structure is visible indicating dynamics in the shearing plane.

Surface Quality

The amount of thermal energy involved in laserassisted turning raises the question as to the extent

of surface damage. The surfaces were, therefore,

evaluated using roughness measurements, SEMphotos and X-ray examination (for determination

of residual stress and micro-extension levels).

Figure 10 shows surface topology and surface

roughness. The surface turned demonstrates the

uniform structure which characterizes a visible

feed marks (Figure 10, top). The values of the

surface roughness measurement (R^: the

arithmetic mean of the departures of the profile

from the mean line, R^: the average height

difference between the five highet peaks and the

five lowest valleys within the sampling length, Rj:

the maximum peak to valley height of the profile

assessment length) are comparable with those

achieved in grinding processes (Figure 10,

bottom).

machined surface

toDOloqy of surface

machining parameters ^

material

SljNij-ceramlcs

(HIPSN)

tool material: PCB cutting speed: 75 ra/mln rate of feed: 6 ram/ml

depth of cut: 0,8-1,2 mm

laser power: 330-380 W

surface roughness

1,32

iii

\>

-1,67

0,0

2,i|7 \m Rt=3,09 pm Rg

1,25

=0,18 xm

Figure 10: Surface quality.

Diffraction analysis: n components of residual stress:

- laser-assisted turning: cross to the turning direction

- traverse grinding: parallel to the grinding direction penetration depth: 8

laser-assisted turning

tool material

(tepth of cut (ai,) feed rate (Vf

laser power (Pl)

CBN

1 m

6 nm/iin

330-380 W

traverse grinding

feed rate (Vfa)

nuifcer of revolutions (nw) Infeed/stoke (Og) Infeed (ae,get >

100 mm/mln

250 rev/ml n 0.02 flin

1 mm

500

MPa

300

200

100

0

naterlal: SljNn - ceranlcs (HIPSN) Ecowresslve residual stress width of diffraction line

X777 50 n/mln 75 m/mln 25 m/s 50 m/s cutting speed Vc

1.0 c

degree c

o

0.6 2

o

0.2 I

0.0

460

Figure 11: Comparison of the compressive resi- dual stress and the width of diffraction

line of machined surface in laser-Jis- sisted turning and grinding, (source:

FhG-IPT/IWM)

Recent investigations (14, 15) show that microextension visible only on X-ray images which is
caused by lattice faults arising in the course of the

machining operation, permits the state of damage

to be reliably characterised. Figure 11 shows the

compressive residual stress and the width of dif- fraction line shown on X-ray images (micro-extension is proportional to the width of difraction

line) of two surfaces which had been turned with

laser assistance and ground.

Both samples were prepared to obtain the same

levels of surface roughness. Lower compressive

residual stress and micro-extension values are obtained in laser-assisted machining operations

than in grinding operation. The difference between

these two operations could be due to laser heating, or annealing by laser heating. Additional tests are necessery to substantiate these results.

Machining Examples

Figure 12: Machining example (HPSN-ceramics).

Figure 13: Machining example (HIPSN-ceramics).

Figure 12 and Figure 13 show examples of the

machining of HPSN- and HIPSN-ceramics

(sintered, rough sample and machined part) with

SEM-photos of the machined surface.

Laser-Assisted Milling of Composite Materials

Functional surfaces are coated with hard iron,

nickel or cobalt (e.g. stellite) based alloys to en- hance the wear and corrosion characteristics of

tools and parts subjected to high stress levels. The

coating, which is carried out using surface welding or laser beam technology, is applied during

initial manufacture and repair operations. Such

composite materials, consisting of a simple b.'ise material and high quaUty coating are becoming

increasingly important as effective cost saving measures. The Co-base alloy (stellite 6) was used as the

work material for milling investigations. Appropriate flat samples were genereted by multi-layer

laser beam cladding up to a thickness of approximately 5 mm.

Hard alloys are usually machined in time-consuming and costly grinding or electro discharge

operations because of their poor machineability.

Their high levels of hardness, strong local variati- ons in hardness and structural inhomogeneity of

the layers and irregular coating pose a number of

problems. External heat supply via laser beam during turning or milling operations on these materials results in reduced material strength (c.f. Figure 2), thereby facilitating cutting (8).

Test Procedure

Milling tests were carried out using a plane-milling machine with a drive power of 8 kW and a maximum spindle speed of 4000 min'^ and an in- tegrated C09-laser. A C02-high-performance la- ser with a maximum output of 5 kW was used to generate the beam. As shown in Figure 14 the beam was deflected onto the workpiece and focused on the surface into an elliptical focal spot in front

of the milling tool via a focusing optic attached to the spindle box on the machine.

CBN (cubic boron nitride) was used as a tool
material in both the conventional and laser-assisted milling operations under examination. The

milling tool selected was a toroidal end mill with

negative cutting edge geometry and a diameter of

32 mm.

1

,

Figure 14: Experimental set-up for laser-assisted

milling.

In the case of laser-assisted machining, there is a significant reduction in force. This change in force

is due to a temperature increase in the entire chip cross section and the resultant material plastification. As a result of the high wear levels sustained

by CBN-tool cutting edges during conventional

machining operations (VB,^,;,^ = 1 mm after 70

mm machining length), the cutting force components rise steeply and finally cause cutting edge

failure. In comparison, significantly less tool wear

occurs in laser- assisted milling operations. The

flank wear of cutting edge (VB„^ax) ^"'Y ^-^^ mm.

Conclusions

To recapitulate, it can be said that the results of

the investigations presented, indicate that laser- assisted hot machining technology offers a number

of advantages.

Results and Discussion

The aim of the investigations was to determine the

degree to which the machinability of the material can be improved by laser-heating. The results obtained from the milling tests will be presented and

discussed in the following. Attention will focus on

the cutting forces and on tool wear in the evaluation of the cutting process.

The force components {¥(. feed force, feed

perpendicular force and Fp: passive force) are compared in Figure 15 for machining length of 70

mm.

too

00£ N cI ^

u

U

i-o

c

3

o 100

I—j lncreffl« of

'—' force for Bllllns

length of 711 ob

f777\

naterlal:

cutting tocl:

stelllte 5

32 m

1

tool aaterlal : CBN

nsQllning

caraiKters: - 100 n/Bln

f

J ' 0,1 nil 8p = 1 mm

a. = 10 ran

laser power: 3,0 KW

Without laser Hlth laser Figure 15: Comparison of the process forces in conventional and laser-assisted milling.

The investigations into laser-assisted turning of si- licon-nitride ceramics show that surfaces with

good grinding quality (R^ < 3 ^m and < 0.5

f,tm) can be obtained. The low cutting forces

permit parts with small diameters to be machined.

As the machining examples show, flexible part

geometry is an additional benefit to be gained

from this technology.

Preliminary investigations into the laser-assisted

milling of stellite 6 coatings have resulted in a re- duction by approximately 70 % in the cutting for- ces occurring during the process and by approximately 90 % in tool wear in comparison with conventional machining operations.

Additional investigations will be undertaken in or- der to evaluate the technological and economic

viability of laser-assisted hot machining. One of

the main objectives of these investigations will be

to determine the optimum process parameters for

the minimization of process forces and tool wear

as well as of the influence exerted on surface

layers.

462
References

1. R. Clausen, "Untersuchungen zur Bearbeitung metallischer Werkstoffe durch Warmzerspanen und Plasma-Abtragen", Dissertation

Technische Universitat Berlin (TU), 1969.

2. W. L. Weingaertner, "Plasmerunterstiitzte

Warmzerspanung", Dissertation RheinischWestfalische Technische Hochschule Aachen

(RWTHV 1983.

3 B.M. Jau, S.M. Copley, M. Bass, "Laser Assisted Machining", 9th North American Manufacturing Res. Conf. Proc , SME, Vol. 19-

22, Nr.5, pp.12-15, 1981.

4. B. M. Jau, "Laser Assisted Machining of

Hard to Machine Materials", Dissertation

University of Southern California, 1981.

5. W. Konig, A. Wagemann, H.-G. Mayrose,

"Laserunterstiitztes Drehen von heiBgepreBtem Siliziumnitrid", Industrie Anzeiger,

Vol.111, Nr.99, pp.34-36, 1989.

6. W. Konig, F. Ferlemann, C. Hermanns, A.

Zaboklicki, "Machining of New Materials

with Advanced Machining Processes - LaserAssisted Machining", 25th ISATA Conference, Florence-Italy, pp. 435-456, 1992.

7. W. Konig, A. Zaboklicki, "Laserunterstiitzte

Drehbearbeitung von Siliziumnitrid-Keramik", be published: VDI-Z, No. 6/July, 1993.

8. W. Konig, F. Treppe, A. ZabokJicki,

"Laserunterstiitztes Frasen", VDI-Z, Vol.

134, No. 2, pp.43-48, 1992.

9. A. Zaboklicki, "Laserunterstiitzte Zerspanung

von Verbundwerkstoffen", VDI-Berichte,

Nr.965.2, VDI-Verlag, 1992.

10. W. Konig, "Fertigungsverfahren, Band 1: Drehen, Frasen, Bohren", VDI-Vedag, 1990.

11. H. Kiso, T. Taguchi, M. Fukuhara, T.

Kimuro, "Machining of Advanced Ceramics

by Turning with Pol ycrystal line Diamond

Tool", Bull. Japan Soc. of Prec. Eng., Vol.21,

Nr.2, pp. 142-143, 1987.

12. H. Salmang, H. Scholze, "Keramik",

Springer-Verlag, 1983.

13. K. Uehara; H. Takeshita, "Cutting Ceramics

with Technique of Hot Machining", Annals of

the CIRP, Vol.35, Nr.l, pp.55-58, 1986.

14. W. Pfeifer: "Influence of Machining-induced

Residual Stresses on the Strength of Ceramics", Proceed, of the 9th Intern. Conf. on

Exper. Mech., Vol.3, pp.1237-1245, 1990.

15. W. Pfeifer, E. Sommer: "Bewertung der Festigkeitseigenschaften oberflachenbearbeiteter

Keramikbauteile", DGM Informationsgesellschaft mbH, Oberursel 1991.

STUDY OF THE FORMATION OF MACRO- AND MICROCRACKS DURING MACHINING OF CERAMICS

G. M. ZHANG, D. K. ANAND, S. GHOSH, and W. F. KO

University of Maryland, College Park, MD

This paper presents an experimental study on the formation of macro- and micro-cracks

during machining of ceramic materials. In this study. Aluminum oxide (AI2O3) was the

test material and polycrystalline diamond tipped carbide inserts were used. The cutting

forces were recorded during each turning process and the surface finish was measured

after machining. An environmental SEM was used to obtain high-magnification images

of macro- and micro-cracks induced by machining. With the assistance of a computerbased vision system, quantification of fracture surfaces with respect to crack nucleation,

growth, and cleavage was attempted. Results from this research provide an insight into

the prevailing mechanisms of material removal during machining of ceramics, and

suggest the development of microcrack-controlled machining technologies.

1. Introduction

The need for high-strength materials has

led to the development of advanced ceramics.

Advantages that ceramics have over other

materials include superior heat resistance, wear

resistance, and corrosion resistance. Increasing

applications for ceramic materials are being

found in industries such as aircraft, automobile,

and micro-electronics.

Although most of the ceramic parts are manufactured to near net shape by pressing and

sintering processes, precision machining is required to achieve a high degree of geometrical

accuracy. Consequently, machining of ceramics

has been a rapidly growing field. A traditional

ceramics machining process is grinding where

ceramic material is removed by the fine cutting

edges of abrasives particles. In general,

productivity of the grinding process is low, and

micro-cracks are formed on and beneath the

machined surface that often cause premature

failure of ceramic components during service due

to brittle fracture. With the ever increasing

number of ceramic materials, new and innovative

processing methods are needed to achieve cost

reduction and quality assurance. Many attempts have been made in this

regard. As reported in (1), a "laser lathe" was

developed where the dual beam principle was

applied to remove ceramic materials in a molten

form. However, the thermal damage induced on

the surface and the sub-surface during laser

machining is one of the drawbacks of the

technique. Another processing method is abrasive jet machining, which uses a stream of

high pressure fluid with abrasives insdes to pierce and wash ceramic materials away (2). But,

the availability of the advanced equipment and

the associated economic factors have limited its application on the shop floor. Recently, a new

method, called "heat assisted machining," which

integrates laser machining and grinding to remove ceramic materials in a ductile regime, is being developed. Product quality and

productivity is expected to be significantly

improved. However, with the wide selection of

abrasive grains in the market to high levels of

performance and dimensional accuracy, grinding

is still widely used in machining ceramics (3).

465
This paper presents a research focused on

basic understanding of the material removal

mechanisms during machining of ceramic

materials. Aluminum oxide (AI2O3) was the test material and Polycrystalline diamond tipped

carbide inserts were used. All machining tests were performed on a CNC turning center. Cutting

forces were recorded during machining and

surface finish was measured after machining. An

environmental SEM was used to examine both macro- and micro-cracks formed on the sample

surfaces. With the assistance of a computerbased vision system, 3-dimensional images of the

fracture surfaces were reconstructed and

quantified with respect to the micro-crack

formation during machining. Contour plots are used to characterize the size and shape of cracks

and their orientation. Effects of three basic

machining parameters, namely cutting speed, feed,

and depth of cut, on the crack formation were

investigated usinga factorial design technique.

Results from this research provide an insight into

the micromechanisms of material removal during

machining of ceramics, and suggest the

development of crack-controlled machining

technologies.

2. Procedures of Experimental

Investigation

Research results over the last few years have

shown that material removal in machining

ceramics is dominated by brittle fracture (4-6).

Evans and Marshall investigated the cutting

mechanisms of grinding ceramics. They pointed

out that the cutting mechanisms (or wear

mechanisms) are due to plastic flow and lateral

cracking (7). They predicted a lateral fracture

threshold loading limit. When the force

developed during machining is below the loading

threshold, plastic cutting occurs. On the other

hand, lateral cracking occurs when the developed

force is above the loading threshold. The

interplay between the dynamic characteristics of a machining process and the micro-fracture

mechanisms of the workpiece material at molecular level calls for a systematic study of the

physics behind the material removal process

during machining of ceramics. This research

focuses on an experimental investigation, which

consists of three major components:

1. Machining of ceramic material for

sample preparation.

2. Profilometric measurements of

machined surfaces.

3. SEM examinations of microstructural

and fractographic feature of the material

being machined.

2.1 Machining of Ceramic Material

Figure 1 shows the experimental setup for

machining aluminum oxide (AI2O3). A CNC

lathe was used. The machined AI2O3 is a

cylindrical bar with initial diameter = 16 mm and

length = 60 mm. The purity is 99.8% and the

grain size is 10 - 12 |im. The material is a strong,

dense recrystalized high-alumina ceramic

material. Some material properties related to machining are listed below.

Hardness

Modules of Elasticity

Compressive Strength

Fracture Toughness

Melting Point

1100-1200 kg/mm2

380 GPa

3000 MPa

4.0 MPa*Vm

2050 OC

Ceramic Bar Cutting Fluid Injector

CNC Sianter Lathe

Chip

oiiector

Strain Indicators

Polycrystalline Diamond

Tipped Carbide Inscr and Dynamomeicr

Data

Acijuisiuon System 0

Microcomputer

Figure 1 Experimental Setup for Machining

Alumina

A dynamometer was attached to the

toolpost and the tool holder was fixed on the

dynamometer. During machining, polycrystalline

diamond tipped carbide inserts, with nose radius

0.25 mm and -7 ^ rake angle, were used to machine the alumina bars. In order to absorb the

heat generated during machining, distilled water

was used as cutting fluid during machining.

Chips are collected in a funnel shaped device with a built-in filter to separate chip from the cutting

fluid. The dynamometer measured the cutting

466
force in both the cutting speed and feed

directions. The data were recorded and stored in a PC-based data acquisition system.

For setting the machining conditions, or

selecting feed, depth of cut, and cutting speed, a

factorial design method with dupUcates was used

(8) Table 1 lists the eight combinations of the

three machining parameter settings and the

measured cutting force data in both the cutting

speed and feed directions. For example, the row

listed as Test 5 in Table 1 means:

(1) Three machining parameter settings are feed = 0.05 mm/rev, depth of cut = 0.77

mm, and cutting speed = 67m/min;

(2) The average measured cutting force

component in the feed direction is 15.4

N (an average of two measurements,

13.9 N and 16.5 N), and the average

measured cutting force component in

the cutting speed direction is 6.5 N (an

average of two measurements, 6.2 N

and 6.7 N); and

(3) The measured roughness average value

from the machined surface, = 2.30

|im.

2.2 Surface Finish Measurements

After machining, surface profiles were

traced for the machined surfaces on a Talysttp

Stylus Instrument. In this research, the

roughness average, R^, was used as the

characterization index for the finish quality

evaluation. The eight measured Ra values are

listed in Table 1. The best fmish quality, R^ =

1.05 |im, was achieved at a combination of feed =

0.012 mm/rev, depth of cut = 0.26 mm, and

cutting speed = 67 m/min. The worst finish

quality, R^ = 2.30 jim, was at a combination of

feed = 0.05 mm/rev, depth of cut = 0.77 mm, and

cutting speed =11 m/min. These results are in

agreement with the established metal machining

theory. Figure 2 presents surface profiles along

the feed direction and taken from four different

samples. It is interesting to note that all the four

patterns of height variation displayed in Fig. 2 are of random nature. No detectable effects of the

tool nose radius and the tool motion pattern can be found with respect to the surface profile

formation. However, these profiles display

Table 1 Machining Parameters Settings and Results

Parameters Settings Force Measurements, N Resultant Surface

Test No. Feed Depth of Cut

Cutting

Speea

Feed Force Tang. Force Force Roughness,

mm/rev mm m/min. Runl Run 2 Run 1 ""Run 2 N [im

Test 1 mi 0.772 67. 8.47 5.^0 9.14 1.35

Test 2 0.012 0.772 11. 22.46 22.68 4.98 4.93 23.11 1.92 Tests 0.012 0.264 67. 3.83 3.71 2.16 2.05 4.32 1.05

Test 4 0.012 0.264 11. 3.24 5.03 1.47 1.99 4.48 1.15 Tests 0.051 0.772 67. 13.94 16.49 6.16 6.67 16.51 2.17

Test 6 0.051 0.772 11. 21.34 23.99 8.72 8.75 24.29 2.30

Test 7 0.051 0.264 67. 4.24 3.93 3.08 2.91 5.07 1.51 Tests 0.051 0.264 11. 5.76 5.14 2.61 2.20 5.96 1.70

467
several deep valleys along the feed direction.

The deepest valleys are found in the case of high

feed and large depth of cut (Fig. 2d). These

narrow and deep valleys suggest occurrence of

brittle fracture during the chip formation.

(mm)

0.035

0.025

0.015

0.005

0.0 0.1 0.2 0,3 0.4 0.5 0.6 (mm)

Ra= 1.05 m

(a) Test 3: f=0.012 mm/rev,

d = 0.264 mm, v= 67 m/min

(mm)

0.035

0.005

0.0 0.1 0.2 0.3 0.4 0.5 0.6 (mm)

Ra= 1.35 urn

(b) Test 1: f=0.0 12 mm/rev,

d = 0.772 mm, v= 67 m/min

(mm)

0.035

0.025

0.015

0.005

0.0 0.1 0.2 0.3 0.4 0.5 0.6 (mm)

Ra= 1.70 urn

(c) Test 8: f=0.051 mm/rev,

d = 0.264 mm, v= 1 1 m/min

(mm)

0.035

0.025

0.015

0.005

0.0 0.1 0.2 0.3 0.4 0.5 0.6 (mm)

Ra = 2.30 ^im

(d) Test 6: f=0.051 mm/rev,

d = 0.772 mm, v= 1 1 m/min

Figure 2 Surface Profiles Measured from the

Machine Surface

2.3 SEM Examination

To understand the process of chip

formation during machining and identify the

relationship between the fracture mode and the

microstructure of the materail, a detailed

examination and analysis of the fracture surface

is required. In this research, an environmental

scanning electron microscope (ESEM) was used

to obtain high-magnification images that provide

details about the geometry of the formed chips

and the topography of the machined surfaces.

Another advantage of using ESEM is that no

conductive coating is required as in the case of

ordinary SEM. This allows a direct identification

of macro- and micro-cracks induced by

machining without the influence of applied

coating.

Figure 3 presents a set of electron

micrographs of the chips collected during

machining. Figure 4 presents a set of

micrographs of the surfaces machined under the

eight combinations of depth of cut, feed, and

cutting speed settings at two levels.

3. Analysis of Experimental Results

The basic philosophy applied for data

analysis is first to identify if there are similarities

between the machining of ceramic materials and

the machining of metals. Afterwards, focus will

be given to analyzing new and important findings.

It is expected that there exist significant

differences in the micromechanisms of material

468
removal between the two types of machining

process.

3.1 Cutting Force Generated during

Machining

Following observations are made by exa- mining the cutting force results listed in Table 1:

1. The cutting force produced during

machining is proportional to the cutting

area. The small cutting forces, such as 4.32 N and 4.48 N, are associated with

the combinations of low feed and depth

of cut settings, and the largest cutting

forces, such as 24.3 N and 23.1 N, with

the combinations of high feed and depth

of cut settings. This observation

follows the law of material strength, i.e., force is proportional to the applied area

at a given stress level, or at a given

loading threshold.

2. As a first approximation, the unit cutting

force, calculated as a ratio of the

measured cutting force to the cutting

area, ranges from 650 N/mm^ to 2500

N/mm2. The low limit of the unit

cutting force happened during

machining at a high cutting speed

setting, a phenomenon that could be

interpreted as the effect of high

temperature on the material removal.

However, such a wide variation of unit

cutting force can only be contributed to the fact that the fracture strength of a

brittle material fluctuates. The

fluctuation can be as much as an order

of magnitude (5). In fact, this

characteristic indicates that ceramic

materials can be removed at a low unit

cutting force if the machining condition

is preferred, pointing out great promise

to search new and innovative machining

technologies.

3. The magnitude of the tangential cutting

force component is much smaller than

that of the cutting force component

along the feed direction, a phenomenon

just opposite to what has been observed

during the machining of metals where

the magnitude of the tangential force is much larger than that of the feed force.

3.2 Mechanisms of Material Removal

The analysis presented in this section

includes a detailed examination and analysis on

fracture mechanisms for clarification of the role

of the microstructure, or understanding the

metallurgical factors that control the material

removal process. The analysis is first divided

into two parts for examining the chip formation

and surface topography formation, respectively.

Then we combine the results from these analyses

and propose a model to describe the mechanisms

of material removals including the chip formation

process. Special efforts are made to illustrate a developed computer vision system for visualizing

and quantifying the fracture surface formed

during machining.

3.2.1 Mechanisms of Chip Formation

Information on the chip formation

mechanisms can be obtained through the

examination of the four micrographs of chips

shown in Fig. 3:

1. The chip fragments shown in Fig. 3a

and Fig. 3b were formed during two

machining processes with depth of cut

settings at 0.77 mm and 0.26 mm,

respectively. It is not surprising to conclude that a high material removal

process, such as the one with a large

depth of cut, will produce large chip

fragments. However, the size of chip

fragments formed during a single

machining process varies significantly.

For example, the chip fragments shown

in Fig. 3b are significantly different in

size. This phenomenon indicates that

special attention must be paid to the

investigation of chip formation during

machining of ceramic materials.

2. The two chip fragments shown in Fig.

3d were taken at a high magnification.

They represent two types of chip

observed in this experimental study.

The unique feature among the two types

of chips is their lateral surfaces, which

are formed by brittle fracture. This

feature indicates that the chip formation

during machining is a macro-scale

fracture process, or cleavage fracture

469
470OR
471
mode. However, there exists a unique

difference between these two types of

chip. The chip on the right side of the

figure has a smooth flat surface. The

parallel lines, or the plowing marks, on

the smooth surface resemble the chip

flow over the rake face of a cutting tool, a phenomenon routinely observed

during machining of metals. Note that

this chip was formed under a machining

condition where the temperature

measured in the cutting zone was about

TSO^C. This suggests that formation of

this type of chip could start from plastic

deformation in the cutting zone, then

progress into plastic flow over the rake

face, and end with separation from the

machined part due to fracture. This

suggestion can be interpreted as the

existence of deformation mode in a fracture process for brittle materials (5,

8). Research in the future direction

should pursue applying the elasticplastic fracture mechanics to characterize the crack tip field in the

chip formation process. Figure 5

illustrates the proposed formation

process of the two types of chip,

namely, the machined chip and the

fractured chip during machining. The

fractured chip is defined as those chip

fragments formed without direct contact

to, or flow over, the tool rake face. This

may provide an explanation to the

random nature of the height variation

pattern observed in taking surface

profiles from the machined surface.

DisUlled

Force

Figure 5 Two Types of Chip Formed during

Machining of Alumina

3. The micrograph at a high

magnification shown in Fig. 3c

depicts the intergranular fracture

developed within a fractured type of

chip formed during machining. The

location of the crack with respect to the chip as a whole is marked in

Fig. 3a. It shows that the crack is about 5 )im in width and 120 fim in

length. It is intersting to note that

this crack was trapped when a fractured chip was formed,

indicating that the formation of

fractured chip fragments is due to the catastrophic failure of the

ceramic material under high stresses (above the loading threshold),

| developed during machining. !

3.2.2 Mechanisms of Surface Topo-I

graphy Formation

One of the ultimate objectives of I machining engineering ceramics is to achieve a high degree of the geometrical accuracy of a designed part. The topography formed on a machined surface plays a key role in this regard.

Based on the previous discussion, the machined

surface is an assemblage of fracture surfaces left by the chip formation. Reliability of a machined \ part during its service period is directly related to the status of macro- and micro-cracks formed on

the machined surface as well as the region just

beneath the machined surface. A new approach,

which integrates profilometry, microscopic

analysis and image processing, is developed to investigate the mechanisms of surface topography

formation (9). ;

1. Macro- and microscopic examination of

the fracture surface. The eight

micrographs shown in Fig. 4 are the

images taken from the surfaces

machined under the eight machining

conditions with a magnification of 500.

The intergranular fracture along the

grain boundaries can be clearly

identified on these machined surfaces,

indicating their preferred, river-like

crack paths, as depicted in Fig. 4. An

important observation is that almost all of the intergranular cracks initiated at triple point junctions (intersections of

the grain boundaries) and dislocation

pile-ups at the grain boundaries. Figure

6 is a micrography of a chip fragment

collected during machining. A

magnification of 4,500 is used to reveal

both microstructural and fractographic

472
X 4,500

1 |im

Figure 6 Microstruclural and Fraclographic feature around a Single Grain

(A Triple Point Junction Is Indicated)

a single grain The

pairs of

2.

feature around

picture depicts the thret

contacting points between the

neighboring grains. It is verv' likely that

the cutting force induces local residual

tensile stresses at these triple points

dunng machining. In combination with

stresses from deformation generated

within the grains due to the cutting

action, crack nucleation is initiated at these locations. As the machining

process goes on, the progressive

development of grain boundary

microcracking leads to crack growtli, or propagation. Figure 6 also illustrates

the cleavages due to microcracking at the triple point junctions. The

progression of intergranular cracks

causes an entire grain to dislodge from

the machining area, leading to intergranular fracture or cleavage

fracture. The process of grain

dislodging forms micro-cracks, and

then macro-cracks, and finally, the

topography of a machined surface.

Profilometric examination. By

examining Fig. 2 and Table I jointedly, a small Pcak-to-Valley value is associated witli a machining process

with a small cutting force. It indicates

that most of the crack initiation, growth,

and cleavage were confined with a very

limited local areas, or "shori-crack"

regions (5), when the cutting iorce was

rather small. As these "short-cracks"

extended, the toughness of alumina built

up and resisted extension of these

developed cracks. Figure 7 illustrates a steadily increasing toughness with

expanding crack size for alumina (5).

On the other hand, the protlle with a large Peak-to-Vallcy value is associated

with a machining process generating a large cutting force. This indicates that

the induced local residual tensile

sti'esses exceeded a critical value, but the

toughness keeps at a constant level

because the crack length has reached a

critical value, say 10 [im illustrated in

Fig. 7 for alumina. Under such

circumstances, accelerated crack

propagation will lead to the occurrence

of catastrophic failure of alumina.

During machining, this could imply the

increase of formed fractured chip

fragments. An excessive amount of

fractured chip fragments could

introduce large irregularities on the

473
machined surface, as illustrated in Fig.

2d.

1 10 10' 10^ icr icp Crack Length, c (^m)

Figure 7 Fracture Toughness as a Function

of Crack Length (5)

Stereophotogrammetric examination

through image processing. Stereoscopy

offers a direct and nondestructive

procedure for determining the elevation

of a fracture surface at selected regions

(10-12). In this research, a software

package, which integrates image

processing, profilometry, and computer

graphics, was developed to visualize the

surface topography formed during

machining, display patterns of crack

propagation, and quantify the size and

shape of intergranular and transgranular

cracks from the ESEM micrographs.

Figure 8 presents a visualized surface

topography in three-dimensional space

and an associated contour map. The

corresponding area of the surface

topogarphy is marked in the

micrograph shown in Fig. 4e(area 5).

The reconstructed fracture surface

topography vividly depicts the surface

texture formation during machining. It visualizes the appearance of surface

cracks and quantifies the relative

positions between the surface

roughness and the surface cracks. For

example, the distance 4.2 |j.m marked in Fig. 8a, obtained through calibration

using the profilometric measurements,

provides a valuable information for

post-machining, such as grinding and

polishing, as well as reliability

assessment of the machined parts. By

Reference Plane

for Ra Measurement

m

(a) Visualization of Surface Topography and

Crack Indications

13 13 i.O SS> 2JD U Ol5 03 SS SS T3

7 (im

PiieU(i2)

(b) Contour Map for Quantifying the

Formed Cracks

Figure 8 Fracture Surface Reconstruction

and Crack Closure

examining the deduced contour map

shown in Fig. 8b where the interval

between the contour lines is 0.5 |im.

The numbers shown outside the frame

are the heights of the location, the

474
lowest point of the surface topography

being taken as 0.0 ^im. On the other

hand, the pixel number, as illustrated in

Figs, 8a and 8b, provides a scale for

calculating dimensions in a horizontal

plane. By adding the width between

two contour lines, a quantitative

information about the size and shape of

a crack can be estimated. For example,

the width marked as 0.4 jim in Fig. 8b

covers 6 contour lines. The information

provides the size of a V-shape crack

having an angle of 1 .6^ with its vertex

4.2 |j.m away from the reference plane

for Ra measurements. It is interesting

to note that the contour map could assist us in identifying transgranular cracks.

For example, the boundaries of a grain

sized 10 \im can be identified on the

map, as shown in Fig. 8b. Within the

grain, a crack path is also identified. It

is marked 0.8 |im. If further

investigation on this spot is pursued, we

might be able to have information

about transgranular cracks developed by

machining. The principal objective of

modem quantitative fractography is to express the geometric characteristics of

the features on the fracture surface in

quantitative terms. The stereophotogrammetric examination through image

processing developed in this research

may provide a powerful tool to establish

the quantitative relationships between

the crack formation and machining

conditions.

3.3 Material Removal Mechanisms

during Machining

From the above discussion, as a material

removal process, machining of ceramic materials

shares certain common characteristics with what

we have observed from machining of metals.

Both machining processes involve cutting force

generation, chip formation, and surface

topographical generation. However, machining of

ceramic materials possesses unique

characteristics. Based on our experimental work.

we present a model to describe the material

removal process with emphasis on the

relationships between the macroscopic and

microscopice fracture behaviors. Figure 9

illustrates the five essential stages in this basic

physical process:

(a) Dynamic Loading and Induced Stress Field

(b) Crack Initiation

(c) Crack Propagation

(d) Chip Formation

475
Two Types of Cutting

bol

Surface

Damage

Sub-surface

Damage

(e) Formation of Surface and

Sub-surface Damage

Figure 9 Five Essential Stages in the Material

Removal

1. Dynamic Loading. When a machining

process starts, the tool approaches to the

part material. At the instant when the

tool touches the part material, the impact

between them induces stresses in the

material and a stress field is formed

(Fig. 9a).

2. Crack Initiation. When subjected to the

induced stresses, crack nucleation

begins immediately. Their sites are locations of stress singularities where

high internal tensile stresses build up.

These locations can be triple-point

junctions, dislocation pile-ups at grain

boundaries, and second-phase particles.

All of these flaws could initiate microcracks (Fig. 9b).

3. Crack Propagation. As the machining

process goes on and stresses continue

to develop in the material, the formed

microcracks grow in the material

immediately surrounding them. Both

intergranular and trangranular cracks

are developed and advanced obstacles

are encountered (Fig. 9c).

4. Chip Formation. The process of crack

propagation is terminated when more

micro-cracks propagate and produce

unstable fracture. The material

separated by fracture can be in a

position either contacting with the rake

face of the cutting tool or away from the

cutting zone, depending on the pattern

of crack propagation. Consequently,

two types of chip will be formed. The

type of machined chip (shadow areas)

3.4

represents the fractured material directly

under the tool action and also subjected

to the temperature influence. On the

other hand, the type of fractured chip

depends on the pattern of crack

propagation, which is governed by the

material microstructure (Fig. 9d).

Formation of Surface and Sub-surface

Damage. Due to brittle fracture, chip

fragments are formed during

machining. At the same time, surface

texture is formed on the machined

surface with cracks because brittle

fracture. These cracks are called surface

damage and, in general, unavoidable

during machining of ceramic materilas.

In addition, the residual effects of brittle facture near the surface layer bury

numerous micro-cracks to a certain

depth, leading to sub-surface damage

(Fig. 9e).

Development of Crack-controlled

Machining Technologies

Based on the comprehensive understanding of material removal mechanisms, we

should have gained guidelines for developing

new and innovative machining technologies to control the crack formation. The following

suggestions are made in this regard.

1. Control of Microstructure of Ceramic

Material. It has been recognized that

machinability can be significantly

improved by controlling microstructure

of metals through heat treatment. The

same philosophy applies to the

machining of ceramic materials as well.

The microstructure of ceramic materials

is central to strength and toughness

properties. By controlling the size and

shape of individual grains, or the

volume fraction of incorporated second

phases, we may be able to balance the

need for avoiding structure failure from

design point of view and the need for

degrading short-crack toughness from

the viewpoint of improving machinabiUty. A typical example would be the

development of Dicor/MGC, a glassceramic material designed for use in

dental restorations (14). The material

has a unique microstructure consisting

of mica flakes of approximately 70

volume percent dispersed in a non- 476
porous glass matrix. The cleavage

along the literal planes of mica flakes

functions as stress concentration, a situation similar to the stress

concentration around the inclusions in

free-machining steels, offering an

excellent machinability.

2. Control of Machining Parameters. It is certain that the tool geometry will play a key role in determining the distribution

of induced stresses in the material being

machined. By limiting the volumetric

ratio of fractured chip to machined chip,

we may be able to manage surface and

sub-surface damage. The ductile region

grinding is a typical example of keeping

at a low level or eliminating the

formation of fractured chip during

machining. Proper selections of the

three cutting parameters (feed, depth of

cut, and cutting speed) may render an optimal machining performance

possible. In the study of environmental

effects on fracture mechanics, it has

been well recognized that stress- corrosion cracking causes intergranular or cleavage separation by a loss of

cohesion (15). Therefore, use of cutting

fluids can improve machining

efficiency. In fact, chemical-assisted

machining has been studied intensively,

showing great promise for crackcontrolled machining (16-17).

3. New Processing Methods. Laser-based

heat assisted machining is being

developed. By controlling the

temperature of the cutting zone, an increase of pre-fracture plastic

deformation could reduce brittle fracture

significandy, leading to a full control of

machining process. On the other hand,

crack arrest through post-processing

would be a typical example to improve

surface quality after machining.

4. Conclusions

In this research, a fundamental study

related to machining of alumina is conducted. It consists of cutting force measurement,

assessment of surface finish quality, and analysis

of micromechanisms of material removal during

machining. The results are summarized as

follows.

1. Machining of ceramic materials is characterized by crack initiation and

propagation, leading to chip formation.

There are two types of chip, i.e., machined

chip and fractured chip, nd cleavage fracture are the three key elements in the machining

process. A model is proposed to describe

the material removal process.

2. By integrating image processing,

profilometry and computer graphics, a computer-based vision system is developed

for expressing the geometric characteristics

of features on the fracture surface through

three-dimensional visualization and contour

mapping.

3. Collaboration between shop floor machinists

and material processing engineers is needed,

to 1). optimize ceramic microstructures, that

provide required resistance to the initiation

and propagation of brittle cracks while

maintaining an acceptable machinability; and

2). develop new and innovative machining

technologies for effectively controlling the

crack formation during machining.

Acknowledgements

The authors acknowledge the support of

the University of Maryland Research Board, the

Department of Mechanical Engineering, and the

Institute for Systems Research at the University

of Maryland under Engineering Research Centers

Program: NSFD CDF 8803012. Special thanks

are due Mr. Myron E. Taylor, Directior of the

Central Facility for Microanalysis, who provided

valuable assistance in conducting the SEM study.

The technical discussions with Dr. Brain Lawn of

the National Institute of Standards and

Technology, and Dr. Isabel Illoyd and Dr. James

Dally of the University of Maryland provided

many insights of this research. The assistance

from Mr. Doug Beale and Mr. Don DeVoe

during the course of this work is also appreciated.

References

1. G. Chryssolouris, J. Bredl, S. Kordas,

and E. Wilson, "Theoretical Aspects of a Laser Machine Tool," ASME Winter

Annual Meeting Proceedings, PED 20,

pp. 177-190, December 1986.

2. M. Mazurkiewicz, "Understanding

Abrasive Waterjet Performance,"

477
Machining Technology. Vol. 2, No.l, pp.

1-3, 1991.

3. K. Subramanian, S. Ramanath, and Y. O.

Matsuda, "Precision Production Grinding

of Fine Ceramics," Proceedings of the

First International Conference o n Manufacturing Technology. Chiba, Japan,

1990.

4. R. C. Bates, "Micromechanical Modeling

for Prediction of Lower Shelf, Transition

Region, and Upper Shelf Fracture

Properties," Part 3, in Fracture

Mechanics: microstructure and

micromechanisms." Edited by Nair, et al, ASM International, pp. 131-168, 1989.

5. B. R. Lawn, " Fracture of Brittle Solids.

" Second Edition, Cambridge University

PresSi Cambridge, in press.

6. J. Cagnoux, and A. Cosculluela,

"Influence of Grain Size on Triaxial

Dynamic Behavior of Alumina," Part 6 in

Dynamic Failure of Materials: Theory.

Experiments, and Numbers." Eds. H.

Rossmanith, and A. Rosakis, Elsevier

Apphed Science^ pp. 73-84, 1991.

7. A. G. Evan, and D. B. Marchall, "Wear

Mechanisms in Ceramics," I n Fundamentals of Friction and Wear. Ed.

D. A. Rigney, Metals Park, Ohio,

American Society of Metals, pp. 439-

452, 1980,.

8. G. Box, W. Hunter, and J. Hunter,

" Statistics for Experimenters. An

Introduction to Design. Data Analysis,

and Model Building ." John Wiley &

Sons, Inc., 1978.

9. J. Landes, and R. Herrera,

"Micromechanisms of Elastic/Plastic

Fracture Toughness," Part 4, in

" Fracture

Mechanics: microstructure and

micromechanisms. " Edited by Nair, et al, ASM International, pp. 111-130, 1989.

11. E. E. Underwood, "Recent Advances in Quantitative Fractography," Part 2, in

Fracture Mechanics: microstructure and

micromechanisms." Edited by Nair, et al, ASM International, pp. 87-109, 1989.

12. E. E. Underwood, "Quantitative

Fractography", Chapter 8, in Applied

Metallography. G. F. Vander Voort, Ed.,

Van Nostrand Reinhold, pp. 101-122,

1986.

13. J. E. Hilliard, Quantitative Analysis of

Scanning Electron Micrographs," Journal

of Microscience. 95, Part 1, pp. 45-58,

1972.

14. D. Grossman, "Structure and Physical

Properties of Dicor/MGC GlassCeramic," Proceedings of the 1991

International Symposium on Computer

Restorations . Regensdort-Zurich,

Switzerland, pp. 103-115, 1991.

15. W. Gerberich, "The Micromechanics and

Kinetics of Environmentally Induced

Fractures," Part 3, in

" Fracture

Mechanics: microstructure and

micromechanicsms. " Edited by Nair, et al, ASM International, pp. 201-208, 1989.

16. G. Zhang, T. Hwang, D. Anand, and S. Jahanmir, "Tribological Interaction in

Machining Aluminum Oxide Ceramics,"

Proceedings of the Navy Tribolog y

Workshop . Annapolis, Maryland, May

11-13, pp. 9-13, 1992.

17. G.M. Zhang, T.W. Hwang, and D.K.

Anand, "Chemo-Mechanical Effects on

the Efficiency of Machining Ceramics,"

Proceedings of 1993 National Science

Foundation Design and Manufacturing

Systems Conference . Charlotte, NC,

January, pp. 421-428, 1993.

10. K. Komai, and M. Noguchi, "Fracture

Surface Reconstruction Technique and

Crack Closure," Part 5, in "Fractography",

Edited by Koterazawa, et al., Elsevier

Applied Science, pp. 161-186, 1990.

478
THE MACHINABILITY OF DIFFERENT DENTAL

MATERIALS FOR CAD/CAM SYSTEMS

K. H. KUNZELMANN and R. HICKEL

University of Munich, Munich, Germany

In spite of increased machinability and optimized physical properties, presently available

materials for dental CAD/CAM restorations suffer from chipping at machined edges.

The aim of this study was to compare glass ceramic materials, sintered leucite

reinforced ceramic, dental hybrid composites and an experimental hydroxiapatite with

regard to the amount of chipping at machined edges. Samples in shape of a computer

generated test-inlay with rectangular edges were machined using a dental CAD/CAM

machine (disc diameter: 4 cm, average grain size: 64 jum, unloaded surface speed: 45

m/s). The machined samples were evaluated with both SEM and profilometer. Samples

with a fine microstructure (hot-pressed, fine grain size of the sinter particles) had

smaller margin defects. Increasing crystallite contents improved margin quality, too. The

average chip size at machined egdes of the tested materials was: phlogopite containing

glass ceramic: 60 //m, leucite reinforced ceramic: 60 /xm, tetrasilicic fluormica glass

ceramic: 80 jim, feldspathic porcelain: 105 ixm.

Introduction

The increasing esthetic demands of patients

and excellent biocompatibility have led to growing interest in using ceramics as restorative

materials in dentistry. Usually the restorations are fabricated by slip-casting, controlled

crystallization of glass or hot-pressing of

precrystallized glasses. A lost-wax process is used to shape the molten glasses or glassceramics. The properties of these materials are highly dependent on the skills of the dental

technician. In addition process-dependant

microdefects reduce the fracture toughness and

increase the time-dependant failure probability

of restorations [1].

machined edges

profiiometTic surface tracings

Figure 1: Computer generated 3D-data set for the

CAD/CAM-unit.

479
Efforts to automate the production of dental

restorations have induced the developement of

CAD/CAM-units to process dental ceramics [2,

3, 4]. Computer-aided manufacturing of dental

restorations helps to reduce the above mentioned limitations by using optimized ceramic

materials with little variation in microstructure.

With already commercially available devices the

prefabricated ceramic blocks are milled with

diamond coated grinding discs to fit into tooth

cavities. In spite of increased machinability and

physical properties presently available materials

for CAD/CAM purposes suffer from chipping

at machined edges. The chipping defects do not

only reduce the accuracy of fit of the restora- tions but may be the predominate cause of

catastrophic failure [5].

The aim of the present study was to optimize

ceramic materials for better stability of machined edges in presently available CAD/CAM

milling devices for dental restorations. For this

purpose experimental ceramic materials reinforced with leucite were compared with commercially available machinable ceramics. In

addition a conventional dental hybrid composite

and an experimental hydroxiapatite material were included to evaluate the amount of chipping at machined edges of other materials which

are commonly used for dental restorations.

Materials and Methods

The tested materials were machined with a dental CAD/CAM milling device driven with a high-speed motor. The cutting tool is a diamond

coated grinding disk with a diameter of 4 cm

(average grain size: 64 /xm). The unloaded

surface speed of the grinding disc is 45 m/s.

Ceramic blocks complete one rotation per

second around the axis and against the grinding

disc. The samples were machined to the shape

of a computer generated test-inlay with rectan- gular edges (Fig. 1). The samples were water

cooled during the milling process. To reduce

surface tension a combined detergent and lubricant was added to the cooling water. ,

The experimental ceramic materials were

prepared by a dental ceramics manufacturer.

The raw materials were fused at 1450°C to molten glass. For increased homogeneity the

melt was quenched, dried, crushed and molten

for a second time. In a following tempering

process the amorphus glass was converted to a

partially crystalline glass ceramic which was

crushed and ground to particles [6]. The particles ware shaped in a sinter process to the

material blocks used in our study . The variables

which were investigated were different crystallisation conditions, cooling rates after sintering to the final shape, sinter conditions and grinding

sizes of the sinter particles.

The materials were coded and all experiments were done without knowing the composition or

the properties of the coded samples. The machined samples were evaluated in the SEM.

The investigation of the experimental materials was done in four test series. In a first screening series the influence of different crystallite contents on machinability was examined

qualitatively with the SEM. Seven materials

with varying leucite contents and varying sinter

conditions were used. Table 1 summarizes the

linear thermal expansion coefficients of the

experimental materials and their sinter conditions. With increasing leucite content the thermal expansion coefficient also increased. According to the manufacturer A - D are highly

crystalline glass ceramics, E is semicrystalline,

F is amorphus glass and G is a feldspathic

porcelain for enamelling to metals.

In the second series the influence of internal

stress caused by the contraction mismatch of the

glassy phase vs. the crystal phase leucite on the

defect size at machined edges was studied. For

this purpose two highly crystalline materials, EST and E-MT, with the composition summarized in Table 2 but with identical sinter conditions were compared [6] . One half of the specimen were cooled with the furnace switched-off

and the other half at a cooling rate of less than

5°C per minute.

480
Table 1: Thermal Expansion Characteristics and Sinter Conditions of

the Experimental Ceramics

Material

(external code)

a„ (lO 'VK)

[-500 °C]

a„ (lO V'K)

[500-600 °C]

Pressure to form

greenware [MPa]

Sinter temperture

l°C]

A 18.8 24,9 2.1 820

B 18.8 24.9 3.3 820

C 18.8 24.9 2.1 850 D 18.8 24.9 3.3 850

E 20.0 23 3.3 920

F 8.3 7.9 3.3 920 G 14.5 14.2 3.3 880

Table 2:

Composition (wt %) of the basic glasses for E-MT and E-ST

Oxide SiO, AI.O, K,0 Na,0 B2O3 CeO, CaO BaO TiO,

Basic glass of E-MT 63.0 17.7 11.2 4.6 0.6 0.4 1.6 0.7 0.2

Basic Glass of E-ST 64.0 14.3 10.1 8.0 0.6 0.9 3.3 1.5 0.3

In the third series, the influence of microstructure on machinability was studied. To

obtain a finer microstructure the samples were

hot-pressed. The sinter process took place at a temperature of 1050°C and simultaneously

applied pressure of 0.5 MPa. The quality of the

machined edges was compared to the samples

of series 1.

Further improvement of the sinter process

was achieved in the fourth series by a finer

grain size of the sinter particles and hot-pressing. During the preceeding experiments the

average grain size of the particles was between

30 - 40 /xm, whereas the materials in series 4

had an average grain size of 7 ixm. To validate

the influence of the crystallite phase a highly

crystalline material (PK-MT) was compared

with a material with a higher content of glassy

matrix (BS-PK).

The results of these optimized materials were compared with already commercially materials

which were optimized for machinability. Two of

them were developed for dental application

whereas the third one is widely use in medical

implant technology.

In the last series, materials with different

micro structure were tested for their feasibility as machinable restorative materials. Dental

composites are successfully used either as plastic

materials which are light-cured in the oral cavity or as inlay systems which are fabricated extraorally and bonded to the teeth with a composite

bonding system. A conventional dental hybrid

composite (20 wt% bifunctional methacrylate,

79 wt% anorganic filler particles with an aver- age diameter < 1,5 /xm, 1 wt% photoinitiator

and pigment) was formed in cylindrical shape

for milling and then light cured as recommended

481
in vacuum with intensive light. As teeth consist

predominately of hydroxiapatite the restoration

of teeth with the same material would be the

most natural way to treat carious lesions. An

experimental hydroxiapatite from bovine bones

(theoretical density: 99,5 %) served as screening sample.

The surface roughness was measured with a profilometer (diamond stylus, tip angle 90°, 5

fim tip radius, ISO 4287/1, X = 0,8). The

roughness of the machined edges, the width and

the depth of the defects at the edges of the

specimen were recorded with a diamond knife

(3 mm broad, tip angle 120°, 5 ixm tip radius,

DIN 4768 and DIN 4768/1, X = 0,8).

Results

Specimens with a higher crystallite contents

exhibited less surface roughness than materials

with a higher glassy matrix contents (Table 3).

SEM evaluation of the samples with a high

leucite contents shows regions with surface

defects of brittle fracture mode besides plastically deformed areas without substantial microcracking (Fig. 2) while samples with little or no

crystallite contents showed only brittle fractures on the surfaces.

All materials suffer from chipping at the

machined edges (Fig. 3). The sizes of the shellshaped defects at the machined edges range up

to 300 fxm with increasing values for higher

contents of glassy matrix. The cooling rates had no influence on the size of the defects. All

samples of series 2 exhibited defects up to 200

fxm along the machined edges. Hot-pressing of

the samples in series 3 resulted in a more

homogenous surface and less machining damage

along the edges with flaws of only 60 [xm. An

additional gain in margin stability was achieved

in series 4 by finer grain sizes of the hotpressed materials. The PK-MT samples had

defect sizes up to 50 fxm and the BS-PK samples up to 100 jim.

The marginal defects of the commercially

available dental ceramic materials ranged from

80 to 100 jwm. The phlogopite containing glassceramic samples had the smallest defect sizes of

the commercially available materials with defect

sizes at the edges up to 60 /xm (Fig. 4). Both

the dental hybride composite material and the

experimental hydroxiapatite showed defects up

to 150 fxm. The fracture mode of the defects of

both materials was of plastic in nature. The

composite had a fine surface structure whereas

the hydroxiapatite was characterized by a porous, fine-grained surface with smooth islands

of highly condensed material (Fig. 5).

The result of the measurements with the

surface roughness analyser are given in Table 4,

using the parameters and R^^^^. The R^-value

is the mean of the maximum peak-to-valley

values from 5 consecutive sampling lengths,

while Rmax is the maximum peak-to-valley value

of all 5 sampling lengths.

Table 3: Surface roughness of test specimen

with different leucite contents. (R^ = average surface roughness ac- cording to ISO 4287/1 (X = 0,8),

SD = standard deviation).

Material R, [^im] SD [nm]

A 1.47 0.07

B 1.66 0.08

C 1.62 0.08 D 1.58 0.1

E 1.83 0.11

F 1.76 0.11

G 1.80 0.18

482
Figure 2: Machined surface of material A. The Surface exhibits brittle fracture mode are besides plastically deformed areas (right window).

483
Figure 4: Small defects at the machined edge of the optimized experimental

material PK-MT.

20kM .0 25kx 4QQy 318

Figure 5:

SEM picture of the surface structure of an experimental hydroxiapatite.

2 0k V 1 . 00kx 10 . Q)' 31 1

484
Table 4: Roughness of the machined edges recorded

with a diamond knife (DIN 4768 and DIN

4768/1).

Material R2[/xm]

A 6,04 7,33

B 5,51 5,86

D 6,58 9,1

E 6,67 8,73

F 9,11 11,96

G 6.64 9,59

E-MT (cooling rate: < 5°C) 9.16 14,35

E-MT (cooling with furnace

characteristic)

6,59 8,44

E-MT (hot-pressed, conventionally ground)

6,51 7,51

PK-MT (hot-pressed, fine ground)6,06 6,58

ground)

5,98 7,09

feldspathic porcelain 5,55 7,19

phlogopite containing glass ceramic5,89 6,18

tetrasilicic fluormica glass ceramic7,13 8,28

Discussion

Material removal under the load of the

cutting tools takes place by microcracking and

chipping of the surface. The induced microcracks result in larger chips if the crack is not

deflected and branched by crystallites. This is confirmed by the increased surfaces roughness

of materials with a higher portion of glassy

matrix and the SEM pictures. With raised

leucite contents plastically deformed areas can be observed with nearly indistinguishable grain

boundaries (Fig. 2). Kelly [5] attributes this as "polishing wear" resulting in declined wear rates with the number of grinding cycles. This is in

accordance with the reduced working times of

the grinding tools for tetrasilicic fluormica glass

ceramic compared to feldspathic porcelain [7]

Although chipping at machined edges was

less for materials with a higher leucite contents

other factors seemed to influence the dimension

of the defects. The contraction mismatch between the thermal expansion of the leucite

cystallites and the glassy matrix can result in

stabilizing internal stress [6, 8]. At machined

edges, however, the internal stress can increase

the size of the defects. A slower cooling rate

after firing the samples can reduce the inner

tensions leading to increased margin quality.

The different tempering methods, however, had no influence on the margin stability of the

edges. Another possible explanation of the

large margin defects was breaking of the margin

comers along the grain boundaries of the frit particles. This was verified by series 3 with the

improved sinter technique, i. e. hot-pressing of

the sinter particles, and series 4 with the finer

grain size of the frit particles. Both series

resulted in improved margin quality. Series 4

additionally confirmed the findings of series 1 that a higher crystallite contents results in less

margin defects, because material BS-PK, which

showed larger flaws, had a lower leucite contents than material PK-MT. The margin quality

of the fine ground and hot-pressed experimental

material is comparable to the phlogopite containing glass ceramic.

The differences between the phlogopite

containing glass ceramic material, the experimental material PK-MT and the tetrasilicic

fluormica glass ceramic may be explained

through their microstructure. The phlogopite

containing glass ceramic, which shows the

smallest defects at machined edges, is a glass

ceramic material with spherical crystallites. The

485
experimental materials are leucite reinforced.

Leucite crystallites are cubical or tetragonal in

shape [9] . In both cases only small edge defects

are observed whereas the tetrasilicic fluormica

glass ceramic had larger defects (Table 4)

because of the interlocking nature of the flat plate-like crystals [10]. Instead of single crystallites, crystallite agglomerates broke away from

the machined edges. Further investigations may

show the influence of the cristallite shape on

fracture toughness, fatigue strength and chipping at machined edges, because on the one hand it can be expected that the interlocked,

plate-like crystallites give better fracmre or

fatigue data while on the other hand the machinability of fine structures may be better with

the less interlocked spherical, cubical or tetra- gonal cystallites.

Esthetic appearance also needs to be considered in the dental context. While the tetrasilicic

fluormica glass ceramic or PK-MT are tooth

colored and can be processed to dental restora- tion in a one step procedure, the phlogopite

containing glass ceramic is too opaque for

esthetic restorations. Better esthetic results can

be obtained with non-tooth-colored materials in a two step procedure if they are used as copings

which can be individually characterized in

consecutive firing processes. This technique

could be transfered even to high-strength alumi- na ceramic containing at least 70 % pure alumi- na in the body. Here the greenware from aluminum oxid could be shaped to a coping with the

CAD/CAM-unit, reinforced with the glassy

matrix and color-characterized afterward

leeding to a highly fracture resistant and

esthetic dental restoration.

Both, composite and hydroxiapatite samples,

demonstrate that they have potential as machinable materials. With these materials the working

time of the grinding tool is nearly tenfold compared to the ceramic materials. Composites can be used now for small, non-stress bearing

restorations. Further experiments are needed

however, to optimize these materials for larger

restorations.

Conclusion I

! The first step to optimize ceramic materials ; for milling conditions similar to the available ' dental CAD/CAM-systems should be to mini- ; mize strucmral defects. The influence of the i reinforcing crystallite system can be determined

in a second step. In dentistry improving of the

esthetic appearance is the third step in the j development of new machinable materials. Since

the used CAD/CAM-unit represents only a fixed i set of tribosystem variables, the development of i future systems should be more flexible regard- | ing their milling devices in order to adapt to the I requirements of new materials. |

I

I

References

[1] J. W. McLean, "The Science and Art of

Dental Ceramics", Oper Dent. Vol. 16,

pp. 149 - 156, 1991.

[2] D. Rekow, T. M. Speidel, and A. G.

Erdman, "Computer-Aided System to Automate Production of Posterior Dental

Restorations", Abstract 1336, J Dent

Res, Vol. 65, p. 317, 1986.

[3] F. Duret, J. L. Blouin, and B. Duret,

"CAD-CAM in dentistry. J Am Dent

Ass, Vol. 117, pp. 715 - 718, 1988.

[4] W. H. Mormann, M. Brandestini and F. Lutz, "Das Cerec-System:

Computergestiitzte Herstellung direkter

Keramikinlays in einer Sitzung",

Quintess Zahnarztl Lit , Vol. 38, pp. 457

- 462, 1987.

[5] J. R. Kelly, H. Liithy, A. Gougoulakis,

R. L. Pober, and W. H. Mormann,

"Machining Effects on Feldspathic

Porcelain and Glass Ceramic:

Fractographic Analysis", Proceedings of

the International Symposium on Computer Restorations , pp. 253 - 273,

1991.

486
[6] G. Beham, " IPS-Empress: Eine neue Keramik-Technologie "

IVOCLAR/VIVADENT Report #6. 1990

[7] W. Mormann, and O. Furrer,

"Efficiency of an Electric Drive in the

CEREC CAD-CAM Unit", J Dent Res , Vol. 71, Spec. Issue, Abstr. # 3, p. 516, 1992.

[8] K. J. Anusavice, P. H. Dehoff, B.

Hojjatie, and A. Gray, "Influence of

Tempering and Contraction Mismatch

on Crack Development in Ceramic

Surfaces", J Dent Res, Vol. 68, No. 7,

pp. 1182 - 1187, 1989.

[9] H. Claus, "Die Bedeutung des Leuzits

fiir die Dentalkeramik", Zahnarztl

Welt/Reform, Vol. 90, No. 6, pp. 44 - 46, 1981

[10] D. G. Grossman, "Structure and

Physical Properties of DICOR MGC

Glass-Ceramic " , Proceedings of the

International Symposium on Computer

Restorations , pp. 103 - 115, 1991.

Acknowledgements

- This study was supported by the SANDERSTIFTUNG.

- The authors gratefully acknowledge Mr. P. Wollwage and Prof. Dr. Holand of

IVOCLAR/VIVADENT for providing the

experimental ceramic formulations used in

this study.

- The phogopite containing glass ceramic was

developed at the Friedrich-SchillerUniversitat Jena, Otto-Schott-Institut fur

Glaschemie.

- The experimental hydroxiapatite was

supplied by MERCK.

487

MECHANICS OF TOOL WEAR DURING MACHINING OF

ADVANCED FIBROUS COMPOSITES

G. SANTHANAKRISHNAN, R. KRISHNAMURTHY, and

S. K. MALHOTRA

Indian Institute of Technology, Madras, India

Advanced fibrous composites, owing to higher specific

strength , elastic modulus and also impact and corrosion

resistance find wide range of applications in aerospace,

automobile and other industrial sectors. To obtain

dimensional tolerance and also to maintain certain

functional features, composites are usually machined.

Owing to their heterogeneous nature and also depending

upon the orientation and nature of reinforcing fibre used,

the composites pose considerable problems in machining.

Pulsating cutting forces, higher order cutting temperature

and chip clogging encountered during machining of

composites result in thermo-mechanical stress induced

spalling, abrasion and chip notching modes of tool wear.

Different types of tool wear and the associated

mechanisms, observed during machining of carbon fibre

reinforced plastics (CFRP) and Kevlar fibre reinforced

plastics (KFRP) composites with plain and coated cemented

carbides are presented in this paper.

INTRODUCTION

Rapid developments in

composites including resin,

metal and ceramic based

composites have opened their

application range. This is

largely due to their high

specific strength, elastic

modulus, corrosion resistance

and good fatigue resistance.

Though composites are usually

moulded, for obtaining close

fits and tolerances and also for

achieving near-net shape,

certain amount of machining has

to be carried out. Due to their

non-homogeneity, composites pose

considerable problems in

machining. The data available

for machining of homogenous

metallic materials cannot be

used for composite machining.

Further each composite differs

in its machining behaviour since

its physical and mechanical

properties largely depend on

type of fibre, fibre content,

fibre orientation and

variabilities in the matrix

material. Among the resin

based composites, carbon fibre

reinforced plastics ( CFRP) and

Kevlar fibre reinforced

plastics (KFRP) find the wider

application in aerospace

industries. In FRP composites,

which is an assemblage of

plastics and fibres in different

orientations, it is difficult to

investigate the mechanism of

material removal during

machining [ 1 , 2 ]

, since the action

of each fibre encountered by the

489
•U\

T3

(U

ty\ c

itu

-*-a

2

Q.

E

500

400

300

200

Tool material • K 101 (Sintered o P20j carbides

A Tic (Cemeots

A CW (Ceramics)

a = 1.0mm _ s = O.lmm/rev.

Dry

Thermocouple=III type

2 3 4

Spindle speed

6 8 10

rpm xlO 2

20

Fig. 1 Relation between spindle speed and

temperature at cutting edge [3]

cutting tool will be varying

from moment to moment. Thus, a

cutting tool in machining a FRP

composite will experience an

environment of pulsating load,

apart from load fluctuation. The

cutting edge will also

experience higher order cutting

temperatures [ 3 ] during

machining. Figure 1 illustrates

the influence of cutting speed

and type of cutting tool on

cutting temperature during

machining of composites. It is

seen that the cutting

temperature increase with

cutting velocity. For each

tool, there exists a critical

velocity around which there is a

rapid increase in cutting

temperature. Since composites

have poor thermal conductivity,

the cutting tools must dissipate

major part of the heat

generated; thereby experiencing

higher cutting temperatures.

Apart from pulsating

forces, and higher temperatures

the cutting edge will also

experience large scale abrasion

due to the rubbing of the fibres

on the machined work surface

over the flank portion of the

cutting tool. Thus, a cutting

tool will be subjected to impact

fatigue, thermal degradation and

abrasion; thereby experiencing

different forms of tool wear.

Figure 2 [4] illustrates the

influence of cutting temperature

on wear and associated causes

for tool wear. From this

figure, it can be seen that most

of wear mechanisms reported

become operational during

machining of composites.

Observations on tool wear during

machining of composites have

illustrated mixed modes of tool

wear.

EXPERIMENTAL

Machining trials have been

carried out on filament wound

CFRP and KFRP tubes in a precision, rigid and high speed

490
a

Cutting speed (temperature)

Workpiece Galling

removal

Removal galling

build up with

carbide particles

Alternating thermal

strains in the cutting edge

(b)

Fig. 2 Schematic representation of (a) Wear os o function

of cutting speed. (b) Causes of Wear W

VDF lathe using plain and coated

cemented carbides as detailed in

Table 1. Work piece material and

dimensional specifications and

layup sequence are detailed in

Tables 2,3 & 4. Apart from

measurement of cutting forces

and surface finish, tool wear

was also observed using optical

and scanning electron

microscopes. Typical observations

are presented in the following

sections

491
TABLE 1. MACHINING CONDITIONS

Cutting direction

Cutting speed

Feed rate

Depth of cut

Tools

Right hand turning

V, m/min

s, mm/rev,

a , mm — 12.7mm square

carbides ( P3 0 , K2 0 ), Coated

(TiN/TiC coated on K-type substrate)

Tool nomenclature :

50 to 200

0.025 to 0.150

1 and 1.5

inserts - Sintered

carbides

Coolant

-6

Dry

-6 75 90 0.8

TABLE 2. WORK PIECE - MATERIAL SPECIFICATIONS

Fibres

a. Carbon rovings - BESFIGHT, 6.4 ii filament dia, 6000

f ibres/ tow

b. Carbon cloth - 0.017 inch (0.432mm) thick, 7.4 oz./yd^.

( 2 3 Ogm/m^

c. Kevlar-49 roving - 4560 DEINER-969, 546 + 16 Tex.

d. Kevlar-49 cloth - 0244, 4 end satin weave, 0.3mm thick,

2 2 0 gm/m^

Resin

a. Epoxy Resin Ciba Geigy LY 556

b. Epoxy Hardner Ciba Geigy HT 972

TABLE 3. WORK PIECE -DIMENSIONAL SPECIFICATIONS

MATERIAL INNER DIAMETER (m) OUTER DIAMETER (m)

CFRP

KFRP

0 . 0975

0. 1020

0.1200

0.1226

TABLE 4. LAY UP SEQUENCE FOR CFRP AND KFRP COMPOSITE TUBES

I Layer ;

II Layer : III Layer : (This layer

thickness was

one layer of cloth with warp placed parallel to

the axis of the tube.

One layer of cloth with warp placed

the axis of the tube.

Hoop winding (90°) with roving

sequence was repeated till the

achieved.

at 45" to

r equ i r ed

492
(a)

Fig. 3 Typical observation of flank wear on (a) P30 and

(b)K20 Cemented carbide tools (CFRP Machining)

RESULTS AND DISCUSSION

Machining of CFRP Composites

Flank wear of plain cemented

carbides

During machining, the

cutting tool flank experiences

severe adhesion due to sliding

against the fibre in the work

surface. This results in the

generation of flank wear. Figure

3 illustrates typical observation of flank wear on tool

grades P30 and K20 (Table 5) in

machining CFRP composites. It is

seen that while P30 tool

experiences abrasive flank wear,

K2 0 tool exhibit relatively

smoother wear surface. The

hardness of P30 tool was 14500

MPa (HV30) , while it was 16000

for K20 tools. Despite the

marginal difference in the

hardness, K20 tool exhibited a

fairly smooth and broader flank

wear. During machining, TiC

particles in the P30 tools can

be dislodged [5] and transferred

to the work surface. On further

sliding, this loosely held TiC

partricles could abrade the

tools, in addition to the

plowing and grooving of the tool

by the carbon fibres. K2 0 tools.

493
TABLE 5. TOOL AND COATING MATERIALS PROPERTIES

PROPERTIES P-30 K-20 TiN Tic

coating coating

nctX.Llllc:so^nvJU^ 1*1 r^d 1X A^ w) nu ou 1X D nU nU nW Z D u u u J

T z

"5 nu nu nu Comp . Strength MPa 4700 5000

Conductivity W/m- K 45 80 46.1 25

Coeff.of Thermal

Expansion 10 m/m- K 6.7 5

which are plain cemented

carbides without the hard TiC

particles, exhibited smoother

flank land associated with

localised small scale

deformation of the contacting

asperities, resulting in

smoother flank wear. This

situation is analogous to

machining of cast iron, where

K20 performs better than P30.

Flank wear of coated carbides

CFRP composites have been

machined with TiN coated (K type

sub-strate) tools. A typical

observation on flank wear of

coated carbide tool is presented

in Figure 4 . It is seen that

the continuous sliding of the

machined work-surface over the

tool flank has displaced the

coating material down to the

flank wear boundary, presenting

typical tear-drop projections.

Tear-drop projections represent

material wiped down to the wear

boundaries by sliding contact.

When the just machined worksurface slides past the

secondary flank of the coated

carbide, localised attrition

between the coating material and

work surface will occur,

resulting in higher contact

temperatures over such spots.

This can facilitate transport of

the coating material

experiencing attrition down to

the flank land boundary.

Comb"- cracking during flank wear

During machining of CFRP

composites, the coated carbide

tools exhibited a pattern of

comb cracking over the flank

land. Figure 5 presents a typical observation of combcracking on TiN coated tools. A

number of parallel cracks

running in a direction normal to

the cutting edge can be seen on

the flank-land. The occurrence

of comb-cracking can be

attributed to alternate thermomechanical stressing of the tool

material during machining of

composites. During machining, a cutting tool usually experiences

higher temperatures at three

distinct zones: depth of cut

line (del) zone, nose, and

secondary grooving zone (Figure

6) . The zones farther away

from the cutting nose undergo

rapid heating and cooling,

experiencing quenching stresses

and cracking [6]. Apart from

this, the orientation of fibres

with respect to the cutting

edge induces force fluctuations.

Thus, the coated carbide would

have experienced thermomechanical stresses, resulting

in cracking of the coating in a

section over the del zone. Once

originated subsequent cracks

occurred in zones closer to the

cutting nose. With continued

machining, comb cracking running

parallel to the cutting edge

494
001

Fig. 5 Comb - cracking observed on TiN coated K-type

substrate tools (CFRP Machining)

495
groove

Fig. 6 Schematic illuslration of

tool wear

also occurred, weakening the

tool-flank and spalling of the

coating. This can be seen in

Figure 5. Spalling of hard and

brittle material is due to decohesion failure and usually

results in sharp spall-boundary.

The occurrence of de-cohesion

failure and consequent spalling

is illustrated in Figure 7.

Flank wear asssociated with

fatigue

Cyclic loading of cutting

zone can promote fatigue of

hard materials of the cutting

tool. Continuation of machining,

even after occurrence of

localised spalling resulted in

higher order chipping of the

coated tool . The observed

chipping of the TiN coated

tools, associated with a dimpled

fracture of the substrate is

presented in Figure 8. The loadreversal during machining of

CFRP composites could have

induced low-cycle fatigue of the

tool material over the cutting

zones resulting in chipping of

the flank portion.

Machining of KFRP Composites

Machining. of KFRP was

observed to be more hazardious

severe, owing to the clogging

and chipping of the cutting

edges by the fussy chip produced

during KFRP machining. This has

resulted in sudden chipping of

the cutting edge, making it

difficult to predict the tool

performance. Figure 9 presents

the types of chips produced

during CFRP and KFRP machining.

Short segmented chips with

broken carbon fibres were

produced during CFRP machining;

while continuous tenacious and

fussy chip in lumps were

produced with KFRP machining.

This has resulted in clogging of

the cutting edge and consequent

chipping. Chipping of P30 tools

Fig. 7 Cohesion failure and consequent spalling of TiN

coated K-type substrate tools (CFRP Machining)

496
Fig. 8 Chipping of TiN coated K-type substrate tool

associated with a dimpled fracture of substrate

Fig. 9 Type of chips produced during FRP machining

(a) CFRP and (b) KFRP

497
^3KV lOOHM 31 . 082

Fig. 10 Notchining over del zone - P30/KFRP Machining

over del during machining of

KFRP composite is presented in

Figure 10. Unlike the case of

carbon fibre, Kevlar is

tenacious and cannot be cut

easily, resulting in producing

fussy chip in lumps. Figure 11

presents bending stress strain

characteristics of reinforcing

fibres [7]. It is seen that,

while carbon is relatively

brittle Kevlar is tenacious,

production of fussy chip in

lumps and continuous rubbing of

chip over the tool material

resulted in spalling of coating

with TiN coated tools as well.

Typical spalling of coating

observed on TiN coated tool is

presented in Figure 12.

Continuous rubbing of fussy chip

in lumps over the cutting zone,

might have caused localised

galling and removal of material

resulting in spalling.

CONCLUSIONS

Observations on machining

of CFRP and KFRP composites have

shown two distinct tool wear

mechanisms; 1) thermo-mechanical

and stress induced tool wear

with CFRP, and 2) dominant chip

notching with KFRP composites,

- Comb-cracking and consequent

spalling by de-cohesion

failure was the dominant

cause of flank wear of TiN

coated carbides.

- P30 carbides experienced

abrasion dominant flank wear,

while K20 carbides exhibited

relatively wider, but smoother

flank wear while machining

CFRP composites.

- All the cutting tools used

for machining of KFRP composites

498
1600

Carbon

Glass

0.5 1.0 1.5 2.0 2.5 3-0

Bending strain %

Fig. 11 Bending stress-strain characteristics of Reinforcing fibres [7]

30HM 35. OS4

Fig. 12 Typical spalling of coating observed with TiN

coated K-type substrate tool ( KFRP machining)

experienced severe chip notching

and galling over the del zone.

- From the types of chips

produced it can be concluded

that for machining of KFRP

composites, a cutting tool

having sharp cutting edge with

as small a nose radius as

possible may be recommended for

effectively cutting the chip,

thereby minimising the

tendency for chip clogging and

notching.

REFERENCES

1. K.Sakuma and M.Seto, "Tool

Wear in Cutting Glass Fibre

Reinforced Plastics (The

Relation Between Fibre

Orientation and Tool Wear)",

499
Bull -of JSME. Vol.26, No. 218,

pp. 1420-1427, Aug. 1983.

2. K.Sakuma and M.Seto, "Tool

Wear in Cutting Glass Fibre

Reinforced Plastics (The

Effect of Physical Properties

of Tool Materials )", Bull . of

JSME. Vol.22, No. 163, pp. 107-

114, Jan. 1979

3. K.Sakuma and M.Seto, "Tool

wear in Cutting Glass Fibre

Reinforced Plastics (The

Relation Between Cutting

Temperature and Tool Wear)",

Bull .of JSME. Vol.24, No. 190,

pp. 748-755, April 1981.

4. W. Schintlemeister , "Cutting

Tool Materials Coated byChemical Vapour Deposition",

WEAR . Vol.100, pp. 153-163,

1984 .

5. G. Santhanakrishnan, "Investigations on Machining of FRP

Composites and Their Tribological Behaviour", Ph. D.

Thesis . IIT Madras, India,

June 1990.

6. R . Krishnamurthy , and A.S.

Choudhyury, "Comb Cracking of

Cemented Carbides", Proc 13th

AIMTDR Conference . Jadavpur

University, India, pp.99-B103,

1988.

7. K.Friedrich, "Ultra High

Strength Materials",

Fotschritt Berichte Per VDI

Zeitschriften, Vol.5, No. 82,

1984/1.

500
THE METALLURGY OF CBN AND ITS WEAR

IN HIGH SPEED MACHINING OF

FERROUS MATERIALS

G. RAI

Smith Megadiamond, Provo, UT

Cubic boron nitride (CBN) tools are increasingly being used in high-speed machining of

materials due to many advantages when using CBN tooling as an alternate to tungsten

carbide or to grinding operations. These are: excellent high-temperature strength, required

to maintain its shape at cutting temperatures; high melting temperature; relative chemical

inertness towards most materials; and a hardness second to only diamond. However, use of CBN in machining of some ferrous-containing materials, which include ferritic steels and

some forms of cast iron, is limited due to higher than expected tool wear. This paper

describes, in general terms, the high-pressure process of synthesizing polycrystalline CBN

tool blanks, the role of second phase material in the microstructure, and its interaction with

CBN grains. The results of an Auger electron microscopy study on diffusion couples

comprising a CBN layer in contact with ductile cast iron, exposed to 1000°C for extended

periods of time, is described. It was found that the concentration of boron in the cast iron

decreased with increasing sputtering time, whereas the nitrogen concentration remained

virtually unchanged. In the CBN layer, however, the carbon concentration at the interface was higher than the bulk concentration in the CBN layer.

INTRODUCTION

The trend toward higher productivity places

a premium on reducing manufacturing time by use of high speed, precision machine tools. A cutting

tool used for these systems must be resistant to fracture and to plastic deformation in order to maintain its integrity. If these requirements are met, then resistance to wear becomes the primary

concern. The maximum cutting speed for the

cemented carbides is limited by their inability to

resist plastic deformation in the machining of steel and other alloys. The carbide family of tools

starts melting around 1250°C and the strength of

carbide falls off rapidly with increasing

temperature. Peak machining temperatures of

about 1000°C are easily attained in practice^^\

resulting in severe plastic deformation at the

cutting edges. The advent of a superabrasive

family of cutting tools, comprising essentially

polycrystalline diamond (PCD) and polycrystalline

cubic boron nitride (PCBN) materials, has made

it possible to further increase the speeds and

feeds.

The use of polycrystalline diamond cutting

tools is somewhat limited to non-ferrous materials

and ceramics due to the affinity of carbon with

most metallic materials, particularly iron based

engineering materials. On the other hand, PCBN

cutting tools have found wide acceptance in

machining a variety of materials ranging from cast iron to superalloys because of their relative

chemical inertness. Many advances have been

made in the development of grades of CBN-based

tools to further reduce the chemical affinity of tool

material with the work piece. The purpose of this paper is to briefly describe the manufacturing

process, the microstructural constituents of PCBN

501
and its chemical interaction with nodular cast iron

as evaluated by a diffusion couple study.

SYNTHESIS OF POLYCRYSTALLINE

CUBIC BORON NITRIDE

The polycrystalline CBN blanks with a tungsten carbide substrate are synthesized by a

high-pressure, high-temperature process. The

starting material, generally, is CBN crystals which

have been cleaned, classified and reduced. The

feed powder, which is a mixture of CBN crystals

and sintering aids, is placed onto a WC-Co disc,

then encapsulated using refractory metals. This

capsule is placed in a heater tube which, in turn,

is surrounded by pyrophyllite, a pressure

transmitting media used in high-pressure systems.

Among the many technologies for achieving

ultra high pressures, cubic press technology has

been found to favor PCD products having curved

surfaces^^\ The cubic press employs the use of

six anvils, as shown in a schematic diagram

(Figure 1), to apply pressure simultaneously on all six faces of a pyrophyllite cube, which houses the

heater tube and the encapsulated CBN material for

sintering. The pressure applied is in the range of

50 to 70 Kbars (5 to 7 GPa) above the determined

critical pressure for HBN to CBN transformation,

as shown in Figure 2. In some instances

hexagonal boron nitride (HBN) is used along with

cubic boron nitride crystals and catalysts, which

could be metals or metal nitride. Hexagonal

boron nitride, in this case, is dissolved into the

solvent or catalyst within the high-temperature,

high-pressure CBN stability region. Cubic boron

nitride then very likely nucleates heterogeneously on the existing CBN sites and grows. Control of

the nucleation and growth processes is used to vary the physical properties of the polycrystalline

composite. The sintering process utilizes liquid

phase to promote intercrystalline CBN bonding.

The liquid phase is provided by melting of the

carbide substrate at the CBN-carbide interface.

This seems to be achieved by the extensive

alloying effect of boron and nitrogen^^^ with the

pseudo-eutectic phase of tungsten carbide cobalt,

which is reported to wet CBN crystals extremely

well. The liquid phase infiltrates the CBN powder

bed by capillary action and fills interstitial positions.

In liquid phase assisted sintering some

shrinkage occurs as a result of particle

rearrangement, which is facilitated by the

lubricating action of liquid and driven by its surface tension, which pulls the particles together. In addition, regions of forced contact between

CBN particles are under mechanical stress, which

locally increases the particles' chemical i potential^'*^ This results in more rapid solution of

j those particle regions into liquid phase, with ' matter being transported by diffusion in the liquid

to be deposited elsewhere. Therefore, the particle centers approach each other and the pores shrink i just as in the solid state sintering process. ': MATTER TRANSPORT IN ULTRA HIGH

PRESSURE SINTERING

Densification by solution-precipitation in

liquid phase sintering is largely governed by j diffusion of material through the liquid phase, i The diffusion of matter can be accomplished by I lattice or grain boundary diffusion, as shown in

Figure 3. Since CBN crystals are the closest

\ packed structures and are comprised of atoms of i different chemical nature, the lattice diffusion

contribution to sintering is very limited. On the

other hand, atom transport by grain boundary

diffusion is the more dominant mechanism for

sintering. The role of pressure in sintering of

CBN grains is essentially to provide a stable I region in which back conversion of CBN to HBN

is prevented. Its effect on the diffusion constant,

whether for boundary or lattice diffusion, is not well understood, but can lead to very high

compressive stresses at grain contacts. According

to the Gibbs-Thompson equation^^, the increase in chemical potential or solid solubility due to stress

is given as:

j aC = C„ [ exp ( a^p, v /kT ) - 1 ] (1)

where aC is the solubility increase, a^ppi is applied

stress, Co is the solubility under no applied stress,

V is the molecular volume of the solid, T is temperature, and k is the Boltzman constant. For

a two to three times increase in effective stress,

solubility can easily be increased by 50 to 70% at 1300°C. The effect of ultra high pressure,

therefore, appears to be two fold: it increases the

solubility of CBN particles in the liquid metal and

m
increases the diffusion rate. Both of these tend to enhance the sintering rate. DEVELOPMENT OF CBN

MICROSTRUCTURE

During the sintering process, the initially

separate CBN crystals are transformed into grains

of a polycrystal that contact their neighbors on

nearly all sides. The void space between particles

is filled with the liquid pool of metal infiltrating

from the molten tungsten carbide interface. At

first these are long channels that extend to the

surface, but later develop more defined shapes.

Sometimes these channels are quite thick and are devoid of any CBN grains. Figures 4a and 4b

show features of these channels, also called metal

"rivers", that are found in the high-pressure

sintering process. Figure 4b is a higher

magnification microphotograph showing details of

precipitation of tungsten carbide grains from the

metal pool formed during cooling. It is necessary

to control formation of these metal rivers, as these

constitute weak regions in the cutting tool leading

to excessive wear.

The microstructure of polycrystalline cubic

boron nitride essentially comprises CBN grains of

1 to 3 micron in size. The metallic second phase

is found along the triple points and also along

grain boundaries, as shown in Figures 5a and 5b

at two magnifications. The size of the metal pool

is larger at the triple points compared with the

metal pool size on the grain boundaries. Figures

5c and 5d are scanning electron micrographs of

specimens subjected to leaching for one hour in a mixture (50% by volume) of nitric and

hydrofluoric acids. The figures show extensive

intercrystalline bonding of CBN crystals, with

voids representing prior presence of metallic

catalyst material. A qualitative energy dispersive

analysis suggested that the metal pools are

essentially W-Co alloys with smaller amounts of

Al. The amount of metal as determined by

density measurement generally varies between 20

to 24% by weight of the CBN layer.

WEAR OF CBN TOOLS DURING

MACHINING

Non-catastrophic tool wear (viz., gradual

wear) during machining or turning operations can be grouped into two regions:

1. Low Speed or Mechanically Activated

Wear. The primary mode of wear in most

ceramic and hard materials falls under this category, where removal of tool material occurs by scouring action of hard phase inclusions or protruding asperties in the work piece^^\ Thermal

and mechanical fatigue may also play some role in

abrasion wear^^. While hardness of the tool

material is the most obvious and significant

consideration in abrasive wear, the role of

toughness of the tool material is not so well

defined. In brittle materials, at a critical load, a sub-surface median vent crack is formed on the

plane of the applied stress. With additional

loading this sub-surface crack grows to a critical size, at which point it becomes unstable and a

catastrophic failure ensues^*\ In tool materials

such as CBN, abrasion can be assumed to take

place when two adjoining cracks intersect. The

removal volume is then related to the length of

cracks in the tool. Laboratory research shows that

abrasion wear of tool materials involves

microfracturing of grains, indicating that

toughness of grains plays a major role in their wear resistance. The wear rate (W) is found to relate to fracture toughness (Kjc) by the following

equation^^''°\ where A is a constant: W = A/(Kic)' (2)

This seems to have been found true for carbide

cutting tools also. It is quite clear that for similar

physical properties of cutting tool material, a tougher tool material will show better abrasion

resistance.

2. High Speed or Thermally Activated

Wear, The physical mechanisms that are responsible for producing gradual wear depend on

cutting temperature. For each combination of tool

and work piece materials, the transition

temperatures between wear regimes are determined by the chemical and mechanical

properties of the tool-work piece system. In low

speed operations, chemical interaction between the

tool and work piece is insignificant. In high speed

machining, where temperatures^^ can rise in excess of 1000°C, the chemical stability of the
tool material with respect to the work piece

becomes a very important contributor to wear. The chemical solubility of the tool material in the

work piece increases exponentially with

temperature. In essence, the tool material

dissolves in the hot flowing chips and the transport

of material from the tool cutting edge to hot chips can be visualized to occur in two steps, namely:

transfer of tool material by solid state diffusion across the interface and transport of dissolved tool

material from interface to chip bulk^'^\

The wear of PCBN tools in the machining of

ferritic materials has been studied. Heath^^^^

observed the reactivity of PCBN with soft nonmartensitic phases of iron at relatively higher

speeds. There are evidences in literature^^'*^ indicating that machining non-hardened steels with

PCBN was not economical due to "diffusion

wear" resulting from high solubility of boron and

nitrogen atoms in ferrite. Aspinwal, et al^'^ reported that wear of PCBN tools is larger when

cutting ferritic, as opposed to pearlitic, cast iron

with surface speeds in excess of 90 m/min. They

also found that relatively small changes in

chemical composition of the work piece may

produce very large changes in wear rate.

The flow of chips across the tool face is reported to produce peak temperatures of 1000°C

in the vicinity of the cutting edge. The peak

temperature seems to be governed by the nature of

chip formation Vv'hich could change from

continuous to discontinuous. It was reported that crater wear in PCBN tools machining hardened

M2 steel was lower than in its soft condition.

Soft M2 steel exhibited continuous chip formation

which became discontinuous when steel was

hardened. In the milling of cast iron at speeds of

1,800 surface meters per minute, the PCBN wear

rate for pearlitic steels was significantly lower

than for ferritic steels. The segmentation of the

pearlitic chips was more pronounced than ferritic

steel^^^\ It seems very likely that wear of PCBN

in machining at high speeds is governed by a chemical dissolution process where tool

constituents, namely the tungsten carbide-cobalt

binder phase and boron and nitrogen atoms,

interact mechanically and chemically with the

work piece. It appears, however, that chemical

interaction characterized by solid state diffusion of

atoms is a more dominant wear mechanism of

PCBN tools.

DIFFUSION COUPLE EXPERIMENTS

Since both B and N exhibit solid solubility in

ferrite, a diffusion study involving PCBN and

ferrite is very useftil in understanding the extent of

matter transport across the interface. For this purpose, ferritic cast iron and PCBN tools of the same size were polished and pressed using an

Instron machine to provide an intimate contact. These were held together tightly with the help of

a vice. The difftision couple was then annealed

under pure hydrogen in a furnace maintained at 1000°C for 40 hours. The sample was allowed to cool in the ftirnace.

AUGER ELECTRON SPECTROSCOPY (AES)

OF CAST IRON AND PCBN

The extent of inter-difftision of CBN and

ferritic cast iron constituents was analyzed using

Auger electron spectroscopy (AES). Both the cast iron and CBN pieces were first cleaned and then

sputtered to remove the outer layer to a depth of

approximately 40 to 50 A. The sputtering rate was determined to be 5 A per second. The

elemental analysis of both the cast iron and CBN

pieces was continued as a function of sputtering

time until no changes in the atomic concentration

of elements could be detected. The atomic

concentration profiles as a function of sputtering

time for different elements in cast iron and CBN

are shown in Figures 6 and 7, respectively. From

these figures the following inferences can be

drawn:

1 . The atomic concentration of B and N does

not change after about 2 minutes of sputtering,

indicating that diffusion is limited to a depth of

less tlian approximately 0.05 micron.

2. The B to N atomic ratio decreases from

the bulk of the tool material to the interface of the

difftision couple. This indicates a faster rate of B

diffusion into cast iron as compared to nitrogen.

3. The atomic concentration of C changes

significantly from the diffusion interface to the

bulk of tool material and remains approximately

constant after a depth of approximately

0.08 micron.
4. The Fe atomic concentration in the CBN

material does not significantly change with

sputtering time, indicating a lack of Fe diffusion

in the tool material.

5. The atomic concentration of nitrogen in

the cast iron piece does not change with sputtering

time, suggesting that nitrogen is probably lost at the interface in the form of gas and only B

diffuses into cast iron.

Since the AES system analyzes constituents over an area of 1 mm x 1 mm, it also includes the

light element analysis in the second phase (WCCo) of the PCBN structure. It is, therefore, quite

likely that diffusion of elements such as carbon

into the tool material may also have occurred

through the WC-Co phase of the structure. During machining using PCBN cutting tools, the

moving chip provides a most efficient sink for

diffusion of tool constituents. Since the tool is always in contact with the newly formed surface

of the work piece, die diffusion rates can be faster

than predicted by the static diffusion couple study.

SUMMARY

PCBN cutting tools provide an excellent

combination of hardness, toughness and relative

chemical stability to most work piece materials.

However, chemical interaction between tool

material and some forms of cast iron appears to be

a major cause of wear. Interdiffiision between the

tool and work piece materials could take the form

where B and N essentially dissolve into the work

piece; the work piece constituents, on the other

hand, dissolve into the tool material mainly

through the second phase (WC and Co).

ACKNOWLEDGEMENT

The author expresses his deep sense of gratitude

to Heidi Allen for her diligence in preparing this

manuscript and to David Miess for his assistance.

REFERENCES

1. M. C. Shaw, "Temperatures in Cutting,"

Proceedings of the Special Prod. Engineering

Conference, ASME, Special Volume, Nov.

1988.

2. H. T. Hall, "High Pressure-Temperature

Apparatus," Metallurgy at High Pressures and

High Temperatures , edited by K. A.

Gaschneider and M. T. Hepworth, Gordon

and Breach Science Publishers, NY, 1964.

3. G. Rai, "Microstructure and Mechanical

Properties of CBN Twinserts," Proceedings of

the Int'l Conference on New Diamond

Science and Technology, pp. 1069-1079,

1991.

4. J. H. Rosolowski and J. E. Burke,

"Sintering," General Electric Corporate

R&D Technical Report, 1988.

5. O. H. Kwon and G. L. Messing, "HIP of

Liquid Phase Sintered Ceramic Composites:

Tailoring Multiphase and Composites

Ceramics," Materials Science Research , edited by R. E. Tressler, G. L. Messing, C.

G. Pantano, and Robert E. Newnham, Vol.

20, pp. 41-53, 1985.

6. J. G. Baldoni and S. T. Buljan, "Ceramics

for Machining," Ceramic Bulletin . Vol. 67,

pp. 381-387, 1988.

7. E. Rabinowicz, Friction and Wear of

Materials . NY: John Wiley, 1966.

8. A. Misra and I. Finnic, "A Review of the

Abrasive Wear of Metals," trans. ASME,

Vol. 104, pp. 94-101, 1982.

9. S. T. Buljan and V. K. Sarin, "Machining

Performance of Ceramic Tools," Cutting

Materials , edited by F. W. Gorsler, AIME,

Metal Park, OH, pp. 335-348, 1981.

10. J. G. Baldoni, S. F. Wayne, and S. T. Buljan, "Cutting Tool Material: Mechanical

Properties - Wear Resistance Relationship,"

trans. ASME, Vol. 29, pp. 347-352, 1986.

11. B. M. Kramer and N. P. Suh, "Tool Wear by

Dissolution: A Quantitative Understanding,"

trans. ASME, J. Engg, for Industry . Vol.

102, pp. 303-312, 1979.

505
12. P. J. Heath, Ultra Hard Materials . ASM

handbook, 9th ed.. Vol. 16, pp. 333-352,

1987.

13. P. Hancock, "Coated Cutting Tools - Influence of Composition and Surface

Integrity in High Speed Machining," Final

Technology Report. Cranfield Institute of

Technology, 1978.

14. D. K. Aspinwal and W. Chen, "Report on CBN Tooling Applications in Machining of

Iron," University of Birmingham, 1987.

15. P. A. Dearnley, "A Metallurgical Evaluation

of Tool Wear and Chip Formation When

Machining Pearlitic Grey Cast Irons With

Dissimilar Graphite Morphologies," Wear. Vol. 101, pp. 33-68, 1985.

16. Z. Palamii, "Cutting Temperature in

Intermittent Cutting," Int'l J. of Machine

Tools and Manufacture . Vol. 27, pp. 261-

274, 1987.

506
I

I

Fig. 1. Anvil configuration of Cubic Press

507
200

180

160

140

120

P 100

(Kb)

80

60

40

20

(B)

WURTZITE BN

CUBIC BN

(ZINCBLENDE)

LIQUID

^^-^ HEXAGONAL BN

(LAYER LATTICE) i

1000 2000

T(K)

3000 4000

15

10

GPa

Fig. 2 Thermodynamic stable regions for conversion and growth

of cubic boron nitride.

508
t

Fig. 3 Transport of matter between cBN particles through grain boundaries

and bulk diffusion. Pores disappear when they are touching grain

boundaries.

509
Fig. 4 Optical photomicrographs of cBN sintered under ultra high pressure

conditions showing (a) channels or rivers of metallic liquid which have formed

(1 OOx), and higher magnification view (b), of the same showing WC precipitation

from liquid upon cooling. (500x).

510
Fig. 5 Scanning electron micrograph (a) atSOOx of polycrystaliine cubic boron

nitride sintered using ultra high pressure technology. Micrograph (b) at 1 0OOx

shows catalyst material which can be seen at grain boundaries and at triple points.

511
c

d

Fig. 5 Scanning electron micrograph (c) of polycrystalline cubic boron nitride

sintered using ultra high pressure showing etched structure atSOOx. Micrograph

(d) is of the same structure now atSOOOx showing intercrystalline cBN bonding.

512
80

0 2 4

SPUTTERING TIME MIN.

Fig. 6 Concentration profile of elements in cast iron as determined

byAES after diffusion.
ATOM

%

024

SPUTTERING TIME MIN.

Rg. 7 Concentration profile of elements determined by AFS

in the PcBN tool material after diffusion.

514
LASER AND ELECTRICAL

DISCHARGE MACHINING

515

CRACK-FREE PROCESSING OF HOT-PRESSED

SILICON NITRIDE CERAMICS USING PULSED

YAG LASER

N. MORITA

Chiba University, Chiba, Japan

This paper describes the effect of crack-free laser processing of hot-pressed

silicon nitride ceramics on fracture strength and residual stress. The fracture strength

of the processed test pieces was determined by three-point bending tests. The residual

stress in the processed surface was characterized by X-ray diffraction measurement.

The strength of the laser- processed workpieces was reduced to 90-80% compared with

that of the ground workpieces because of subsurface layer damage, which included the

heat-affected layer and the residual stress layer. In particular, the latter greatly

affected the fracture strength because the higher compressive residual stress layer

generated during grinding was released by laser processing. It is concluded that the

newly developed crack-free processing technique has great industrial possibilities.

1. Introduction

Advanced fine ceramics have many excellent

physical and chemical properties. However, their hardness and brittleness make them difficult to machine. The focused laser beam is one of the tools which can be used to machine fine ceramics'""''". However, the laser machining method has some diss- dvantages such as the generation of recast layers'^*

and the development of thermal -stress- induced

cracks'^'. These defects decrease the fracture strength

of ceramics, and also worsen other good characteristics'^'. This is probably the primary reason why laser applications for machining ceramics have been

limited*«'''^'. This paper describes the feasibility of a crackfree processing method for hot-pressed silicon nitride (Si3N4) ceramics using an yttrium aluminum garnet

(YAG) laser. The influences of the laser waveforms,

such as pulse duration, peak output power and pulse

repetetion frequency on the generation of cracks and

recast layers during laser drilling are evaluated. The

fracture strength and the residual stress are measured

for the purpose of evaluating the crack-free surface.

2. Influences of Laser Waveforms on the Generation of Cracks and Recast Layers

2. 1 Experimental procedure

Two types of YAG laser oscillators, a pulsepumped YAG laser with pulse durations of 120 ^.s to 7 ms and a continuous-pumped YAG laser, were used in the present experiments.

Moreover, in order to obtain the laser pulses with

durations in the order of magnitude of 10 us, we set up

a 100 mm-diameter aluminum disk with a slit, as

illustrated in Fig. 1, and rotated it at 10 000 revolutions per minute using a DC motor so that the pulsed YAG

laser beam was partly chopped by the slit through

which it passed. The chopped laser beam with durations of 20 us, 65 us and 110 ^ls can be obtained by

using disks with slit widths of 1 mm, 3 mm and 5 mm,

5i7
respectively.

In addition, the laser pulses with durations of 160 ns to 700 ns were oscillated by switching the AO

(acoustooptic) -modulators mounted on the continuous-pumped YAG laser. These short pulses generated

by AO-modulator operation are generally called Qswitched pulses.

The pulsed laser waveforms which were obtained

in the manner described above are shown in Fig. 2. Table 1 shows combinations of laser parameters

examined in the experiment, such as pulse duration

( / ), output energy per pulse (E), the number of irradiated pulses (A'^) , and the peak output power iP).

The number of irradiated pulses was controlled so

that the total output energy Et{ = ExN) was

maintained constant at 24.4 J, which was based on the

energy of the longest pulse duration of 7 ms. The peak

output powers were made nearly fixed at 4 kW,

regardless of pulse durations, except that of the Qswitched pulse.

The hot-pressed silicon nitride ceramics (Si3N4)

which were sintered using Y2O3 and AI2O3 as sintering

aids, were investigated. They were ground into 10 x

15 X 5 mm size by the resin-bond diamond wheel of

# 400 grains. The properties of Si3N4 ceramics are

indicated in Table 2. Si3N4 ceramics are one of the

ceramic materials which are very sensitive to constructive defects such as cracks and voids. In the

observation of recast layers and cracks, a scanning

electron microscope (SEM) and an X-ray microanalyzer (XMA) were used.

2. 2 Results

Si3N4 ceramics decompose into silicon vapor and

(-Rotation disk

Mirror Lens Pulsed YAG laser

Lens Laser beam

DC motor

Workpiece

Slit width : 1,3,5 mm

nitrogen gas directly at about 2 150 K under the

condition of slow heating at atmospheric pressure. Rapid laser heating and vaporizing may very possibly

produce the liquid silicon phase around the laser- processed surface. Figure 3 shows the state of the

crack development in the laser-drilled hole. The

crack propagates from a recast layer (which looks

white in the SEM photograph) into the matrix.

Table 1 Combinations of laser parameters examined in the experiment

t 160ns 500ns 2Pps 65^is llQiJS 120ps 800ps 7ms E lOnJ 3.7mJ 75(nJ 230mJ 380mJ 670nvJ 3.2J

P 15.4 A.O 3.6 3.6 3.5 5.6 h.O 3.5

<uf

8.

a.

El ;e2;

Pulse duration:

Energy :

Table 2 Properties of Si3N4 ceramics.

Specific heat (J/kgK) 712.3

Density (kg/m^) 3260

Thermal conductivity (W/nrK) 29.3

Poisson ' s ratio 0.27

Young's modulus (GPa) 303.9

Thermal expansion 3.2

coef f icient(xlO-6/K)

(a)Q-switching pulse (b)PuIse chopped by

rotation disk

t=160ns

I

t i

1

20fjs/div.

Fig. 1 Experimental equipment to obtain laser pulses

with durations of the order of 10 ^s

0.2|js/div.

(c)Normal pulse

t=0.8ms

0.2ms/div.

Fig. 2 Typical pulsed laser waveforms

518
Figure 4 shows the characteristic X-ray photographs of the recast layer. The recast layer contains

richer silicon and yttrium elements than dose the

L Heat-affected layer

Recast layer

Fig. 3 State of the crack development in the laser-drilled hole

(b)Si (c)Y

Fig. 4 Characteristic X-ray photographs of the recast layer

Si(Reference)7^Si02(Reference)

'v-Recast layer

Crystal: PET Si-Kp spectra

6.768

Wavelength A

Fig. 5 Silicon X-ray emission spectrum observed in the recast layer

matrix. Figure 5 illustrates the silicon A'/? X-ray emission spectrum observed in the recast layer. X-ray

analyses identified the recast layer as Si, not as Si.N, or SiO^. From these observations, we see that the

recast layers are mainly composed of silicon. We believe that the recast layer is formed

because the silicon vapor generated by decomposition

of SijNj is saturated in the hole and condenses into

liquid during laser heating.

Figure 6 shows the dependence of the crack

length and the recast layer thickness on the peak

output power. The increase in the peak output power

causes the crack propagation and thick recast layer

formation. It is necessary to maintain the peak output

power as low as possible for crack-free machining

ceramics.

1

60

50

w 1 i

40

OOc

01 u

i-H 01>, 30

20

CO '

Recast 10

0

Pulsed YAG laser Ma tec ial : S i3NAA^

Lens focal length f:50 mm y^.kl^ J Defocus length def:0 ram E=3.21J^^

O Crack A Recast layer

,1.96 J ^ ^

' A^^''"'"^ Pulse duration t:0.8 ms Shot number: 10 shots

3 4 5 5 7 B

Peak power :P kW

Fig. 6 Dependence of the crack length and the recast

layer thickness on the peak output power

70

60

50

40

30

20

10

0

Pulsed YAG laser Material :Si3N4

, Lens focal length f:50, 75mm Q

Defocus length def:0 mm

Hole-,

Crack A\~\J O Crack ? ,A

A Recast layer J

10- 10^ 10-5 10"^ 10-^ Pulse duration: t s

Total pulse energy Et:25 J

10-

t 160ns 500ns 20ps 65ps llOps 12Qus 800ps 7ms E lOmJ 3.7nU 75mJ 230mJ 380mJ 670mJ 3.2J

N 2500 6600 326 106 65 37 8 1 P 15.

4.0 3.6 3.6 3.5 6.9 4.0 3.8

Fig. 7 Influences of the pulse duration on the crack

length and the recast layer thickness

519
Figure 7 illustrates the influences of the pulse

duration on the crack length and the recast layer

thickness. The crack length and the recast layer thick- ness are reduced with a decrease in the pulse duration.

The crack length is smaller than the recast layer

thickness when the pulse duration is less than 65 \is. This indicates that shallow cracks are developed

within the recast layer without the growth of cracks

into the matrix. In particular, no cracks were obser- ved in the recast layer when Q-switched pulses with a duration shorter than 500 ns were used. This suggests

that the thermal stress distribution induced by a short

laser pulse is only localized in the laser-heated shal- low area. Figure 8 shows the dependence of the crack

length and the recast layer thickness on the pulse

repetition frequency. Repetition frequency higher than

10 kHz thickens the recast layers and propagetes the

cracks suddenly into the matrix. This result suggests

that the short laser pulses with high repetition fre- quency establish a thermal stress distribution similar

to that induced by a long laser pulse. The critical repetition frequency for crack- free processing shifted

to the side of frequency lower than 10 kHz when the

higher peak output power was used. Therefore, in order to achieve crack-free machining of SisNi ceramics, it is necessary to process them

using a laser beam pulsed for a duration of under 500 ns at a repetition frequency of less than 10 kHz. It appears that the crack development results from the

thermal shock and the thermal expansion mismatch

between the recast layer and matrix.

3. Fracture Strength and Residual Stress of

Crack-Free Surface

3. 1 Experimental procedure

Two techniques were applied to the evaluation of

laser-processed surface integrity. One technique

involved measurement of the fracture strength by the

bending method, and another was residual surface

stress measurement by an X-ray diffraction technique. The X-ray method has previously been suc- cessfully used to measure surface stress in Si3N4

ceramic material""*. Hot-pressed silicon nitride (Si3N4) ceramics,

which were mixed and sintered with yttria (Y2O3) and

alumina (AI2O3) as sintering aids, were examined in the experiments. The properties of the Si3N4 ceramics

are shown in Table 2. Standard test pieces were prepared for the threepoint bending tests provided in JIS R-1601. They

were finished to 3x4x40 mm size and 0.8 ^m surface

roughness (Rmn^) by a surface grinding process under

the conditions shown in Table 3 with resin-bond

diamond wheels.

In the next process, one side of the standard test pieces was finished by a crack-free laser milling

process*"'under the conditions shown in Table 4 using

a continuous-pumped and Q-switched YAG laser beam. The laser milling process is illustrated in Fig.

9. The removal depth of the laser milling was fixed at

x: .H

00 4J

25

20

15

10

5

Q switching pulse Material :Si3N4

Peak power P:6.6 kW Shot number: 3000 shots

Lens focal length f:75 mm

Defocus length def :0 mm O Crack A Recast

t=400ns

Pulse frequency kHz

Fig. 8 Dependence of the crack length and the recast layer thickness on the pulse repetition frequency

Table 3 Surface grinding process conditions

Process Machining conditions

Rough Wheel 200 IQt mm, #200, Resin

grinding Depth of cut 1 2 pm

1st Feed rate 15 m/min(X),4 mm/pass(Y)

Rough Wheel <^200 lot mm, #A00, Resin

grinding Depth of cut 5 ^im 2nd Feed rate 8 m/min(X),2 mm/pass(Y)

Finish Wheel i))200 lOt mm, #600, Resin

grinding Depth of cut 2 pm

Feed rate 5 m/min ( X) , 0 . 5mm/pass ( Y

Wheel rotation rate : 2850 r.p.m.

Table 4 Crack-free laser milling process conditions

jnditions

Depth of cut (mm) 0.1 oTT

Pulse repetition

frequency (kHz) 320 1000 Surface roughness

Rz ipm) 3 5 Feed rate (mm/s) 2.80

Scan width (pm/pass) 50 Assisting gas (kgf/cm^) Air(2 .5)

Laser beam

r-1

\ ! Removal

\ I Scan , ^, '// ^ depth

Fig. 9 Laser milling process

520
0.1 mm and 0.3 mm by controlling the pulse repetition

frequency to keep it at the levels of 320 Hz and 1 000 Hz. respectively. Under these conditions, no cracks were observed in the very thin recast layers which were formed as a reduction product on the laser- milled surface. The bending strengths of the standard

test pieces and laser-milled ones were analyzed by the

Weibull statistics method.

The test pieces were ground into 10x15x5 mm

size for the X-ray residual stress measurement under

the same conditions shown in Table 3 . The residual

stress in the ground and laser-milled surface was

measured under the X-ray diffraction conditions given

in Table 5 . In the present experiments, strain normal

to the (411) lattice planes of /3-Si3N4 was measured

using Cr-Ka radiation. In this case, high resolution

against depth is attained because the X-ray penetration depth for the (411) lattice planes is small, 6^11

ptmTM. We used the sin^ <p method for the determination of residual stress.

3.2 Influences of crack-free laser processing on

fracture strength

It is necessary to investigate the influences of crack-free laser processing on fracture strength to determine its practicability. Figure 10 shows the

Table 5 X-ray residual stress measurement conditions

x-ray diffraction conditions

X-ray diffraction method

Characteristic X-ray

Diffraction plane

Diffraction angle (deg)

Filter

X-ray-irradiated area (mm^)

Tube voltage (kV)

Tube current (mA)

Divergent angle (deg)

Scanning speed (deg/min)

Parallel beam method

Cr Ka spectrum

3-Si3N4 (411) 125.58 Vanadium foil 15y(2~3)

50400.5 0.45

Table 6 Average fracture strength and Weibull modulus

obtained from the results shown in Fig. 2 Machining method Average fracture strength(kgf /mm2) Weibull modulus

(A) Grinding (#600) 93.9 18 7 (B) Laser machining (O.lmni 82.6 17 2 depth) after (A)

(C) Laser machining (0.3mm depth) after (A) 7A.1 ' 24 8

Table 7 Annealing conditions

Weibull plots of fracture strengths of the three point

bending test pieces with various depths of laser

milling. Table 6 shows the average fracture strength

and Weibull modulus obtained from the results shown

in Fig. 10. The average fracture strength of laser- milled test pieces is smaller than that of the ground

ones, 88% for 0.1-mm-depth milled test pieces and

79% for 0.3-mm-depth milled ones, respectively. The

strength reduction of 10 to 20% seems to depend on

both the heat-affected layer and residual stress layer. In the next step, lapping tests and annealing tests were carried out to determine the reason for the

99.9

99.0

90.0

in 50.0

10.0

5.0

1.0

(a)Grinding(#600)—

(b) Laser machining

(O.lmm depth)

(c) Laser machining

(0 . 3mm depth)

- after (a) 1

3-point bending test Load , , i'

(a)

Machined surface i 1

2/

^« -

30 40 50 60 70 80 90 100 150 Fracture strength (kgf/mm2)

Fig. 10 Weibull plots of fracture strengths of the threepoint bending test pieces with various depths of

laser milling

99.9

99.0

i I r

-(a)Grinding(#600)-

90.0

50.0

10.0

5.0

- after (b)

Test piece Annealing condition

1 Ground (Standard) in Ar gas at 1000 "C for 12 hours

Laser i 1 1 ed in N2 gas at lOOO'C for 10 hours Fig.M

1.0

(b) Laser machining

(O.lmm depth) after (a)'^

(c) Lapping(5pm depth)

j

•P/

(b) (c)

(a)

3-point bending test Load * /

Sample

Machined surface

J/

30 AO 50 60 70 80 90 100 Fracture strength (kgf/mm^)

Weibull plots of fracture strengths of lasei test pieces and lapped ones

150

521
strength reduction after laser-milling. Annealing

conditions are shown in Table 7. Figure 11 shows the Weibull plots of fracture strengths of laser-milled test pieces and lapped ones. The increase in fracture

strength after lapping, which was about 20 MPa (2 kgf/mm^), was very small. Also, the strength of

sufficiently lapped test pieces did not recover to the

previous level. The small increase in strength by

lapping shows the existence of subsurface damage

layers such as brittle thin recast layers with microcracks. The difference in fracture strength between

the ground test pieces and lapped ones suggests that compressive residual stress is generated in the test- piece surface after grinding.

It is well known that the residual compression

stress layer established by diamond grinding, which is about 20 \im thick, improves mechanical properties

such as fracture strength and toughness"". Figure 12 shows the Weibull plots of fracture strengths of the

ground testpieces and annealed ones after grinding.

The decrease in strength by annealing explains the

release of residual compression stress because the

lapping treatment after annealing did not make the

strength recover.

It seems that not only the recast layers, but also the removal of the residual compression stress layer causes the decrease in fracture strength in laser

milling. Figure I3 shows the Weibull plots of fracture

strengths of the laser-milled test pieces and annealed

ones after laser milling. The average strength is almost constant after annealing. This means that residual stress does not develop in laser processing.

3^3 Influences of crack-free laser processing on

residual stress The residual stress after machining has direct

effects upon the mechanical properties of materials.

Table 8 shows the results of the experiment on the

effect of surface finishing on residual stress. A negative value of stress means compression. The residual

stress is compression independent of the nature of the

finished surface. The residual stress at any laser- milled depth is smaller than that of the ground surfaces.Table 9 shows the results of the experiment on

the effect of surface finishing on both average fracture

strength and Weibull modulus. The difference in strength between the ground test pieces and annealed

ones after grinding is about 48 MPa (4.9 kgf/mm^)

This value reflects the residual stress in the ground

surface, as shown in Table 8, (A). This proves that the decrease in strength after annealing depends on

the release of residual surface compression stress. The difference in strength between the laser-milled

test pieces and annealed ones after laser milling is close to the residual stress in the laser-milled surface, as shown in Table 8, (B) and (C).

Table 8 Results of the experiment on the effect of sur- face finishing on residual stress Measured surface Residual stress (kgf/mm2)

(A) Ground (#400) -9.82+1. 14 (B) Ground (#400) & laser machined (0.1mm depth)

-2.34±0. 98

(C) Ground (#400) & laser machined (0.3ram depth)-1.19+0. 98

99.9

99.0

90.0

50.0

10.0

5.0

1.0

\ \ \ (a)Grinding(#600)

.(b)Annealing' after

grinding

3-point bending test Load . , ^

^Sample

Machined surface

30 AO 50 60 70 80 90 100 150 Fracture strength (kgf/mm2)

Fig. 12 Weibull plots of fracture strengths of the ground

test pieces and annealed ones after grinding

99.9

99.0

90.0

50.0

10.0

5.0

1.0

•(a)Grinding(#600)

(b) Laser machining

(O.lmm depth)

after (a)

(c) Annealing after

laser machining

3-point bending test Load Sample

Machined surface

30 40 50 60 70 80 90 100 150 Fracture strength (kgf/mm^)

Fig. 13 Weibull plots of fracture strengths of the laser- milled test pieces and annealed ones after laser milling

522
Table 9 Results of the experiment on the effect of sur- face finishing on both average fracture strength

and Weibull modulus

M3ctiLni.ng niGttiod Av€ir3g6 £ir3ct.uir€ strength(kgf /mm2)

Uowt; i Ki 1 1 1

J. uu 1 1 modu 1 u s

7a) Grinding (#600) 93.9 18 7 (B) Annealing in Ar gas 89.0 16 1 after (A)

(C) Laser machining (0.1mm 82.6 17 2 depth) after (A)

(D) Annealing in N2 gas 78.7 29 3 after (C)

(E) Lapping (5pm depth) 84.9 12 2 after (C)

Figure 14 shows the results of the SEM observation of the laser-processed and ground surfaces. It was found that the recast layers are mainly composed

of porous silicon oxide (SiOa) on the laser-processed

surface, with no cracks. The surface roughness value was smaller than 6 jim /?max. From the experimental results described above, we may conclude that crack-free processing using a Q

4

Hill 1 1 IliHlMffTriiVfai

( a ) Ground surface

( c ) Lapped surface

Fig. 14 Results of the SEM observation of the laser- processed and ground surfaces

-switched YAG laser makes the fracture strength 10 to 20% lower because of the brittle recast layer and

the release of residual surface compression stress after grinding. The combination of crack-free processing and diamond grinding will be more effective in the improvement of surface integrity.

Laser beam

Fig. I? Schematic illustration of the laser-milling process

a) Slope shaping

b) Sinusoidal plana shaping

Fig. 16 Cross sections of samples which were milled b> controlling the pulse repetition frequency ( a

linearly and (b) sinusoidally.

4. Applications

We have applied crack-free laser processing to three-dimensional shaping. Figure 15 shows a sche- matic illustration of the laser-milling process'*^'. The

pulsed laser beam is scanned on the test piece surface

by NC-controlled stage feed. The pulse repetition

frequency was changed for the required machined

depth on each position of the test pieces. Figure 16 shows cross sections of samples which were milled by

controlling the pulse repetition frequency ( a ) linearly

and ( b ) sinusoidally'"^.

The laser-turning process was developed for cylindrically shaped test pieces. Figure 17 illustrates the laser- turning process. The laser beam was

523
Laser beam

Fig. I7 Laser-turning process.

Fig. 16 Screw-cutting sample with 8 mm diameter

irradiated in a tangential direction to the rotational

axis to facilitate control of the machining depth.

Figure IB shows a screw-cutting sample of which the diameter was 8 mm"*'.

5. Conclusions

We have investigated the influences of heataffected layers and residual stress on fracture

strength for the purpose of evaluating the practicability of our newly developed crack-free processing

technique for ceramic materials. The conclusion

reached is that the 10 to 20% decrease in fracture

strength after laser processing depends on both brittle recast layers and the release of residual compression

surface stress. The residual compression surface

stress induced by diamond grinding improves the fracture strength. Therefore, the removal of residual

compression stress after laser milling causes the decrease in fracture strength. We think that this crack-free processing method has great industrial

possibilities in combination with the traditional machining process.

References

( 1 ) Chryssolouris, G., Bredt, J. and Kordas, S., The

Winter Annual Meeting of ASME.PED- Vol. 17

(1985),p.9.

( 2 ) Chryssolouris, G. and Bredt, J., Proceeding of International Symposium on Machining of Advanced Ceramic Material (1987), p. 1.

( 3 ) Chryssolouris,G. and Bredt, J., Interceram, Vol.37. No. 2(1988), p. 43.

( 4 ) Wallace, R. J., Bass, M'. and Copley, S. M., J. Appl.

Phys., Vol. 59, No. 10(1986), p. 3555.

( 5 ) Wallace,R.J. and Copley,S.M.,Advanced Ceramic

Materials, Vol. 1, No. 3(1986), p. 277.

( 6 ) Paek, U. C. and Gabliano, F. P., IEEE Journal of Quantum Electronics, Vol.QE-8, No.2(1972)p.ll2.

( 7 ) Sakamoto, H., Morita, N., Watanabe, T. and

Yoshida Y., Proceedings of the First International Conference on New Manufacturing Technology

Chiba(1990),p. 571.

( 8 ) Morita, N., Watanabe, T. and Yoshida Y., Trans.

Jpn. Soc. Mech. Eng., (in Japanese) , Vol. 56, No.522-C(1990), p. 498.

( 9 ) Morita, N., Ishida, S., Fujimori, Y. and Ishikawa,

K., Pulsed Laser Processing of Ceramics in Water,

Applied Physics Letter, Vol. 52, No. 23(1988),

p. 1965.

(10) Walls. D. J., Evans, A. G., Marshall, D. B. and

James, M. R., J. Am. Ceram. Soc, Vol. 69, No.

1(1986), p. 44. (11) Morita, N., Sakamoto, H., Watanabe, T. and

Yoshida Y., The Winter Annual Meeting of ASME, MD-Vol. 16(1989), p. 9. ('2) Tanaka, K., Kurimura, T., Matsui, E. and Akiba,

Y., Materials (in Japanese) , Vol. 36 (1987) , p. 817.

((3) Lange, F. F.. James, M. R. and Green, D. J., J. Am.

Ceram. Soc, Vol.66, No. 6 (1983).

(14 ) Morita, N., Watanabe, T. and Yoshida, Y., CrackFree Processing of Hot- Pressed Silicon Nitride Ceramics using Pulsed YAG Laser, JSME Int. J., Ser. Ill, Vol.34, No. 1 (1991), p. 149.

(15) Morita, N., Watanabe, T. and Yoshida, Y., CrackFree Processing of Hot- Pressed Silicon Nitride Ceramics using Pulsed YAG Laser, JSME Int. J., Ser. Ill, Vol.3 , No. 1 (199 ), p.

524
ELECTRICAL DISCHARGE MACHINING OF

ADVANCED CERAMICS

N. M. FAULK

Tennessee Technological University, Cookeville, TN

Finishing advanced ceramic materials into intricate shapes with the required tolerance

is often extremely difficult using traditional machining methods. This has led to the

development of more sophisticated machining methods which do not depend on mechanical

contact. Electrical discharge machining (EDM) is such a techniques and it successfully overcomes many of the difficulties faced by other techniques. Since there is no direct contact,

this method is ideal for hard, brittle, and refractory materials. There are a few material

requirements, such as a minimum electrical conductivity, but for machinable materials very

good tolerances of below 1 /xm have been achieved.

Research has successfully demonstrated that several ceramics such as TiB2, B4C, and

several composites containing nitrides, carbides, and borides can be shaped to an ultra-smooth

finish. Cut and uncut surfaces were examined, indicating the cutting mechanism, the surface

quality, and contaminants.

Introduction to EDM

Electrical Discharge Machining (EDM) is derived from conventional diesinking, first developed as a practical metalworking method by

Lazarenko and Lazarenko (1) after World War II. Much theoretical and experimental work has been

conducted since that time in order to identify the

basic processes involved. Recent research has

started to extend this metal-finishing technique to processing ceramics and composites (2-3). These

studies have shown that EDM can be successfully

applied to ceramics, including monoliths, cermets,

and ceramic-matrix composites (CMCs) if the

electrical resistivity is below 300 Q • cm (4). Electrical discharge machining has been

developed into an outstanding technique for

precision machining. Diesinking and wire EDM

machines allow intricate cutting and shaping of

appropriate solid materials to nearly any threedimensional size or shape. Tolerances below 1 /im

(40 /liin.) have been achieved (4,5); EDM can produce a mirror finish, even on ceramics, giving

a surface roughness of less than 0.3 (12 /^in.)

(6,7). This technique is especially important for

machining hard, brittle, and high-melting-point

materials, which cannot be machined by

conventional techniques such as grinding.

The rate of material removal is influenced

mainly by the electrical and thermal properties of

the workpiece (the metal or other material), but not

at all by the hardness of the material. This is an advantage for ceramics, which are often limited by

their hardness and/or low toughness. In EDM, the

workpiece and the shaping tool are the electrodes

separated by a liquid dielectric. Because there is no

direct physical contact between the electrodes, no

mechanical stress is placed on the workpiece.

Depending on the application, there are two

types of machines used for EDM operation. The

diesinking machine, also known as the ram-type,

plunge, or vertical erosion machine (Fig. 1), can be

used for tapping, cutting holes, and helical

machining, and is capable of cutting or sinking very

complicated shapes into a workpiece. The workpiece is the cathode and the shaping tool (die) is the

525
anode in this type of machine. It is desirable to remove material from the workpiece, but not from

the tool, in order to use the die repeatedly for the same application. Heavy hydrocarbons and

kerosene are most commonly used as the dielectric

fluid in this conventional form of EDM.

After EDM

Figure 1 . Schematic diagram of the

die-sinking machine.

Wire-cutting machines (Fig. 2) evolved

from research on diesinking machines (5). Instead

of using a shaped electrode as the tool, wire-cutting

machines use a thin metallic wire under tension to cut like a jigsaw. Some erosion of the wire is allowable because the wire is continuously wound

between spools as it travels through the material,

but excessive wire erosion can cause wire breakage.

This breakage is the dominant practical and

economic problem of wire-cut EDM. In this type

of machining, water is the usual dielectric, and

brass-, steel-, or molybdenum-based wire is used

for the electrode. The polarity is switched for wirecutting machines: The workpiece is the anode, and

the wire is the cathode. Since the early 1980s,

most of the research and development work in the

EDM area has focused on wire-cutting (5). Refinements include increased cutting speeds,

extended taper angles, and independent

programming of the upper and lower wire guides so

that operators can cut shapes that have different

geometries at the top than at the bottom. Also in very limited use as an attachment to the wire

machine is a rotary spinning or indexing fixture. This process, also referred to as rotary EDM,

allows for better flushing and for cutting circular

parts that would otherwise be impossible to cut.

1. Wire (cathode)

2. Workpiece (anode)

3. Wire Guide Table

4. Workpiece Table

5. Machine Body

6. Wire Spool

7. Used-Wire Spool

8. Wire Tension Adjusters

Figure 2. Schematic diagram of the

wire-cutting machine.

526
The Electrical Discharge Process

Electrical discharge machining is based on

electric discharge erosion. Figure 3 illustrates how

this discharge process works, with discussion

following the work of Lee and Feick (8), Foster

(5,9), and Eubank and coworkers (10,11). The

three necessary components, that is, the two

electrodes and the dielectric fluid flowing in the gap

between them, can be seen in Fig. 3. Note that the

dielectric field contains ultrafme solid impurities.

An electric field generated by applying

approximately 200 volts DC across the gap, with an average width of 40 fim (1.6 mils), generates a magnetic field between the workpiece and the

electrode. At first, the two pieces are insulated by

the dielectric, so no current flows across the gap.

The resulting electric field, however, causes the

micron-size particles to be suspended and to form a bridge. This causes the breakdown, or

deionization, of the dielectric. The voltage (V) falls to about 25 V, and the current (I) rises to a constant value set by the operator. The plasma channel, seen forming in panel 3 of Fig. 3, grows during the

pulse time (t), or 'on-time". The intense heat

produced causes a vapor bubble to form around this channel. The surrounding liquid dielectric restricts plasma growth, thus concentrating the input energy,

V«I«t, into a very small volume. Energy densities

up to 3 J/mm^ (50 Btu/in.^) result, which cause the

plasma temperature to reach nearly 40,000 K

(72,000°F), and the plasma pressure to rise as high

as 300 MPa (2700 atm) due to inertial and viscosity

effects (10).

During this on-time, the high-energy plasma

affects both electrodes in different ways, depending on the properties of the materials involved. For

most applications, the plasma melts both electrodes

(9), as explained below. As the current flow halts,

the bubble implodes, drawing the molten material

into the gap. The dielectric fluid flushes and

solidifies this molten material and then carries it away, while also cooling the electrode surfaces.

The cycle is then repeated.

The pulse duration is one operating

condition that greatly influences the material

removal rate (MRR). The optimum pulse duration

depends on which type of EDM machine is being

used. For diesinking machines, optimum times of

10 to 100 us are common, whereas for wire-cutting

machines, times are less than 2 /is (3). The

difference is explained by the following discussion.

Because of bombardment of fast-moving electrons

at the start of the pulse, the anode melts rapidly

first, but then begins to resolidify after a few

microseconds. Recall that for the diesinking

machine the anode is the shaping tool. This

resolidification is believed to be due to the rapid

expansion of the plasma radius at the anode, which

causes a decrease in the local heat flux at the anode

surface. This expansion is referred to in the

literature as channel widening (12) or expanding

circle heat source (11). Melting of the cathode

occurs one to two orders of magnitude later because

of the lower mobility of the positive ions. Thus,

for the diesinking machines, time must be allowed

for the resolidification of the die, but for the wirecutting machine, much lower times may be used to allow high erosion rates of the anodic workpiece.

Mechanisms Involved

Two mechanisms for ceramic erosion have

been reported during EDM: melting and

evaporation, and thermal spalling. Each mechanism

is discussed below, along with examples of the

materials that each erodes.

Melting and evaporation of the electrodes is the most common method for erosion and the one

traditionally explained and modelled (9-14). For

metals and most ceramics, the plasma melts both

electrodes. As the current flow ceases, a violent

collapse of the plasma channel and vapor bubble

causes superheated, molten liquid on the surfaces of

both electrodes to explode into the gap, where the

liquid dielectric resolidifies the molten material.

Analysis of the debris was carried out on a WC-Co

system by Gadalla and Tsai (15-16), and they

concluded that, in some cases, the plasma

temperature causes evaporation of material from the

electrode. Their studies clearly show that material

is removed in both the liquid and gaseous state; perfectly spherical particles are apparently

resolidified particles from the gaseous state, whereas irregularly shaped particles are solidified

from the liquid state (17).

Not all the molten material can be removed

after the spark due to surface tension, tensile

strength, and bonding forces between the liquid and

solid. Crosby reports that for carbides cut with a combination of diesinking and wire-cutting EDM,

the recast layer has a depth of about 5 /^m (0.2 mil)

(18).

527
•

••0

1. Build-up of an electric field. 2. Formation of a bridge by conductive particles. 3. Beginning of discharge due to emission of negative particles. 4. Flow of current by means of negative and positive particles. 5. Development of discharge channel due to a rise in temperature and pressure. Formation of vapor bubble.

6. Reduced heat input after drop in current. Explosion-like removal of material.

7. Collapse of vapor bubble.

8. Residues: material particles, carbon, and gas.

Figure 3. A sequential set of eight panels showing how EDM works

528
Pure ceramics eroded by melting include

B4C, SiC, and graphite; CMCs such as AUOj-TiC,

Sialon-TiN, Si3N4-TiN and SiC-TiB2, and cermets,

including WC-Co and Si-SiC (15,19). Table 1 contains information on most of these materials.

Many of the CMCs listed in this table have a conductive phase added to decrease the resistivity

(2). Consider AUOj-TiC; its polished surface can be seen in Fig. 4, where the light phase is the TiC.

The EDM surface with resolidified melt-formation

droplets is shown in Fig. 5. Surface roughness

measurements show the roughness to be < 30 /xm

(< 120 mils) (19). This reflects, in part, the large

grain size of the material.

Table 1 . Materials Used in This Study

Material Composition* Resistivity

fi-cm

1. TiBo N/A - 1.8e-7

2. B,C N/A 0.5

3. Al^Oj-TiC 70/30 3.5e-3

4. AU03-TiC 70/30 l.Oe-2

5. Si-SiC 5/95 1.43e-4+

6. Si-SiC 2/98 1.47e-4^

7. Sialon-TiN 80/20 7.0e-4

8. Si3N4-TiN 75/25 1.5e-3

9. SiC-TiBj 80/20 1.0

10. BN-TiB^ 65/35 ; 1.2e-3^

11. BN-TiBj 75/25 1.4e-3^

* First value corresponds to first phase listed.

^ Resistivity calculated assuming linear change with

composition.

Thermal spalling is usually defined as the

mechanical failure of a material, without melting,

due to the internal stresses created by heating

and/or cooling, which overcome the bond strength.

This occurs as a material expands and contracts during the sudden temperature changes associated

with EDM, resulting in tension and compression

sufficient to cause tensile or compressive failure,

respectively.

To calculate the spalling temperature,

equations that show the thermal stress, created due

to a temperature gradient can be used (20). During

heating, the specimen is exposed to high

temperatures in the plasma; fluxes on the order of

10'° to 10" W/nf (3.2 X 10' to 3.2 x 10^° Btu/h/ft^) are reported for EDM (21). Due to heat transfer

from the plasma to the workpiece, a compressive

force perpendicular to the surface of the workpiece

is created. Moreover, because the surface attempts

to expand, but cannot, due to the constraint from

the bulk of the specimen, compressive stress is created parallel to the surface. When the stress exceeds the ultimate compressive strength of the

material, failure by thermal stress is assumed.

During cooldown (flushing), the reverse occurs; quench rates of 10^ to 10' K/s are observed (21).

The compressive forces created by the plasma

disappear and tensile forces are created in the

specimen as the surface attempts to contract. Failure occurs if the stress exceeds the tensile

strength.

Materials with high thermal expansions and

low thermal conductivity are subjected to higher

stresses due to more severe gradients. If a material

has a low strength or Young's modulus, as well, the

tendency to spall is enhanced. When the sparks are

localized in the same spot, the material may also

break by fatigue after a limited number of cycles.

An additional property that plays a role in whether

or not a material spalls is the melting point. If the

melting point is low enough, the material simply

melts or evaporates, as discussed above, and no

stresses build up to cause fracture. Petrofes (19)

reports that materials with melting points below

2800°C (5100°F) are cut by the melting

mechanism.

An example of a ceramic that erodes by

spalling is pure TiB^. Figure 6 shows the polished,

uncut surface of this material, whereas the cut surface is seen in Fig. 7. Note the absence of melt

droplets or a recast layer as was seen in Fig. 5. Instead, fracture of crystals and flaking are observed in this high-melting-point ceramic. The

liquid droplets seen on the surface are not from the

surface melting but rather from the brass wire (2). Energy-dispersive x-ray spectrometry (EDS) proves

that these droplets contain copper and zinc only.

Other than these droplets, there is no recast layer on the machined surface. The small holes present

in the grains themselves are due to dislocations

initially present in the crystals, which the thermal

process of EDM revealed by thermal etching.

529

Figure 6. Polished surface of TiB. (Material 1, Table 1).

Figure 7. EDMed surface of TiBo.
Application of EDM to Advanced Ceramics

Studies have been performed to answer

several questions that are important from an

industrial point of view, such as the degree of

contamination of the surface; the effects of EDM on

the hardness, surface stresses, and strength; the

material removal rate attainable; and the effect of

the grain size of the material. These are considered

separately below and summarized in Tables 2

and 3.

Table 2. Summary of EDM on Materials Studied

Material

Melting Removal

Temperature Cutting Rate

Material (°C) Mechanism (mm^/min)

1 2920(m) Spall 8.28

2 2450(m) Melt 0.60

3 201 8(m) Melt 2.42

4 201 8(m) Melt 2.52

5 1415(m) Melt 3.15

6 1415(m) Melt 2.00

7 1878(d) Melt 4.40

8 1878(d) Melt 2.33

9 2700(s) Melt 3.52

10 3000(m) Spall 12.2

11 3000(m) Spall 10.4

m = melting, d = decomposition, s = sublimation

I. Gadalla and Petrofes studied the surfaces

created by applying EDM to several samples of

advanced ceramics (22). Elemental mapping was

performed on many of the samples in order to determine the extent of contamination on the surface

by the material or impurities in the dielectric, which

in the cited case was tap water. The brass wire that was used (63/37% copper/zinc) caused some

contamination. Copper alloy resolidified droplets

with a lower concentration of zinc and a fine

dispersion of zinc exist over the entire surface. The

zinc dispersion is expected due to its higher vapor

pressure. Viewed by EDS, the surface subjected to EDM showed traces of sodium, potassium, and

chlorine from the tap water added to the deionized

water (the dielectric fluid used) to adjust its conductivity.

II. The Vicker's microhardness of surfaces

prepared by EDM was compared to that of diamond sawn and polished surfaces (19); see Table 3. It was concluded that for those materials that were

machined by the melt mechanism, EDM produced

surfaces with lower hardnesses than those obtained

with a diamond saw. This conclusion implies that the internal stresses created by EDM on quenching

the liquid phase are less than those created by

mechanical erosion. On the other hand, the

hardness obtained for spalled materials by EDM

was greater than that obtained by a diamond saw. This indicates that the hardness obtained by the

spalling mechanism is comparable to mechanical

erosion. In some cases, regardless of the erosion

mechanism, polished surfaces gave the highest

hardness due to the high mechanical stresses created, capture of debris by a softer matrix

material, and/or possible redistribution of phases.

III. Surface residual stresses were

determined by Gadalla and Tsai for surfaces of

WC-Co composites subjected to EDM and were

found to be anisotropic (15). In another study (10),

it was found that flaking due to spalling causes

residual compressive stresses, but the contracting

due to solidification of melted surfaces causes

residual tensile stresses.

IV. The decrease in strength of a ceramic

machined by EDM is its biggest drawback. Saito

(23) reports that even when no cracks are observed on the machined surface, the value of the bending

strength is between 40 to 50% of its initial value for

several ceramics and CMCs. Nakamura does

report, however, that when no cracks are observed on the surface, the bending strength can be

recovered by removing entirely the roughness with

a fine diamond grinding wheel of 400 grit or more

(24) .

532
Table 3. Vicker's Microhardness Testing'

for Machined Surfaces

Material Diamond EDMed Polished

Saw*

HV HV HV

1 1366 ± 200 3253 ±210 2971 ±240

2 2569 ± 680 1995 ± 600 2800 ±100

3 1734 ± 300 1306 ± 330 2163 ± 110

4 2193 ± 270 1521 ±280 2349 ±220

5 2111 ± 510 1633 ± 680 N/A*

6 2914 ± 120 2341 ±210 N/A*

7 1099 ± 70 1197 ± 310 1721 ±240

8 1737 ±210 1353 ± 350 2373 ± 420

9 2700 ± 490 1672 ± 460 2598 ± 450

10 265 ± 70 561 ± 200 630 ± 200

11 300 ± 70 343 ± 70 375 ± 180

^ All tests performed at 1kg load 30 sec, except

materials 10 and 11. Softness of these materials

required 300 g load 15 sec.

* Low-speed diamond saw: 300 rpm, 300g load,

low diamond concentration blade.

* Material too porous to get a reading on the

indentation.

V. For a wide variety of advanced

ceramics, including pure ceramics, cermets, and

CMCs, Petrofes reports cutting rates of 10 to 50

mm'/min (0.02 to 0.08 in.'/min), but does not claim that these numbers are optimum (19). Saito

reports rates of 100 to 200 mm-/min (0.2 to 0.4

in.^/min) for conductive Sialon (contains TiN) and

Si3N4 using optimum cutting conditions (23).

VI. Gadalla and Petrofes report that larger

grain sizes of the nonconductive phase (for

composites) cause higher roughness and a higher

removal rate because the loss of a single grain

leaves a large crater and removes more material,

compared to a small grain (22). They also report

the case of an AljOj—TiC composite of grain sizes

20 and 30 /xm (0.8 and 1.2 mils), respectively, in which the x-ray diffraction pattern of the machined

surface showed much broader peaks than did its polished surface. This indicated that the melted

material produced by EDM precipitated finer

crystals with high residual stresses on quenching.

Expanding Use of EDM for Ceramics

Advanced materials that are difficult or impossible to machine by conventional abrasive

methods can be machined into complex shapes

using EDM, it they have low enough electrical resistivity (< 300 n»cm). Because the tool and the

workpiece are separated by a dielectric field,

extremely hard and brittle materials can be cut. The EDM process takes place by melting,

evaporating, or fracturing the surface, depending on

the melting point of the material.

Several materials have been successfully

cut, and as awareness of the requirements for EDM

grows, more ceramics manufacturers are adding

conductive phases to their products to make them

machinable. Use of this technique is expected to grow as operators begin to see ways to incorporate

EDM into production.

References

1. B. R. Lazarenko and N. I. Lazarenko,

"Machining by Erosion," American Mach. . Vol. 91, No. 26, p. 120-121, 1947.

2. N. F. Petrofes and A. M. Gadalla,

"Electrical Discharge Machining of

Advanced Ceramics," American Ceramics

Soc. Bull. . Vol. 67, No. 6, p. 1048-1052,

1988.

3. N. F. Petrofes and A. M. Gadalla,

"Processing Aspects of Shaping Advanced

Materials by Electrical Discharge

Machining," Adv. Mater. Manuf. Proc . Vol. 3, No. 1, p. 127-153, 1988.

4. R. F. Firestone, "Abrasionless Machining

Methods for Ceramics: The Science of

Machining and Surface Finishing II," National Br. Stand. Spec. Publ. . No. 562,

B. J. Hockey and R. W. Rice, Ed., p. 261-

281, 1979.

533
5. R. F. Foster, "Evolution of Wire EDM,"

Tool. Prod. . Vol. 47, No. 12, p. 90-93,

1982.

6. J. C. Quinlan, "What's New in EDM,"

Tool. Prod. . Vol. 52, No. 12, p. 40-49,

1987.

7. L. Houman, "EDM Surfaces: The

Persistent Problem of Measurement,"

Carbide Tool J. . Vol. 14, No. 4, p. 24-26,

1982.

8. D. W. Lee and G. Feick, "The Techniques

and Mechanisms of Chemical, Electrochemical, and Electrical Discharge

Machining of Ceramic Materials;" The

Science of Machining and Surface

Finishing . National Bur. Stand. Spec.

Publ., No. 348, p. 197-211, 1972.

9. R. F. Foster, "Advanced Wire Cutting

EDM," Carbide Tool J. . Vol. 14, No. 4, p. 22-23, 1982.

10. D. DiBitonto, M. R. Patel, and P. T. Eubank, "Theoretical Models of the

Electrical Discharge Machining Process,

Part I: A Simple Cathode Erosion Model,"

J. Appl. Phvs. . Vol. 66, No. 9, p. 4095-

4103, 1989.

11. M. R. Patel, M. A. Barrufet, P. T. Eubank, and D. D. DiBitonto, "Theoretical

Models of the Electrical Discharge

Machining Process. Part II: The Anode

Erosion Model," J. Appl. Phys. . Vol. 66,

No. 9, p. 4104-4111, 1989.

12. C. van Osenbruggen, "High-Precision

Spark Machining," Philips Tech. Rev. . Vol. 30, No. 6, 7, p. 195-208, 1969.

13. S. M. Pandit and K. P. Rajurkar, "A New

Approach to Thermal Modelling Applied to Electro Discharge Machining," No. 81-

WA/HT-3, American Society of Mechanical

Engineers, 1981.

14. A. Erden, "Effect of Materials on the

Mechanism of Electric Discharge

Machining (E.D.M.)," Trans. ASME. Vol.

105, No. 2, p. 132-138, 1983.

15. A. M. Gadalla and W. L. Tsai, "Machining

of WC-Co Composites," Adv. Mater.

Manuf. Proc . Vol. 4, No. 3, p. 411-423,

1989.

20. R. W. Davidge, Mechanical Behavior of

Ceramics . Cambridge University Press, p. 118-122, 1979.

21. A. E. Berkowitz and J. L. Walter, "Spark

Erosion," J. Mater. Res. . Vol. 2, No. 2,

p. 277-288, 1987.

22. A. M. Gadalla and N. Petrofes, "Surface of

Advanced Ceramic Composites Formed by

Electrical Discharge Machining," Mater.

Manuf. Proc . Vol. 5, No. 2, p. 253-271,

1990.

23. N. Saito, "A Review of Special Machining

of Fine Ceramics," Sumitomo Elect. Tech.

Rev. . Vol. 25, p. 1-19, 1986.

24. M. Nakamura, K. Kubo, and S. Hirail,

Electr. Mach. Tech. . Vol. 8, No. 2, p. 1- 12, 1984.

534
EFFECT OF GRAIN SIZE ON WIRE ELECTRICAL

DISCHARGE MACHINING OF

POLYCRYSTALLINE DIAMOND

S. Z. WANG, K. P. RAJURKAR, and J. KOZAK

University of Nebraska-Lincoln, Lincoln, NE

Application of ultra-hard cutting tool materials such as polycrystalline diamond (PCD) is continuously expanding. However, shaping of PCD blanks is a long, labor-intensive and

costiy process. Wire electrical discharge machining (WEDM) has the potential to be an

effective and economical technique for the production of tools from PCD blanks. This

paper reports the effect of grain size of PCD, discharge frequency and discharge energy

of WEDM on machining feed rate and quality of machined surfaces of PCD blanks.

Surface roughness and notch depth (which is often observed at the diamond WC interface) were measured to study the cutting edge quality of the PCD blanks.

INTRODUCTION

Polycrystalline diamond (PCD) tool blanks have

been developed to machine special alloys,

composites, and abrasive nonmetallic materials.

Due to the unique properties of PCD tool blanks,

they are widely used in the automotive industry

(for machining pistons, brake cylinders,

transmission cases, and engine blocks), in the

aerospace industry (for machining engine

components), the jewelry industry (for flycutting

and engraving), in the woodworking industry (for

machining particle board, fiberboard, and

composite materials), and in the textile industry

(for machining water jet cutting nozzles) (1).

PCD is essentially a synthesized, extremely

tough, and intergrown mass of randomly

orientated diamond crystals. It is produced by

sintering together the carefully selected diamond

particles at a very high temperature (approx. 1500

°C) and ultra-high pressures (about 6 GPa). PCD

materials are extremely tough and strong. They

can be regarded as a composite material which

combines the hardness, abrasion resistance,

compressive strength, and high thermal

conductivity of diamond with the toughness of

tungsten carbide. They are available in several

grades, based on the average diamond grain size

in the PCD layer (2).

Shaping of PCD tool blanks is challenging due to inherent properties of high strength and high

toughness. Diamond grinding is one of the most

commonly used techniques (3). The structure of

the PCD layer is very homogeneous, and the

pertinent high wear-resistance causes problems even in grinding. In order to remove the material

from the PCD blank, the diamond layer of the

grinding wheel must be renewed by a trueing or

dressing operation frequentiy, resulting in rapid wear of the wheel. During grinding, depending on the PCD grade and contact force machining

rates are 0.5 to 2.0 mmVmin and the G ratio

(ratio of workpiece removal volume and grinding

wheel wear volume) is 0.005 to 0.02 (4). Thus,

the grinding wheel wear rate is 50 to 200 times

higher than the workpiece removal rate. Hence,

grinding is suitable only to a limited extent for

the production of PCD profile tools. Although

diamond in the PCD layer is not electrically

535
conductive, PCD sintered material manufactured

with cobalt as a solvent catalyst has adequate

electrical conductivity suitable for electrical

discharge machining (EDM).

The removal of material in EDM is achieved with

the erosive effect resulting from electric sparks

between two electrodes. Sparks of short duration

(0.1 - 2000 yis) are generated in a liquid dielectric

gap separating tool and workpiece. In wire

electrical discharge machining (WEDM) a wire of

diameter ranging from 0.05 to 0.25 mm is used

as a tool electrode and the dielectric (deionized

water) is injected into the machining gap.

WEDM produces complex 2- and 3 -dimensional

shapes in hard materials with a high degree of

accuracy (0.005 - 0.05 mm), without the use of a complicated and expensive EDM profile

electrode. Because of the use of relatively

inexpensive tool (wire) and non-contact nature of

the process, WEDM offers a lower cost

alternative with inherent higher accuracy.

Although EDM are becoming an attractive

alternative for PCD cutting, the relevant

published information is extremely Umited. The

machining feed rate and the surface roughness are influenced by wire speed. The machining rate has been found to be maximum for a wire speed

of 4 - 5 m/sec (5). The surface roughness also

increases as the wire speed increases. One study

(6) indicates that higher removal rates can be

achieved with a high generator voltage and pulse

frequency. The optimum cutting edge waviness

is apparently achieved with a discharge current of

4 - 6 A and a low generator voltage of 120 V.

Recently, authors have reported the effect of

pulse frequency, charge current, and pulse on- time on erosion rate and the machining

performance of cutting blanks with different layer

thicknesses of PCD and tungsten carbide (7).

The influence of frequency and energy of electric

discharge on the material removal rate (i.e., cutting speed) of PCD blanks, and the mechanism

of removing diamond grains from the matrix

based on the effect of thermal stresses on the

diamond grains and cobalt phase during electrical

erosion have also been studied (8). The

machining feed rate has been found to increase

with discharge frequency and discharge energy.

The wire electrical discharge machining rate of

the PCD layer is abnost 3 - 5 times lower than

that of the tungsten carbide substrate. The

difference in values of thermal expansion of

diamond and cobalt plays an important role in the

diamond grain removal during WEDM of PCD.

This paper reports the result of an investigation

of WEDM of PCD blanks with different grades.

The goal of this research is to study the influence

of different diamond grain sizes on the material

removal rate (i.e., cutting speed). The surface

roughness and "notch" depth have also been

measured to study the cutting edge quality of the

PCD layer.

EXPERIMENT

PCD blanks of three different grades were machined by a 5-axis CNC wire electrical

discharge machine. The three grades correspond

to 2 jim, 10 pm, and 25 jam diamond grain sizes

in the PCD layer. The blank consists of a 0.7

mm thick PCD layer and a 2.48 mm thick

tungsten carbide substrate. '

'

'

The discharge energy with a transistor-controlled

RC relaxation circuit is determined by the value

of the capacitor which is in parallel with the

machining gap and the value of discharge

voltage. The no-load operating voltage is adjustable from U = 80 to 200 V with variable

current of I = 0.5 to 128 A and a pulse frequency

f = 1 to 250 kHz. The manually set value of

pulse frequency controls the pulse off-time by

means of transistors and relative trigger circuits.

Because of the RC circuit, the discharge period is very short (usually less than 4 ps), and therefore,

stable arcing does not occur during machining

with this system.

A stratified (zinc coated brass) wire of 0.25 mm

diameter was used as the negatively polarized

tool electrode. The wire speed and the wire

tension were kept at 6 m/min and 1.5 N

respectively. The operating voltage (U) was set

at 120 V, charge current (I) was set at 48 A, and
pulse on-time (tp) was set at 3.2 ps. The capacity

(C) and pulse frequency (f) were varied between

1 - 3.3 jiF and 8-20 kHz respectively.

An on-line data acquisition system and a trigger

circuit were used to record the machining feed

rate and the actual number of sparks. The

collected on-line data was used to obtain the

average machining feed rate and the average

discharge frequency. The discharge frequency

which indicates the actual number of sparks on

the workpiece surface is a more realistic

parameter than the pulse frequency which is the

setting parameter of the machine (9). The

material removal rate (MRR) was determined by

weighing the workpiece before and after

machining during each experiment and

calculating the result based on the period of

machining time. The machined surfaces were

studied by scanning electron microscopy (SEM)

and surface roughness profilometer. A

toolmakers microscope was used to measure the

"notch" depth between the PCD layer and

tungsten carbide substrate.

For all three grades of PCD, higher capacitance

and higher discharge frequency result in higher

values of machining feed rate. The linear trend

of increase of machining feed rate (Vf) with

discharge frequency can be explained on the

basis of the following relations developed in

(7,8).

The average feed rate of machining (V^) (shown

in Figure 2) in mm/min and/or the cutting speed

(VJ in mmVmin defined as V^, = Vf • H are a function of discharge frequency (fj and

capacitance (C),

and,

where

A = k

(1)

(2)

30 f//

C

EXPERIMENTAL RESULTS AND ANALYSIS

Machining Rate

The experimental results of the machining feed

rate (Vf) as a function of discharge frequency (f^) and capacitance (C) are shown in Figure 1.

3.5- C = 3.3;iF. C = l>iF

1000 2000 3000 4000 5000 6000 7000

Discharge Frequency, Hz

Figure 1 Feed Rate of Machining (Vf) vs. Discharge Frequency (fj)

k = constant of proportionality

Ud = peak voltage, V

H = height of workpiece, mm

d^ = diameter of wire, mm

Sb = gap between wire and workpiece,

mm

Thus, the discharge frequency (f^) and

capacitance (C) seem to be the parameters

influencing the erosion rate and, therefore, the

Wire

dw

"JU^

Worlqaece

1

Vf t

Figure 2 WEDM Operation [8]

537
quality of the machined surface.

The higher machining feed rate of 2 pm grain

size as compared to 10 pm and 25 pm grain sizes

may be due to a finer and more uniform

distribution of electrically conductive cobalt

binder (6), and lower bond strength between

particles and the binder. The large grain sizes

seem to result in a better diamond intergrowth

under high pressure and high temperature

synthesis process. The difference in coefficients

of thermal expansion of diamond grain and cobalt

may also be contributing to easier dislodging of

smaller grains than larger grains (8). The

increase of machining feed rate (for all grades)

with discharge frequency is due to the lower offtime (i.e., non-productive or non-machining time)

corresponding to higher discharge frequency.

The spark generated crater size depends on the

discharge energy including pulse on-time. The

off-time provides a pause between two successive

discharges for removing the debris and flushing

the dielectric fluid.

The experimental results also indicate that there

is no significant difference between machining

feed rates of 10 pm and 25 pm grain sizes. The

regression equations corresponding to lines shown

in Figure 1 for the different grades of PCD

blanks are,

Grain Sizes:

2 pm

10 pm

25 pm

V, = 4.33-10-'C°-^Y

= 3.29-10-'C°^y^

(3)

V, = 3.41-10-'C'''*Y.

The calculated values of the volume of material

removal (V^,) for different grades PCD blanks are shown in Figures 3 and 4.

10 15 20

Grain Size, micrometer

25 30

Figure 3 Volume Removal (V„) vs. Grain

Sizes (2, 10, and 25 jim)

60

•= 13'

-3 12

11

I

lo^

« 9

73>o

E

4)

S

7-

"""^ ^

C = 3.3>iF

Grain Sizes: 2 ;im

* 10;im

^^-p?^

C = 3.3;iF ° 25}\m.

C = l;iF

1000 2000 3000 4000 5000

Discharge Frequency, Hz

6000 7000

Figure 4 Volume Removal (V^) vs. Discharge

Frequency (f^)

where C is in pF, f^ is in Hz, and Vf is in

mm/min. For parameter values: = 110 V, C

= 1-3.3 pF, H = 3.18 mm, Sb = 0.25 mm, and

d^ = 0.25 mm.

The volume of material removal (Vn^) per

discharge can be estimated using the following

relation (7),

V (4)

The material removed or eroded per discharge is a function of the discharge energy which in turn depends on the value of the capacitance. Hence,

a higher capacitance value results in higher

volume of PCD removed per discharge. The

grain size as mentioned earlier plays an important

role in the erosion process. However, it seems

that the change in volume eroded per discharge

is insignificant beyond a critical value of grain

size.

538
The off-time represented by discharge frequency

plays a less dominant role in WEDM erosion

mechanism than in die-sinking erosion

mechanism. The stability and dielectric

deionization phenomenon and the avoidance of

arcing are greatly influenced by the duration of

the off-time in die-sinking EDM. In wire EDM,

the motion of wire, the flow path of dielectric

fluid and the additional effect of gravity in

removing debris result in more or less uniform

effect of off-time on the erosion process.

Therefore, it is expected that the off-time or discharge frequency will have an insignificant

effect on erosion per discharge. Results shown in

Figure 4 seem to confirm the above mentioned

physical phenomena. However, a slight decrease

in erosion per discharge with higher discharge

frequency or lower off-time similar to die- sinking

EDM can be seen for all results except the blanks

with 10 |im and 25 pm grain sizes at 3.3 pF. The

results for 25 pm grain size PCD blank show a consistent distribution around an average value of

9.5 X 10"^ mmVper discharge, where as a rapid

increase in the erosion rate per discharge with

discharge frequency is evident in case of 10 pm

grain size PCD blank. The wire broke at discharge frequency values higher than 2.5 kHz

at 3.3 pF with 10 pm and 25 pm grain sizes PCD

blanks. A further investigation of 10 pm grain

size PCD blank erosion under different machining

parameters needs to be carried out to understand

this behavior.

values over the entire PCD surface layer revealed no consistent trend with machining parameters.

It seems that machining parameters used in this

investigation did not affect the surface roughness

values. The PCD layer was not large enough for

recommended 0.762 mm cut-off (wavelength) for

which the traverse length is 5.08 mm and,

therefore, a cut-off of 0.254 mm (a traverse length of 1.524 mm) was used in measuring the

roughness values. This constraint on the cut-off

may have contributed to the lack of emergence of

any specific trend of surface roughness values.

Therefore, two separate sets of measurement were

taken. The first set corresponds to the zone 1 at the entry of the wire in the cutting region. The

second set corresponds to the zone 2 extending

from the approximately middle of the workpiece

cutting length to the exist of the wire (Figure 5).

The results of surface roughness (RJ are shown

in Figure 6.

Wire

Direction ^—Zone 1 — ^—Zone 2 —

Tungsten Cairbide Substrate

Qualitv of Machined Surface

The quality of the cutting edge and corresponding

microgeometrical features influence the overall

performance and economy of the cutting process.

A longer tool life, lower cutting forces and

vibrations, and better surface finish require a properly finished and micro-defect free PCD

cutting tool edges. Therefore, the effects of

discharge frequency (f^), discharge energy

(capacitance C), and grain size on surface

roughness and notch depth have been

investigated.

The measurement of average surface roughness

Figure 5 Surface Roughness Measurement

Zones

Figure 6 shows that the surface roughness of the

I?GD layer with 2 pm grain size is slightly less

than with 10 pm and 25 pm grain sizes. From

the calculations, the average roughness for 2 pm

grain size is 2.5 pm and 3.1 pm for 10 pm grain

size. There seems to be a definite difference

between surface roughness values of the two

zones. The surface roughness (RJ of the first zone of the PCD layer is approximately 50%

times more than that of the second zone. The

difference in surface characteristics may be

539
5.5

U 5i

(L>

g 4.5

3

O

c/3 2.5

Grain Sizes:

, a ji

Zone 1

2)im

* 10 pm

25 ;im

Zone 2

1000 2000 3000 4000 5000 6000 7000

Discharge Frequency, Hz

Figure 6 Surface Roughness (R,) vs. Discharge Frequency (f^) on Two Zones

explained by analyzing the change of discharge

frequency during machining. A typical discharge

frequency record is shown in Figure 7.

Initially when the wire enters the cutting region,

the process is unstable and following a transition zone becomes fairly stable. The initial instability, longer off-time providing sufficient

time for resolidification of molten but not ejected

workpiece material, and the movement of wire

(i.e., non-uniform distribution of sparks and

related energy) tend to generate a rougher surface

in the first zone. The stability of the process as well as shorter off-time (higher discharge

frequency) helps in continuous and uniform

600

80 100

Time, s Figure 7 Discharge Frequency (fj)

Machining Workpiece (10 lum grain size)

of

distribution of sparking which results in better

surface roughness in zone 2.

The machined surfaces were also studied by

scanning electron microscope. The SEM

photographs (Figure 8) show the significant

differences among the three grades of PCD

blanks. The machining conditions were set at I = 48 A, tp = 3.2 ps, C = 3.3 pF, and f = 12 kHz.

The discharge frequency values corresponding to 2 pm, 10 pm, and 25 pm grain sizes PCD blanks

are 2773 Hz, 2388 Hz, and 2370 Hz respectively.

The smoother surface generated by WEDM of 2

pm grain size workpiece shown in Figure 8(a) as compared to that of 10 pm and 25 pm grain size

workpieces shown in Figure 8(b,c) may be due to a finer and more uniform distribution of

electrically conductive cobalt binder and higher

discharge frequency or lower off-time as mentioned above. A comparison of Figures 8(a),

8(b), and 8(c) also shows that the larger voids

(cavities) are left on the surface of PCD layer

with large grain size. These grains are either

removed during the erosion process due to the

difference in coefficient of thermal expansion of

diamond and cobalt and/or subsequent shock

wave impact resulting from the collapse of the

plasma channel.

A "notch" formation at the transition zone between the PCD layer and the tungsten carbide

substrate during WEDM process can be seen in

Figure 9, similar to observations have been

reported in (8). It is believed that processes

taking place during the production of PCD

composites lead to a higher cobalt concentration

in this transition or interface zone. The cobalt

being electrically conductive is eroded easily

from this high cobalt density region. The grain

size has a significant effect on the notch depth

and surface characteristics of PCD composites as shown in Figure 9(a-c). The notch depth is smaller for 2 pm grain size PCD composite than

that of either 10 pm or 25 pm grain size

composites. Figure 10 shows the variation of

notch depth measured by a toolmakers

microscope (accuracy 0.001 mm). The uniform

and relatively smaller density of cobalt in the

PCD composite transition zone seems to result in

540
1^-

k%

(a) 2 pm (b) 10 pm (c) 25 pm

Figure 8 SEM Photograph for the Three Grades of the PCD Layer (x750)

(a) 2 pm (b) 10 pm (c) 25 pm

Figure 9 SEM Photograph Showing the Difference in the Notch Depths (xl50)

lower notch depth with 2 pm grain size as compared to those of 10 pm and 25 pm grain

sizes. The higher energies (i.e., large capacitance

values) expectedly lead to a deeper notch. The

pulse frequency does not seem to have a

significant effect on notch depth, although the

trend at 12 kHz with a capacitance of 3.3 pF is not consistent with the rest of the results.

5A1

Figure 10 Notch Depth (K^) vs. Grain Sizes (2,

10, and 25 ^m)
SUMMARY

This paper reports the investigation of the effect

of discharge frequency and discharge energy (i.e., capacitance) on the machining feed rate, the

volume of material eroded per discharge, and

cutting edge quality during wire electrical

discharge machining of polycrystalline diamond

blanks of different grain sizes. The machining

feed rate and the volume of material removed per

discharge decrease as the diamond grain size

increases from 2 pm to 10 pm and have no

significant change between 10 pm and 25 pm

diamond grain size. The surface roughness and

the notch depth were measured to investigate the

quality of the PCD layer. The average roughness

of 2 pm grain size is slightly less than that of 10

pm and 25 pm grain sizes. No significant

difference between the surface roughness values

of the PCD layer with 10 pm and 25 pm grain

sizes was observed when the discharge frequency

and the discharge energy are changed. However,

it was observed that the surface at the entry of

the wire is rougher (by almost 50% more) than

the surface near the exit of the wire. The notch

depth shows that 2 pm grain size as better cutting

edge quality than 10 pm and 25 pm diamond

grain sizes.

ACKNOWLEDGEMENTS

The support of the National Science Foundation

(Grant #DDM-92 15298) and the State of

Nebraska (Nebraska Research Initiative Fund) is gratefully acknowledged. Authors are thankful to Diamond Abrasives Corporation, USA for

supplying the SYNDITE PCD materials used in

this study.

REFERENCES

1. "J & M Diamond Tool Makes Great

Strides in Machining Polycrystalline

Diamond," EDM Digest, Vol. Xm, N. 1, 1991, pp. 14-15.

Diamond and PCD," Diamonds in

Industry , DeBeers Industrial Diamond

Division, 1991, pp. 3-5.

3. Kenter, M., "Effect of Process Parameters

When Grinding PCD-3" Industrial

Diamond Review, Vol. 52, N. 553, June

1992, pp. 313-318.

4. Wyss, R., and Pollak, E., "Machining

Concept for PCD Tools," Swiss Quality

Production, 1990, pp. s70-s75.

5. Fengguo, C, Yansheng, M., Yuanfang,

W., and Jinjuan,' F., "A Machining

Technique of Superhard Materials PCD &

CBN Using Wire EDM," Proceeding of

International Symposium for ElectroMachining (ISEM- 10) , 1992, pp. 359-363.

6. Spur, G.,Puttrus,M., and Wunsch,U.W.,

"Wire EDM of PCD," Industrial Diamond

Reviews, Vol. 48, June 1988, pp. 264-266.

7. Wang, S. Z., Kozak, J., and Rajurkar, K.

P., "Study of Wire Electrical Discharge

Machining of Polycrystalline Diamond,"

To be published in Transaction of the

NAMRI/SME, May 1993.

8. Kozak, J., Rajurkar, K. P., and Wang, S. Z., "Material Removal in WEDM of PCD

Blanks," Submitted to Joumal of

Engineering for Industry, 1993.

9. Rajurkar, K. P., and Wang, W. M., "OnLine Monitor and Control for Wire

Breakage in WEDM," Annals of the

CIRP, Vol. 40/1, 1991, pp. 219-222.

2. Jennings, M., "Wire Drawing with Natural

542
LASER MACHINING OF SILICON NITRIDE CERAMICS

A. G. SOLOMAH

SAC International Ceramics Ltd., Mississauga, Ontario, Canada

Silicon nitride ceramic materials were machined using a high power CO2 laser beam. The

laser beam was operated under continuous or pulsed mode at different power levels.

Argon, oxygen and nitrogen were used as assist gases. The as-cut surfaces of silicon

nitride materials were found to possess different surface flaws, mainly porosity, which

affect the subsequent surface grinding processes. It was concluded that oxygen is the most

suitable assist gas during laser cutting under continuous or pulsed mode operation. Laser

cutting of silicon nitride can not be used alone to finish such ceramic materials, however

it can be used as a complementary process to conventional diamond grinding especially

in high speed cutting of complicated shapes and geometries.

INTRODUCTION

Silicon nitride and carbide are prime candidate

materials for ceramic components in advanced heat

engines due to their superior properties, such as

strength, toughness, wear resistance, light weight

and oxidation resistance at high operating

temperatures, e.g. 1300 C (1). However,

manufacturing components from such hard and

tough ceramic materials requires sophisticated

processing techniques to produce "near-net" shapes

after densificadon. The final product must be

machined to achieve the required dimensional

tolerance and acceptable surface fmish to increase

the reliability of such components under

anticipated operating conditions.

Diamond tooling and machining are the accepted

techniques to machine such hard ceramic materials

(2). However, diamond machining is very

expensive and it is a very slow process, especially

when complex shapes are required. Machining

ceramics with diamond tooling accounts for about

60-70% of the cost of the final product. Cutting

and drilling are considered the most time

consuming steps in diamond machining of

advanced ceramic materials, especially when

complex geometries are required.

Laser machining can be a very cost-effective

technique to reduce the cost of ceramic machining,

especially for line cutting, drilling and profile

cutting. Several researchers in Japan and Europe

have investigated the use of high power CO2 lasers

to cut silicon nitride, zirconia and alumina

ceramics (3-7). Success has been achieved in

machining these ceramic materials under controlled

conditions. The economics of such a new

machining process are not available yet. However,

we have performed a preliminary energy savings

analysis comparing diamond tooling and the laser

in cutting silicon nitride materials, which showed

a saving of 60% (see Appendix).

MATERIALS AND EXPERIMENTAL WORK

Hot-pressed silicon nitride ceramic materials

containing about 3 wt% MgO and other non-
specified amounts of different sintering additives were used to carry out the preliminary work to investigate the feasibiHty of cutting with a high

power (1.5 kW) carbon dioxide laser. The density

of the hot-pressed billets was > 99% of the

theoretical density. The hardness, H, and fracture

toughness, Kjc were determined using indentation

techniques and their values were 16 GPa and 4.5

MPa.m'^^, respectively. The flexural strength, a, was 560 Mpa, as measured using a three point

bending test.

The laser was a single mode CO2 laser which has

the capability of producing a high quality

continuous output beam at 10.6 |im over a power

range of 50 to 1500 W. A lens focussed the

output to a spot as small as 100 |im, corresponding

to a maximum power density of more than 10^

W/cm^. The laser could also be pulsed at a

repetition rate up to 5 kHz. Motion of the

workpiece under the focussed laser beam was

accomplished by means of an x-y positioning table

together with a programmable motion controller.

Speeds of up to 15 cm/s were achievable with

travel of 30 cm in each direction. In the present

study, the laser beam was operated in two different

modes, continuous wave and pulsed; the power

ranged from 600 to 900 W in each case, while the

frequency in pulsed operation varied from 150 to 900 Hz. The focal length was fixed at 12.7 cm and

the working distance between the nozzle and

workpiece was 150 |im. Oxygen and argon were

used as assist gases. The thickness of the silicon

nitride workpiece was 6 mm and the cutting speed

was 0.25 cm/s.

The as-cut surfaces were characterized using visual

inspection for colour changes and macrocracking,

while their microstructure and microcracking were examined using scanning electron microscopy

(SEM) provided with an energy dispersive x-ray

analyzer (EDAX).

Conventional diamond grinding was utilised to remove the as-cut surface in order to achieve a surface finish quality similar to that produced by

diamond machining. A 220 grit diamond wheel

was used to remove the damage layer followed by

320 grit to remove aa additional depth of 40 |im to reach the bulk material.

RESULTS AND DISCUSSION

1. Cutting using continuous wave laser

Surface quality changed drastically when the

power level changed from 1200 to 600 W. With

oxygen assist gas, smoother surfaces were

achieved when lower power levels were utilized,

especially near the bottom of the workpiece. High

power levels caused severe shattering and fracture,

which can be explained oh the basis of thermal

shock due to the lower heat conductivity of Si3N4

generating high levels of differential thermal

strains between the top and the bottom of the

workpiece. Lowering the power level reduced this

problem and produced a smoother surface. The

cutting mechanism responsible for this process is the oxidation of Si3N4 according to:

2Si3N4 + 7O2 —->6SiO + 8NO (1)

6SiO -1- 3O2 —-> 6Si02 (2)

The thickness of the oxidised layer, or silica scale,

was measured as 150 to 350 |im, using optical

microscopy.

Microstructural examinations using SEM revealed

that the porosity changed as the laser beam

penetrated throughout the thickness. Figures 1(a)

and 1(b) show that the porosity decreases in

density and in pore size with penetration depth.

The shiny spots shown in the bottom of Fig. 1(a) are metallic silicon inclusions, as revealed by

EDAX analysis. It is thought that at the top layer

of the workpiece oxidation takes place rapidly,

generating SiO and NO gases which are further

converted into higher oxygen content, thus

consuming more oxygen. This situation can reduce

the oxygen content of the gas mixture which

would then be less oxidising in the deeper layers,

thereby producing fewer gas bubbles and a lower

porosity. This is in agreement with the presence of

metallic silicon inclusions deeper in the workpiece

which do not exist in the top layers.

544
Fig. 1(a) SEM micrograph showing the morphology

of the silicon nitride edge as cut with 600 W

continuous laser power and oxygen assist gas.

Fig. 1(b) Porous layer in Fig. 1(a) shown under

higher magnification (490X vs 60X).

Under an argon atmosphere, metalHc shiny

surfaces were produced but fracture occurred at all laser power levels. Since argon is an inert gas, cutting was carried out by a decomposition/evaporation process, according to the following reaction:

Si3N4 >3Si + 2N2 (3)

The temperature of the above reaction is > 1785 C.

The penetration of molten silicon was impeded

through the entire thickness and therefore fracture

occurred. Visual and optical examination have

confirmed that the first three mm of the thickness

has the best surface finish achieved by the laser

beam under any of the experimental conditions. It

is expected that such laser cutting can be utilized to produce a high quality surface finish, which

require no further finishing, for thin silicon nitride

materials, e.g. wafers and substrates of thickness

less than three mm.

2. Cutting using pulsed laser

When a thermal pulsing mode was utilized,

significant reduction in fracturing was observed

due to the diminished thermal shock compared to condnuous mode laser operation. Under oxidizing

cutting condidons, surface finishes similar to those

of the condnuous mode were obtained. No

significant differences in surface finish were nodced with changing frequency (150 to 900 Hz).

However, microstructural examinations using

SEM have revealed that porosity is more uniform

and finer in pore size than that found during the

continuous cutting mode. Figure 2 shows such

topographical features which are distinct from

those shown in Figure 1.

545
Fig.2 SEM micrograph showing the morphology of

the cut edge for silicon nitride cut with 600 W of

pulsed laser power and oxygen assist gas. CONCLUSIONS

1. High power CO2 laser cutting of silicon nitride

ceramic materials can be successfully achieved

under oxidising conditions using either continuous

wave or pulsed operation.

2. The oxidised surface layer has a thickness of

150 to 350 }im, which requires further grinding

using conventional diamond tooling.

3. Thin wafers or substrates made of silicon nitride

ceramic materials can be successfully machined

using inert assist gases, e.g. argon, yielding a

surface finish requiring no further grinding.

4. The mechanisms involved in the laser cutting of

silicon nitride ceramics depend on the assist gas, which controls the cutting and surface finish

quality.

REFERENCES

1. R.N. Katz, " Recent Development in High

Performance Ceramics," in Materials Technology - 1976, A.G. Chynoweth and W.M. Walsh, Jr., Eds., Am. Inst. Phys., Pub., pp. 236-252 (1976).

2. R.H. Chand, K.P. Castello and T.M. Payne,

"

Finish Machining of Technical Ceramics - An

Engineering Approach," in Proc. of SME

Conference on Advanced Ceramics '90, Rosemont,

Illinois (1990).

3. J. Yamamoto and Y.Yamamoto, "Laser

machining of Silicon Nitride," in Proc. LAMP '87 , Osaka, Japan, pp. 297-303.

4. M. Naeem, M.E. Preston and J.R.Tyrer,

"Machining of Engineering Ceramics With A High

Power C02 Laser," Adv. Manuf. Eng. , Vol. 2, 27-

31 (1990).

5. R.F. Firestone, "Laser and other non-abrasive

machining methods for ceramics," in Advanced

Ceramics '87 Conference, 17-19 February 1987 , MR87- 112-6.

6. I.J. McColm and N.J. Clark, Forming, shaping

and working of high-performance ceramics. Chapter 5, Chapman and Hall Publisher, London,

England (1986).

7. T. Araya and T. Umino, "Present Status of C02

Laser Processing," J. Mater. Eng. , 15, 299-306

(1991).

546
APPENDIX

Energy Consumption of Laser vs Diamond

Machining

To demonstrate that laser cutting of ceramic

materials is an energy efficient process compared

to conventional diamond machining, tests were

performed on the cutting of a 6 mm thick plate

(100 mm diameter) of silicon nitride using both

techniques.

With diamond cutting, a total cut length of 144 cm

was made in 128 min using a surface grinder,

which consumed 2.3 kW of electrical power. Thus

the electrical energy consumed is 4.9 kW-hr.

Using the carbon dioxide laser at a power level of

600 W, the same length of cut was made in 9.45

min. Such a laser consumes 12.5 kW of electrical

power, giving electrical energy consumption of

1.97 kW-hr. Therefore, the laser machining would

require only 40% of the energy consumed by

diamond maching.

547

TECHNOLOGY IN PULSED ND:YAG LASER DRILLING

OF AN AL-LI BASED/SIC METAL MATRIX

COMPOSITE

T. M. YUE and W. S. LAU

Hong Kong Polytechnic, Hong Kong

C. Y. JIANG

Nanjing Aeronautical Institute, Nanjing, China

The behavior of laser micro-drilling of an Al-Li/SiC metal matrix composite was studied using

a pulsed Nd:YAG laser. The concept of critical intensity was employed to predict the profile

and the diameter of the drilled hole. The dimension of the hole was found to be related to the

relative position of the laser focal spot to that of the workpiece. In the case of through hole

drilling of thin sheet material, the diameter of the hole can be best controlled by the beam

intensity. Moreover, a moving laser source, by nature of its tunnelling action, can improve on

the uniformity of the dimension of the processed hole, this is particularly valid for thick

materials.

1. INTRODUCTION

The unique properties of Al-based/SiC metal

matrix composites (MMCs) are, undoubtedly, of great

attractions to many engineering applications.

However, they have not been widely utilized until quite recently because of the high production cost and

technical difficulties encountered in their manufacture.

To-day even with the advanced manufacture routes,

though they can be produced in a wide variety of

grades and forms (1-4), most of these materials are not produced to finished net shapes. Thus postproduction shaping becomes essential to generate

usable engineering components.

It is well accepted that Al-based/SiC

composites are difficult to machine; and apparently the

only sensible tooling material for this kind of exotic

materials is polycrystalline diamond(PCD) (5,6).

Nevertheless, in the realm of micro-machining, laser technology seems to provide the best alternative. The

"non-contact" nature and the extreme high intensity of

the beam precludes the problem of tool wear and also minimizes any unacceptable microstructural changes.

At present, information uncovered in the open

literature concerning laser machining of MMCs are scarcely available. Previous study has shown that pulsed Nd:YAG laser can offer a real possibility in cutting Al-Li/SiC composites in that the quality of the

cut was largely dictated by the beam intensity (7).

This research aimed to deepen the knowledge on the behavior of laser drilling of Al-Li/SiC materials so as

to establish a model capable of predicting the profile

of the "machined" hole.

2. MATERIAL

The material subjected to the present

investigation was a newly developed Al-Li(8090)/SiC

MMC which has already attracted much attentions in the aerospace industry owing to its extremely high

specific mechanical properties. The material was in 2.2mm sheet supplied in the annealed condition. The

nominal composition of the matrix material is presented in Table 1. The matrix is reinforced with

17%(wt.) SiC particles having an average size of

3^iTi. Figure 1 shows a typical microstructure of the composite material.

Table 1 Major elements in 8090 alloy

Element Weight (%)

Li 2 .2-2 .

Mg 0 . 5-1 .

Cu 1 . 0-1 .

Al balance

549
Fig.l Typical microstructure of the MMC material

Since, ' '

and

i(M=j^ (3)

/(0,0) R^z)

Now the threshold radius which describes a circle of

r = ri(z), having its intensity = I„ can be written as:

3. THEORETICAL BACKGROUND

In spite of the complex and unsteady

conditions exist during laser processing in that the

material is melted, vaporised, and expelled in a very

short time span. It has been shown that in laser

drilling of various materials the depth of the

"machined" hole is strongly related to the critical intensity (I J* of the laser beam and for the 8090/SiC

composite material the I,, was found to be 7.5x10'

W/m= (8). Assuming a Gaussian TEMqo mode condition

I, is given by:

r,(z) = [^2(0)+(e^)2]in

/(0,0)i?-(0) (4)

\ /Ji?2(O)+(0z)2]

where.

/(0,0):

2PLP IE

(5)

(Plp= the pulse power; E = pulse energy; tj = pulse

duration)

/[r,(z),z] =/(0^)exp[—^] (1)

The notations of the various parameters are defined as shown in the following diagram:

According to equation 4 the profile of the

machined hole will depend on E and t, provided the

physical parameters of the laser are fixed. And from

equation 4 both the maximum value of rt(z), i.e. r^^,and its corresponding position, Z^^, can be

obtained by taking the derivative of equation 4 with

respect to z:

Schematic diagram of the Laser optic system

is the threshold intensity which initiates vaponzation

-R\0)/Q (6)

4. EXPERIMENTAL PROCEDURES

Laser micro-drilling was conducted on a pulsed Nd:YAG laser with a maximum time-average

power of 120W. The beam parameters were varied in the experiment with the pulse energy(E) ranging from

1 to 5J/pulse; pulse duration(tj)= 1ms; with the pulse

550
frequency(fp) fixed at 23Hz. All of the experiments were carried out using argon as the flushing gas

operating at a pressure of 0.2 MPa.

To validate equation 1, i.e., the profile of the

machined hole, blind-holes were drilled in the plane of

the rolling direction using five different beam

intensities: 1000W/cm% 1500W/cm-, 2000W/cm%

2500W/cm' and 3000W/cm\ Based on the outcome of

these results, the relationship between the maximum

diameter of through holes that can be achieved with

different beam intensities was studied by a defocusing

method (locating the surface of the material at the

position corresponding to Z^J.

5. RESULT

The blind-holed specimen was sectioned using

the EDM wire-cut and the profiles are revealed in Figure 2. Figure 3 compares some of these results to those as predicted by equation 3. In general, the

experimental results agreed well with the theory,

except at the position near the entrance of the blind

holes; the results will be discussed in a later section.

As in the case of through hole drilling, the

relationship between the beam intensity and the hole

diameter at the beam entrance and exit surfaces are presented in Figure 4 and Figure 5 respectively and

the top surface of the drilled specimen is shown in Figure 6. The measured diameter at the entrance side was significantly larger than the predicted value,

whereas for the exit surface they were reasonably in good agreement.

Fig. 2 Blind hole profile as a function of intensity

(I, from nght to left: 1000-3000W, steps 500W)

Q.OO -

theoretical -.50

experimental

-1.00

-1.50 s

-2.00 -

-2.50

-3.00 M

-3.50 -

-4.00 1 1 1 1 1 1 1 -2.00 -1.50 -1.00 -.50 0.00 0.50 t.OO 1.50 2.00 r^(2) (mm)

Fig. 3a Blind hole profile at intensity = 2000W

— theoretical i\

1 1 1 —I— 1 1 1 -2.00 -1.50 -1.00 -.50 0,00 O.SO 1.00 1,50 2.00

r^(z) (mm)

Fig, 3b Blind hole profile at intensity = 3000W

Fig. 6 Through hole diameter as a function of intensity

(I, from left to right: 1000-5000W steps lOOOW)

551
0.9G

0.80

. 0.70

0-60 +

i 0.50 a

„ 0.40

OJoI 0.30

0.20

theoretical

expenoieriTal

2000 3000 4DO0 5000 Intensity .1^. (Ul

Fig. 4 Intensity vs. top surface hole diameter

0.60

ro 0.40 - Q

o 0.30 •• X

theoretical experimental

0 1000 2000 3000 4000 5000 6000 Intensily .1^, 'Ul

Fig. 5 Intensity vs. exit surface hole diameter

6. DISCUSSION

(a) Blind hole drilling

The good agreement between the experimental

profile and that as predicted by equation 2 suggests

that in blind hole drilling the contour of the hole can be best described by a single parameter namely, the

critical intensity, I,. The results, both theoretical and

experimental, also show that a spear-head profile is obtained in the absence of relative movement between

the workpiece and the laser beam along the Zdirection. The fact that the dimension at the entrance of the hole is larger than the predicted value is possible due to the melting effect occurring at the

surface. As the laser beam reaches the surface of the

specimen, areas next to the vaporized zone would be

melted and expelled with the aid of the argon jet.

Such skin effect is also time dependent, i.e. if more

pulses are required to accomplish the hole then the

surface regions next to tht vaporized zone would be

exposed longer to sufficiently high radiation to melt

the material. As a consequence, the tapered entrance would become larger and deeper. However, such

552

effect was much less pronounced inside the blind hole

because of the shielding effect of the top surface

stopping most of the laser radiation of less than that of

the critical intensity from entering the hole. In addition to this, due to the geometry of the hole any

melted phases would be less likely to escape. Built on

this argument, the profile of the blind hole at the

surface could be better described by the melt zone

contour based on calculations of heat conduction

models rather than using the critical vaporizing

criterion.

(b) Through hole drilling

In theory, through hole drilling is possible

provided the thickness of the sheet does not exceed the

overall length of the profile. But again the shape of

the hole would not be uniform and will depend on the

position of the workpiece relative to the focal point of

the laser beam according to equation 4. For instance,

if the bottom surface of the workpiece is at the

corresponding Z^^ position then an inverted-taper

hole would form (Figure 7). Nevertheless, it is possible to improve the uniformity of the hole by a moving laser source travelling in and out of the

specimen. Figure 8 presents the result of such a case, apparently the profile of the hole machined with the

laser tunnelling action was better than the one without.

Therefore, for thick materials if a moving laser source

is employed or as in the case of thin sheets, by

placing the top surface of the material at the Z^^

position the diameter of the hole can be estimated

using equations 6 and 4. Results shown in Figures 4

and 5 confirmed that the size of the hole was related

to the intensity of the beam and it is believed that by

further refinement of the theory of co-operating with

the melted boundary condition a more representative

solution can be obtained.

Based on equation 4, it is also possible to control the dimension of the through hole by a defocusing technique, in that the position of the

specimen can be placed at different levels along the

beam path. However, when using such technique to machine thin specimens, corrections must be made to account for the melted regions.

The idea of employing the concept of critical beam intensity to predict the profile of the "machined"

hole can, certainly, offer some practical information on laser micro-machining of MMC materials. The

approach is unique in the sense that it alleviates many

of the uncertainties such as metal reflection, laser-gasmaterial interactions, etc. which arise during laser processing.
laser drilling direction

Fig. 7 Specimen placed at resulted an

inverted taper hole

O.Smm

Fig. 8 Left: with a stationary laser source; Right: with a travelling laser source

CONCLUSIONS

(1) The profile of the laser machined holes can be

determined using the critical intensity

criterion; however, the model required

modification when applied to the entrance

region.

(2) The profile of the hole would vary depending on the relative position of the workpiece to that of the laser focal spot.

(3) The maximum diameter of the machined

through holes was related to the critical intensity.

(4) In using a moving source technique, the

uniformity of the through hole can be

improved.

REFERENCES

1. A.Niklas, L.Froyen and L.Delaey,

"Comparative Evaluation of Extrusion and

Hot Isostatic Pressing as Fabrication

Techniques for Al-SiC Composites" , Materials

Science and Engineering . A135 pp. 225-229,

1991.

2. A. R.E. Singer, "Metal Matrix Composites

Made by Spray Forming" Materials Science

and Engineering , Vol.A135, pp. 13-17. 1991.

3. T.M.Yue and G.A.Chadwick, "Squeeze

Casting of Light Alloys and Their

Composites-A Review", Proceedings of the

6th International Manufacturing Conference

with China. 10-12 March 1993, Hong Kong.

4. D.R.Pank, J.J.Jackson, "Metal Matrix

composite Processing Technologies for Aircraft Engine Applications", ASM/TMS

Materials Week ^92 . 2-5 November 1992,

Chicago, Illinois USA.

5. M. Brun, M. Lee and F. Gorsler, "Wear

Characteristics of Various Hard Materials for Machining SiC-Reinforced Aluminium Alloy",

General Electric Corporate Research &

Development , Report No.84CRD327, 1984.

6. G.A. Chadwick and P.J. Heath, "Machining

Metals Matrix Composites", Metals and

Materials . Feb., pp. 73-76, 1990.

7. W.S. Lau, T.M. Yue, C.Y. Jiang and S.Q.

Wu, "Pulsed Nd:YAG Laser Cutting of

SiC/Al-Li Metal Matrix Composite",

Proceedings of the Machining of Composite

Materials Symposium . ASM/TMS Materials

Week, Chicago USA, 2-5 Nov. 1992, pp.29-

34.

8. C.Y. Jiang. W.S. Lau, T.M. Yue and L. Chiang, "On the Maximun Depth and Profile

of Cut In Pulsed Nd:YAG Laser Machining",

to be published in the CIRP Annals 1993 .

ISBN 0-16-041820-8

J^-U. S. GOVERNMENT PRINTING OFFICE : 1993 O - 341-931 (82673)
it
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