I


NIST Special Publication 766

NBSIBAM 1986 Symposium on

Advanced Ceramics

Edited by

S. M. Hsu

Ceramics Division

Institute for Materials Science

and Engineering

National Institute of Standards

and Technology

Gaithersburg, MD 20899

and

H. Czichos

Bundesanstalt fur Materialprufing

Unter den Eichen 897

D-lOO Berlin 45

German Federal Republic

May

U.S. Department of Commerce

Robert A. Mosbacher, Secretary

National Institute of Standards and Technology

Raymond G. Kammer, Acting Director

NOTE: As of 23 August 1988, the National Bureau of

Standards (NBS) became the National Institute of

Standards and Technology (NIST) when President

Reagan signed into law the Omnibus Trade and

Competitiveness Act
Library of Congress U.S. Government Printing Office

Catalog Card Number: 89-600737 Washington: 1989

National Institute of Standards

and Technology

Special Publication 766

Natl. Inst. Stand. Technol.

Spec. Publ. 766

159 pages (May 1989)

CODEN: NSPUE2

For sale by the Superintendent

of Dociunents

U.S. Government Printing Office

Washington, DC 20402
PREFACE

The National Institute of Standards and Technology (NIST) (formerly NBS)

and the Bundesanstalt fur Materialprufung (BAM) share many similar missions

and responsibilities in materials characterization, standards, and

measurement sciences. Consequently, scientific interactions have been

conducted on an informal basis between the organizations for many years.

In 1984, these activities were increased through the auspices of the

Versailles Project on Advanced Materials and Standards (VAMAS)

.

Discussions with Professor H. Czichos of BAM at the Gordon Research

Conference on Tribology in 1984 and the Wear Conference in 1985 indicated

that more frequent, structured, exchanges of scientific and technical data

between our organizations would be valuable. Specifically, Professor

Czichos suggested an annual seminar series similar to the well-recognized

Leeds-Lyon Conference on Tribology. Enthusiastic response from NIST and

BAM culminated in a 2 -day seminar on ceramics at BAM in April 1986. A

second joint seminar was held at the NIST in April 1987 and addressed

corrosion and wear. A third seminar on polymers was held in September 1988

at BAM.

My participation in both seminars has convinced me of the value of these

exchanges and I hope that these meetings which further international

collaboration and understanding will continue.

This volume, a result of the first seminar, contains valuable data and

scientific information which is of particular use to the ceramics technical

community

.

STEPHEN M. HSU

Chief, Ceramics Division

National Institute of Standards and Technology

(formerly the National Bureau of Standards)

ABSTRACT

Advanced ceramics offer many advantages that other materials do not

possess. They have high strength, dimensional stability, are chemically

inert, lightweight, wear resistant, and have desirable properties in

electrical, optical, and thermal applications. At high temperatures, they

are the only class of material with reasonable properties. As such,

advanced ceramics is increasingly used in electronic, structural, wear,

optical, and electrical applications and devices.

Worldwide production of advanced ceramics is growing rapidly. Since

ceramics are based on alumina and silica, the most abundant minerals on

earth, effective utilization of ceramics carries implications into the next

several centuries.

One of the major technical barriers to widespread use of ceramics is the

inability of industry to manufacture reliable ceramics reproducibly and

economically. Advanced ceramics are sensitive to small defects introduced

during processing and generated during use. The identification of key

parameters and subsequent measurements of these parameters are crucial to

control the quality of ceramics. The National Institute of Standards and

Technology (NIST) (formerly National Bureau of Standards, NBS) , and the

Bundesanstalt fur Materialprufung (BAM) to a large extent, share a similar

mission to assist their respective industries to achieve reliable

manufacturing through standards. This symposium therefore provides timely

exchange of technical information on a very significant subject area.

V

CONTENTS

Page

A Critical Assessment of Requirements for Ceramic Powder

Characterization; A. L. Dragoo, C. R. Robbins, and

S. M. Hsu, NBS . 1

Microstructural Modification in Ceramics Resulting from

Diffusion Induced Boundary Migration; J. E. Blendell, NBS 19

Structural Reliability of Ceramics at Elevated Temperatures;

E. R. Fuller and S. W. Wiederhorn, NBS 29

An Ultrasonic Study of Ceramics from the Green State to

Fully Fired; A. L. Dragoo, NBS 42

Nondestructive Characterization of Ceramic Materials: Thermal Waves

G. S. White, NIST, E. H. LeGal LaSalle,

Laboratoire d'Optique Physique, ESPCI 59

Computer-aided Ultrasonic Testing of Non-Oxide Ceramics;

A. Hecht, E. Neumann, P. Rose, BAM 79

Computertomography ; J. Goebbels, H. Heidt, A. Kettschau, and

P. Reimers, BAM 97

Nondestructive Characterization of Ceramic Materials: Small Angle

Neutron Scattering, K. A. Hardman-Rhyne and E. R. Fuller, NBS 101

Introduction to the Versailles Project on Advanced Materials and

Standards (VAMAS) Technical Working Area: Wear Test Methods;

S. Becker and J. Lexow, BAM Ill

Fretting of Ceramic Materials, D. Klaffke, BAM 126

Tribology of Advanced Ceramics, S. M. Hsu, NBS 144

vii

A CRITICAL ASSESSMENT OF REQUIREMENTS FOR CERAMIC POWDER

CHARACTERIZATION

A. L. Dragoo, C. R. Robbins and S. M. Hsu

National Bureau of Standards

Gaithersburg, Maryland

Abstract

The detailed characterization of ceramic powders is very important for the

reproducible manufacture of advanced ceramics that will perform reliably

in service. The basic issues are: what to measure, how to measure it and

how to assure quality in analytical measurements by all laboratories. "Vfhat

to measure" involves understanding the relationship between powder

characteristics and ceramic microstructures . "How to measure it" requires

the development of measurement methods and the determination of

repeatability and reproducibility. Standard Reference Materials (SRM) are

required to assure the quality of measurements and the comparability of

measurements between different laboratories and techniques. From

considerations of the distributed nature of powder and ceramic propei'ties,

new SRM' 3 are proposed which will certify distributed properties. As an

example of such SRM's, technical requirements are developed for the

production of an SRM with a certified particle size distribution for

ceramic powders. Factors which enter the certification of a particle size

distribution for a ceramic powder include the approximate mathematical

representation of the distribution, weighting of the distribution by the

measurement technique, the particle shape distribution, and statistical

variances introduced by the powder, the sampling methods, and the methoa of

measurement

1
I. Introduction

In their study conducted for the National Bureau of Standards, Charles

River Associates Inc. (CRA)^ identified three key barriers to the

commercialization of ceramic materials for advanced technological

applications: production cost, reliability in service, and reproducibility

in manufacture. For the case of ceramic components for heat engines CRA

saw reliability as a major technological barrier and improved

characterization of ceramic powders as providing "important leverage in

overcoming the reliability and reproducibility problems. "2 Improved

characterization of ceramic powders requires not only improved measurements

but also improved understanding of how powder characteristics are related

to properties and performance of resulting ceramic materials.

Ceramic powders and ceramic microstructures share a common feature

that is important for the improved characterization of both powders and

microstructures: i.e., due to local variations on the scale of grain size,

characteristics are distributed rather than single-valued in nature. The

distributed nature of characteristics requires that a measured

characteristic must be represented, at least, by two statistics: by a

mean and a standard deviation for the material variation. These two

statistics often may not be sufficient; for example, if the distribution is

multimodal or higher skewed.

The importance of the distributed nature of ceramic characteristics is

met, for example, in the fact that the strength of a material under brittle

fracture is determined by a limiting flaw which may be a crack, void, pore

cluster, or large grain. Limiting flaws are the larger outliers within the

population of flaws. Flaws result from a variety of factors including

2
variability in the starting powder. For example, hard agglomerates in

ceramic powders have been noted as the cause of gross voids in a variety of

ceramics sintered from active powders.^'®

What are the characteristics of powders which most influence the

microstructure of a resulting ceramic? And, how can the relevant

properties be measured best? Table 1 lists many of the physical and

chemical characteristics of powders. It is evident from a survey of the

physical characteristics, in particular, (1) that some properties are

related, such as specific surface area, porosity and particle size; and (2)

that characterization of a feature, such as agglomerates, involves several

characteristics. Further, a qualitative characteristic such as the

"hardness" of an agglomerate requires description in terms of both chemical

and physical characteristics which are measureable. Thus, a single

characteristic does not appear to be generally sufficient to predict the

sinterability of a powder.

Techniques, such as sedimentation and electrical zone sensing, which

employ a macroscopic amount of material for measurement of particle size

and porosity provide a measure of the distributed aspect of those

characteristics. As an example of the considerations required in the

development of a standard for a distributed characteristic of a powder,

requirements for particle size SRM will be considered below in Sections III

to VI. The ability to measure a distributed characteristic on a

macroscopic quantity of powder is not the case for many other

characteristics, especially for chemical characteristics. Information

regarding the distribution of a chemical characteristic is most often

obtained by microscopic techniques, such as analytical microscopy.

3
Improvement in the measurement of physical and chemical characteristics

of powders requires development of improved measurement techniques and

sample preparation procedures that have increased precision and accuracy.

Precision is "the agreement within a set of observations or test results

no obtained as are directed in a method. ^ Each observation or test must be

carried out in a manner to assure the independence of each result.

"Repeatability" is defined here as that aspect of precision involving

comparison of results within a single laboratory and with a single

technique. "Reproducibility" is defined here as that aspect of precision

involving comparison between laboratories or between different techniques.

Accuracy is "the degree of agreement between the true value of the

quantity ... and the average of many observations. ^ The terms "bias" and

"systematic error" are used here to refer to the deviation between the true

and observed values. Standards are required to measure the bias between

techniques and in the case of absolute measurements, to establish the

amount of bias within a technique.

Standard Reference Materials (SRM) are characterized and certified by

the National Bureau of Standards for some chemical or physical properties.

The issues of precision and accuracy are addressed in the characterization

of these materials. Kaye^^ has drawn attention to the need for the

National Bureau of Standards to play a greater role in providing standards

for particle size and shape for the fineparticle industry. With the

expectation that ceramic powders will provide some of the future

fineparticle standards, the NBS Ceramics Division is engaged in a

4
long-range program of developing and certifying Standard Reference

Materials (SRM's) intended to improve the assurance of chemical and

physical quality of ceramic powders.

II. Existent Particle Size Standards

A variety of reference materials are presently available for the

calibration of particle size measurements. Table 2 summarizes the particle

size standards which are available from governmental organizations.

Table 3 lists a variety of particle size references which are available

commercially. Spherical polymer particles with very narrow size ranges

have been certified for single particle size values. SRM 1690 consists of

polystyrene spheres which were manufactured on board a space shuttle to

achieve improved roundness. Mulholland et al.^^ described the

certification of SRM 1690, a nominal one micrometer diameter standard.

Although latex and polystyrene spheres have excellent dimensional and shape

characteristics, limited quantities and costs may restrict their use to

applications such as primary calibrants in a laboratory and as reference

materials where only small quantities are required. la addition, these

materials are transparent to x-rays and cannot be used to calibrate

instruments which use x-ray attenuation. Finally, these materials provide a

single-point calibration.

The reference materials based on glass beads and quartz powder or sand

have certified size distributions obtained by sieving, sedimentation and

electrical sensing zone techniques. Although glass beads can be obtained

with nearly spherical shapes, it is difficult to prepare glass materials

that are bubble free. The development and certification of quartz powder

5
and sand as Certified Reference Materials (CRM's) for distribution by the

Community Bureau of Reference of the European Economic Community were

described by Wilson^^^

The introduction of new synthesis methods for ceramic precursor

powders, such as reactions in flames and plasmas, and the use of alkoxides

and other organometallic precursors in aqueous and nonaqueous liquid

systems, has begun to provide a variety of submicrometer sized particles

which often have very regular shapes. The growing availability of these

powders poses challenges to physical characterization methods used to

measure them and to the standards used to assure the quality of these

powders. However, these new powders offer the opportunity of new source

materials for improved standards for physical, as well as chemical,

characteristics of powders. Here, we briefly examine some of the

requirements for such a ceramic powder if it is to provide the basis for a

Standard Reference Material for particle size measurements. A standard for

submicrometer sizes and with a size distribution approximating an idealized

distribution, such as the well-known lognormal distribution, is of special

interest because it will enable comparisons of different methods of size

measurement to be compared with a greater degree of confidence.

III. Empirical and Idealized Particle Size Distributions

Estimates of a mean size x, and a standard deviation, o, always can

be obtained directly from the experimental data. However, if a particle

size distribution can be represented satisfactorily by an ideal or model

distribution, then estimates of other statistics can be obtained. For

example, it is possible to calculate distribution means^^^

6
X max X max

- f xqj(x)dx/ f ^1^5^)

Jx min Jx min

and arithmetic means

X max X max 1 /k

^k.O - f x'<qo(x)dx/ f qo(x)dx (2)

Jx min Jx min

where the index refers to the basis of the distribution: i.e.,

j - 0 for a number distribution

J » 1 for a length distribution

j - 2 for an area distribution

J - 3 for a volume distribution.

M-I^j is the first order moment with respect to an Jth type distribution;

M|^^Q is a kth order moment with respect to a number distribution. Since

different measurement techniques impose different distributions on a

population of particles, a functional representation of a particle size

distribution enables .results of various techniques to be compared.

The lognormal distribution which was introduced over a century ago by

Galton^'^ and McAlister^^ been found to model observed particle size

distributions found both in nature and from industrial processes^ ^"^^. In

addition to the lognormal distribution there are several other

distributions which may be considered. The Rosin-Ramraler distribution^ ^'

was found to describe the properties of powdered coal. The Pearson Type V

distribution^^ has been used as the assumed fiinctional form for particle

size measured by a commercial photon correlation spectrometer^^.

7
For the lognormal distribution the percentage, cl(t), of particles

between z and z + dz, where z - inx, is expressed as

r (z-z)'^-|

d<<) - q(z)dz » exp - dz. (3)

1 r (z-z)2-

exp

The statistics z and of the distribution can be expressed in terms of

the geometric mean, Xg, and the geometric standard deviation:

z = inxg (4a)

(J2 = in Og (4b)

The geometric mean is equal to the mode and median values if the number

frequency distribution is plotted with respect to the logarithm of the

particle size. The lognormal distribution has the property that the

standard deviation is independent of the choice of basis. Simple

logarithmic relationships exist between the geometric means for the higher

order distributions and Xg and Og. Also, various arithmetic means and

moments of lognormal distributions with respect to different bases can be

related to these statistics^^.

Procedures for fitting an empirical distribution to a lognormal

distribution and for testing the distribution were given by Kottler^'^"^''^.

He examined the calculation and use of the log-probit plot. Although such

a plot appears linear if the data exhibits lognormal behavior, the data

cannot be fitted by a simple linear regression procedure. From the

cumulative probability

t

p - ; q(t)dt (4)

—00

where t is the normal deviate, t = (x - x)/a, and, where q(t) is given by

8
eq (3) I a change, 6P, is related to 6t by

6P - q(t) 5t . (5)

It is t which is proportional to Inx. To estimate the goodness of fit by

a calculation of Chi-squared for the errors in the normal deviate, it is

necessary to use the Mueller-Urban weights

W = NCq(t)2/P(l-P)] (6)

where N is the number of particles in the population. The standard error

bands about the P' « P(computed) versus log x line are given by P* ± S',

where

S' - CP'(1-P')/N]^/2 (7)

Minimization of the value of Chi-squared is complicated by the fact that

the errors in P are cumulative, and therefore, not Independent. Kottler

described a method of successive approximations by which Chi-squared can be

minimized.

IV. Particle Shape

Ceramic powders, with a few exceptions, do not consist of spherical or

nearly spherical particles. Kaye^*-* has defined an ideal reference material

as "usually composed of dense, smooth, spherical fine particles, the

characteristics of which can be measured directly with different

characterization procedures that employ different physical measurement

technology." However, the condition of sphericity may be relaxed

(1) if the particle shape is regular so that a physical dimension can

be related to a measured equivalent diameter;

(2) if a numerical shape characteristic is narrowly distributed for a given particle size;

(3) if the shape characteristic does not vary significantly over the

range of particle sizes of interest.

9
The relaxation of the sphericity condition requires the determination of

one or more particle shape factors. Development of shape factors will

require the use of microscopy and image analysis to obtain related

measurements of particle size and shape. King28»29 discussed the use of

image analysis and the development of particle size statistics from image

data.

V. Methods for Particle Size Measurements

The concept of particle size, although intuitively meaningful, is

operationally ambiguous since the measured value depends significantly on

the method of measurement because the measurement imposes a particular

selectivity on the particle size distribution. Variance in the measurement

is sometimes superimposed on the inherent variance of the powder. Biases

may be introduced by the experimental procedure and by the approximations

which relate the instrumental response to the particle size. At the

present time, a definitive method does not exist for measuring a particle

size distribution. A definitive method is "based on sound theoretical

principles and which has been experimentally demonstrated to have

negligible systematic errors and a high level of precision"30. achieve

high accuracy in a definitive method, especially for measurement of a

physical property, the measurement must be referenced to the standards for

the fundamental units. Determination of a particle size distribution and

associated shape factors will require a suite of methods which may include

microscopy and image analysis, sedimentation with ultracentifugation and

electrical sensing zone measurements.

10
Wood^'' has discussed various methods of particle size analysis from the

standpoint of the dynamic range of the method and the basis, or weighting,

factor, j, of the measured distribution, see Eq. (1) and (2). Fig. 1

summarizes a variety of size analysis methods with respect to dynamic range

and weighting factor. Wood distinguishes between a primary weighting

factor which describes the manner in which classes of particles are counted

and a secondary weighting factor which is a function of equivalent diameter

basis and the measurement principle. An analytical technique which

determines the equivalent diameter (ED) from projected area has a secondary

weighting factor J - 2; a technique which determines ED from a volume

measurement has J « 3; and a technique which determines ED from a

surface-to- volume ratio has j = 2.5.

It is important that the method used to measure the distribution must

not yield a large standard error, S, see Eq. (8), over the range of

certification. As Kottler^^ noted, S is proportional to N"'/^ where N is

the total number of particles measured. For a technique which measures the

relative concentration on a mass basis, such as sedimentation, N is

replaced effectively by the mass of particles in the volume over which the

measurement is carried out. This consideration implies that image

analysis methods may require the analysis of on the order of 10,000 images

to obtain a satisfactorily limit of uncertainty.

VI. Certification, Packaging and Stability of SRM's

The development and certification of an SRM requires the performance of

the following tasks:

identification of the measurement need;

determination of the characteristic to be certified;

identification of potential source materials;

11
small, actual differences between measurement methods which become

discernible with improved precision;

• underestimate of the standard deviation for a method.

VII. Conclusion

As ceramic comp>onent3 find a wider application in the national and

international economies, the need for stcindards for quality assurance in

their manufacture and use will increase. Since the quality of the starting

powder is an essential factor in attaining reliability and reproducibility

of ceramic components, it follows that new Standard Reference Materials will

be needed for the measurement of particle size. Since the particle size of

ceramic powders is generally a distributed property, SRM's with certified

particle size distributions are expected to be necessary for the

development and use of improved ceramic materials.

Since SRM's are developed to meet the measurement needs of industry and

commerce, it is important that new SRM's for particle size measurements

reflect user needs with respect to material, size range and possible

auxiliary properties. Consequently, the suggestions of and collaboration

with the user community are welcome.

Acknowledgements

The partial support of this work by the Office of Automotive Technology

Development, Department of Energy, is gratefully acknowledge.

References

1. J. R. H. Black, S. L. Blur:, S. H. Kalos, and G. C. Tassey, "An

Economic and Technical Assessment of Advanced Ceramic Materials,"

Am. Ceram. Soc. Bull. , 64 [1] 39-MI , 50 (1985).

2. Charles River Associates Inc., Technological and Economic Assessment

of Advanced Ceramic Materials, Vol. 1; Planning Report 19.

NBS OCR 8A-470-1 . National Bureau of Standards, August 1984.

3. C. A. Bruch, "Sintering Kinetics for the High Density Alumina

Process," Am. Ceram. Soc. Bull., 41 [12] 799-806 (1962).

12
H. J. S. Reed, T. Carbone, C. Scott, and S. Lukasiewicz, "Some Effects of

Aggregates and Agglomerates in the Fabrication of Fine Grained

Ceramics," p. 171 in Processing of Crystalline Ceramics. Edited by

H. H. Palmour, R. F. Davies, and T. M. Hare. Plenum Press, NY, 1978.

5. W. H. Rhodes, "Agglomerate and Particle Size Effects on Sintering

Yttria-Stabilized Zirconia," J. Am. Ceram. Soc. , 64 [1] 19-22 (1981).

6. A. L. Dragoo and L. P. Doraingues, "Preparation of High-Density

Ceria-Yttria Ceramics," J. Am. Ceram. Soc , 65 [5] 253-59 (1982).

7. F. F. Lange and M. Metcalf , "Processing-Related Fracture Origins: II,

Agglomerate Motion and Cracklike Internal Surfaces Caused by

Differential Sintering," J. Am. Ceram. Soc. , 66 [6] 398-406 (1983).

8. F. W. Dynys and J. W. Halloran, "Influence of Aggregates on Sintering," J. Am. Ceram. Soc , 67 C9] 596-601 (1984).

9. Standard Terminology for Statistically Related Terms, ASTM Standard

D4392-84

10. B. H. Kaye, "Standard Finepart icles - Who Needs Them?" Chemtech,

40-43, Jan. 1980.

11. G. W. Mulholland, A. W. Hartman, G. G. Hembree, E. Marx and

T. R. Lettieri, "Development of a One-Micrometer Diameter Paricle Size

Standard. Reference Material," J. Res. Natl. Bur. Stand. (U.S.) , 90

[I] , 3 - 26 (1 985).

12. R. Wilson, "Reference Materials of Defined Particle Size Certified

Recently by the Community Bureau of Reference of the European

Community," Powder Tech ., 27, 37-43 (1980).

13. DIN 66141, "Darstellung von Korn-(Teilchen)groessen-verteilungen,"

February 1974.

14. F. Galton, "The Geometric Mean, in Vital and Social Statistics," Proc.

Roy. Soc. London , 29, 365-67 (1879).

15. D. McAlister, "The Law of the Geometric Mean," Proc Roy. Soc. London , 29. 367-76 (1879).

16. J. Aitchison and J. A. C. Brown, The Lognormal Distribution with

Special Reference to Its Use in Economics. Cambridge University

Press, Cambridge, 1957.

17. K. Mahmood, "Lognormal Size Distribution of Particulate Matter," J.

Sedimentary Petrol. , 43 [ 4] 1 1 61 -66 ( 1 973 )

18. S. K. Kurtz and F. M. A. Carpay, "Microstructure and Normal Grain

Growth in Metals and Ceramics. Part 1. Theory," J. Appl. Phys. , 51

[II] 5725-44 (1980).

13
19. P. Rosin and E. Rammler, "The Laws Governing the Fineness of Powdered

Coal," J. Inst. Fuel. 7, 29-34 (1927).

20. P. Rosin and E. Rammler, "Grinding Properties of Coal," Zeraent, 20

210-13, 240-41 , 317-21 , 343-46, 363-65 (1 931 ).

21. J. K. Ord, "Pearson System of Distributions," pp. 655-59 in

Encyclopedia of Statistical Sciences, Vol. 6. Wiley, New York, 1985.

22. L. G. Bunville, "Commercial Instrumentation for Particle Size

Analysis," pp. 1-42 in Modern Methods of Particle Size Analysis.

Edited by H. G. Barth. Wiley, New York, 1984.

23. T. Allen, Particle Size Measurement, 3rd ed.; pp. 157-61. Chapman and

Hall, London, 1981 .

24. F. Kottler, "The Distribution of Particle Sizes, Part I," J. Franklin

Inst ., 250, 339-56 (1950).

25. F. Kottler, "The Distribution of Particles Sizes, Part II,"

Franklin Inst ., 250, 419-41 (1950).

26. F. Kottler, "The Goodness of Fit and the Distribution of Particle

Sizes, Part I," J. Franklin Inst ., 251, 499-514 (1951).

27. F. Kottler, "The Goodness of Fit and the Distribution of Particle

Sizes, Part II," J. Franklin Inst ., 251, 617-41, (1951).

28. R. P. King, "Determination of the Distribution Size of Irregularly

Shaped Particles from Measurements on Sections or Projected Area,"

Powder Technol. , 32, 87-100 (1982).

29. R. P. King, "Measurement of Particle Size Distribution by Image

Analyzer," Powder Technol. , 39. 279-89 (1984).

30. J. K. Taylor, Handbook for SRM Users. NBS Special Publication

260-100. National Bureau of Standards, Sept. 1985.

31. B. C. Wood, "Selection Criteria for Particle Size Analysis Equipment,"

Powder Bulk Solids , 35-39 (April 1986).

32. T. Allen and A. A. Kahn, "Critical Evaluation of Powder Sampling

Procedures," Chem. Engr. , CE108-CE112 (May 1970).

33. T. A. Hatton, "Representative Sampling of Particles with a Spinning

Riffler. A Stochastic Model," Powder Technol ., 19. 227-33 (1978).

34. R. C. Oppenheim and J. A. Hersey, "Physical and Chemical Changes in

Powders on Storage," J. Powder Bulk Solids Technol ., 1 (3), 3-9

(1977)

.

14
Table 1. Characteristics of Ceramic Starting Powders

Physical Characteristics Chemical Characteristics

Grains (Primary Particles)

Size

Shape

Agglomerates (Secondary

Particles

)

Size

Shape

Porosity

Amount (See Density)

Size

Shape

Particle Contact

Coordination

Strength

Bulk Composition

Major Elements

(1 to 100$)

Minor Elements

(10 ppm to U)

Trace Elements

( < 10 ppm)

Inorganic Species

(for example,

sulfates and

nitrates

)

Organic Species

Water and other

volatiles

Phases

Density

Specific Surface Area

Permeability

Compactibil ity

Flowability

Surface Composition

Major Elements

Minor* Elements

Trace Elements

Inorganic Species

Organic Species

Water

Phas es

15
Table 2. Particle size standards certified by governmental organizations.

Particle size is in units of um.

Material Source Use Property

SRM Wb, Magnification NBS

Scale

SRM 474, Line Width

SRM 1688, Latex Spheres

SRM 1689, Latex Spheres

SRM 1690, Polystry. Spheres

SRM 1691, Polystry, Spheres

SRM 1692, Polystry. Spheres

SRM I960, Polystry. Spheres

SRM 1003, Glass Beads

SRM 1004, Glass Beads

SRM 1017a, Glass Beads

SRM 1018a, Glass Beads

CRM 066, Quartz Powder BCR

CRM 067, Quartz Powder

CRM 068, Quartz Sand

CRM 069, Quartz Sand

CRM 070, Quartz Powder

CRM 130, Quartz Sand

CRM 1 31 , Quartz Sand

CRM 132, Quartz Sand

CRM 165, Latex Spheres BCR

CRM 166, Latex Spheres

CRM 167, Latex Spheres

SEM(s)

OM(s)

In preparation

In preparation

Diam

In preparation

Stokes Diam

Sieves

Sieves

Sieves

Stokes Diam

Stokes Diam

Vol Diam

Stokes Diam

Stokes Diam

Vol Diam

Vol Diam

Vol Diam

Particle

Counting

,

Light scattering

1.2.3,5,50

1 to 12

or planning

or planning

-0.9

~0.3

or planning

'10.0

3 to 5

34 to 120

100 to 310

225 to 780

0.35 to 3.50

2.4 to 32.0

160 to 630

14 to 90

1.2 to 20

50 to 220

480 to 1800

1400 to 5000

2.223±0.013

4. 821 ±0.019

9.475±0.01

8

16
Table 3. Examples of commercially available particles for size references,

Polymer Spheres

Polystryrene (PS)

Polyvinyl-Toluene (PVT)

Polyvinyl-Divinyl Benzene (DVB)

Polystryrene-Butadiene (PSB)

Spherical Pollens

Paper Mulberry

Ragweed

Bermuda Grass

Walnut

Corn

Uniform Glass Microspheres Test Dusts

AC Fine

Arizona Road

17
18
MICROSTRUCTURAL MODIFICATIONS IN CERAMICS RESULTING

FROM DIFFUSION INDUCED BOUNDARY MIGRATION

J. E. Blendell

Ceramics Division

Institute for Materials Science and Engineering

National Bureau of Standards

Gaithersburg, MD 20899

INTRODUCTION

The recent observations of Diffusion Induced Grain Boundary

Migration (DIGM) in ceramics have caused a reevaluation of the stability of

interfaces under conditions where diffusion is occurring. We will review

what DIGM is, show typical microstructures that occur, explain why this

phenomenon is important to ceramics processing and give some examples of

how DIGM affects the properties of ceramic materials.

DIGM is the migration of grain boundaries due to the diffusion of a solute in the boundary.^ The boundaries are otherwise stable; that is,

they will not migrate at these temperatures without the introduction of the

solute. In DIGM experiments this stabilization is usually accomplished by

annealing the samples at high temperatures to cause grain growth. When the

solute is introduced into the system at a lower temperature, the boundaries

migrate. The area swept out by the moving boundaries has been alloyed.

The alloying has occurred to a much larger extent than can be accounted for

by lattice diffusion at the experimental temperatures. DIGM has also been

observed during dealloying, in which the solute diffuses out of the sample.

For a critical review of DIGM see Handwerker.^

There is an analogous process. Liquid Film Migration (LFM) , that occurs in solid- liquid systems under similar conditions to DIGM except that

a liquid film separates the grains, rather than a grain boundary. The

microstructures observed in LFM are nearly identical to those observed in

DIGM. The differences are due to the shape of the grains in a solid- liquid

system. For example, in the MgO system, with a liquid present the grains

tend to be spherical, and when no liquid is present the grains tend to be

polyhedral

A typical microstructure produced by DIGM is shown schematically in

Figure 1, along with many of the other features commonly observed in DIGM.

The original position of the boundary is shown as a dotted line, the final

19
positions are solid lines and the cross-hatched regions have been alloyed

with the solute. The differences between DIGM and grain growth are striking. During DIGM boundaries move away from their centers of curvature

as they migrate, where as in grain growth the boundaries move towards their

centers of curvature. The area of the boundaries increases during DIGM in

contrast to grain growth where the grain boundary area decreases. Also

DIGM occurs at temperatures much lower than those at which grain growth

occurs. During grain growth, the boundaries tend to remain smoothly

curved, or locally planar if facetting occurs. Boundaries that are migrating during DIGM often show sinusoidal shapes with the same boundary

migrating into both adjacent grains. As shown in Figure 2, a range of

boundary shapes are observed in the same sample although there is usually a tendency for one boundary shape to dominate.

One important problem about DIGM has been the identification of the

driving force. There exist many driving forces for boundary migration,

such as curvature, phase transitions, temperature gradients and electric

field gradients. When these have been eliminated by experimental design,

DIGM is still observed during solute diffusion. While there is a reduction

in the free energy of the system by the formation of the alloy, this energy

reduction is not sufficient to drive boundary migration.^ For example, in

isotope self diffusion, the same free energy of mixing is dissipated, but

there is no coupling of the energy of mixing to a force driving the motion

of the boundary. This is because the isotope atoms are chemically

identical and the boundary has no interaction with the isotope. Thus the

boundary does not "know" that isotope self diffusion is occurring. Another

way to think about it is that for a boundary to move there must be a force on it. The motion of the boundary can be prevented by applying an opposite

force. If the free energy decrease during isotope mixing could drive the

motion of the boundary, then there exists a force that could prevent the

boundary motion and therefore, isotope mixing. As there is no force that can be applied to prevent, or reverse isotope mixing, the free energy

decrease due to isotope mixing can not apply a force on the boundary. The

non- ideal terms in the free energy might, in theory, apply a force on the

boundary, but the occurrence of DIGM both for positive and negative

deviations from ideal solution behavior, and both for alloying and

20
dealloying, with identical microstructures and velocities, demonstrate that

these excess thermodynamic terms do not provide the driving force.

It is now well established that the driving force for DIGM is the

coherency stress due to the difference in lattice parameter between the

constrained alloyed region and the unalloyed bulk.'' All other proposed

mechanisms have been eliminated either by thermodynamic arguments or by

experiment. Most importantly, the critical experiment of demonstrating

that the migration velocity goes to zero as the difference in lattice

parameter between the diffusion zone and the unalloyed bulk goes to zero,

has been performed.''

COHERENCY STRESS MODEL

DIGM arises when there is diffusion of a solute down a grain boundary

and into the lattice adjacent to the grain boundary. When the solute is of

a different atomic size than the solvent, this diffusion causes the lattice

to be locally stressed. When there is a difference in stress between the

grains, atoms will migrate from the higher stressed side, across the

boundary to the lower stressed side. The boundary will thus migrate into

the region of higher stress. The velocity of a boundary is determined by

the difference in stress across the boundary and the rate of diffusion of

atoms across the boundary.

The stress necessary for DIGM will exist even when lattice diffusion

is effectively "frozen out", that is, at low temperatures. A atomic jump

of only one lattice spacing into the lattice from the grain boundary is

needed to create the coherency stress. Even at low temperatures there is

always some mixing of solute atoms from the grain boundary into the

lattice, although it may be very limited in depth.

The stress difference across the boundary can arise for several

reasons. Because of the anisotropy of the elastic modulus, even for cubic

materials, there will be a difference in stress for different

crystallographic orientations. The strain is determined by the change in

lattice parameter with composition and is isotropic for cubic materials.

Alternatively, if the dislocation content of one grain is higher, or one grain is oriented so that the stress can cause dislocation generation, then

stress relief can occur in that grain. Thus there will always be a difference in stress across a grain boundary which will initiate migration.

21
For the case of two solid grains separated by a liquid (LFM) , the

compositional change due to the stress can be determined from the free

energy, f , as a function of composition using a common tangent

construction.^ This is shown in Figure 3. is the composition of

the stressed material and c^g is the equilibritim composition. At the

stressed grain, material of composition c^^^^^j.^^^^ will dissolve in the

liquid, migrate across the liquid film (down the composition gradient) and

precipitate on the unstressed grain at composition c^^.

MICROSTRUCTUEAL EFFECTS

DIGM has been observed^ in the MgO system with additions of NiO or CoO

as shown in Figure 4. Boundaries are induced to migrate, increase their

curvature and cause alloying at much higher rates than when diffusion is

not occurring.

LFM has also been observed^ in the same system, using V2O5 as the

liquid phase. Typical microstructures are shown in Figure 5. The

composition profiles have been determined® for the liquid case and are shown in Figure 6. The conditions of temperature and time used in these

experiments were calculated using the coherency stress model.

The microstructures observed are very different from typical

structures seen in ceramics. The interlocking nature of the grain

boundaries is a feature of boundaries that have moved due to DIGM. This

feature can lead to enhanced creep resistance by increasing the diffusion

distance. Another unique result of DIGM is a reduction in the grain size.

Since all the grain boundaries migrate into the adjacent grains, a first

order model predicts that there will be a reduction in the grain size by

approximately a factor of 14 (the number of neighbors). Thus it may be

possible to reduce the grain size of a ceramic by solute diffusion after

sintering. This would provide advantages in increased strength and in

addition, could significantly change the electrical properties. In

metallic systems DIGM has been observed to switch from occurring in the

bulk, to a surface phenomenon as the temperature is increased.^ This can

be used to modify the surface of a material to produce a surface coating

without changing the bulk microstructure or properties.

However there are situations where DIGM would not be advantageous.

For fiber reinforced composites, DIGM can lead to increased bonding between

22
the fiber and the matrix. This will reduce the effectiveness of fiber

pull-out as a toughening mechanism. For ceramics used as nuclear waste

storage materials, boundary migration due to the presence of compositional

inhomogeneities can occur. Radioactive material could be leached out at

significantly higher rates than predicted. Since DIGM can occur at very

low temperatures and long times, DIGM will not be observed in leach tests

performed in the laboratory.

SUMMARY

DIGM is the motion of grain boundaries due to the diffusion of a solute down the grain boundary and into the lattice. The areas swept out

by the boundary are alloyed while areas not swept remain unalloyed. DIGM

occurs at temperatures much lower than those where grain growth occurs

The driving force for DIGM is the coherency stress due to the

difference is atomic size between the solute atoms and the solvent

material. The coherency stress model has been very accurate in predicting

the temperature and times necessary for observing DIGM and LFM. Other

proposed mechanisms have been eliminated either by thermodynamic arguments

or by experiment.

The changes in the microstructure due to DIGM can have a major effect

on the properties of materials . Interlocking boundaries may increase the

high temperature creep resistance of ceramics, change the toughness for

grain boundary fracture, but will degrade the properties if a specific

microstructure is required as in many composites. Also, the ability of

ceramics to be effective nuclear waste storage materials may be reduced by

the dealloying that can occur. Thus DIGM is a important process for

determining the long term performance of ceramic materials.

23
REFERENCES

1. M. Mllert and G.R. Purdy, Acta Met, 26, 333(1978).

2. C.A. Handwerker, Diffusion in Thin Films , ed. D. Gupta, Noyes

Publications, (1987).

3. R.W. Balluffi and J.W. Cahn, Acta Met, 29, 493(1981).

4. W. Rhee and D.N. Yoon, Acta Met, 35, 57(1987).

5. C.A. Handwerker, J.W. Cahn, J.E. Blende 11 and D.N. Yoon, Diffusion in

Solids: Recent Developments , ed. M.A. Dayananda and G.E. Murch, - Metallurgical Soc . , Warrendale, PA (1985).

6. J.E. Blendell, C.A. Handwerker, C.A. Shen and N.D. Dang, Ceramic

Microstructures '86 . ed. J. Pask and A.G. Evans, Plenum Press, New

York(1987)

.

7. C.A. Handwerker and J.E. Blendell, to be published.

8. C.A. Handwerker, R. Marinenko, D.E. Newbury and J.E. Blendell, to be

published.

9. J.E. Blendell, C.A. Handwerker, W.A. Kaysser and G. Petzow, Surfaces

and Interfaces in Ceramic and Ceramic -Metal Systems , ed. J. Pask and

A.G. Evans, Plenum Press, New York(1981)

.

24
Figure 1. Schematic of microstructures observed in DIGM. Boundaries

that were initially Straight have migrated away from their

centers of curvature and alloying has occurred only in the

regions swept by the boundaries . The regions behind the

migrated interface have epitaxial relationships with the

matrix grains as indicated by the letters.

4

~1 /

Figure 2

.

Mlcrostructure observed in Ircn foils that have been exposed

25
A Co Ccoherent ^coherent B

Figure 3. Free energy, f, as a function of composition in a binary

system, A-B. Shown is the common tangent construction for

determining the composition of an unstressed planar

solid, Cgq, in contact with its equilibrium liquid and for a coherently stressed solid, c^^j^^^^^^, in contact with its

equilibrium liquid.

Figure 4. DIGM in MgO-NiO system. Surface of MgO sample exposed to NiO

at 1280°C for 16 hrs. The surface relief due to solute

addit ion in the regions swept by the grain boundaries can be

clearly seen. Facetting is seen on several grain surfaces.

26
Liquid Film Migration in MgO-V2 05 as a result of Co diffusion

at 1250°C for 20 min. Boundaries have migrated away from

their centers of curvature and alloying, as indicated by the

lighter gray, has occurred. The grains were initially

spherical

.

27
Figure 6. Quantitative Electron Microprobe analysis showing Co

concentration in the grains due to alloying. The original

grains are pure MgO as indicated by the darkest gray. The

regions of intermediate gray indicate alloying with Co.

28
STRUCTURAL RELIABILITY OF CERAMICS

AT ELEVATED TEMPERATURES

E. R. Fuller, Jr. and S.M. Wiederhorn

Ceramics Division

Institute for Materials Science and Engineering

National Bureau of Standards

Gaithersburg, MD 20899

ABSTRACT

Current methodologies for assessing the structural reliability of

ceramic materials at elevated temperatures are reviewed. Typically these

methods have been based on fracture mechanics concepts, such as those used

to describe low- temperature , brittle fracture, in which failure is assumed

to ensue from the subcritical growth of a pre-existing flaw to a critical

size. However, recent high- temperature data for ceramics suggest that this

view may be over simplified, and that flaw or damage accumulation may be an important aspect of high- temperature reliability. The need for

incorporating these concepts into a high- temperature reliability

methodology is discussed.

INTRODUCTION

If strength were the sole criterion for structural performance, hightemperature structural ceramics, not metallic alloys, would be the

materials of choice for a number of high- temperature structural

applications [1], e.g., gas turbine engines. This observation was based on

the fact that the strength of the better high- temperature structural

ceramics is greater than that of the alloys currently being used in turbine

engines. Accordingly, in figure 1, the strength of a grade of yttriadoped, hot-pressed silicon nitride is compared with that of Waspaloy, and

several other alloys used for high- temperature applications [2,3]. The

figure indicates that this silicon nitride has a potential 400 C use- advantage over Waspaloy, as indicated by the temperatures for which the

strength of both materials falls precipitously. If a way could be found to

utilize this high- temperature strength of ceramics, then the economy of a myriad of high- temperature structural applications would be improved

considerably

.

But to be used in heat- temperature applications, ceramics not only

have to be strong, but they also have to be tough. Reliability of design

also requires that the variance of strength be low to avoid failure as a consequence of subcritical crack growth, or momentary overloads in service.

Demands such as these are difficult for ceramic materials to meet, since

they traditionally lack the low variance in strength, the high fracture

toughness, and the thermal shock resistance of metals. Nevertheless,

high- temperature resistance to chemical attack, and a high mean strength at

elevated temperatures suggest a considerable potential for these materials

in heat engines. This potential for high- temperature structural

29
.'()0 400 600 800 10OO 1200 TEMPERATURE. C

Figure 1. Comparison of the fracture strength of a yttria-doped silicon

nitride [3] with the yield strength of three commercial alloys [2].

applications has been the incentive during the past fifteen years for most

of the work to improve the mechanical behavior of structural ceramics.

At room temperature, ceramics are brittle materials. Their strength

is determined by microstructural flaws, such as voids, inclusions and

cracks, which may be present in the material. At high temperature,

ceramics exhibit some ductility. However, since the number of degrees of

freedom for plastic flow is generally not sufficient for arbitrary shape

changes, ceramics still behave in a semi -brittle manner. In addition,

time -dependent plastic flow, such as occurs during creep, results in the

generation of cracks and cavities that eventually lead to stress rupture of

ceramics. These two characteristics of ceramics, flaw sensitivity at low

temperature and flaw generation at high temperature, largely determine the

mechanical behavior and structural reliability of these materials.

In this paper we review the techniques that have been developed to

improve the reliability of structural ceramics, and we relate these

techniques to the mechanical and chemical behavior of the materials at

elevated temperature. Four important high- temperature phenomena related to

structural reliability are noted: 1) The enhanced chemical reactivity of

materials at elevated temperatures, which can lead to structural

modifications and flaw generation; 2) the existence of a threshold stress

or stress intensity factor for creep crack growth; 3) a predominantly

strain-controlled creep-rupture regime for stresses below this crack growth

threshold, where reliability is governed by distributed damage processes;

and 4) the redistribution of applied stresses as a result of creep and

damage processes

.

30
STRUCTURAL CHANGES AND FLAW GENERATION

At elevated temperatures a number of new processes can occur that

influence structural reliability. For example, kinetic processes, such as chemical reactions, diffusion, and viscous flow, occur more rapidly, and

thermodynamic processes can result in phase transformations and new phase

equilibria. Even in the absence of an applied stress, strength degradation can occur as a consequence of specific chemical reactions between the

ceramic and the environment, or because of physical transformations that occur as a consequence of phase instability. For example, the reaction of

inclusions in magnesia- doped, hot-pressed silicon nitride with the oxygen

in the air can resulted in the formation of surface corrosion pits that

replaced machining damage as the source of failure [4-6]. Other examples

include: second-phase precipitation along grain boundaries, as in yttriadoped silicon nitride [7], devitrification of a grain-boundary glass, as in

vitreous bonded alumina [8] , and chemical modification of the grains

themselves, as in zirconia toughened ceramics [9-10]. Specific chemical or thermodynamic reactions are usually not considered in reliability

methodologies, and techniques for predicting failure as a consequence of

these processes have not been developed. Because of this, materials

exhibiting these degradation mechanisms must be either improved by ceramic

processing, or else not used under conditions where the degradation can occur

.

When stresses are present, the high- temperature strength of a material

may degraded either by the growth of cracks from flaws, as at room temperature, albeit usually by different processes, or by the generation, growth

and linkage of creep cavities or cracks. At the same time such strength

degradation processes are occurring, other processes may occur that reduce

the severity of flaws and thereby improve the strength. The resulting

strength will thus depend on the relative kinetics of the competitive

processes. This competition is manifested by a threshold for crack

growth, above which creep crack growth is the predominant failure mechanism

and below which stress -enhanced flaw generation is the important mechanism.

These concepts are elaborated further in the following sections.

CRACK GROWTH THRESHOLD

Although the growth of cracks and flaws can be an important factor in

the lifetime of structural materials at elevated temperatures, the

influence of creep processes, sintering and chemical reactions on the

fracture process cannot be ignored. These processes can cause effective

"blunting" or "healing" of flaws or cracks , whereby giving rise to a crack

growth threshold, i.e., a threshold stress intensity factor, K^^ • below

which cracks do not propagate. However, when the applied stress intensity

factor, Kj , exceeds this threshold, Kj > K^^ > crack propagation results in

the same type of delayed failure as is observed at low temperatures,

although crack propagation in this instance may be accompanied by a cracktip creep deformation zone. Below this threshold, propagating cracks

arrest as a consequence of creep or other high- temperature processes that

reduce the flaw severity. This threshold becomes an important hightemperature design parameter, because the material's response is

drastically different above and below the threshold. Above the threshold,

31
design by crack growth models of failure [11-15] is appropriate; below the

threshold, damage generation models are appropriate [16-18].

The existence of a high- temperature threshold for crack growth has

been suggested by a number of authors [5,16-23], based on both strength and

crack propagation data. The most convincing evidence for such thresholds

comes from strength studies on specimens containing small indentation

cracks. Such specimens are subjected to a static stress until failure occurs and the time- to- failure is recorded. If failure does not occur in a specified time, the specimens are intentionally broken in a controlled

manner to measure their strength. "Blunting" or "healing" of the

indentation crack results either in failure from some site other than the

indentation, or in a substantial increase in the strength. Techniques such

as these have been used by Tressler and co-workers [21-22] to determine the

crack growth threshold in silicon carbide, and by Quinn [20] to map out the

crack growth regime for magnesia-doped, hot-pressed silicon nitride.

Recent data by Grathwohl [16] and by Dalgleish et al. [18] clearly

indicate a boundary between those failure mechanisms controlled by crack

growth, and those mechanisms controlled by flaw generation. Grathwohl'

s

data, for example, show that different creep rupture responses are obtained

for smooth and indented specimens, see figure 2. Specimens that failed

from the indentation cracks follow a continuous stress -rupture curve,

typical of a crack- growth dominated stress -rupture behavior. The stress

10° 10^ 10^ 10^ lO'* 10^

TIME TO FAILURE (min)

Figure 2. High- temperature creep rupture data for a hot-pressed silicon

nitride (Annawerk GmbH) at 1200°G. At a given applied stress, the timeto-failure is plotted both for as-machined specimens (open circles) and for

precracked specimens (filled circles). The precracked specimen at 100 MPa

did not fail from the precrack. From Ref. [11].

32
rupture curve for the smooth bar specimens, however, shows a distinct break

at a critical value of the applied stress. The high stress portion of the curve is identical to that of the indented specimens and therefore, is

believed to be the result of crack growth. The break in the curve represents a possible "blunting" threshold, indicative of a transition from

crack- growth dominated failure processes to creep rupture and/or flaw

nucleation processes , although Grathwohl interprets this break as a consequence of the stress relaxation in the flexural specimen from creep.

CREEP RUPTURE

At elevated temperatures, creep is the process that limits lifetime of

most ceramics. Because ceramics are inherently brittle, even at elevated

temperatures, lifetime is not usually limited by component extension, that occurs as a consequence of creep, but rather by the formation of cavities

and cracks, that are generated by the creep process. High- temperature

processes (sintering, plastic deformation, chemical reaction) that anneal

cracks and reduce their severity also generate new flaw populations that

are now the sources of failure. This dual role of high- temperature

processes in determining failure is illustrated in figure 3 for a crack in

siliconized silicon carbide which has been loaded at 1300°C [24]. Although

the crack has been blunted by creep deformation and its growth arrested,

continued deformation in the vicinity of the original crack tip results in

the formation of cavities along shear bands. The linkage of these cavities

accounts for the eventual failure of this material. Blumenthal and Evans

[25] have observed a similar behavior for a commercial grade of hot-pressed

aluminum oxide at elevated temperature.

Figure 3. Blunted crack in siliconized silicon carbide at 1300°C.

Continued deformation causes the crack to generate damage along shear bands

emanating from the blunted crack tip [24]

.

33
Creep -rupture at elevated temperatures usually involves both damage

generation, such as the nucleation and growth of cavities, and the linkage

of these cavities to form cracks that cause failure. Depending on the

specific damage mechanism, any of these processes may be the rate limiting

step for failure. In an attempt to understand high- temperature failure

mechanisms of ceramics and their dependence on microstructure , a considerable amount of theoretical work has been conducted to model the

failure process [26-33] . Failure mechanisms are usually classified

according to material type: those having a two-phase structure, and those

having no second phase at the grain boundary. Usually these mechanisms

assume that the time-to-failure depends only on one of the critical

processes given above. The results are expressed in terms of a parametric

relationship between failure time, temperature, applied stress and strain

rate, as was done by Monkman and Grant [34] for metals. Ideally the

constants in these parametric relationships can be expressed in terms of

physical properties, so that, in principal, the failure time can be related

both to basic material properties, such as viscosity and diffusivity, and

to microstructural characteristics, such as grain size and volume fraction

of second phase. To date, however, quantitative comparisons between theory

and experiment have not proven satisfactory, suggesting the need for

additional work to understand the fundamentals of creep rupture in

ceramics

.

Accordingly, current understanding of creep -rupture in ceramic

materials at elevated temperatures is largely empirical. The small body of

data that is available suggests that the failure of structural ceramics

follows parametric relationships, such as the Monkman-Grant relationship

[34], in which the time-to-failure, tf , has an inverse relation to the

minimum strain rate, *, so that, tf* = constant. This relationship between

strain rate and time- to- failure was first demonstrated for ceramics by

Kossowsky et al. [35] for a grade of magnesia-doped, hot-pressed silicon

nitride, figure 4. Since two distinct curves are obtained for the data in

helium and the data in air, this study also indicates an influence of

environment on the creep-rupture time. More recently, Grathwohl [16]

confirmed these findings with flexural tests for another grade of silicon

nitride. Grathwohl also measured the strain during creep and found that

the strain-to-failure was constant for his material, as would be expected

from the Monkman-Grant relation. This finding supports the idea of a strain- limited rupture criterion. Recent studies by Johnson et al. [36]

and by Wiederhorn et al. [8] suggest that the Monkman-Grant coefficient,

tf%, for some ceramic materials is not a constant, but rather depends on

the applied stress and/or the temperature.

The development of a data base for lifetime prediction at elevated

temperatures requires the evaluation of Monkman-Grant coefficients for

material-environment systems of interest. As these coefficients must

dependent on both the material's microstructure and the environment, and

since no unifying concepts have emerged from creep-rupture data to date, it seems unlikely that a universal coefficient will be found for this

equation. However, once the Monkman-Grant coefficient, and its stress and

temperature dependence are known (or for that manner, the parameters for some other parametric relationship) , the relationship can be used for the

prediction of failure time. These equations replace those derived for the

crack- growth dominated lifetime predictions at low temperatures.

Strain- limited failure criterion, such as the Monkman-Grant

relationship, implies that lifetime at elevated temperatures is determined,

34
5x10 r ^

I I M 1 1

1

1 ~n I I I I M 11 n—I I

I 1 1 1

1

1

10 ^ 10 ^ _

10 ^2 Steady State Creep Rate z^{U M

Figure 4. Dependence of creep-rupture life on the steady state creep rate

in tension for a hot-pressed silicon nitride (HS-130). The solid and open

symbols are for specimens tested in air and in helium, respectively. From

Ref. [35].

not by the rate of crack growth, but by the rate of creep. Accordingly,

improvements in lifetime can be achieved through the use of processing

techniques to reduce the creep rate. Since the dependence of creep rate on micros true tural variables is well understood, the use of creep theory to

improve lifetime at high temperature is a feasible procedure.

CREEP REDISTRIBUTION OF STRESS

Thus far, the stress state has been assumed to be invariant in time.

However, when a component is subjected to stress gradients, such as in a flexural test specimen, this assumption is only valid if the creep rate is

linearly proportional to the stress. If the creep rate is not linearly

proportional to stress, then the stress distribution will change in time,

and depending on the processes controlling creep, steady state creep may

not be established. Several steady state solutions for this stress

redistribution have been obtained for power law creep of flexural beams,

examples of which are shown in figure 5. This figure illustrates three

35
COMPRESSIVE STRESS

COMPRESSION SIDE ^^c/^ol^

1 2

Re = n, = 2

s-10 7 nc=n,^2; S-1

^

/\ A 1

-elastic

1 TENSILE STRESS TENSILE SIDE

((Jt/Oe,)

Figure 5. Calculated steady-state stress distributions for a flexural beam

in creep. The creep rate in compression and tension are given,

respectively, by the equations: = A d^^ and \^ - k (Sa)"*'

.

separate cases for the steady state creep stress distribution. The

"elastic" solution is obtained for a linear relation between strain rate

and stress. The sigmoidal- shaped curve is obtained for power law creep

with equal rates of creep in tension and in compression. However, when the

creep rate is more rapid in tension than that in compression, an asymmetric

stress distribution, such as the one illustrated, results. The S=10 in

figure 5, corresponds to a factor of ten in the stress dependence of the

creep rate. For this last case, the neutral axis moves towards the

compressive surface of the component [37-40]. To fully understand creep

that is asymmetric in stress, creep data in both compressive and tensile

loading must be obtained.

An example of nonlinear creep behavior is illustrated in figure 6 for

a grade of siliconized silicon carbide which is known to nucleate cavities

when subjected to a tensile load at high temperatures [41]. Cavities in

the tensile portion of the bend bar do not support the applied load, a consequence of which is an internal reduction of the cross -sectional area

of the test specimen on the tensile side of the flexure bar. The local

stress on the material located near the cavities is higher than the average

stress, and thus, the material creeps faster in this region than in the

compressive section which contains no voids. The net result of this

behavior is a shift of the tensile stress field towards the compressive

section of the flexure specimen and a reduction of the average tensile

stress in the specimen. The shift of the neutral axis can be demonstrated

experimentally from a microstructural examination of the specimen

36
COMPRESSIVE

SURFACE

0.2 mm

Figure 6. Cavity formation within the cross-section of a siliconized

silicon carbide flexure specimen which has been tested at IBOCC and an initial applied stress of 300 MPa [41]. Note that cavity formation on the

tensile side of the specimen has proceeded to beyond the original neutral

axis of the specimen, indicating a shift in the neutral axis.

37
subsequent to testing. As illustrated in figure 6, just prior to failure

cavity generation has occurred all the way across the tensile section and

past the geometric center of the specimen. This shift of the stress

distribution can also be demonstrated by placing parallel rows of

indentations into the side surface of test bars [41]. After deformation,

the position of the plane of neutral strain can be shown to have moved

towards the compressive side of the specimen, suggesting a similar movement

of the neutral stress axis. The transient solution to this problem has

recently been analyzed by Chuang et al . [40].

Regardless of the exact process controlling creep, the fact that the

stress in a component may change with time has to be factored into the

design of ceramic components to obtain accurate predictions of lifetime.

The reduction of the tensile stress as a function of time, for example,

gives rise to a stress threshold for failure, as the stress in the tensile

surface decreases [16-17]. Our understanding of bending stresses and their

role in the creep rupture process is still at an early stage of development

and work is needed to elucidate material behavior in multi-axial stress

fields. Although some steady-state solutions of creep in flexure have been

obtained, transient state creep analyses are also needed. Such transient

state creep solutions have recently been obtained by Chuang et al. [40] for

beams in flexure, but other stress state configurations will be needed for

practical lifetime predictions. In addition, practical analyses for creep

rupture of ceramic materials will require creep measurements to be made in

both tension and compression to characterize the material behavior in

"pure" stress states, before more complex biaxial stress states can be

understood. Measurements in the tensile configuration are especially

needed, because, at present, very little definitive work has been conducted

on structural ceramics using tensile testing.

In addition to the discussion above, the relation between damage

mechanisms and creep rates will have to be studied more fully to develop a fundamental understanding of the factors that determine the reliability of

ceramic materials at elevated temperatures. The incorporation of damage

processes into theories of creep have been pursued by only a few authors

,

see for example ref. [42]. Two facts emerge from these studies. First,

damage and creep are interrelated: creep is affected by the damage process,

and visa versa. Second, each damage mechanism is controlled by a distinct

set of equations that determines the relation between the creep and the

damage processes, and ultimately, the rupture time. In this regard, Ashby

and Dyson [42] have shown that by combining different damage mechanisms

with damage theory, the time- to -failure for metallic materials can be

expressed in terms of a Monkman- Grant- type equation. The dependence of the

Monkman-Grant coefficient on stress and temperature depends on the

particular mechanism responsible for creep - rupture . The value of the

coefficient is inversely related to the tolerance of the material to strain

at failure. Analyses, such as that presented by Ashby and Dyson [42], have

yet to be applied in a systematic way to ceramic materials.

SUMMARY

Concepts of reliability are well established for room- temperature

applications of ceramic materials. The assumption that the failure

originates from pre-existing flaws is supported by a considerable body of

experimental data. Based on this assumption, structural reliability

38
dictates control of the flaws present in a structural material and control

of the external conditions (stress, environment, and temperature) that lead

to the growth of flaws and eventually to component failure. There is a substantial body of theory to deal with room- temperature failure, which is

currently being applied in both low- and high- temperature structural

applications

.

At high temperature the situation is somewhat different. If the

applied stresses are greater than a threshold stress for crack growth, then

failure also depends on pre-existing structural flaws, in which case, the

techniques developed for failure prediction at room temperature can also be

applied to high- temperature behavior. However, in these cases crack

propagation may be accompanied by crack- tip creep processes, which will

modify the methodology somewhat. In addition, the time -dependence of the

stress, both the applied stress and the creep-modified stress state, now

becomes an important variable.

When the applied stresses are less than the threshold for crack

growth, pre-existing flaws may become benign with failure occurring from

flaws that are generated by the high- temperature processes. Concepts of

reliability in this creep -rupture regime must be based on these flaw

populations, which implies an understanding of the flaw generation, flaw

growth and linkage processes. Empirical descriptions of lifetime, however,

are available for this regime of material response: failure time and

minimum creep rate appear to be related by a Monkman-Grant- type

relationship. To obtain a deeper understanding of the failure process,

theoretical studies are required to relate the coefficients for these

Monkman-Grant relations to the physical process occurring during creep

rupture. Also, a larger data base is needed to test the theories that have

been developed.

When non-uniform stresses are present in a component, such as the

bending stresses in a flexural test specimen, the applicability of current

lifetime prediction schemes are questionable, because of the redistribution

of these stresses with time. Stress no longer becomes a design constant,

but rather becomes part of the lifetime analysis. As indicated above,

creep -modifications to the stress state also influence the manner in which

high- temperature data is obtained, i.e., shifts in the neutral axis of

flexural test specimens with time. Our understanding of these processes is

in an early stage. Although theoretical analyses of the creep problem in

bending are available for steady state creep, and are just becoming

available for transient state creep, the dependence of creep on damage

accumulation mechanisms has still to be clarified for practical ceramics.

Finally, to fully understand the behavior of materials that are subjected

to complex, time -dependent stress states, creep data are required in "pure"

stress states (for example, both compression and tension as a minimum), and

most likely, in mixed stress states as well.

Acknowledgement: The authors are pleased to acknowledge the support of the

U.S. Department of Energy, Fossil Energy Materials Program (ERF) and U.S.

Department of Energy, Advanced Heat Engines Project (SMW)

.

39
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Deutsche Kermische Gesellschaft , 1986, pp. 911-929.

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E.M. Lenoe, R.N. Katz and J.J. Burke, Plenum Press, New York, 503-533

(1983)

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Mater. Sci. 21, 810-824 (1986).

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Science and Technology of Zirconia, II, edited by N. Claussen, M.

Ruble and A.H. Heuer, Am. Ceram. Soc, Columbus, OH, 371 (1984).

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12. J.E. Ritter, Jr., in Fracture Mechanics of Ceramics, Vol. 4, edited by

R.C. Bradt, D.P.H. Hasselman and F.F. Lange, Plenum Press, New York,

667-686 (1978).

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(1973) . 14. S.M. Wiederhorn and E.R. Fuller, Jr., Mat. Sci. and Engr. 71, 169-186

(1985)

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Orlando, FL, January 24-26, 1977, in ERDA Rep. CONF-770110, pp. 231-

246 (Energy Research and Development Administration)

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16. G. Grathwohl, in Proc. 2nd Intl. Conf. on Creep and Fracture of

Engineering Materials and Structures, edited by B. Wilshire and D.R.J.

Owen, Pineridge Press, Swansea, UK, 565-577 (1984).

17. A.G. Evans, Mater. Sci. Eng. 71, 3-21 (1985).

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575-581 (1985).

19. M.H. Lewis, B.S.B. Karunaratne, J. Meredity and C. Pickering, in Creep

and Fracture of Engineering Materials and Structures, edited by

B. Whilshire and D.R.J. Owens, Pineridge Press, Swansea, 365-379 1981.

20. G.D. Quinn, in Methods of Assessing the Structural Reliability of

Brittle Materials, edited by S.W. Freiman and CM. Hudson, ASTM Spec.

Tech. Publ. 844, 177-193 (1984).

21. E.J. Minford andR.E. Tressler, J. Am. Ceram. Soc. 66, 338-340 (1983).

22. R.E. Tressler, E.J. Minford and D.F. Carroll, in Creep and Fracture of

Engineering Materials and Structures, edited by B. Wilshire and

D.R.J. Owen, Pineridge Press, Swansea, 551-563 (1984).

23. E.R. Fuller, Jr., S.M. Wiederhorn, N.J. Tighe, and L. Chuck, Mater.

Sci. Eng. 71, 209 (1985).

40
24. L, Chuck and E.R. Fuller, Jr., to be published.

25. W. Blumenthal andA.G. Evans, J. Am. Ceram. Soc. 67, 751-759 (1984).

26. R.L. Tsai and R. Raj, ActaMetall. 30, 1043-1058 (1982).

27. J.R. Porter, W. Blumenthal and A.G. Evans, Acta Metall. 29, 1899-1906

(1981) . 28. C.H. Hsueh and A.G. Evans, ActaMetall. 29, 1907-1977 (1981).

29. A.G. Evans andA.S. Rana, ActaMetall. 28, 129-141 (1980).

30. M.D. Thouless andA.G. Evans, J. Am. Ceram. Soc. 67, 721-727 (1984).

31. M.D. Drory andA.G. Evans, J. Am. Ceram. Soc. 67, 727-731 (1984).

32. J.E. Marion, A.G. Evans, M.D. Drory and D.R. Clarke, Acta Metall. 31,

1445-1457 (1983).

33. A.G. Evans and W. Blumenthal, in Fracture Mechanics of Ceramics,

Vol. 6, edited by R.C. Bradt, A.G. Evans, D.P.H. Hasselman and F.F.

Lange, Plenum Press, New York, 423-448 (1983).

34. F.C. Monkman and N.J. Grant, Proc . ASTM 56, 593-620 (1956).

35. R. Kassowsky, D.G. Miller, and E.S. Diaz, J. Mater. Sci. 15, 601-610

(1980)

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759-763 (1984).

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38. T. Fett, Res Mechanica, in press.

39. T.-J. Chuang, J. Mater. Sci. 21, 165-175 (1986).

40. T.-J. Chuang, S.M. Wiederhorn, and C.F. Chen, to be published in

Proc. 3rd Intl. Conf. on Creep and Fracture of Engineering Materials

and Structures, Swansea, UK, April 1987.

41. S.M. Wiederhorn, L. Chuck, E.R. Fuller, Jr. and N.J. Tighe, in

Tailoring Multiphase and Composite Ceramics, Materials Science

Research, Vol. 20, edited by R.E. Tressler, G.L. Messing, C.G.

Pantano, and R.E. Newnham, Plenum Press, 755-773 (1986).

42. M.F. Ashby and B.F. Dyson, in Advances in Fracture Research, Vol. 1,

edited by S.R. Valluri, D.M.R. Taplin, D. Rama Rao, J.F. Knott and R. Dubey, Pergamon Press, New York, 3-30 (1984).

41
AN ULTRASONIC STUDY OF CERAMICS FROM THE GREEN STATE

TO FULLY FIRED

A. L. Dragoo

National Bureau of Standards

ABSTRACT

Ultrasonic measurements were investigated as a means for detecting hard

agglomerates in alumina samples during processing from green state to

fully dense. A powder consisting of hard agglomerates was prepared by

calcining a spray-dried alumina powder. Longitudinal and shear wave velocities were measured in green state, bisque fired (1000 °C, 20 h) and

fully dense (1465 °C, 12 h) alumina samples. Ultrasonic signals were coupled into the green state samples by means of an elastomer which

required minimum contact pressure. The sound velocity measurements were compared with measurements of density and pore size distribution. For

the green and bisque state samples significant differences were observed

between samples formed from hard agglomerates and those formed from an uncalcined ("normal") spray- dried powder. Dense state samples obtained

from both types of powder yielded equivalent sound velocities and

densities. Mercury porosimetry measurements indicated that small pores were removed or sealed off in passing from the green to the bisque state,

and that pore size was substantially reduced in passing from the bisque

to the dense state. Green state samples from normal and calcined powders

showed differences in pore size distributions at finer pore sizes, i.e.,

below 20 nra.

42
AN ULTRASONIC STUDY OF CERAMICS FROM THE GREEN STATE

TO FULLY FIRED.

A. L. Dragoo

National Bureau of Standards

Gaithersburg, Maryland 20899

USA

The objective of the present study was to investigate the use of sound

velocity measurements as a means for the determination of the effect of

soft and hard agglomerates on the sintering of alumina compacts. The

effects on the density and elastic moduli were investigated for compacts

at three stages of sintering: green-state or unfired, bisque or partially fired, and fully fired. In addition to ultrasonic

characterization of the samples, the density and porosity of the samples were determined.

EXPERIMENTAL

The procedure for the preparation of samples is illustrated by the

flowchart in Fig. 1. 500 g of 99.99 percent alumina powder, having the

cumulative particle size distribution, shown in Fig. 2, as determined by

x-ray sedigraph, and 3.18 g of MgCNOg )

• (6H2O) were dispersed in 1 L of

water containing small amounts of polyvinyl alcohol (PVA) as a binder and

glycerol as a plasticizer. The slurry was spray-dried to form

agglomerated spheres. A micrograph of the spray-dried powder is shown in

Fig. 3. The spray- dried powder was divided into equal portions, and one portion was calcined at 1000 °C for 12 h to form "hard" agglomerates.

Calcination resulted in partial sintering of the material within the

spherical particles but not between them.

Two sets of compacted samples were prepared. The first set was formed

from the uncalcined powder. These samples are termed "normal samples"

and are designated by an "N." The second set was prepared from the hard

agglomerate powder. These samples are termed "calcined samples" and are designated by a "C." Disks, approximately 3.2 mm thick and 28.6 mm in

diameter, were formed by uniaxial compaction with pressures up to 206

MPa. Each of the two sets of samples were subdivided into three groups

of samples: green-state samples; bisque-fired samples which were sintered

43
at 1000 °C for 20 h; and dense samples which were sintered at 1465 °C for

12 h.

The samples were characterized by ultrasonic measurements, density

measurements and mercury porosimetry. Sound velocity measurements were

carried out on pairs of samples from each of the groups for a total of

12 samples.

Sample densities were determined for the green samples by a dry-weight

dimensional method and for the sintered samples by a liquid immersion

technique. The densities are listed in Tables 1 and 2. Accuracy of the

density values for the green samples was estimated to be better than ±1

percent and for the sintered samples, ±0.1 percent. The density of two

dense samples were also determined with a helium pycnometer.

Mercury intrusion porosimetry was used to measure the pore size

distribution and the bulk and structural densities for the three types of

ceramic samples. The results of those measurements are summarized in

Table 2. The bulk density is defined by the external volume of the

material and includes both open and closed pores; the structural density

includes the material plus only the closed pores in the measured volume.

The three estimates of the mean pore diameter are discussed below.

Ultrasonic Techniques

The ultrasonic techniques used in this study have been described in

detail elsewhere by Jones et al.^ and only will be summarized here.

Through- transmission, using separate transmitting and receiving

transducers, and pulse-echo-overlap techniques were used to make time-off light ultrasonic measurements. The through- transmission configuration

is illustrated schematically in Fig. 4,

The shear-wave transit times were measured by means of two echo-overlap

procedures: through- transmission overlap for green and bisque samples and

pulse -echo overlap for the dense samples. For the green and bisque

44
samples, a novel elastomer material,^ capable of supporting shear waves

at MHz frequencies, was used to dry couple the transducers to the

samples. A single transducer with a viscous coupling agent was used for

the dense samples.

The longitudinal -wave transit times were measured by a throughtransmission technique for the green samples and by pulse -echo -overlap

for bisque and dense samples. Green and bisque samples were dry coupled

to the transducer by covering the surface of the sample with a natural

latex elastomer (surgical glove material) followed by applying silicone

oil between the elastomer and the transducer. This coupling method

yielded a better longitudinal -wave signal than that obtained with the

shear-wave coupling material. Dense samples were measured with a direct

oil coupling.

Mercury Porosimetry

The measurement of pore size by mercury Intrusion assumes cylindrical

pore structures to establish a relationship between the applied pressure

and the pore radius. Pore size distributions are obtained for fraction

of total volume for per size interval and for fraction of total surface

area per size interval. Median pore diameters are calculated for each

distribution. These median values are listed in Table 2 under the

columns labelled "Med. (V)" and "Med. (A)", for volume-based and surface

area-based distributions, respectively. In addition, the measurement

returns an estimate of the mean pore size from the ratio of total volume,

V, to total surface area. A; i.e., (4V/A) . For very narrow pore size

distributions the three estimates of the mean pore diameter are nearly

equivalent. For broad or multimodal distributions the three estimates

are disparate, with the volume distribution yielding the highest

estimate

.

RESULTS AND DISCUSSION

The relationship between velocities to sample density for shear-waves and

^Proprietary material provided to NBS for test and

evaluation by B. Boro Djordjevic, Martin-Marietta Laboratories,

Baltimore, MD.

45
longitudinal waves are shown Fig. 5 and 6, respectively. For both types

of waves, there is nearly an order- of -magnitude increase in sound

velocity from the green to the dense state. The error bars in the

figures are a combination of between sample variations, dimensional

uncertainties, and ultrasonic transit- time errors. The accuracy of the

velocity measurement technique was determined by making reference

measurements on an aluminum alloy and on single -crystal quartz and was found to be bettwe than 1 percent for the dense samples.

As shown in Fig. 5 and 6 and by the elastic moduli in Table 1, the

ultrasonic measurements are clearly different for normal and calcined

samples in the green and bisque states. For the green and bisque states,

the velocities for calcined samples are markedly less than those for

normal samples, whereas dense state samples were similar with respect to

both density and sound velocities. The sound velocities are a function

both of the density of the sample and its processing history. For green

state samples, different amounts of binder in the normal and calcined

samples may be significant. Kupperman and Karplus^ showed that greenstate ceramics of different densities, as a result of different forming

pressures, manifest an increase in velocity with density which is

comparable to that observed here. However, in this work a single forming

pressure was used, with the calcined samples yielding slightly lower

densities than the normal samples.

Elastic moduli and Poisson's ratio were calculated for each sample,

assuming sample isotropy.'^ The results are given in Table 1. The

propagated errors in the calculation of Young's moduli are, for worst

case, ±2, 8 and 3 percent for green, bisque and dense states,

respectively. For the shear moduli, the worst case errors are ±2, 7 and

2 percent; for the bulk moduli, ±3, 13 and 4 percent. Between sample

variation contributed the largest portion of the uncertainty in each

case

.

Sample (elastic) homogeneity and isotropy were measured ultrasonically on

one sample in each sample pair. Longitudinal wave homogeneity was evaluated from measurements at the center and 3 peripheral positions.

46
The standard errors were 2, 1.5 and 0.3 percent for green, bisque and

dense states, respectively. Such an increase is expected since

homogeneity is presumed to improve as sintering proceeds. Sample

isotropy was evaluated by rotating the shear polarization direction at

the center of each sample in increments of 45° from an arbitrary initial

orientation. The anisotropy was within measurement precision for shear waves

.

The porosimetry results in Table 2 were obtained from measurements on a single specimen of each type which were broken into 2 to 4 pieces. The

relative errors in the mean pore diameters are about 20 percent for the

green and bisque samples, with the exception of Med. (A) and (4V/A) values

for the green calcined samples which are about 80 percent, and for the

dense samples, about 40 percent. The uncertainties for the bulk and

structural densities are about 3 percent for all cases. These estimates

of the uncertainties do not include a contribution from sample - to - sample

variation.

The calcined and normal samples differ significantly only in the case of

the green state samples. For these samples the calcined and normal state

samples differ in their porosity, pore size distributions and mean pore

size. The calcined sample showed slightly lower bulk densities,

indicating more porosity. The three estimates of the mean pore size for

this sample are slightly larger than those for the normal sample. The

calcined sample showed a tendency to delaminate near the midline of the

sample. The large uncertainties in Med. (A) and (4V/A) for the calcined

sample are associated with significant variation in the observed pore

size distribution at small pore diameters.

The normal green sample exhibited surface area distributions with strong

modes near 100 and 5 nm; i.e., toward the upper and lower extremes of the

size range. Since the importance of the small pores is reduced in the

volume distribution which emphasized the mode near 100 nm,

Med. (V) » Med. (A) , normal green state

is obtained in Table 2. The normal green state sample yielded more

reproducible pore size distributions than the calcined sample.

47
For the bisque and dense states, the porosity characteristics of calcined

and normal samples became nearly identical. The surface area-based pore

size distributions for both types of bisque fired samples showed a very

strong mode centered near 100 nm and a weak mode centered between 5 and

10 nm. The dense state samples showed a single mode centered near 10 nm.

The porosimetry results indicate that small pores were eliminated or sealed off during processing from the green state to the bisque state.

From the bisque state to the dense state, it is evident that pore size is

reduced. Comparing the porosimetry results and the ultrasonic results,

the sound velocity measurements are expected to be affected primarily by

the largest pores, whereas porosimetry measurements can detect both large

and small open pores. On the other hand, sound velocities are expected

to be more sensitive to the continuity and rigidity of the material

structure

.

From the standpoint of sintering process, from green to bisque state, the

small increases in sound velocities and elastic moduli suggest that the

material structure became more rigid and continuous, probably as a result

of the loss of binder and the formation of necks between particles.

Also, fine pores were eliminated or sealed off. From bisque to dense

state, pore size was markedly reduced and the identity of the hard

agglomerates was lost.

CONCLUSIONS

Shear and longitudinal sound velocities were measured and elastic moduli

determined for green, bisque and dense state alumina samples. Dry

coupling of shear waves into green state compacts was accomplished with

an elastomer which required minimum contact pressure. Ultrasonic

measurements were able to detect the presence of hard agglomerates in

green and bisque state samples but not in the dense state samples. This

results suggests the possible use of ultrasonic measurements to monitor

the presence of hard agglomerates in green state materials.

48
Mercury porosimetry measurements indicated that small pores were eliminated or sealed off during the bisque-firing and that the pore size was substantially reduced for sintering to full density.

Acknowledgements

The author recognizes the major contribution of M. P. Jones who, as a graduate student at Johns Hopkins University, worked in collaboration

with C. R. Robbins, of the Ceramic Powder Characterization Group, and

with G. V. Blessing, of the Ultrasonic Standards Group at NBS , to prepare

many of the samples and to conduct the ultrasonic measurements. The

porosimetry measurements by C . R. Robbins are also recognized. The early

suggestion of this research problem and discussions with W. Coblenz

(present address: Norton Company, Worcester, MA) were important in the

formation of research program. This work was supported by the Office of

Nondestructive Evaluation at the National Bureau of Standards.

49
DIMENSIONAL DENSITY, ELASTIC MODULI AND POISSON'S RATIO

(P. RATIO) OF NORMAL (N) AND CALCINED (C) SPRAY -DRIED ALUMINA

POWDERS

Sample

State

Density (g/cm'^ ) Dimensional Young

Elastic Moduli (GPa)

s Shear Bulk p. Ratio

Green-N 1. 93 2.6 1.7 0 6 -0.23

Green-C i . Q C OO 0.6 1 .

1

u i -0.72

Bisque-N 2. 12 29.0 13.0 11 0 0.07

Blsque-C 2. 00 11.0 6.6 2 0 -0.14

Dense -N 3. 91 400.0 162.0 248 0 0.23

Dense -

C

1465

3. 92 405.0 164.0 256 0 0.23

50
SAMPLE DENSITIES AND PORE DIAMETERS FOR SPECIMENS FORMED FROM NORMAL (N)

AND CALCINED (C) SPRAY-DRIED ALUMINA POWDERS

Sample

State Dimens

.

Density

Bulk

(g/cm^)

Struc

.

He Pycn.

Pore

Med. (V)

Diameter (urn)

Med. (A) (4V/A)

Green-N 1 .93 1.87 3.57 0.075 0.018 0.039

GreenC1,88 1.82 3.44 0.101 0.062 0.070

Bisque -N 2. 12 2.04 3.71 0.106 0.091 0.069

Bisque -

C

2.00 1.92 3.66 0.108 0.092 0.071

Dense -N 3.91 3.74 3.80 0.008 0.006 0.008

Dense -

C

1465

1500

3.92 3.72 3.80 3

3 .952.931

0.008 0.005 0.008

Dimens. = density determined from dimensional measurements on weighted specimens.

Bulk - bulk density obtained from Hg intrusion porosimetry.

Struc. — structural density obtained from Hg intrusion porosimetry.

He Pycn. = density obtained using He pycnometer.

Med. (V) = median pore diameter based on distribution of pore voliime.

Med. (A) = median pore diameter based on distribution of pore area.

(4V/A) = average pore diameter calculated from total pore volume (V) and

total pore area (A)

.

51
1. Jones, M.P.; Blessing, G.V. ; Robbins, C. R. Dry-Coupled Ultrasonic

Elasticity Measurements of Sintered Ceramics and Their Green States.

Mater. Eval . 44(7): 859-62; 1985.

2. Kupperman, D. S.; Karplus, H. B. Ultrasonic Wave Propagation

Characteristics of Green Ceramics. Am. Ceram. Soc. Bull. 63(12): 1505-

09; 1984.

3. See, e.g., Kolsky, H. Stress Waves in Solids. New York: Dover

Publications; 1963.

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Figure 1- -Flowchart for the preparation of alumina samples containing

soft and hard agglomerates

.

53
Figure 2- -Cumulative particle size distribution for alumina starting

powder as obtained with x-ray sedigraph. Cumulative percent finer is

plotted versus the Equivalent Spherical Diameter (ESD) , in /xm, calculated

from Stokes Law for settling of spheres.

54
SPRAY-DRIED PARTICLES

Figure 3- -Micrograph of spray- dried alumina powder.

55
o

o ^

to

• D. > e u <H a o o O 0) o X

00 0) a

n

a

Ba

CO

I

Figure 4- -Block diagram of apparatus used in the through- transmission

pulse -echo technique.

56
6

6

>-UO

_J

u

>

q:

CE

u

X

6

5

3

2

I

0

I N

BISQUE

GREEN.- N

C

DENSE I N,C

0

DENSITY (g/cm^)

Figure 5 - -Ultrasonic shear wave velocity vs. sample density for the

green, bisque and dense states, with both normal (N) and calcined (C)

processed spray-dried powders.

57
\

E

6

UO

_J

u

>

z

o

1

1

10

9

8

7

6

5

4

3

2

10

DENSE I N,C

BISQUE

GREEN

I N

N

0

DENSITY (g/cm3)

Figure 6 - -Ultrasonic longitudinal wave velocity vs. sample density for

the green, bisque and dense states with both N and C type powders.

58
NONDESTRUCTIVE CHARACTERIZATION

OF CERAMIC MATERIALS: THERMAL WAVES

G. S. White

Ceramics Division

Institute for Materials Science

and Engineering

National Bureau of Standards

Gaithersburg, MD 20899

and

E. H. LeGal LaSalle*

Laboratoire d'Optique Physique

ESPCI

10 Rue Vauquelin

75005 Paris, France

ABSTRACT :

Investigation of thermal wave techniques for application to the

nondestructive evaluation (NDE) of ceramic materials is in its early stages.

Thermal waves have the advantage of having an adjustable thermal diffusion

length, which determines the probe length, and of being a surface and near

surface probe. Research at NBS has shown this technique to be sensitive to

some forms of surface polishing damage in vitreous silica and to variations in

porosity of alumina.

INTRODUCTION :

Although many techniques for nondestructive evaluation (NDE) have been

developed, and are in use, for testing metals, the corresponding progress for

ceramics has been limited. The principal reason behind this lack of progress

in ceramic materials is the brittle nature of these materials; because of

their low fracture toughness (on the order of 1 to 10 MPam^ ^ ^ compared to -100

59
MPam^''^ for metals), catastrophic flaws are measured in tens of micrometers

rather than in millimeters, as for metals. In addition, defects which result

in failure of ceramics usually lie on the surface. Therefore, NDE techniques

which measure bulk properties are often inappropriate for finding defects in

ceramics. In this paper, application of thermal wave techniques to the NDE of

ceramic materials is discussed and related to the concerns listed above. Two

examples relating to recent NBS research are discussed.

THERMAL WAVE PROCEDURE :

The concept of thermal waves as a material probe has developed over the

past few years-"^ with contributions from several research groups. The

following picture has been reached. When a periodic heating source is applied

to a material, the resulting periodic heat flow can be described as highly

damped travelling waves.

S = Aexp(ja>t-(l+j)x/M) , (1)

where S is the thermal wave signal, x is the distance traveled, and n =

(2K/Cpw)^^^ is the thermal diffusion length which is a function of the thermal

conductivity, K, the heat capacity, C, the density of the specimen, p, and the

angular modulation frequency of the heat source, u>. As the strong damping

component in eqn. 1 indicates, the effective depth to which such waves travel

in the material depends upon the modulation frequency of the heating source;

60
the lower the frequency, the deeper the penetration. For most applications,

the practical penetration range of the thermal waves can be considered ~fi.

A number of detection schemes exist which measure certain aspects of this

heat flow, and two of them, the photoacous tic effect and the mirage effect,

will be qualitatively described here. The photoacoustic effect signal (Fig.

1) results from thermally induced pressure variations in the atmosphere above

the heated specimen which has been placed in a sealed chamber. The pressure

variations drive a microphone and the resulting signal is proportional to the

temperature of the specimen^ . The mirage effect (Fig. 2) , takes advantage of

the modulation of the index of refraction gradient above or in the heated

region of the specimen to deflect a probe laser beam and to produce a signal

proportional to the gradient of the temperature'^ . Generally, the probe beam

is passed through the heated atmosphere above the specimen surface without

touching the specimen itself. In both of these detection techniques, the

heating can be caused by any modulated source; e.g. laser, electron-beam,

flashlamp , etc.

Because the thermal behavior can be described as a damped travelling wave

phenomenon, the signal can be divided into two independent components:

magnitude and phase. The magnitude is a measure of the strength of the signal

and, therefore, reflects optical as well as thermal properties of the

specimen. The phase component is a measure of the delay between the time heat

is applied and the signal is detected and is much less sensitive to optical

features'" .

As eqn. 1 implies, the interrogation range of the thermal waves depends not

only upon the modulation frequency of the heat source, but also upon thermal

parameters of the specimen under study. Table I compares n at two different

frequencies for three commonly used ceramic materials: vitreous silica.

61
alumina, and silicon carbide. There is an order of magnitude change in when

the specimen is changed from vitreous silica to silicon carbide, and this

change is independent of any experimental parameters. In principle, the

values of /i can be maintained from material to material by variations in w

although, in practice, the frequency capabilities of the modulator may not

have the necessary range.

Although the understanding of thermal waves has greatly improved, only

recently have there been systematic attempts to use this understanding in the

NDE of ceramic materials. Because thermal waves are essentially a surface and

near surface effect and because the thermal diffusion length can be modified

over a large range by changing the heating modulation frequency, thermal waves

seem naturally suited to this application.

THERMAL WAVE APPLICATIONS:

Residual Polishing Damage : There is concern over the possible creation of

residual damage resulting from surface preparation of optical components.

Because thermal waves have proven useful in the interrogation of metallic and

semiconductor layered specimens-'^ • ^

, it is a reasonable step to use them to

look for near surface residual damage in glass. The possible types of

residual damage envisioned include scratches, subsurface microcracks , and

residual stresses. Since the penetration depth of the thermal waves depends

upon the modulation frequency of the heat source, the presence of a damaged

layer with thermal properties different from those in the bulk would result in

a change in the frequency dependence of the signal as the thermal waves cross

the boundary between the damaged and undamaged material. The size of this

change would depend upon the magnitude of the change in (K/pC)^^^. Therefore,

62
one of the most direct applications of thermal wave techniques to the

detection of a residual damage layer would be to measure the frequency

dependence of the thermal wave signal.

Specimens, 1.2 x 1.5 x 0.21 mm, were cut from polished vitreous silica

slides. The specimens were abraded with a 180 grit diamond wheel and then

polished in steps using 30, 15, and 1 yum diamond polishing compound and .3 jim

alumina powder. At each step, specimens were set aside for thermal wave

investigations.

The experiments consisted of measuring the photoacoustic or mirage signals

at 5 random spots on the specimen surfaces as a function of heating modulation

frequency, using both photoacoustic and mirage detection techniques. The

frequency range was limited by the mechanical chopper to 20 Hz < w < 250 Hz.

The signal from both the photoacoustic cell and the mirage system beam

deflection detector was filtered in a vector lock- in amplifier. Both phase

and magnitude were measured at each of the positions on the specimen surface.

Measurements were made following each of the polishing steps described

above. While the actual signal magnitudes varied as a result of the polishing

procedure, the results described below, obtained from the specimens which

underwent final polishing, are representative of the general trends observed

at each polishing step.

In Figure 3a, the signal magnitude is plotted as a function of modulation

frequency using the mirage effect detection technique. In Figure 3b, the

corresponding phase data are displayed. In both cases, the data are plotted

versus the square root of the frequency. The frequency range covered

corresponds to thermal diffusion lengths, and, hence, to penetration depths,

between 25 and 100 micrometers. The line through the data points in Fig. 3a

is an empirical exponential fit and shows that no discontinuity, corresponding

63
to a damage layer, is observed in the magnitude signal. The phase data, shown

in Fig. 3b, do show a discontinuous slope at about f = 30 Hz, corresponding to

a penetration depth of about 80 /xm. However, because 30 Hz was at the low end

of the frequency capabilities of the apparatus, this result is most likely an

artifact; see below.

The photoacoustic phase data, Fig. 4b, covering the same frequency range,

give no indication of the discontinuity. Similarly, the photoacoustic

magnitude data (Fig. 4a), which theory predicts should be a linear function of

1/f for isotropic thermal insulators^ , does not suggest the presence of a

damage layer. A straight line fit connects all the data for both the phase

and magnitude measurements. Therefore, the preponderance of the thermal wave

data has failed to give any indication of a damage layer in the depth range of

25 to 100 fj,m.

However, while no damage layer has yet been observed, photoacoustic scans

have shown the presence of scratches remaining from the surface treatments.

Magnitude and phase line scans (f = 200 Hz, penetration depth - 30 /Ltm) across

the surface of an as -received Si02 specimen (Figures 5a and 5b) give no

indication of any defects. After grinding and polishing, a randomly

positioned set of photoacoustic magnitude scans mapped out the remaining

scratch shown in Figure 6. The lines in Figure 6 are spaced 50 /zm apart,

indicating that the scratch is a large defect and should easily be observed by

optical scattering NDE tests although a casual inspection by eye did not

detect the defect prior to the photoacoustic measurements. Subsequent

investigation of the polished surfaces under an optical microscope did detect

a number of residual scratches. The large signal to noise ratio in the figure

suggests that the photoacoustic effect has the sensitivity to detect much

smaller flaws in a systematic surface scan of these specimens.

64
Porosity ; The porosity of ceramic components has long been a quality control

variable. Because /i is a function of the effective K as well as of p, both of

which are functions of the specimen porosity, thermal waves can be expected to

be sensitive to variations in porosity. Therefore, a set of alumina specimens

of varying porosity has been made and thermal wave phase and magnitude

photoacoustic measurements were made as a function of frequency.

The specimens were uniaxially pressed to 103 MPa and fired in air. The

firing procedure consisted of placing the specimens in the furnace at 800 C,

increasing the temperature to 1600 C in approximately an hour (see Fig, 7)

,

following which the specimens were withdrawn after various lengths of time.

The specimens were then sectioned on a diamond saw and the cross - sections were

potted together in epoxy for polishing, optical measurements, and

photoacoustic measurements. The polishing procedure was the same as that

outlined above for the surface damage tests. Because each of the specimens

experienced the same surface preparation, any surface roughness effects in the

thermal wave measurements are expected to be related only to variations in

porosity. In addition, because no binder was used, no surface absorption

effects related to partial binder burnout in the different specimens will

occur.

Before the photoacoustic measurements were made, the percentage porosity of

the specimens was measured using standard pore counting techniques and is

expressed in percentage of the surface area taken up by pores

.

Figure 8 shows magnitude measurements for the specimen removed from the

furnace as soon as 1600 C was reached. Measurements were made at 5 positions

on the specimen and the vertical lines show the resulting standard deviation.

For this specimen, the porosity was measured as 30.8 ± 5.4%. Figure 9 shows

65
the same type of data for a specimen which was held at 1600 C for 118 hr. For

this specimen, the porosity was 3.0 ± 2.6%. There are marked differences

between the two sets of figures, suggesting that thermal waves do detect the

difference in porosity between the two samples. Analytical interpretation of

these signals is difficult because of the influence of unknown surface

roughness, optical absorption, and light to heat conversion efficiencies,

especially as these parameters vary with porosity changes. However,

variations in these parameters are expected to be directly related to

variations in porosity, since other parameters have been held constant.

Therefore, the thermal wave results suggest that, at the least, the thermal

wave signal could be empirically calibrated to give porosity values for a

given material and processing treatment. Alteriiatively , thermal waves could

be used to calculate thermal diffusivities as a function of porosity. Some

work in this direction has already begun on other materials^

.

SUMMARY :

The development of thermal wave techniques as NDE tools has reached the

level at which practical tests need to be considered. One of the most useful

applications of thermal waves would be to the NDE of ceramics, materials for

which there are currently very few NDE tools.

Two applications of thermal wave NDE of ceramics have been discussed.

Feasibility of using photoacoustic and mirage thermal wave techniques to

detect both damage resulting from preparation of optical surfaces in vitreous

silica and porosity in alumina has been investigated using photoacoustic

frequency dependence measurements. While residual surface scratches were

shown to be detected easily, no indication was found of a subsurface residual

66
damage layer in silica. On the other hand, thermal wave signals were found to

be sensitive to porosity changes in alumina.

ACKNOWLEDGEMENTS :

I gratefully acknowledge the support given this work by the Office of

Nondestructive Evaluation at N.B.S.

67
Table I:

Material Frequency (Hz) u (urn)

Si02 20 93

2000 9

AI2O3 20 285

2000 29

SiC 20 1000

2000 100

68
REFERENCES

1. G. Birnbaiim and G.S. White, in Nondestructive Testing. Vol. 7 (R.S. Sharpe

,

ed.), Academic Press, Inc. London, pp. 259-365 (1984).

2. A. Rosencwaig and A. Gersho, J. Appl. Phys . 47, 64-69 (1976).

3. L. Aamodt and J. Murphy, J. Appl. Phys. 52, 4903-4914 (1981).

4. A. Rosencwaig and G. Busse, Appl. Phys. Lett. 36 (9), 725-727 (1980).

5. A. Rosencwaig, J. Opsal and D.L. Willenborg, J. de Phys. 44 (10), C6-483-

C6-489 (1983).

6. B.K. Bein, S. Krueger and J. Pelzl, Can. J. Phys. 64, 1208-1216 (1985).

* Work done while a visiting scientist at N.B.S.

69
C02 BEAM

WINDOW (ZnSe)

SAMPLE

TO LOCK-IN

MICROPHONE

1. Schematic diagram of a photoacoustic cell. A modulated heating beam from

a CO2 laser passes through the ZnSe window into the sealed chamber where it is

incident upon the sample. The heated surface causes a pressure variation in

the air above the sample which is detected by the microphone.

70
HEAT SOURCE

POSITION

SENSITIVE

DETECTOR

2. Schematic diagram of the mirage effect. A modulated source heats the

sample surface. The surface then reheats the atmosphere, causing an index of

refraction (n) gradient above the heated sample surface. A HeNe probe beam,

passing through the heated atmosphere is deflected at the same frequency as

the sample is heated. The resulting deflection, <I>, is detected by a position

sensitive detector.

71
400.0 -

4^

300.0

200.0

'2 100.0 -

0.0 I I I I I I I I I I

I I I I I I I I I

I I I I I I I I I II I I I 0.00 5.00 10.00 15.00

I I I I I

{ 20.00

sqr. rt. f

-50.0

CO

a; -60.0 -

'-70.0 -

a. -BO.O -

90.0 I I I I I I I I I I

I I I I I I I I I

I I I I

I

I I I I I I I 1 I

I 0.00 5.00 10.00 15.00 20.00

sqr. rt. f

3. a) Magnitude signal of the mirage effect for a deflection of the probe

beam normal to the sample surface. The sample has been ground and then

polished.

b) Corresponding phase data for the same sample position and time as

shown in Figure 3a.

72
600.0

C/3

. 400.0

a;

200.0 -

0.0 I I I I I I I I I I I

I I I 1 I I I I !

I

I I M I M I I I

I I I I I I I I I

I M I I I I I I I

I

I I I I I I I I I

I 0.00 0.01 0.02 0.03 0.04 0.05 0.05

1 /frequency

•50.00 n b

-60.00

a;

0^-70.00

CO

-80.00

-90.00

0

"1 TTT I I n I I

I r~r~[—i—r~i—n—

|—n—r~i—i—n—r—

i—|—n—r~r-i—n

r~r~\ 5 10 15 20

Sqr. Rt. f

4. a) Photoacoustic magnitude data for a specimen which has been ground and

polished.

b) Photoacoustic phase data for a specimen which has been ground and

polished.

73
800.0 n

600.0 I I I I

I I I I M I I I

I

i I I I I i I I I I

I I I I I I I I I

I I I I

I I I II 0 2000 4000 6000 8000 10000

Position (microns)

-70.00

q b

^-72.00 -\ m

0)

U

tiO-74.00

0)

0)

PL.

76.00 -

-78.00 -

-80.00 I I I I I M I I

I M I I I I I I !

I

I I I I I I I I I

I

I I I I I I I I I

I

I I I I I I I T-r-| 2000 4000 6000 8000 10000

Position (microns)

5. a) Photacoustic magnitude scan across the surface of an as-reccived

sample. The lack of structure indicates no detectable surface damage.

b) Photoacoustic phase scan across the same sample surface as in Figure

5a. Again, no indication of damage is observed.

74
6. Photoacoustic image of sample after grinding and polishing, showing the

presence of a remaining scratch.

75
1800 n

Time (min.)

7. Furnace temperature plotted as a fmction of time cfciring the firing

treatment of the alumina specimens.

76
1000 n

800-

0 I

I I I I I I I I I

I

I I I I I I I I I II I I I I I I I I

I

I I I I I I I I I

I

I I I I I I I I I

I

I I I I I I I M I 0.00 0.01 0.02 0.03 0.04 0.05 0.06

l/f (l/Hz)

8. Photoaccxtstic ma^itude measurements of the alumina specimen with 30.87.

porosity as a fLnction of heating frec^jency.

77
1200

•rH

800-

0)

^ 400

tie

cd

0

0.

I I M I I I I I

I

I I I I I I I I I

I

I I I I I I I I I I

I I I I I I I I I

I

I I I I I I I I I II I I I I I I I I I 00 0.01 0.02 0.03 0.04 0.05 0.06

l/f (l/Hz)

9. Photoacoustic ma^itucte measuremmtB of alumina specimEn with 351 pctf-cjBity

as a function of heating frequency.

78
Computer-aided ultrasonic testing of non-oxide ceramics

A. Hecht, E. Neumann, P. Rose

Introduc tion

Non-oxide ceramics are mainly used at high- temperatures or for

wear- resi Stan t parts. Typical materials are silicon-carbide

(SiC) and silicon-nitride (Si3N4). Due to the brittleness of

ceramics, even very tiny defects (<50 urn) may cause failure.

Consequently high-frequency ultrasonic testing (UT ) should be

practiced for the detection of defects which are smaller than

e. g. 50 ym. First attempts have been made by developing a 120 MHz-UT-equipment [l ] Since the transmission-factor of the

couplant/material- interface becomes very frequency dependent

at such high frequencies, problems might occur to get an ultrasonic impulse passing through the interface if the surface is

not polished [2]. Also the material's attenuation has to be

considered: Usually, the attenuation of sintered ceramics is

very low even at frequencies above 100 MHz, but can reach very

high values for cast-ceramics at lower frequencies (20 MHz).

This paper presents results of 25 MHz and 50 MHz ultrasonic

testing of reac tion- bonded sil icon- infiltrated silicon-carbide

(SiSiC) containing approx. 12% free silicon, and of hotpressed silicon-nitride (HPSN). Some of the UT results are

compared with low-energy radiographic images.

Ultrasonic test equip-aent "A3AKUS"

All devices of "A3AKUS" , which was developed as a multipurpose UT-equipment [3-5], are displayed in Figure 1. These

devices are commercially available and may be combined in

various arrangements to i--_alize the different UT-technique s

,

e. g. B-scan, C-scan, etc. [s].

79
UT of ceramic materials was carried out first by using a 25 MHz

focussing transducer in immersion technique with the specimen completely under water. Figure 2 shows typical echo

traces of UT in immersion technique, where a front surface echo

and a backwall echo can be distinguished. To detect small

defects, the amplification has to be increased. If the

transducer is placed directly over a flaw, an additional

echo occurs. Its amplitude can be measured by the ultrasonic analyzer's peak level detector after having gated out a distinct depth-zone of the specimen. The start point of

this depth-zone is influenced by the length (duration) of

the front surface echo. This length is approximately 4 mm in

Figure 2, if the materials thickness is 10 mm. Tne latest

stop point of the evaluable depth-zone is fixed directly

by the beginning of the backwall echo.

An important criterion for the selection of an appropriate

focussing transducer is the focal distance (distance between

transducers front surface and focal point on the acoustical

axis). As the ratio between the sound velocities of ceramics

and couplant water is approx. seven, the focal distance in

ceramics is approx. a seventh of its value in water. Consequently, transducers with a long focal distance should be

selected for the inspection of thicker parts, if available.

All measurements presented in this paper are UT C-scans,

where the specimen 'is scanned and the amplitude A measured

within a fixed dep^h-zone. Amplitude resolution is 8 bit.

The data are then displayed as A(x,y), where x,y are spatial coordinates. Figure 3 shovs the exper -^ntal set-up for

a C-scan recording by "ABAKUS" . An additional characteristic of this S2t-up is a spectrum analyzer used as a bandpass filter with variable bandwidth and center-frequency.

This is a measure to suppress the electrical noise and consequently increase the flaw detectabil ity . The bandpassfiltered signals are then recorded by a transient- recorder and

the maximum amplitude evaluated by fast software.

80
computer-aided equipment is a useful tool for detecting, monitoring and evaluating defects. Digital image analysis should

be an additional means to compare results of ultrasonic and

radiographic testing. Detected defects can be evaluated additionelly by use of focussing transducers of higher frequency

and taking microscopic C-scans for flaw sizing as well as

regarding carefully the A-scans at the defect positions. This

might be a suitable method for flaw-type characterization.

Acknowledgement

This work was carried out at the German Federal Institute for

Materials Testing ( Bundesanstal t fur Ma ter ialpriif ung (3AM)),

Berlin. The destrurtive preparation of the HPSN-specimen was

carried out at Elektroschmelzwerk Kempten (ESK). The authors

thank all persons who contributed to this report.

Authors

A. Hecht is with BASF Aktiengesellschaft , ZEW/BFZ-L443

,

6700 Ludw: g shaf en , German Federal Republic. E. Neumann and

P. Rose are with Bundesanstal t fiir Ma ter ialpriif ung (BAiM) , Fachgruppe 6.2, Unter den Eichen 87, -1000 Berlin 45, German

Federal Republic.

81
the ultrasonic transducers sound field. In some cases the

coincidence is good, in others not so good. This is most

likely caused by the different depth-sensitivity of the

focussing ultrasonic transducer.

Further examinations were made by taking microscopic ultrasonic C-scans of interesting regions with a 50 MHz focussing

transducer. Figure 11 shows the ultrasonic image of an inclusion 1 mm below the surface that was detectable by both UT

and RT. As the defect appeared as a white spot on the radiographic image, one could be certain that it was an inclusion.

Figure 12 shows the ultrasonic A-scan at the defect position of

maximum amplitude. No single echo. occurs, but a long sequence

of various echoes caused by circumferential and other types of

waves. This is in accordance with [9], where the echo traces

resulting from spherical inclusions could be predicted.

Figure 13 shows the A-scan of a flaw 0.5 mm below the

surface, that was detectable only by UT. Only one distinct

flaw echo occurs. Consequently this could be no spherical

shaped defect. The ultrasonic footprint- shaped image of this

flaw is shown on the left side of Figure 14. Its planar size

was estimated with 0 . 8 x 0 . 35 mm by a 6 dB-drop method (this

method can be applied because of the very small focussing

diameter). A micrograph oZ the defect (see Figure 14, right)

was obtained after grinding 0.5 mn-. of the specimen's

thickness. The defect was a flat shaped region of porosity

with the dimensions 0.75 x 0.36 x 0.1 mm and was not detectable by RT because of the very small size in depth direction.

Concl usion

As most of the defects in the inspected specimens were much

bigger than 50 um, satisfactory results could be obtained even

at a test:- frequency as low as 25 MHz. However further basic

work is necessary to cope with all the problems arising. A

82
area on the right side is caused by a Si-vein, which is

only a few pm thick but of large extension. All these defects are displayed in Figure 7, where "A" is an optical

image and "B" and "C" are SEM-images. SiC appears dark and the

free silicon white.

Figure 8 shows a photograph of that cross- section of the

SiSiC-plate where the Si-vein is located (see top view,

Figures 4 - 6). In this side-view the Si-vein appears as a thin white line. In addition two areas of different density

are visible, caused by different Si-concentrations in the

SiSiC microstruc ture . Such areas can be detected nondestruc tively by computer X-ray tomography [8]. Figure 8

points out, that only those parts of the Si-vein are detectable by UT in normal incidence, which are located perpendicular to the incident direction. Additional UT in obliqi:e

incidence would be necessary to detect the complete Si-veln

area in the examined specimen.

Ultrasonic testing of HPSN

The inspected specimen was a 9 mm thick crescent- shaped

plate, the length of the secant line is 110 mm. Again the specimen was inspected twice by taking one C-scan from each side.

Figure 9 shows the superposition of these two C-scans as well

as a photograph taken from a radiographic image (3 0 kV tubevoltage). Only few white spots caused by absorptive incl'isions

are visible, but the original radiographic image displays many

more flaw indications.

In order to compare the ultrasonic and radiographic results

only the small marked ares in Figure 9 are considered. These

areas are shown magnified in Figure 10. The flaw images from

the UT C-scan above are also projected as lines into the

radiographic image below. UT inc c ations are of a bigger size

than the indications from radiographic testing (RT) due to

83
All the measured C-scans A(x,y) are stored in the computer/

where mathematical transformations or other signal analyses

may be carried out after the measurements have been taken [s].

The data can also be transferred into a digital imageprocessing system, which is normally used for X-ray film

evaluation [s] in the BAM. This image- processing system

allows to combine, transpose, or superimpose different UT

C-scans, and even superimpose UT C-scans with radiographic

images. It also offers different measures of image enhancement. All figures containing C-scans in the next chapter

are photographs from the image- processing system screen.

Ultrasonic testing of SiSiC

Properties of SiSiC are characterized in [?]. The inspected

specimen was a 10 mm thick plate containing several natural

defects. The C-scan in Figure 4 displays all the flaw indications (echoes, see Figure 2 ) in a depth zone of 4 - 10 mm.

(This limited range results from the "dead zone" of the front

surface echo in Figure 2). The echo amplitude A is represented

by 255 grey-scale values in accordance with the 8 bit resolution

of the used transient- recorder

.

If the specimen is turned around (and the coordinate system too), the depth zone of 0 - 6 mm can also be inspected

(see Figure 5). Both Figures 4 and 5 look very similar, as

most of the flaws are located in the overlapping depth zone

of 4 - 6 mm. Figure 6 is the superposition of Figures 4

und 5, and displays all detected flaws in the complete

volume of the specimen.

Destructive follow-ups showed, that all the spherical

shaped indications are caused by voids or local Si-enrichments, which cannot be distinguished yet. The straight line

in the middle of Figure 6 is the ultrasonic image of a "worm-hole" , accompanied by Si-enrichments. The larger

84
Abstract

Plates of different ceramic materials have been inspected with

a computer aided ultrasonic test equipment at frequencies of

25 and 50 MHz. The measured ultrasonic C-scans were compared

with radiographic images by transferring all data into a digital image processing system. The detectabil ity of different

types of defects is discussed. Flaw sizes can be estimated by

taking microscopic C-scans with focussing high-frequency

transducers, and flaw type characterization can be carried out

by recording particular features of the ultrasonic A-scan.

85
Re f erence s

[l] Goebbels, K. , Reiter, H. , Hirsekorn, S., Arnold, W. : "Non-Destructive Testing of High-Temperature and HighStrength Ceramics". Proceedings of the 12^^ Int. Conference in Science of Ceramics, Saint-Vincent, Italy,

June 27 - 30, 1983, pp. 483-494

[2] Generazio, E. R. : "The Role of the Reflection Coefficient

in Precision Measurement of Ultrasonic Attenuation".

Materials Evaluation 43 (1985) 8, pp. 995-1004

[3] Hecht, A., Thiel, R. , VoB , U. : "A Computer Controlled

Equipment for Different Ultrasonic Measurement

Techniques". Proceedings of the 3

^'^ European Conference

on NDT, Florence, October 15 - 18, 1984, pp. 145-155

[4] Neumann, E., Hecht, A., Schnitger, D. , V06 , U.

,

Gieschler, W. : "Computer Aided Ultrasonic Testing of

Fibre Reinforced Plastics". Proceedings of the 3^^

European Conference on NDT, Florence, October 15-18,

1984, pp. 35-47

[5] Giittner, F. , Limpert, R. , Steinhoff, U. : Ul traschallpriif ung von Hochleistungsverbundwerkstof f en" . Zeitschrift fur Werkstofftechnik 17 (1986)4, S. 135-147

[6] Heidt, H. , Rose, P., Raabe, ?. , Daum, W. : "Digital Image

Processing for Radiography in Nuclear Power Plants" . Proceedings of the 7^'"^ Int. Conference on NDE in the

Nuclear Industry, Grenoble (France), 1985, pp. 317-319

[7] Willmann, G. , Heider, W. : " Sil iziuraca_bid (SiSiC), ein

neuer Werkstoff fiir den Apparatebau" . Zeitschrift fiir

Werkstofftechnik 14 (1983), S. 135-140

86
Goebbels, J., Heidt, H. , Kettschau, A. ^ Reimers, P.:

"Computer X-ray Tomography of Ceramic Materials".

Proceedings of the 2"*^ Int. Symposium on Ceramic Mate

rials and Components for Engines, Lubeck, April 14-17

1986

Sachse, W. : "The Scattering of Elastic Pulses and the

Nondestructive Evaluation of Materials" . Materials Evaluation 35 (1977) 10, pp. 83-106

87
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Block <3iagram of the used multi-purpose ultrasonic

testing equipment "ABAKUS"

88
Backwall

Echo

®-

Materials _ Thickness'

Ultrasonic

Echo Trace

30 dB

+ Flaw Echo

Gated

Depth -Zone

@: Evaluable

Time Interval

Fig. 2: Ultrasonic testing of ceramic plates; typical echo

traces (A-scan)

89
ULTRASONIC ANALYZER

T/R

0

; Gated OutiOutput Peak

©

OSCILLOSCOPE

A ® B ®

SPECTRUM ANALYZER

Working as a bandpassf) iter ullh

variable center frequency and bandtflilh 0 Out 0

> To

DIGIT. STORAGE OSCIL. "TRANSlEifiraRDER"'

A^'

delayed I 0 In

Computer

To CAMAC or IEEE-488 IB

Sped men

Fig. 3: Experimental set-up for the recording of C-scans by

" ABAKUS"
Ultrasonic C-scan of SiSiC-plate (100 x 100 x 100 mm)

with 25 MHZ focussing transducer; depth zone 4 - 10 mm,

resolution 0.25 mm

5: Ultrasonic C-scan of SiSiC-plate (100 x 100 x 100 mm)

with 25 MHz focussing transducer; depth zone 0-6 mm,

resolution 0.5 mm; specimen turned around in Lmmersion tank, data A(x,y) transposed

91
Fig. 6: Superposition of both C-scans of Figures 4 and 5;

resultant depth zone 0 - 10 mm

Single Pore and

Si - Ennchment

Lcxal

Si - Enrichment

500 \im 200 Jim—

Fig. 7: Different types of flaws in SiSiC detectable with

25 MHz focussing transducer

92
ULTRASONIC TESTING A - SCANS

8: Detection of a Si-vein in SiSiC by ultrasonic testing

with 25 MHz focussing transducer

SUPERIMPOSED UT C-SCflNS^ 25: MHZ:

X-RflV IMflGE^ 30 KU

9: Comparison of different NDT techniques at HPSN-spec imen

above: superimposed C-scans

below: radiographic image, 30 kV

note: Copying of radiographic images leads to much

worse image quality

93
Fig. 10: Section of interest (taken from Figure 9)

above: ultrasonic C-scan

middle: radiographic image after image-processing

(detected inclusions as dark spots)

below: radiographic indications encircled by larger

ultrasonic indications from C-scan above

94
Fig. 11: Microscopic ultrasonic C-scan (2.5 x 2.5 mm) of defect

detectable by UT and RT; 50 MHz focussing transducer,

resolution 10 um

Fig. 12: Echo trace (1 usee full vertical scale) of larger

indication of Figure 11

above: only front surface echo

below: additional sequence of flaw echoes, caused by

circumferential waves (see [9])

95
Fig. 13: Echo trace (1 usee full vertical scale) of defect

detectable only by UT

above: only front surface echo

below: additional single flaw echo

Fig, 14: Flat inclusion detectable only by UT

left: Microscopic ultrasonic C-scan (2.5 x 2.5 mm)

with 50 MHz focussing transducer; resolution 10 ym

right: unetched micrograph (mirror image) after

grinding and polishing

96
Computer tomography

J. Goebbels, H. Heidt, A. Kettschau, P. Reimers

Introduction

Computed X-ray tomography (CT) offers the great advantage to

produce maps of the local X-ray absorption inside an object. In

comparison to that conventional radiography gives only a projection image which means that the interpretation of such images

is often complicated or impossible especially at complex object

geometries. Together with the good density resolution this was

the reason for the great success of CT in medicine since about

15 ,2ars. In this paper an insight will be given into the possibilities of CT to localize and characterize defects in cercnic

components .

Development of BAM-Computertomogr aph

In general the medical tomographs are not suited to NDT applications. Therefore at BAM a group started with the development

of a tomograph about 1978. After preliminary investigations with

an experimental scanner, a very flexible machine was developped

and installed in 1984.

Constructed for a wide range of industrial objects the tomograph

cannot represent the most favourable solution for NDT of ceramics but the images made by this system can give the information

needed to construct a CT-scanner for this special application.

The BAM-scanner has the following specifications:

- Mechanical system: 3 axis scanner table

Reproducibility and accuracy of the linear axis: 30 pm

of the angle position: 0,01 ° Maximum diameter of object: 1 m Maximum weight: 1000 kg

- Detector system: 32 detectors ( Ph o t omu 1 t i p 1 i e r s with

plastic scintillators)

Collimator: 1.2 mm width, 2 to 8 mm heigth

Single collimator: 0.2 mm width, 2 mm heigth

97
- Radiation sources: X-ray, 420 kV

60-Co , 1.33 MeV (12 TBq

)

Experimentals

The quality of a tomograph and of the images taken with it is

mainly determined by three parameters:

- spatial resolution

- density or contrast resolution

- artefacts or nonlinear ities which depend on object geometry.

A treatment of this group of effects is out of the scope of

this paper

.

The first two images are examples for the good density resolution of the CT-method.

Fig. 1 shows the hot pressed Silicon Nitride plate described in

the paper of Hecht et al. The density resolution is 1.2 °a . At

the outer range of the plate a region with a low_r density (the

width is about 3.5 mm) is clearly detected.

Fig. 2 shows four tomograms of a Si impregnated Silicon plate.

Two regions with different density can be distinguished. The

histogram exhibits two peaks which correspond to a density difference of 2.7 %

.

Whereas density variations can be analysed quantitatively such

an analysis is difficult for defects with dimensions smaller

than the spatial resolution. It depends on the focal spot size,

the size of collimators and the distance between source, object

table and detector system, or simply the magnification of the

system.

Fig. 3 and 4 show the possibility of this method to detect

defects even smaller than the spatial resolution and to disting

huish between voids and inclusions.

98
Fig. 3 shows a gas turbine rotor with an inclusion in the rotor

blade at 90'. The density of this inclusion is for minimum 19 %

higher than the average density.

In the last figure are shown two adjacent tomograms of a Sintered Silicon Carbide rotor each with two different grey level

representations. The upper images show the - 20 % range the

lower the - 5 % range of the average density. The lower part of

the two torn-: grams shows one respectively two voids (about 150 um

diameter) in the wheel nave.

Conclusions

Computed X- ray tomography i s a u s e f u 1 tool to d

e

tec t density

variations and defects like voids and inclu sions i n ceramic

components

.

For a quantitat i v e d e scription of s

u

ch d e f ects a well suited tomograph with impr ov ed spatial r eso luti on is need

ed .

99
BPiM I HPSN - PLATE I CT<86e4lS.1386>

Fig. 1: Tomogram of a hot pressed

Si 1 icon Nitride plate

Fig. 2: Four tomograms of a

Si impregnated Silicon plate

B^M I Rotor of a Gas Turbine | CT <868415.125?>

Fig. 3: Tomogram of a gas turbine

rotor - Inclusion in the rotor

blade at 90°

Fig. 4: Two tomograms of a Sintered

Silicon Carbide rotor.

Voids in the lower part of

the wheel nave

100
NONDESTRUCTIVE CHARACTERIZATION OF CERAMIC MATERIALS:

SMALL ANGLE NEUTRON SCATTERING

K. A. Hardman-Rhyne and E. R. Fuller, Jr.

Ceramics Division, Institute for Materials Science and Engineering

National Bureau of Standards, Gaithersburg, MD 20899

ABSTRACT

The use of small angle neutron scattering (SANS) is discussed as a nondestructive technique for characterizing ceramic materials. SANS is seen to be a method for quantifying distributed microstructural properties

of ceramics that have been identified by alternate techniques. As such,

SANS is a complimentary technique to other characterization methods. Two

scattering regimes are emphasized: 1) diffraction for microstructural

phenomena of less than 100 nm; and 2) beam broadening for microstructural

phenomena from 100 nm to approximately 20 /xm.

INTRODUCTION

Small angle neutron scattering (SANS) techniques are used to study

microstructural phenomena in the range of 1 to lO'' nm in size. Since these

techniques cover a wide range of sizes, they are particularly useful for

studies of ceramic processing and distributed damage in ceramics. While

many metal and alloy systems have used SANS techniques, only recently have

experiments been conducted on ceramic materials. This is not surprising

considering the difficulties inherent in analyzing SANS data on these

materials. Often ceramics have several microstructural components present

in the material. These include: residual pores from the sintering process,

inclusions or impurities from starting materials, second phases, and

microcracks or cavities from temperature and/or applied stress effects.

All these effects produce small angle neutron scattering. It is

important either to eliminate all effects except the one of interest or to

identify the effects through complementary techniques, such as electron or optical microscopy. While these complementary techniques can identify

defects, voids, and second phases, SANS can quantify these effects

throughout the bulk of the materials in a nondestructive way due to the

general nature of neutrons

Neutrons are an excellent nondestructive probe of microstructural

features. Because they primarily interact with the nucleus of the atoms,

neutrons are highly penetrating without disturbing the material. This

penetration power allows the bulk of the material to be interrogated,

whereas x-rays and other techniques are more sensitive to surface

phenomena. Another strength of neutron scattering is its dependence on the

chemical elements present in the material, through a quantity called the

coherent scattering length. This scattering dependence varies in an

101
unsystematic way from one element to another, so that differences between

elements with similar atomic numbers can be detected [1].

This paper will discuss NBS research using SANS to characterize

ceramic materials. The main intention is to emphasize the advantages and

limitations of SANS techniques and to spark interest in further SANS

research on ceramics. The next section briefly describes the SANS

instrument at NBS. The following section on SANS experiments is divided

into two parts: 1) the diffraction scattering regime, which is applicable

to small particles or defects in the range of 1 to 100 nm; and 2) the beam

broadening scattering regime for defects and particles on the order of 0.1

fj,m to 20 /zm. Theoretical aspects of small angle neutron scattering are not

specifically discussed, but the interested reader is referred to the

following articles [1-4].

SANS INSTRUMENT

The SANS instrument at the National Bureau of Standards is described

in detail elsewhere [5]; however, a diagram of the major components is

shown in figure 1. The wavelength. A, can be varied from 0.4 to 1.0 nm by

selecting the appropriate speed of a rotating helical-channel velocity

selector. This is particularly important in beam broadening experiments

because the wavelength dependency of the neutron scattering is a necessary

part of this analysis. Longer wavelengths are also useful in diffraction

measurements, when larger size particles or voids (>500 nm) are being

examined and when multiple Bragg scattering from the crystal structure of

the material is to be avoided. To obtain longer wavelengths in a reasonable time period, a cold source is needed in the reactor to lower the

neutron thermal equilibrium temperature, and shift peak intensities to

larger wavelengths. The SANS facility at NBS is currently adding a cold

source, which will increase the neutron flux at the sample significantly

and will result in shorter measurement times.

Two types of collimating apertures define the beam direction and

divergence. One type consists of a pair of cadmium pin hole irises, one after the velocity selector and another before the sample chamber.

h 4.5 m 3.5 m H

LSI 11/23

1 6-Bit Minicomputer

Figure 1. Diagram of the NBS small angle neutron scattering facility.

102
The other collimation system is for higher resolution measurements and

consists of a set of channels in cadmium masks which effectively converge

the neutron beam to a point at the center of the detector. A multiple

sample chamber is available, which is computer controlled and can be used

under vacuum. Single samples can be studied as a function of temperature

from 12 to 1600 K. Horizontal and vertical field electromagnets are also

available. Sample sizes are usually 10 to 25 mm in diameter and 2 to 30 mm

thick. Uniform thickness is essential for analyzing the results.

The scattered neutrons are detected on a 640 mm X 640 mm positionsensitive proportional counter which is divided into 128 colximns and 128 rows and has a spatial resolution of 8 mm in each direction. A dedicated

minicomputer processes the signals from the detector and stores the data.

The angle between the incident beam and the scattered beam is the

scattering angle, e. The magnitude of the scattering vector Q is

(47r/A) sin[ e/2 ] , which is approximated by 2ne/\ in the small angle limit.

SANS EXPERIMENTS

Diffraction and beam broadening measurements are quite different and

are usually obtained in different configurations of the SANS instrument.

Diffraction experiments probe for microstructural phenomena in the range of

1 to 100 nm. The SANS detector is usually located directly behind the

sample so that the incident beam is centered in the middle of the detector.

The experiment generally requires 6 to 12 hours and a beamstop is used to

eliminate scattering from the incident beam. Sample thickness is kept

small (2 to 6 mm) to minim.ize multiple scattering effects. Absorption and

incoherent neutron scattering should be reduced for best results in most

SANS experiments with ceramic materials.

The beam broadening effect is wavelength (A) dependent with the most

sensitive results occurring at A greater than 0.7 nm. Typical experiments

require only 3 minutes to 2 hours depending on: A; the void or particle

size; the density; and the sample thickness. The incident beam is set off

to one side of the detector, away from the beamstop, as the exact center of

the neutron beam is needed to analyze the data. This can be done without

harming the detector because the peak intensity is greatly reduced in the

broadened state. Cadmium foil can be used to reduce further the scattered

beam intensity over the entire detector. The samples should be measured at

three wavelengths or more for best results. Also thicker samples, 5 to 30

mm, are desirable to increase the number of scattering events.

Diffraction Regime

Most SMIS experiments are in the diffraction region and in ceramics

are concerned with inhomogeneities such as voids, cavities, microcracks

,

precipitates, sintered porosity, inclusions, nucleation, and growth of

second phases. In principle, it is possible to determine quantitatively

particle size, shape, size distribution, surface area and other

microstructural values. Examples of SANS experiments with ceramic

materials include the following: formation and growth of heterogeneities

in glass by A. F. Wright [6]; creep cavitation in sintered alumina [7-8]

and in silicon carbide [9-10] by R. A. Page and co-workers; growth and

coarsening of pure and doped Zr02 by A. F. Wright, S. Nunn and N. H. Brett

103
[11]; microcracks in sintered YCrOg by E. D. Case and C. Glinka [12]; Fe

and W inclusions in hot-pressed SigN^ by N. J. Tighe, K. A. Hardman-Rhyne

,

and Y. N. Lu [13]; and powder characterization and ceramic processing by K. A. Hardman-Rhyne, K. G. Frase , andN. F. Berk [14-16]. The study

characterizing microcracks in sintered YCrOg [12] will be briefly discussed

here as it illustrates the pertinent features of SANS characterization.

To simplify interpretation of SANS data, systems with only two

components, such as a matrix material and either precipitates or voids, are preferred. However, one way to avoid the complexities of analyzing a multiple component systems is to run a control sample. This was done in

the YCrOg microcrack experiment [12]. YCrOg is a material which is

sintered around 1750 "C, thereby establishing its microstructure (porosity,

grain size, etc.), but which undergoes an apparent phase transition around

llOO'C. Thus, if the material is quenched from above 1100 °C, microcracks

as well as pores are present in the final material. However, if the

material is annealed at 1050 °C, the microcracks can heal and only the

pores remain. Thus, by subtracting the neutron scattering data of a healed

YCrOg sample from that of a microcracked sample, a difference scattering

distribution, called Iqipf. can be obtained, which is due only to the

microcracks in the material. These difference spectra can then be analyzed

to obtain quantitati.ve information regarding the microcracks.

SANS data can be analyzed in two extremes of scattering angle, Q. For

small Q, the so-called Guinier region [17], scattering is assumed to occur from sharp -edged, randomly oriented scattering centers of an given

geometry. The scattering intensity in this regime is generally given by

I - Vp2 exp[-RG2QV3], (1)

where Vp is the volume of a particle (or scattering center) and R^ is the

particle's radius of gyration with respect to its center of gravity.

However, a slightly different functional form for the Guinier region is

obtained when the scattering from microcracks is modeled after randomly

oriented thin disks of thickness 2H and diameter 2a [18]:

I ~ 2V(Vpnp/Qa)2 exp[ -H^qVS] , (2)

where V is the sample volume seen by the neutron beam, rip is the number of

microcracks per volume, or the microcrack number density, and QH < 1 « Qa.

This low Q fit to the data (Id iff) ^® seen in figure 2. This low Q

scattering regime directly provides information about the microcrack

opening, 2H, by plotting the logarithm of the intensity times versus . For the other extreme of scattering behavior, the high Q or Porod

region, the scattering intensity has a Q" ^ dependence [18]. The Porod

region is more sensitive to smaller dimensions of the scattering centers

and results in a characteristic Porod length which measures the total

surface area to volume ratio, if absolute intensities can be determined.

For the microcrack data, the total surface area of microcracks was estimated from the high Q data by normalizing to the scattering from water.

Figure 3 shows the logarithmic function of Idiff versus the scattering

vector Q. The solid line is a Porod [18] functional fit which describes

the data at large Q (Q > 0.03 A" ^) and can be expressed in the following

form:

I ~ (27r/V)(S/QM, (3)

104
qI 1 1 i I ' I I I.I. 010 012 ,014 .016 .018 020 .022 SCATTERING VECTOR Q (A"')

Figure 2. The low Q scattering data from microcracks in a 7 nun thick

specimen of YCr03 . The solid line is a least- squares fit to a scattering

function for randomly oriented thin disks of thickness 2H. From Ref. [12].

Figure 3. The net scattering due to microcracks from a YCrOg specimen.

The solid line is a least-squares fit to a Porod law. From Ref. [12].

105
where V is the sample volume seen by the neutron beam and S is the total

surface area of the scattering centers. Since H « a, S/V ~ 7r<a^>np , where

<a^> is the average square of the microcrack radius.

Since the SANS data alone do not provide a complete characterization

of the microcrack sizes, a and H, and microcrack number density, np , these

measurements were complimented by elasticity measurements on the specimens.

The decrement in elastic modulus with microcracking gives a measure of

<a^>np, where <a''> is the average of the radius cubed, thus allowing the

mean microcrack radius and microcrack number density to be calculated.

These microcrack parameters have been calculated and agree well with

similar parameters in the literature [12]. They are as follows:

microcrack number density (4.7 X 10^ cm" ) , volume fraction (2.6 X 10" •^),

and crack opening displacement (25 nm) . The mean microcrack radius of

5.7 nm corresponds well with the measured grain size of 6 /im, and is

consistent with a model of localized stress induced microcracks

.

Further SANS diffraction experiments of distributed damage due to

stress and temperature are expected as this technique becomes familiar in

ceramic research. Although optical and electron microscopy can identify

small defects in advanced ceramic materials (<10 /im) , SANS can quantify the

size, shape and distribution of these defects in the bulk of the material.

These results, coupled with failure tests, can help in understanding and

improving the structural reliability of advanced ceramic materials.

Beam Broadening Region

Porosity is a critical aspect in the densification process of a sintered ceramic material. To elucidate the extent of such porosity, a quantitative study with SANS has been conducted at NBS to determine average

pore size. Rather than restricting the SANS measurements to the typical 1 to 100 nm size regime of SANS diffraction, we have explored the neutron

beam broadening region by extending the SANS characterization into the tens

of micrometers size regime [4,14]. This extension of SANS technique to

larger sizes is an important development, because it allows a greater

overlap of SANS characterization with other NDE techniques. In addition,

as this new regime of scattering analysis has developed, it has found use

in numerous other ceramic processing applications, e.g., powder

characterization [14-15]. The technique is illustrated here by a characterization of the porosity in a YCrOg material.

Two samples of YCr03 were fabricated from pure powders by isostatic

pressing at 207 MPa (30,000 psi) ; and one of the samples was sintered. The

density of the "green" compact (the unfired ceramic) was approximately 57%

of theoretical density and that of the sintered material was approximately

94%. The starting ceramic powder, with approximately 30% of theoretical

density, was also examined. Since beam broadening measurements are wavelength dependent, SANS experiments were taken at six or seven of the

following wavelengths: 0.485, 0.545, 0.625, 0.700, 0.800, 0.900, 0.950,

1.00 nm. The observed scattering revealed a striking difference between

the samples as illustrated in figures 4(a) and 4(b), which are SANS spectra

for the sintered and "green" compact specimens, respectively. The sintered

material [figure 4(a)] shows little wavelength dependence. In contrast,

the "green" compact reveals dramatic beam broadening that is strongly

wavelength dependent [see figure 4(b)]. This dependence is illustrated in

figure 5 by plotting the normalized intensity versus scattering vector Q

106
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Figure 4. SANS spectra at three wavelengths (a) for a sintered compact of

YCrOg and (b) for a "green" compact of YCrOg . Plotted is the scattering

intensity versus a linear column slice (detector row number) through the

center of the scattering plane.

107
Figure 5. Normalized scattering intensity versus scattering vector, Q, for

SANS from a "green" YCrOg compact at five wavelengths,

for five wavelengths.

The direct beam is wavelength independent with respect to the

scattering angle, e and is defined in part by the instrumental collimation.

Beam broadening data can resemble a Gaussian distribution at low Q values,

where the full width at half maximum, Ae can be determined by the Gaussian

standard deviation parameter as shown below:

I = I, exp[-Q2 72^32 ] (4)

For full width a half maximum AQ = 2.355 Cq . Since Q = Ine/X, the angular

variation, Ae , is given by 0.3748 Xoq . The Ae value contains both the beam

broadening scattering and that due strictly to the direct (unscattered)

beam, thus the contribution from the direct beam, e^, , must be subtracted

from the experimentally determined value, e, to obtain Ae

:

Ae = [e2 - eb^]^/^. (5)

Although the qualitative aspects of the data clearly demonstrate a strong effect of ceramic processing on the neutron scatterers population in

these materials, quantitative measures of the particle or void size, shape

and size distribution are less straightforward. Moreover, the neutron

phase shifts are well within an intermediate range of values for which

neutron scattering is not expected to be analyzable by multiple refractive

108
0.06

0.01 ' ' ' 1 04 0.6 0.8 1.0 X (nm)

Figure 6. Full width at half maximum, Ae (in radians), versus the

wavelength, A, for voids in a "green" YCrOg compact. The squares are the

data and the circles are the Ae values derived from theory.

behavior alone. Therefore, a generalized beam broadening theory [4]

relevant for this region and multiple refraction has been developed to

quantitatively analyze the SANS data for densified ceramics and other

distributed defects in this size regime. Radius, void (or particle)

density ratio and shape factor can be obtained from this theory, which can

be expanded to consider particle packing, polydispersivity and various

shapes of particles and voids other than spheres. Excellent agreement of

data and theoretical values for Ae can be seen in figure 6 . The average

radius void size in the YCrOg "green" compact material is 0.17 ^m and has a void density ratio of 0.42 compared to the overall density ratio of 0.43.

The general beam broadening theory and SANS technique allows us to

study the densification process in a nondestructive way. It is being

extended to study the sintering of spinels (MgAl2 0^) as a function of

temperature [16]. In-situ as well as other ceramic processing experiments

are expected to develop fully the capabilities of this new approach in

SANS. These techniques should also have applications in a number of other

material disciplines, such as magnetic broadening effects, pores in metal

alloys and colloidal chemistry.

109
REFERENCES

1. Kostorz, G., Treatise on Materials Science and Technology . Vol. 15,

ed. , G. Kostorz, Academic Press, New York, 5-8 & 227-289 (1979).

2. Weertman, J. R. , Nondestructive Evaluations: Microstructural

Characterization and Reliability Strategies , eds . 0. Buck and

S. M. Wolf, American Institute of Mining, Metallurgical, and Petroleum

Engineers, New York, 147-168 (1981).

3. Herman, H. , Non-Destructive Evaluation of Materials with Cold Neutron

Beams, report to Naval Air Systems Command, Contract No. N00019-77-M0418, Washington, DC, December 1977.

4. Berk, N. F. ; Hardman-Rhyne , K. A., J. Appl. Cryst. 18, 467-472 (1985).

5. Glinka, C. J., AIP Conference Proceedings No. 89, J. Faber, ed.

,

Neutron Scattering- 1981 , 395 (1981).

6. Wright, A. F., Neutron Scattering 1981 AIP Conference Proceedings,

ed. , J. Faber, 359 (1982)

.

7. Page, R. A.; Lankford, J.; Spooner, S., J. Mater. Sci. 19 [10], 3360-

3374 (1984).

8. Page, R. A.; Lankford, J.; Chan, K. S.; Hardman-Rhyne, K. A.; Spooner,

S., J. Am. Ceram. Soc. 70 [3], 137-145 (1987).

9. Page, R. A.; Lankford; Spooner, S., ActaMetall. 32 [9], 1275-1286

(1984)

.

10. Page, R. A. and Spooner, S., J. Mater. Sci. 21 [4], 1417-1422 (1986).

11. Wright, A. F. ; Nunn, S.; Brett, N. H. , in Proceedings of the

Conference on Zirconia, Stuttgart, West Germany, 1983.

12. Case, E. D.; Glinka, C. J., J. of Mater. Sci. 19 [9], 2962-2968

(1984)

.

13. Tighe, N. J.; Hardman-Rhyne, K. A.; Lu, Y. N. , Ceram. Eng. & Sci.

Proc. 6, 835-849 (1985).

14. Hardman-Rhyne, K. A.; Berk, N. F. , J. Appl. Cryst. 18, 473-479 (1985).

15. Hardman-Rhyne, K. A.; Frase, K. G.; Berk, N. F. , Physica B 136B [1-3],

223-225 (1986).

16. Frase, K. G.; Hardman-Rhyne, K. A., to be published.

17. Guinier, A., X-ray Diffraction . Chapter 10, W. H. Freeman and Company,

San Francisco, 1963.

18. Porod, G., Acta Phys . Austriaca 2, 255 (1948).

110
N3S/3a:^ - Synipcsium on

I'lateria'.s Research and Testing

Introduction to the Versailles Project on Advanced Materials and Standards (VAT-IAS)

Technical Working Area: Wear Test Methods

3y: S. Becker and J. Lexow

In June 1982 at the Zconcmic Sun.-nit of Versailles the

Working Group "Technology, Growth and Z::nplGyment " prcpcsea a multilateral collaboration on science and

technology. 0n3 of the recommendations concerned rtdvanced Materials and Standards'.'

3y now, several Technical V/orking rarties have "been

launched. The Technical Working Area "Wear Test Methods"

has the follo'ving aims:

• Improvements of the reproducibility and comparability

of wear testing by developing internationally agreed wear

test methodologies.

• Characterization of the wear behaviour of "advanced''

materials in comparison with conventional materials.

As a first step a Round Robin on the wear behaviour of a ball-on-disc sliding test system consisting of four pairings

of c<,-Al20^ ceramic and steel (AI3I521C0) under agreed test

conditions like load, sliding velocity, temperature and sliding

distance wss conducted. By May '86 about two thirds of the

participants had already delivered their test results.

The evaluation of the reports is now under way. It is

rather difficult because of the many influencing factors

that have to be considered.

For the different material combinations typical wear patterns

were observed including for example material transfer (steel/

steel and steel/ceramic) and very minor wear for the Al20.^ couple.

The viewgraphs 4-11 show some of the BAI4-results

The final results of the Round Robin will be communicated

to the participants and to the public as soon as the

evaluation has been completed.

Ill
112
INSTRUCTION SHEET

FOR VAiMAS- PROJECT ROUND ROBIN

WEAR TEST .'-lETH'^S

Please Perform Tests at the folloving Agreed Test Procedure *

1. Test System

-Stationary ball (10mm diameter) against rotating disc (40mm

outer diameter, 32mm wear track diameter)

-rotation in the horizontal plane

-direction of rotation of disc to be indicated by each laboratory

-Ball, disc and debris are to be collected and protecti/ely sto

red in plastic containers.

-Holders for disc and ball are to the discretion of each laboratory.

-Please report (if possible) vibrations (e.g. vibration amplitudes and frequency distribution) of test rig at stated

location

-Report stiffness-data of the test rig (if available).

2. Materials -if-si £p^[^ce^

Disc: AISI 52100 \l ' , .^-^ alpha-Al203 K, .

' VA

Sic ) I SiAlON ) to be supplied later iest^(^rUc^ Si3N4 «;i-,N^ ) ——r Ball: AISI 52100

alpha-Al203 /A I y//A

AISI 52100 coated with Ni as per N.B.S. ) in prep<tration

3. Atmosphere

Laboratory air (50±10% rel. humidity; 23±1°C)

(If a different atmosphere is applied it will have to be documented and reported with the test results.)

4. Lubricant

No lubricant will be used in the first stage of the round

robin tests.

5. Operating Variables

-Motion: Continuous unidirectional sliding

-Velocity: 0.1 m/s

-Normal Load: 10 N

-Temperature: 23±1°C

-Sliding distance: 1 km or until f=1.5 or until displacement due to wear = 1 mm

-Nximber of tests: 3-5

*If other conditions are used please indicate them.

p . t . o ,

113
6. Preparation of Surfaces

-Specimen are to be used as received, i.e. no mechanical surface finishing is necessary.

-Surfaces are to be cleaned immediately prior to each test.

a) Washing in freon ( CI2FC-CF2CI ) is preferred to washing in

ethylalcohol since the water content in the alcohol may

corrode the metallic surfaces

b) Drying in warm air

c) Rinsing with hexane

d) Drying in a drying oven at HQOc, 30 min

-Chemicals of pure quality are to be used.

-Samples are to be stored and trans-.orted in desiccators.

7. Measurements

a ) Wear

(Please define whether wear of ball, wear of disc, or total

wear of both ball and disc are measured.

)

-Continuously measured and recorded

-Specimens weighed before and after each test

-Wear scar diameter on ball to be measured with an optical

microscope

-Wear of ball and disc

-Profilogram results of both surfaces before and after the te

b) Friction:

(Please define whether the friction force or the friction

torque are measured.

)

-Submit a simplified graph giving the fluctuations at the

beginning and at the end, minimum and maximum deviations

during the test etc..

c) Debris:

-All debris is to be collected (See 1).

8. Pos t-iMeasurement Handling

-Indicate the sample number with a water resistant felt pen

on the side of the disc which was not subjected to wear.

-Mark the worn area on the ceramic ball with a water resistant felt pen after each test; first with red, second with

blue, third with green.

-After the first and second test turn the ceramic ball into

a new position without contaminating the sample.

-Put the samples back into the bags immideately after the

test and the examinations in order to avoid mixing.

9. Examination

Surfaces and debris are to be examined by optical color photo

graphy and SEM techniques. Photographs at standard magnifications 50x, lOOx, 200x, 500x, lOOOx

10. Report

All results should be reported to Prof. Czichos by

January 31, 1986.

Reports should in particular comprise:

-Information on the test setup

-Photographs of surfaces

-Graphs of wear and friction

-Additional information on deviations from agreed conditions or other information concerning the test.

114
r 23%

50% r.h.

-- ION

'N V -- 0,1 m/s

V4_y s

-. ikm

Material pairings Measurements

disc

steel alumina • Friction force

steel kit 1 kit 2 • Wear

( system, ball, disc

alumina kit 3 kit 4 • Wear surfaces

( SEM, profilometry )

BAM

Berlin Wear test conditions

Becker

86 52 1321

E A

Berlin Evaluation of wear

Becker

86 52 1322

115
02 mm

01 5

transfer

1 01 >

mm' oQj 02

Al,03 steel A[,03

1

0.003 1 0.00007

steel steel Al,03 AI2O3

o peratin g voriobles -motion: sliding

- velocity O lm/s

- normal load: ION

- temperoture: 22t]'Z

- sliding distance: 1 km

BAM Becker

Berlin Wear volume of balls and discs 86 52 1324

116
m CO

117

119
CD

120

CD

m

(ZD

I

CD

uurlQSZ

122
123
Fretting of Ceramic Materials

by Dr.-Ing. Dieter Klaffke

Modern engineering ceramics are of world wide interest in

raany fields of technical applications.

The tribological behaviour of ceramics is not yet well known,

investigations of friction and wear behaviour should accompany

the further development of ceramic materials.

The fretting wear test (oscillating sliding motions with small

stroke) is a useful tool for the study of tribological behaviour

of materials; wear processes are concentrated on a small part

of specimens so that wear volumes are to be evaluated with good

accuracy.

A tribometer for fretting tests has been developed which

enables to measure the friction force, the linear wear and

the contact resistance during the tests. Fretting tests have

been performed with different ceramic materials against ceramics

and steels.

In fretting tests with steel (ball : AISI 52 100) against

different ceramics the oxide-ceramics (alumina, partially

stabilized zirconia) show much less wear thansilicon nitride and

silicon carbide.

In fretting test with ceramic cylinders against different

steels with different thermochemical surface layers, again

the oxides show less wear than the silicon-based ceramics.

The lowest wear of the steel occurs if it is protected by

a vanadium carbide layer.

Wear processes in the case of ceramic against ceramic are

of pronounced heterogeneous character, showing peaks in the

time behaviour as well of the linear wear as of the coefficient

124
of friction, caused by stochastically removing of hard wear

particles. The coefficient of friction in the steady state is

in all cases higher than 0.4 and for some combinations of

ceramics as high as 0.9.

The humidity of surrounding atmosphere can have a pronounced

influence on wear and friction behaviour. In dry air (5 % r.h.

the coefficient of friction is in the range 0.6 to 0.8 for

SiC against itself and drops down to values below 0.2 for

9 5 % r.h. The wear volumes are about one order of magnitude

greater in air of low humidity than at high humidity.

Specially in dry air wear processes show a pronounced two

range behaviour, starting with high wear rate, than changing

to mild wear at a transition point which seems to depend very

sensitively on the relative humidity.

125
Fretting of Ceramic Materials

D. Klaffke

NB5/BAM-5ymposium on

Materials Research and Testing

Subject: Ceramic Materials

Berlin, 22.04.86

126
stationary ball. I <t>10mm. \ Al SI 52100.

800 HV. ]

i

oscillating disc.

'^^0 mm X 5 mm .

lab. air

BAM Characteristic of test system Klaffke

Berlin (fretting wear) 85 52 1 105

Fig. 1

preferred test conditions

Ax = 0,2 mm

F„ = 20 N

f = 20 Hz

T = 22°C

Fig. 2 Fretting test specimen (cylinder)

127
Fig. 3 Fretting test specimen (flat)

Fig. 4 Fretting test specimen (SiC)

128
DMS-Messung

(Reibkraftj Pad- Ausrichtung

Biegegelenk

Normalkraft

(Cewicht)

Btattfedern

Kreuzfedergelenk

Oberprobe

Pleuelstange

AX = 0,2 mm

^ 20 N

N ^ ...J2- '0

f : 20 Hz

Motor Exzenter

BAM

Berlin Tribometer fur SchwingungsverschleiB

Klaffke 86 52 1247

Fig. 6 Coefficient of friction and linear wear as function

of time in a fretting test; steel/steel

129
Werkstoff nahere Bezeichnung Form f^a Mm! Rz Farbe

drucklos gesintertes Zylinder 0.7 4.3 hell- Siliziumnitrid Ebene 0.25 2.0 grau

MgO-teilstabilisiertes Zylinder 0.8 5.0 geib

Zirkonoxid Ebene 0.2 2.6 Aluminiumoxid mit Zylinder 1.6 9.8 weilJ 98,5 % AI2O3 Ebene 1.0 8.0

Silizium infiltrier- Zylinder 0.2 1.0 dunkeltes Siliziumcarbid Ebene 2.6 18.9 grau

gesintertes Silizium- Zylinder 0.1 0.9 mittel- carbid Ebene 0.2 1.7 grau

Si3N4

Zr02

AI2O3

SiSiC

SSiC

BAM

Berlin

Zusammenstellung der untersuchten keramischen

Werkstoffe

Klaffke 86 52 1249

Fig. 7 Ceramic materials used in fretting tests

130
Fig. 8 Wear scars (SEM) on steel balls after fretting tests

against ceramics

131
Zr02

AI2O3

50 fjm ["250 pm

SiSiC

Ax=0.2mm; =20N: f=20Hz; N=^.8-10^; lab.air 22''C

BAM Profilograms of worn ceramic flats Klaffke

Berlin (bail/flat system), perpendicular to sliding direction 85 52 1 108

Fig. 9

50

250

5L

250 ZrOz

—f

2.5L

250 ^hOs

20V^

250 ^'^'C

—1

\ h 1^ ^w:

Ax =0.2 mm; = 20 A/; f = 20 Hz.- N=i.8- 10^.- T= 22°C.- lab. air

BAM Profilograms of worn ceramic flats Klaffke

Berlin (ball/flat system), perpendicular to sliding direction 85 52 1 109

Fig. 10

132
y. ball

0,5 mm^

0.5 mm3

f ^V, flat

1

m

Zr02 AI2OJ SiSiC

Balls :. P 10 mm

,

AlSl 52100.

800 HV.

f =20 Hz

Ax= 0.2 mm

=20N

N =1.8-10^

T =22°C

lab. air

BAM

Berlin

Component wear of ball/flat (steel/ceramiclsystem; determined from profilograms

Klaffke 85 52 1107

Fig. 11

Stahl Warmebehandlung

HV0,2

(HV0,4)

Ra Mm

90MnV8 Vergiiten 400 .03 .28 90MnV8 Vanadieren 880 .13 1.03 90MnV8 Chromieren (1120) .03 .22 42CrMoS4 Vergiiten 350 .03 .26

42CrlVloS4 Borieren (1500) .09 .71

BAM Klaffke

Berlin Zusammenstellung der untersuchten Stahlwerkstoffe 86 52 1250

Fig. 12 Steel used in fretting tests against ceramic materials

133
0.1

0,05

0,05

mm''

0,1 > ^ Moo - 1/1 U1

:e Lj i_i

e

CO CO* > > c c r 2: o r*j <3

CT* -T ^ CT»

0,1

0,05

0,05

0,1

ZrOj

LX - 0,21 mm

= 20 N

^ = 20 Hz

yV = 1,2-10* T = 22°:

>c

BAM

Berlin VerschleiSvolumina an Keramik- und Stahlkorper

Klaffke 86 52 1251

Fig. 13 Wear volumes of ceramic and of steel specimens after

fretting tests

0 5 10 h 15

/ BAM

Berlin

Reibungszahl und linearer VerschleiBbetrag

bei SchwingungsverschleiS von Si3N4-Zyiinder

gegen Si3N4-Ebene

Klaffke 86 52 1253

Fig. 14 Coefficient of friction and linear wear in a fretting

tests; SiN3N4 / Si3N4

134
20

601-

0,8

0,6

0,4

0,2

LX - 0,21 mm

= 20 N

/ = 20 Hz N = 1,2-10* T = 22°C

10

t

15

BAM

Berlin

Reibungszahl und linearer VerschleiBbetrag bei SchwingungsverschleiS von Zr02-Zylinder

gegen Zr02-Ebene

Klaffke 86 52 1252

.g. 15 Coefficient of friction and linear wear in a fretting

test; Zr02 / Zr02

20

T 40 ^m -

60 L

B A M

Berlin

Reibungszahl und linearer VerschleiBbetrag

bei SchwingungsverschleiS von Al203-Zyiinder

gegen Al203-Ebene

Klaffke 86 52 1254

g. 16 Coefficient of friction and linear wear in a fretting

test; AI2O3 / AI2O3

135
Fig. 17 Coefficient of friction and linear wear in a fretting

test; SSiC/SSiC

0,5 mm

BAM REM Aufnahme der VerschleiEspur in der Klaffke

Berlin Al203-Ebene nach SchwingungsverschleiS 86 52 1258

gegen SSiC-Zyiinder

. 18 SEM micrograph of wear scar on the AI2O 3 plane,

counterbody; SSiC - cylinder

136
BAM REM Aufnahme der VerschleiBspur in der Klaffke

Berlin SSiC-Ebene nach SchwingungsverschleiB 86 52 1257

gegen 513 N4-Zylinder

Fig. 19 SEM micrograph of wear scar on the SSiC plane

counterbody Si3N4 - cylinder

0.1

^5 0,05

0

c3 0,05

0,1

S '3 N

0,1

mm^

^^^ = 0,21 mm

= 20 N

A = 20 Hz

/V = 1.2-10' T = 22°:

u,Ui

z

J)

0

mm''

0,1

try

, 0 _~ IT! M < 1/1

:£* o~ cT S:i

i/i t--i •< 1/1

BAM

Berlin

VerschieiBvolumina an Si3N4- und Zr02-Zylinder

und Keramik-Ebenen nach SchwingungsverschleiS

Klaffke 86 52 1256

Fig. 20 Wear volumes at ceramic cylinders and ceramic planes

after fretting tests

137
V bctU

0,7 mm"

0,1 - mm

HfPSN

;

>

*

*

u

t

f

f

t w t

f

*

\00%

AX = ^ mm

A/ = GOO. 000

FiR 21 Influence of humidity on wear volumes at steel balls and

ceramic planes for three Si3N4 modifications

138
Fig. 22 Wear scars and profilograms of wear scars after fretting

steel ball / SiC-flat; at different hvmiidities

139
V baCL

0,1- mm'

0 -

0.1 - mm

' Wy^ fiat

100 Cr- 6 C

5iC

^

^ SchLi'cker^uPi

= 0,3<5 mm

V =600.000 ^ = /C) cp>5

Fig. 23 Wear volumes of steel and ceramic specimens after fretting

test at different humidities; three modifications of SiC

140
<5h t

Fig. 24 Coefficient of friction and linear wear for two modifications

of SiSiC in dry air

0,1 -

0

cyUnder

mm'

mm'

I &

FK

100%

ScSic floii

N =^00.000 ^ ^ rocps

r=^2J2°C

F'K • Feinkorn

fine, ^raln

Coar5e grain

Fig. 25 Wear volumes of ceramic specimens after fretting test at

different humidities; fine grained and coarse grains silicized SiC

141
Fig. 26 Wear scar profilograms on fine grained SiSiC specimen after

fretting test in air with different humidities

Fig. 27 Same as above for coarse grained SiSiC

142
T. h =5%

—'—

ZZI371UJ_.:

ri 1 "T^l

&\ 1 1

1 ., ' -

i = 1

^

! -1 ' 1 i

i

. -1 1 T ' M 1 1

1 U L.

-1 i '-a-rt-r—

;

-! 1 1- 11 I.I

1

-i -i-j_^_j_,J_i.3.L

o

30'

30 +

IP

1o

1

i

: • : i

:

i

Figs. 28-30 Influence of normal force on fretting wear behavior of fine

grains SiSiC against itself at different humidities

143
Tribology of Advanced Ceramics

S. M. Hsu

National Institute of Standards and Technology

Gaithersburg, MD 20899

ABSTRACT

Research in materials science of advanced ceramics has been developing

rapidly. As materials change and new materials are constantly being

developed, the evaluation of these materials for tribological applications

becomes a critical issue. This paper examines several important

considerations in friction and wear testing of ceramic materials. Detailed

test sample characterization is required to define the materials pair.

Sample preparation, selection of load and speed, environmental control, and

contamination effects are crucial to the understanding of wear mechanisms.

The tribological property of Alumina and water was studied. Water has

been found to reduce friction and wear of the alumina system. A

tribochemical film was found on the surfaces of water lubricated alumina

wear surfaces under boundary lubrication conditions. Results suggest that

alumina, under contact stress and high flash temperatures, undergoes phase

transformation from alpha to gamma alumina. Gamma alumina subsequently

reacts with water to form hydroxides of alumina. Depending on

temperatures, boehmite [AlO(OH)] and bayerite [Al(0H)3j will be formed.

Wear debris analysis suggests the presence of bayerite. Independent wear

test results found that boehmite reduces friction. Phase diagram and

decomposition sequence data suggest that boehmite could transform to

144
Considerations in ceramic wear testing

The unique properties of ceramics require that care must be taken when

conducting wear tests. The two critical conditions for wear testing are

speed and load. Under normal four-ball wear test conditions (52100 steel

specimens and hydrocarbon lubricants) , the typical speed and load ranges

are 0.1 to 1 m/s sliding speed and 40-240 Newtons of machine loading. For

ceramics, due to large differences in the hardness of various ceramics, to

achieve the same contact stresses, loads have to be adjusted significantly

for valid comparison.

The contact pressure is dependent upon the geometry of the contact and

the material properties. The geometry of the four-ball contacts provide a

distribution of force based on the applied load of N = 0.408L, where L is

the applied machine load, and N is the normal load on the surface of single

ball. Material parameters, however control the load distribution on the

surface. The initial contact area is described by the Hertz equation given

in Equation [ 1 ]

37rN(ki+k2)RiR2 1/3

a =

4 (R1+R2)

where R]^ = R2 = radius of ball (6.35 mm = 0.25 inch)

N = normal load on the surface of a ball

a = radius of Hertz constant

Subscripts refer to materials 1 and 2. For the case where they are

the same:

1 - .2

ki = =

TT E

145
where u

E

= Poisson's ratio

= Young's modulus

The basic result of these material parameters on the Hertzian contact

diameter is that harder materials deform less under stress and produce

higher contact pressures. A graph of calculated maximum contact pressures

(assuming hemispherical pressure distribution) for alumina and 52100 steel

is given in Figure 2 using values of E and u of 206 GPa and 0.30 for 52100

steel and 403 GPa and 0.236 for polycrystalline alumina, respectively. The

curves show that assuming elastic contact, the maximum contact pressures

for alumina are 50% higher than steel given the same applied load. The

pressure provided by steel at twenty kilogram load is experiences by

alumina at a bit less that six kilograms. These calculations seirve as a

conservative estimate of the contact stresses. One would expect,

therefore, that initial machine loads selected for 4-bali wear tests on

ceramics should be lower than loads usually utilized for 52100 steel if one

wishes to test at the same contact pressure.

The other issue critical in ceramic wear testing is the temperature in

the contact junction. Materials with lower thermal conductivity do not

allow heat to dissipate as rapidly and give higher local contact

temperatures . The difference in thermal conductivity and thermal

diffusivity between ceramics and metals therefore causes higher local

contact temperatures for ceramics under the same test conditions. A recent

study by Munro^ , who modeled the adiabatic temperature distribution within

test specimens of the four-ball wear test predicted high temperatures

different for alumina specimens versus steel specimens (Figure 3) . Under

1A6
conditions of 500 rpm, 40 kg load, and coefficient of friction of 0.075

(lubricated) , temperatures of ~160°C were predicted for the contact

junction using steel specimens. Temperatures of ~400°C were predicted for

alumina, a difference of 240°C. In addition, the thermal gradients in the

alumina were much sharper.

These two examples of contact pressure and temperature differences

between alumina and steel specimens illustrate the need to analyze the

actual contact conditions without assumptions. Therefore, data

interpretation should take these factors into account.

Cleaning Procedures

Interest in the use of ceramics has led to an increasing number of

tribological studies being conducted on ceramics. In these studies, there

is a vast array of different methods used to prepare specimens for testing

ranging from rinsing with ethanol-^ or acetone^ to baking^ and ion

cleaning^. Indeed, in many cases, the cleaning procedure affects the

tribological test results.

Unlubricated four-ball tests at low to moderate load exhibited unusual

friction traces (Figure 4) . These specimens had been cleaned

ultrasonically in hexane , then acetone, and rinsed with a stream of

nitrogen gas. The friction traces in Figure 4 indicate that a contaminant

(residual hydrocarbon or possible water) is present on the surface during

these tests which affect the results. At low load (10 kg) the contaminant

protects the alumina surface and low friction and wear are observed. At

moderate loads (15 kg) low friction and wear are observed. At moderate

loads (15 kg) low friction is observed for the initial stage of the test

with a transition to high friction during the test (after approximately 2

147
minutes) . The high load (20 kg) test also starts ar lower friction but

after only a few seconds jumps to the high friction level. The nature of

these results suggest as one explanation a surface contaminant is present

that reduces friction and wear in these tests.

ESCA was used to analyze the surface to characterize the contaminants.

Figure 5 shows an ESCA spectrum on alumina surface. The lower trace

represents the ascleaned surface while the upper trace represents the

surface after approximately a 30A layer of material was removed by argon

ion sputtering. Quantification was accomplished by comparing the ratio of

the carbon peak height to the oxygen peak height. The results of ESCA

analyses on several cleaning procedures is given in Table 1. Organic

solvents seem to leave a high degree of carbon containing contaminant (a,b)

unless followed by a detergent and deionized water rinse (c) , or by baking

at high temperature (d,e). The difference between the as-prepared results

and the sputtered results gives an indication of the tenacity of the

contaminant. A strongly bound contaminant shows very little difference in

C/0 ratio due to sputtering (a) while a weakly adsorbed contaminant shows a

large difference (d, e). The organic solvents apparently leave organic

residues on the surface (a, b) . The Micro/DI water cleaning in addition to

the organic solvents (c) effectively removes most of these residues by

dissolving them into solution. The bakeout removes the organic

contaminants by oxidizing them to CO and CO2 which are then weakly adsorbed

on the surface (d, e) . Additional friction experiments confirmed the role

of surface contaminant levels on friction using these cleaning procedures

(Figure 6) . Specimens cleaned with hexane/acetone exhibit a low

coefficient of friction which is relatively constant throughout the test.

148
The more rigorous procedure (including an overnight bakeout at SSO'C)

produced an immediate high coefficient of friction.

These results show that surface cleaning procedures can have a

pronounced effect on the tribological performance of ceramic materials.

Based on these results on alumina, a general cleaning procedure is

recommended: 1) Ultrasonic cleaning with hexane/toluene (9:1) to remove

residual polishing oils; 2) Ultrasonic cleaning with acetone to remove

polar surface contaminants and residual hexane and toluene; 3) Ultrasonic

cleaning with a laboratory detergent (micro at 2% concentration in

deionized water) to remove residual organics ; 4) Several ultrasonic rinses

with deionized water to remove residual detergent; and 5) High temperature

bakeout when feasible.

Alxomlna Wear under Different Fluids

Alumina was tested under dry and liquid lubricated conditions using a

four-ball wear tester at 600 rpm (0.23 m/s) and loads ranging from 5 to 60

kg. These loads result in estimated mean Hertzian contact pressures of 1.6

to 3.7 GPa for alumina specimens. The data is plotted in Figure 7 as log

of the wear scar diameter versus log of the load. Unlubricated tests were

conducted under dry air (<10 ppm water) atmosphere, and produced very rough

friction traces with very high levels of friction and wear. Wear tests

conducted under identical conditions, with paraffin oil and water as

lubricants, produced lower friction and wear. The friction traces for the

water lubricated alumina tests were smooth and resulted in steady state

coefficient of friction levels of 0.31, representing a friction reduction

of 50% below that of the unlubricated case. Wear was reduced by a factor

of three as measured by the wear scar diameter.

149
Examination of the water lubricated worn surfaces using scanning

electron microscopy (SEM) (Figure 8) indicated the presence of a "film" in

the contact region. The film had the appearance of a coherent compact of

material with the ability to flow plastically and be smeared out in various

regions of the contact. Further analysis of the film using FTIR indicated

the presence of 0-H bonds. These observations suggested that tribochemical

reactions might be taking place between the water and the alumina.

Tomizawa and Fischer previously had found tribochemical reaction between

water and "inert" silicon nitride^.

To elucidate the nature of the tribochemical reaction products, tests

were performed to react water with two alumina samples: alpha and gamma

alumina. Low temperature tests were conducted at room pressure and 100°C

for 24 hours . Higher temperature tests were conducted in a bomb reactor at

approximately 1.54 MPa (223 psia) and 200°C for 60 minutes. At the end of

the test, the powders were analyzed for reaction products using

thermogravimetric analysis (TGA) and x-ray powder diffraction (XRPD)

.

Alpha alumina showed no reaction at either temperature. Gamma

alumina, however, reacted at both temperatures. The TGA temperature scans

of gamma alumina and the two reaction products of gamma alumina and water

are presented in Figure 9. The low temperature reaction product thermally

decomposed at 272°C and was identified using XRPD as bayerite [Al(0H)3]

.

The higher temperature reaction produced a material that decomposed at

higher temperatures (510°C) and was identified using XRPD as boehmite

[AlO(OH) ]

.

It has been shown that chemical reactions occur between gamma alumina

and water (and at reasonable rates), however, no reaction was observed

between alpha alumina and water. A transformation of alpha alumina to

150
delta alumina was observed by Mines, Bradt , and Biggers . They found that

when alpha alumina is subjected to abrasion in unlubricated wear tests,

particles of delta alumina were formed. Delta alumina is one of the

transition alumina's and in many cases is very difficult to distinguish

from gamma alumina.

Analysis of the wear debris was conducted to provide direct evidence

of the nature of the tribochemical reaction products. Wear tests were

conducted using a specially prepared gold coated test cup to ensure that no

water- iron interactions would interfere with the analyses. Small amounts

of debris were painstakingly collected and analyzed using XRPD and TGA.

The x-ray powder diffraction (XRPD) spectrum indicated that most of the

wear debris was composed of alpha alumina, however, a small amount of

material was observed that appeared to be bayerite. Thermogravimetric

analysis (TGA) showed a large weight loss at 249 °C and a smaller weight

loss at 325°C. These decomposition temperatures are consistent with those

of the trihydroxides and, taken with the other analyses, indicate that

bayerite is present in the wear debris.

Given all of the data, the following picture emerges for the

tribochemical mechanism of water lubricated alumina. High temperatures and

shear stresses in the contact junction produce a phase transformation from

alpha to transition alumina. Transition alumina reacts with water at high

pressure and temperature to form an aluminum hydroxide. Wear debris

analysis, however, indicated that bayerite is present at the end of the

test. It appears, therefore, that boehmite transforms to bayerite as it

emerges from the high pressure, high temperature environment of the wearing

contact to the low pressure low temperature environment outside of the wear

151
scar. The result is a layer lattice material that lubricates the contact

j unction

.

Conclusions

Sample cleaning procedures have been demonstrated to alter the results

of unlubricated friction tests conducted with dense polycrystalline

alumina. A simple cleaning procedure is offered that uses organic solvents

to remove residual polishing oils and greases followed by a detergent

solution to remove residual organics. A series of rinses with high purity

deionized water then removes residual detergent and results in relatively

clean surfaces. The samples should be baked out at a moderately high

temperature to remove moisture or oxidize any remaining contaminant.

It has also been demonstrated how material properties of elastic

modulus and thermal conductivity affect the test severity with regard to

surface stress and temperature, respectively.

Wear tests conducted with water on alumina showed significant friction

and wear reduction. Detailed analysis of the wear debris and coupled with

static reaction experiments have resulted in the following sequence of

events

:

1. Contact conditions induce a phase transformation from alpha to

transition alumina.

2. Chemical reaction occurs between water and transition alumina

under high temperatures and pressures to produce aluminum

hydroxides

.

3. These aluminum hydroxides are layer lattice structures and serve

to lubricate the contact junction.

152
References

1. Kingery, W. D. , Bowen, H. K. , and Uhlmann, D. R. , Introduction to

Ceramics . 2nd Edition, John Wiley & Sons, New York, 1976.

2. Munro , R. G., "Temperature Considerations in the Study of Surfaces

Using a Four-Ball Apparatus," J. Appl. Phys . , 57, 11, p. 4950-4953.,

June 1, 1985.

3. Steijn, R. P., "Sliding and Wear in Ionic Crystals," J. Appl. Phys.,

34, (2), pp. 419-429 (1963).

4. Aronov, V., andMesyef, T., Trans. ASME , 108, PP • 16-21 (1986).

5. Scott, H. G., "Friction and Wear of Zirconia at Very Low Sliding

Speeds," Wear of Materials, pp. 8-12 (1985).

6. Miyoshi, K. , Buckley, D. H., and Srinivasan, M. , "Tribological

Properties of Sintered Polycrystalline and Single-Crystal Silicon

Carbide," American Ceramic Society Bulletin, 62., (4), pp. 494-500

(1983)

7. Mines, J. E., Jr., Bradt , R. C, and Biggers, J. V., "Delta Alumina

Formation during the Abrasive Wear of Polycrystalline Alumina," Wear

of Materials 1979, pp. 540-550.

153
Table 1. Residual Contaminant (Carbon Containing) on Alumina after

Various Cleaning Procedures as Measured by ESCA

Contaminant Level

as Measured by c/o Peak Ratios

Number Cleaning Procedure As-Prepared Sputtered

a Hexane-Toulene/

Acetone/DI

0.94 0.92

b Hexane/Pyridine 0.58 0.20

c Hexane-Toulene/

Acetone/Micro/DI

0.19 0. 14

d Micro/DI/Bake at 0.17 0.05

e Hexane - Toulene

/

Acetone/DI + Bake

at 693°C for

1 1/2 Hours

0.11 0.06

Each solvent was used by immersing the specimens in them and

agitating ultrasonically for approximately 1 minute.

DI - 18 mO, resistance deionized water.

Micro = the laboratory detergent utilized (2% in DI water).

154
CERAMIC

(BRITTLE)

METAL

(DUCTILE)

Failure

Strain (i)

Fig. 1 Typical stress-strain behavior for metals and ceramics at room temperature

155
UJ

DCW

W

UJ

DC

X

<

MPa

5,000

4,000

3,000 -

2,000

1,000 -

LOAD, kg

Fig. 2 Comparison of calculated contact pressures for 52100 steel vs

,

dense polycrystalline alumina

156
500 -

0 150 300 450

TIME, s

Source : Munro (6)

Fig. 3 Calculated contact junction temperature as a function of time

for lubricated alumina and steel four-ball tests

157
I

CDO

I—o

U-co

600

500

400

300 -

200

100 -

— COF 0.643

COF : 0.522

COF 0 168

Conditions : 600 rpm, 0.25 l/min dry air, 10 min duration, alumina specimens

Fig. 4 Friction traces for unlubricated alumina four-ball tests at

various loads

158
Fig. 5 Typical ESCA analysis of an alumina surface

159
0.7

Pyridine/Dl/Micro/DI/580 °C Bake Out

Acetone/He X ane

Fig. 6 Effect of cleaning procedure on the coefficient of friction for

"unlubricated" alumina on alumina tests

160
LOAD, kg

Fig. 7 Comparison of wear as a function of load for unlubricated

,

paraffin oil lubricated, and water lubricated alumina constant'

condition tests

161
10 0 ^Jm

Fig. 8 SEM photomicrograph of lower specimen wear scar water- lubricated

four-ball test

162
Fig. 9 TGA analysis of gamma alvimina and gamma alumina/water

reaction products

163

NBS-n4A IREV. 2-8C1 U.S. DEPT. OF COMM. BIBLIOGRAPHIC DATA

SHEET (See in struct/on sj

1. PUBLICATION OR

REPORT NO.

NIST/SP-766

2. Performing Organ. Report No. 3. Publication Date May 1989

4. TITLE AND SUBTITLE

NBS/BAM 1986 Symposium on Advanced Ceramics

5. AUTHOR(S)

S. M. Hsu and H. Czichos, Editors

6. PERFORMING ORGANIZATION (If joint or other than NBS, see instructions)

NATIONAL INSTITUTE OF STANDARDS AND TECHNOLOGY

{lormerly NATIONAL BUREAU OF STANDARDS)

U.S. DEPARTMENT OF COMMERCE

QAITHERSBURG, MD 20899

7. Contract/Grant No.

8. Type of Report & Period Covered

Final

9. SPONSORING ORGANIZATION NAME AND COMPLETE ADDRESS (Street. City. State, ZIP) same as item 6

10. SUPPLEMENTARY NOTES

Library of Congress Catalog Card Number: 89-600737

I I Document describes a computer program; SF-185, FlPS Software Summary, is attached.

11. ABSTRACT (A 200-word or less factual summary of most si gnificant information. If document includes a si gnifleant bibliography or literature survey, mention it here) Advanced ceramics offer many advantages that other materials do not

possess. They have high strength, dimensional stability, are chemically

inert, lightweight, wear resistant, and have desirable properties in

electrical, optical, and thermal applications. At high temperatures, they

are the only class of material with reasonable properties. As such,

advanced ceramics is increasingly used in electronic, structural, wear, optical, and electrical applications and devices.

Worldwide production of advanced ceramics is growing rapidly. Since

ceramics are based on alumina and silica, the most abundant minerals on

earth, effective utilization of ceramics carries implications into the next several centuries.

One of the major technical barriers to widespread use of ceramics is the

inability of industry to manufacture reliable ceramics reproducibly and

economically. Advanced ceramics are sensitive to small defects introduced

during processing and generated during use. The identification of key

parameters and subsequent measurements of these parameters are crucial to control the quality of ceramics. The National Institute of Standards and

Technology (NIST) (formerly National Bureau of Standards, NBS), and the

Bundesanstalt fur Materialprufung (BAM) to a large extent, share a similar

mission to assist their respective industries to achieve reliable

manufacturing through standards. This symposium therefore provides timely

exchange of technical information on a very significant subject area.

12. KEY WORDS (Six to twelve entries; alphabetical order; capitalize only proper names; and separate key words by semicolons)

BAM; ceramics; fretting, NIST; powder; sintering; technical exchange; wear

13. AVAILABILITY

Unlimited

I I For Official Distribution. Do Not Release to NTIS

Fy] Order From Superintendent of Documents, U.S. Government Printing Office, Washington, D.C. 20402.

QI] Order From N?-ional Teclinical Information Service (NTIS), Springfield, VA. 22161

14. NO. OF

PRINTED PAGES

159

15. Price

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