N I S T 

r

e

c

o

m

m

e

n

d

e

d

p r a c t i c e g u i d e 

Fractography of

Ceramics and 

Glasses 

George D. Quinn 

This publication is available free of charge from: 

http://dx.doi.org/10.6028/NIST.SP.960-16e2 

Special 

Publication 

960-16e2 

 

 

 

NIST Recommended Practice Guide 

Special Publication 960-16e2 

Fractography of 

Ceramics and Glasses 

George D. Quinn 

Materials Measurement Laboratory 

This publication is available free of charge from: 

http://dx.doi.org/10.6028/NIST.SP.960-16e2 

May 2016 

U.S. Department of Commerce 

Penny Pritzker, Secretary 

National Institute of Standards and Technology 

Willie May, Under Secretary of Commerce for Standards and Technology and Director 
 

 

 

 

 

 

Certain commercial entities, equipment, or materials may be identified in this 

document in order to describe an experimental procedure or concept 

adequately. Such identification is not intended to imply recommendation or 

endorsement by the National Institute of Standards and Technology, nor is it 

intended to imply that the entities, materials, or equipment are necessarily the 

best available for the purpose. 

National Institute of Standards and Technology 

Special Publication 960-16e2 

Natl. Inst. Stand. Technol. 

Spec. Publ. 960-16e2 

640 pages (May 2016) 

CODEN: NSPUE2 

This publication is free of charge from: 

http://dx.doi.org/10.6028/NIST.SP.960-16e2 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Preface 

Preface to the Second Edition 

Fractography is a powerful but underutilized tool for the analysis of fractured 

glasses and ceramics. It is applicable to fractures created under controlled 

conditions in the laboratory and to component failures in service. Fractography 

can identify the cause of failure and can even provide quantitative information 

about the loading conditions. 

The goal of this Guide is to make fractographic analysis of brittle materials less 

an art and more an engineering practice for scientists and engineers. This guide 

emphasizes practical approaches for problem solving and failure analyses. 

Fractographic analysis is to large degree pattern recognition. This Guide 

includes a wealth of illustrations to aid fractographers in pattern recognition. 

The Guide also includes an extensive bibliography and a tabulation of 

published case studies so that the reader can pursue topics of specific interest. 

This Guide is a free public service of NIST for the benefit of American 

industry, and the worldwide ceramic and glass communities. The demand for 

this book has been very strong since it was first published in 2007. The first 

print run of 3,000 copies was distributed around the world and the feedback has 

been very positive. Users say that they appreciate the content, the careful 

layout and organization, and the wealth of colorful illustrations and examples. 

This book is often used as a supplement to courses on fractography of ceramics 

and glasses. I am grateful to the National Institute of Standards and 

Technology for printing the book on high-quality paper and in color. Although 

the book is also available electronically, most users prefer the printed form so 

that it is handy on their desks or laboratories 

I became interested in the history of fractography of brittle materials while 

writing the first edition of this Guide in 2003 to 2007. Chapter 1 lists review 

articles that I wrote about the history of fractography after the first edition was 

printed. I learned much while doing these. For example, in the conclusions in 

his 1942 Journal of Applied Physics paper,23 Frank Preston made a profound 

statement about whether pieces from different fractured parts would match, and 

he also summarized the art of fractographic analysis: 

“From the beginning of time to its end, no two cracked surfaces will fit 

each other, unless originally they were part of the same piece. Cracks are 

as distinctive as fingerprints. ... Now the important thing about cracks is 

that they must be propagated. They do not originate all over the final 

fracture surface, but at one tiny spot, and from thence are propagated out 

iii 
 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

to the rest of the area. This fact and the fact that the telltale marks or 

“fingerprints” of the process are left on the surface, provide the 

groundwork for the subject of Fracture Diagnosis, which in the last 

decade has become an important minor branch of physical science.” 

This guide emphasizes practical approaches for problem solving and failure 

analyses, but it also distills the state of the art in our “important minor branch 

of physical science.” I get gratification from the free dissemination of all the 

know-how developed over decades by many conscientious researchers, 

materials scientists, and engineers. I salute them and have done my best to 

credit their contributions by liberal use of citations and the extensive 

bibliographies in this Guide. 

This edition is about 15% larger than the first edition. It corrects some of the 

errors, including the description of vicinal illumination. There are 300 new 

illustrations; many of which are from my colleagues around the world. All 

told, there are approximately 1000 illustrations and tables in this second 

edition. In Chapter 1, Introduction, and Chapter 2, Resources, I added a little 

about the History of Fractography and the new documentary standards. The 

fracture mirror size procedures in Appendix D were adopted in an ASTM 

standard 1678 in 2009, and the text in this Guide has been adjusted 

accordingly. Chapter 3 on Equipment has updated microscopy techniques, a 

rapidly evolving field. A major correction was made to the description of 

“vicinal illumination,” which was described erroneously in the first edition. 

Chapter 4 on Patterns has new material on sketches and drawings, edge 

chipping, thermal stresses and thermal shock, compression fracture, and 

laboratory-scale tests on medicinal vials and dental crowns. Chapter 5 on 

Fracture Surfaces has new information on mechanical fatigue, slow crack 

growth mechanisms, scarps, glue chips, and a new marking: “step hackle.” The 

term “microstructural hackle” has been refined. Chapter 6, Origins, has 28 

additional pages with dozens of new figures, and new information on bubbles, 

processing and firing flaws, cusps and geometric sharp points. There is a new 

definition of “Griffith flaw.” Chapter 7, Quantitative Analysis, has a number of 

additions including a figure showing flow lines of stress around bubbles or 

flaws; new text about “far-field stress;” new text about the accuracy of the 

Newman Raju stress intensity factors and how to extend them to a/c >1; and 

Bansal’s approximation. Chapter 9, Composites, has new examples of highlyfilled resin matrix composite examples from the dental materials field. Strictly 

speaking, they are not ceramic or glasses, but they behave brittlely and have 

many of the same fractographic features. There are several new cases studies 

in Chapter 10, including an AlON coarse-grained IR dome; ceramic dental 

crowns, and a single crystal sapphire spacecraft refractor. Chapter 11, 

iv 
 

 

 

 

 

 

 

 

 

 

Preface 

Conclusions, has a new paragraph about the revolution in microscopy as well 

as some thoughts about the fractographer as an expert system. A new beautiful 

fractograph from Dr. J. Quinn is added to the conclusions. 

Acknowledgements 

The author derived much inspiration from the late Professor Van Derck 

Fréchette of Alfred University. Van made many important contributions to the 

science of fractography of brittle materials, not the least of which was his 1990 

seminal book: Failure Analysis of Brittle Materials. Van solved many 

practical, vexing, and spectacular cases and trained hundreds of fractographers 

over his long and distinguished career. 

The author has had a continuing and fruitful collaboration with Professor James 

Varner at Alfred University. Dr. Varner and the author continue the Van 

Fréchette legacy by training a new generation of fractographers. 

The author thanks Dr. Jeffrey Swab, of the U.S. Army Research Laboratory; 

Aberdeen, Maryland and Mr. Michael J. Slavin (formerly of ARL) for their 

help in crafting a framework for fractographic analysis that was used to create 

the U.S. Army Military Handbook (MIL HDBK) 790 in 1990. This MIL 

HBDK was the precursor to American Society for Testing and Materials 

standard (ASTM) practice C 1332 adopted in 1996. This was the first formal 

standard for fractographic analysis of ceramics and brittle materials in the 

world. Many of the concepts in this Guide stem from our collaboration on 

those standards. Dr. Swab and I learned much about how fractographers work 

from a Versailles Advanced Materials and Standards round robin on ceramic 

fractography that we organized in 1993-1994. 

The author has learned much from discussions with and readings of the prolific 

writings of Mr. Roy Rice, Prof. Jack Mecholsky, and Dr. Stephen Freiman. 

They have significantly advanced the state of the art of materials science 

through their fractographic work. I note with sadness the passing of Mr. Roy 

Rice in 2011. He was one of the most respected fractographers in the world 

and from whom I learned much. His books on mechanical properties and 

microstructure of ceramics and composites are very valuable. Roy gave me 

many of the single crystal fractographs in Chapter 8 of this Guide. 

Mr. Are Tsirk, an enthusiastic knapper who introduced many of us to the 

fracture of lithics, also succumbed to a long illness in February 2015 within 

weeks of the publication of his own book on fractures and fractography of 

v 
 

 

 

 

 

 Fractography of Ceramics and Glasses 

knapping. He gave me some marvelous photos of scarps and parabolic wake 

hackle images in his final months. We will miss his knapping demonstrations 

at fractography conferences. 

Dr. Roger Morrell of the National Physical Laboratory has been an inspiration, 

particularly with his astute practical applications of fractography to failure 

analysis and his own Guide to Fractography. 

I also acknowledge the insightful work of Dr. Janet B. Quinn formerly of the 

American Dental Association, who made her own contributions to the science 

of brittle materials fractography and sometimes presented challenging 

fractographic problems to her spouse. About one year after the first edition 

appeared in 2007, just in time for her own course on Dental Materials 

Fractography, she suddenly became ill and died. At the time, she was applying 

fractography to ceramic, glass-ceramic, and composite dental restorations such 

as crowns and bridges. Janet and her dear friend Dr. Susanne Scherrer of the 

University of Geneva have taught a new generation of dental materials 

fractographers around the world how to use fractography effectively. 

vi 
Table of Contents 

vii
 
 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

TABLE OF CONTENTS 

Preface 

1.	 Introduction 1-1
 

2.	 Resources 2-1
 

3.	 Tools and Equipment 3-1
 

3.1 Low Power Optical Examination and Component Fracture
 

Documentation 3-2
 

3.2	 Photographing the Overall Component: Cameras, Cellphones,
 

and Simple Digital Microscopes 3-7
 

3.3 Stereo Binocular Microscope 3-11
 

3.4 Stereoptical Microscope Camera Port and Digital Cameras for 

Microscopes 3-14
 

3.5	 Digital image Processing and Focus Stacking 3-17
 

3.6	 Discussion Stereomicroscope 3-19
 

3.7	 Digital Microscopes 3-19
 

3.8	 Illumination Sources 3-20
 

3.9	 Coatings 3-26
 

3.10	 Size Measurements 3-30
 

3.11	 Compound Optical Microscope 3-32
 

3.12	 Replication Equipment 3-35
 

3.13	 Scanning Electron Microscope 3-43
 

3.14	 Stereo SEM Imaging 3-56
 

3.15	 Field Emission SEM 3-58
 

3.16	 Environmental SEM 3-59
 

3.17	 Transmission Electron Microscope 3-59
 

3.18	 Atomic Force Microscope (AFM) 3-60
 

3.19	 Optical Profilometer 3-60
 

3.20	 Confocal Scanning Light Microscope 3-62
 

3.21	 Stress Wave and Ultrasonic Fractography 3-64
 

3.22	 High Speed Photography 3-64
 

3.23	 X-ray Topography 3-65
 

3.24	 Dye Penetration and Staining 3-66
 

3.25	 Polariscope 3-70
 

3.26	 Other Equipment 3-71
 

3.27	 Other Resources 3-72
 

3.28	 The Future 3-72
 

viii 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Table of Contents 

4. General examination and fracture patterns 4-1
 

4.1 Specimen Reconstruction 4-2
 

4.2 Crack Branching Patterns 4-3
 

4.3 Crack Branching Angles 4-6
 

4.4 Crack Branching Distances 4-9
 

4.5 Fragmentation Patterns 4-10
 

4.6 Laboratory Test Fracture Patterns 4-11
 

4.6.1 Tension strength 4-13
 

4.6.2 Flexural strength 4-14
 

4.6.3 Biaxial flexural strength 4-19
 

4.6.4 Laboratory tests: additional tips 4-22
 

4.7 Component Fracture Patterns – General 4-22
 

4.8 T Intersections and Crossing Cracks 4-26
 

4.9 Invisible cracks 4-28
 

4.10 Plates and Windows 4-28
 

4.11 Tempered Windows 4-30
 

4.12 Thermally Induced Plate and Window Fractures 4-32
 

4.13 Bottles and Pressure Vessels 4-33
 

4.14 Torsional Fractures 4-39
 

4.15 Chipping 4-40
 

4.16 Laboratory Test Fractures and Component Fractures 4-44
 

4.17 Controlled Component Fractures 4-44
 

4.18 Finite Element Analysis 4-46
 

4.19 Characteristics of Some Common Fracture Modes 4-47
 

4.19.1 Mechanical overload fracture 4-47
 

4.19.2 Thermal fracture and thermal shock 4-47
 

4.19.3 Impact or contact fractures 4-49
 

4.19.4 Corrosion or oxidation 4-49
 

4.19.5 Residual stresses 4-49
 

4.19.6 Time dependent fracture 4-51
 

4.19.7 Compression fracture 4-51
 

5. Fracture Surface Examination 5-1
 

5.1 Introduction 5-1
 

5.2 Fracture Mirrors, A Fractographer’s Best Friend 5-3
 

5.2.1 Fracture mirrors in glass 5-3
 

5.2.2 Fracture mirrors in ceramics 5-16
 

5.2.3 Fracture mirrors, special cases 5-20
 

5.3 Hackle 5-31
 

5.3.1 Coarse hackle 5-31
 

5.3.2 Wake hackle 5-33
 

5.3.3 Twist hackle 5-36
 

ix 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

5.3.4 Shear hackle 5-42
 

5.3.5 Corner hackle 5-43
 

5.3.6 Step hackle 5-43
 

5.4 Wallner Lines 5-45
 

5.4.1 Introduction 5-45
 

5.4.2 Primary Wallner lines 5-47
 

5.4.3 Secondary Wallner lines 5-50
 

5.4.4 Tertiary Wallner lines 5-53
 

5.4.5 Wallner Lambda lines 5-58
 

5.5 Arrest Lines 5-58
 

5.6 Scarps 5-61
 

5.7 Glue Chips 5-63
 

5.8 Transgranular and Intergranular Fracture 5-66
 

5.9 Stable Crack Growth 5-69
 

5.9.1 Slow crack growth (SCG) at ambient temperature 5-69
 

5.9.2 Stable extension from local residual stress or R-curve effects 5-73
 

5.9.3 High temperature slow crack growth in ceramics 5-74
 

5.9.4 High temperature creep fracture 5-77
 

5.9.5 Mechanical fatigue from cyclic loading 5-81
 

5.10 Organizing all the Information -Fractographic Montages 5-84
 

6. Origins of Fracture 

6.1 Origins, Flaws, and Defects 6-1
 

6.2 The Spatial Distribution of Flaws 6-1
 

6.3 Are Flaws Intrinsic or Extrinsic? 6-4
 

6.4 Matching Fracture Halves 6-5
 

6.5 External Surfaces 6-6
 

6.6 Volume-Distributed Flaws 6-7
 

6.6.1 Pores and Bubbles 6-7
 

6.6.2 Porous regions 6-14
 

6.6.3 Porous seams 6-14
 

6.6.4 Agglomerates 6-15
 

6.6.5 Inclusions 6-17
 

6.6.6 Compositional inhomogeneities 6-21
 

6.6.7 Large grains 6-22
 

6.6.8 Grain boundaries 6-25
 

6.7 Surface-Distributed Flaws 6-27
 

6.7.1 Surface voids or pits from processing 6-27
 

6.7.2 Pits from environmental exposure 6-28
 

6.7.3 Handling scratches and polishing scratches 6-29
 

6.7.4 Sharp object impact or contact 6-33
 

6.7.5 Blunt object impact or contact 6-34
 

x 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Table of Contents 

6.7.6 Grinding and machining cracks	 6-40
 

6.7.7 Chips	 6-53
 

6.8 Other Flaws	 6-54
 

6.8.1 Processing and firing cracks in ceramics	 6-54
 

6.8.2 Artificial or controlled flaws, glasses and ceramics 6-59
 

6.8.3 Other glass origins	 6-61
 

6.9 Fracture Oddities	 6-69
 

6.10 Contaminants	 6-71
 

6.11 Combined or Hybrid Flaws	 6-73
 

6.12 Baseline Microstructure Origins	 6-75
 

6.13 Geometric Sharp points or Cusps	 6-76
 

6.14 Partially-Exposed Flaws	 6-77
 

6.15 Glaze flaws	 6-78
 

6.16 Microfault Pockets 6-80
 

6.17 Polished Microstructural Sections 6-82
 

6.18 Flaw Size	 6-84
 

6.19 What to Report	 6-84
 

7. Quantitative Analysis	 7-1
 

7.1 Introduction	 7-1
 

7.2 Stresses from the Extent of Fragmentation 7-2
 

7.2.1 Annealed parts	 7-2
 

7.2.2 Fragmentation of tempered parts	 7-5
 

7.3 Stresses from the Branching Distances	 7-7
 

7.3.1 Annealed plates and bars	 7-7
 

7.3.2 Tempered plates	 7-8
 

7.4 Stresses from the Mirror Size	 7-10
 

7.4.1 Mirror analysis and constants	 7-10
 

7.4.2 How to measure mirror sizes	 7-13
 

7.5 Fracture Mechanics Analysis of the Flaw Size	 7-17
 

7.5.1	 Introduction to fracture mechanics 7-17
 

7.5.2	 The Newman-Raju Y factors for semielliptical surface flaws in
 

bending 7-30
 

7.5.2.1	 The Newman-Raju formulas 7-30
 

7.5.2.2	 Graphical curves for the Newman-Raju formulas 7-35
 

7.5.2.3	 Newman-Raju formulas for deep (a/c > 1) semielliptical
 

surface cracks 7-36
 

7.5.2.4	 Accuracy of the Newman-Raju formulas and alternatives 7-36
 

7.5.2.5	 Bansal’s approximation for semielliptical surface flaws 7-37
 

7.5.3	 Irregularly-shaped surface cracks 7-38
 

7.5.4	 Three-dimensional, blunt, and inclined flaws 7-39
 

7.6 Relationship of KIc and A	 7-40
 

xi 
 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

7.7 Mirror to Flaw Size Ratios	 7-41
 

7.8 Comparing Measured to Calculated Flaw Sizes	 7-43
 

7.9 Crack Velocities from Wallner Line Analysis	 7-46
 

7.10 Slow Crack Growth	 7-51
 

7.11 R- curve Behavior	 7-55
 

7.12 Indentation Mechanics	 7-61
 

7.13 Fractal Analysis	 7-65
 

7.14 Estimation of Residual Stresses	 7-68
 

7.14.1 Introduction	 7-68
 

7.14.2	 Estimates of residual stresses from fragmentation 7-68
 

7.14.3	 Estimates of residual stresses from fracture mechanics
 

analysis of flaws 7-68
 

7.14.4	 Estimates of residual stresses from fracture mirror size
 

analysis 7-69
 

7.14.5	 Estimates of residual stresses from indention crack sizes 7-70
 

7.14.6 Estimates of residual stresses by other means 7-70
 

7.15 Weibull Analysis	 7-71
 

8. Single Crystals	 8-1
 

8.1 General	 8-1
 

8.2 Preferred Cleavage Planes	 8-2
 

8.3 Fractographic Techniques	 8-5
 

8.4 Fracture Surface Markings	 8-8
 

8.5 Origins	 8-21
 

8.6 Other Sources	 8-32
 

9. Ceramic and Glass Composites	 9-1
 

9.1 Particulate, Whisker, or Self-Reinforced Ceramic Composites 9-1
 

9.2 Fiber-Reinforced Composites	 9-2
 

9.3 Dental composites	 9-6
 

10. Case Studies 10-1
 

Case 3 Conflicting Carbide Data (Silicon carbide flaws and
 

Case 9 Modeler’s Match (Fracture origins in MEMS scale
 

Case 1 Ruptured Rotor (Ceramic gas turbine rotor) 10-1
 

Case 2 Busted Barrel (Silicon carbide machine gun liner) 10-7
 

slow crack growth) 10-11
 

Case 4 Vulnerable Vials (Broken medicinal vials) 10-14
 

Case 5 Damaged Dome (AlON IR dome) 10-16
 

Case 6 Suffering Setter Plate (Silicon carbide furnace plate) 10-18
 

Case 7 Ruptured Radomes (Fused silica missile nosecones) 10-20
 

Case 8 Maligned Machinists (Bend bars made by different shops) 10-23
 

xii 
 

 

 

 

 

Table of Contents 

SiC micro tensile specimens) 

Case 10 Fractious Fractographers (A fractography round robin) 

Case 11 Perilous Prostheses (Four ceramic dental crowns) 

Case 12 Mangled Margins (Twenty dental crowns) 

Case 13 Ruined Refractors 

10-25 

10-28 

10-30 

10-43 

10-46 

11. Conclusions 11-1 

Appendix A Bibliography 

Appendix B Fractographic Case Studies 

Appendix C Fracture Mirror and Branching Constants 

Appendix D Guidelines for Measuring Fracture Mirrors 

Index 

xiii 

 

 

 

 

 

 

 

 

 

Introduction 

1. Introduction 

Why did it break? What was the origin of failure? Did it break as expected or 

from an unexpected cause? Was there a problem with the material or was the 

part simply overloaded or misused? Why did this part break and others not? 

What was the stress at fracture? Was the laboratory strength test successful or 

was there a misalignment? These are common practical questions and the 

fractographer can often give straightforward definitive answers. 

The curse of brittle materials is that they are prone to catastrophic fracture. 

Brittle fracture is fracture that takes place with little or no plastic deformation. 

Nature has partially compensated for this shortcoming by furnishing clear 

fracture patterns and fracture surface markings that provide a wealth of 

interpretable information. Indeed, in many respects, fractographic analysis of 

ceramics and glasses is easier and can produce more quantitative information 

than the fractographic analysis of metals or polymers. 

Figure 1.1 shows broken glass and ceramic rods. Using the techniques 

described in this Guide, fractographers are able to determine that the rods were 

broken in bending, from surface flaws, and even determine the fracture stress. 

(a) (b) 

Figure 1.1 Fractured glass (a) and ceramic (b) rods. 

This Guide takes a broader view of fractography than merely the examination 

of the fracture surfaces. The author wrote the definition that is in ASTM 

standard 13221

: 

fractography, n – means and methods for characterizing fractured 

specimens or components 

1-1 
 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Fréchette2

 also took broad view and described the science of fractography as 

the study of fragments and their interpretation in terms of material properties 

and conditions leading to failure. 

Examination of the fracture surfaces of broken pieces is an important element 

of fractographic analysis, but much can be gleaned by first looking at the sizes, 

shapes, and breakage patterns of the fragments. In some cases, the breakage 

pattern is all that is necessary and examination of fracture surfaces is 

unnecessary. For example, even novices can recognize an impact site on an 

automobile windshield without examining the fracture surfaces. A simple 

visual examination can tell the observer whether a component fracture was 

thermally or mechanically driven, whether the stress was large or small, and 

whether the stresses were uniaxial or multiaxial. 

It is surprising how many instances there are in the literature where undue 

attention was spent on small fracture surface regions that were probably not 

even in the area of fracture initiation. It behooves one to look at the overall 

breakage pattern first, before one jumps to a microscopic examination of the 

fracture surfaces. 

Most people recognize that fractography is a valuable tool for failure analysis, 

but fewer appreciate its value in routine mechanical testing or support of 

materials processing. Although thousands of ceramic or glass items and test 

specimens are broken daily, only a tiny fraction is examined fractographically. 

A wealth of information is lost about the causes of fracture and the nature of 

the material. The flaw type may be just as important as the fracture stress in a 

strength test. Rice rued this state of affairs in 1977:3 

"The most significant experimental procedure that can aid the understanding of 

mechanical properties is a study of fracture surfaces, especially to identify 

fracture origins.... It is indeed amazing the number of mechanical properties 

studies conducted that were extensively concerned directly or indirectly with 

the size and character of flaws and microstructure from which failure 

originated in which no attempt was made to experimentally observe and verify 

the predicted or implied flaw character." 

Some of the reluctance to apply fractographic analysis is that it is unfamiliar to 

most engineers and scientists. Although there are a few notable exceptions, 

fractography is not commonly taught in schools and is often learned gradually 

and autodidactically by experience over many years. Some regard 

fractographic analysis as a subjective practice that can only be applied by 

masters. Others deem it as too interpretive. The reality is that what may seem 

1-2 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Introduction 

mysterious to some is in fact objective and quantitative to an experienced 

fractographer. a Interpretation is an essential step, but brittle materials often 

leave unequivocal markings on fracture surfaces that even a novice can 

interpret with confidence. The markings are the direct consequence of crack 

perturbations during propagation. An important element of fractographic 

analysis is pattern recognition. Certain types of fracture leave telltale fracture 

patterns on the fracture surfaces, or have telltale breakage patterns or fragment 

shapes. For example, hackle lines radiating from a fracture mirror are features 

that even novices can identify. These lines lead an observer back to the origin. 

Other markings may be more subtle and can be overlooked by the casual or 

inexperienced observer. With a little practice and experience, any material 

scientist or engineer should be able to analyze brittle fractures and make an 

interpretation, or at least recognize what pieces should be brought to the 

attention of a more experienced fractographer. The novice may be tentative at 

first, but fractography is a cumulative learning experience. The more 

fractography one does, the easier it becomes. 

A good starting place to gain experience is by examining test coupons broken 

in laboratory conditions. Flexural strength specimens are a common starting 

point. Pattern recognition skills can be reinforced by seeing multiple examples 

in a set of specimens. Component failure analysis is often much more difficult, 

especially if only a single example is available or the fracture occurred under 

unknown loading conditions. 

What are the skills that a fractographer needs? Fractographers ideally should 

have knowledge of materials science, microscopy, mechanics (stresses and 

strains), and an aptitude for problem solving and troubleshooting. Puzzlesolving skills can be helpful as well as an ability to be able to sketch or draw as 

discussed in section 4.7. They should be comfortable with microscopy, since 

many of the features to be studied are smaller than the unaided eye can discern. 

Some knowledge of materials science is necessary since strength-limiting flaws 

and crack propagation behavior are controlled by processing and the 

microstructure. 

Cracks propagate in response to stresses and strains, so fractography often 

comes quite naturally to mechanical engineers. That is not to say that materials 

scientists, physicists and geologists do not make good fractographers, but they 

will have to learn the fundamentals of stress analysis. At a minimum, a 

a This is reminiscent of Arthur C. Clarke’s Third law, which is: “Any sufficiently 

advanced technology is indistinguishable from magic.” Profiles of the Future: An 

Inquiry into the Limits of the Possible, A. C. Clarke, Bantam, 1973. 

1-3 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

fractographer should be comfortable with concepts of uniaxial tensile stresses 

(stresses primarily in one direction, such as in a direct tensile strength test 

specimen), uniaxial bending stresses (such as in a beam loaded in bending), and 

biaxial stresses (whereby there are tensile stresses in two different directions 

such as in a pressurized flat plate or a pressurized bottle). Some important 

stress configurations are illustrated in Figure 1.2. 

σ1 

σ2 

σ2 

σ1 

τ

τ

σ1 

σ2 

σ1 

σ1 σ1 

σ2 

σ1 

σ2 

σ1 

σ2 

Uniaxial Biaxial Equibiaxial Torsion 

Tension Tension Tension σ1 = σ2 σ2 = -σ1 

Figure 1.2 Uniaxial or biaxial stress states may be represented by forces 

(arrows) acting on the surfaces of tiny differential volume elements. They are 

shown here as 3-dimensional (top row) or 2-dimensional (bottom row) views. 

The lengths of the arrows are in proportion to the stress magnitude. σ1 and σ2 

are normal stresses and τ are shear stresses. In some cases, the stress state is 

triaxial, and a third stress direction, σ3, may be added to the diagrams above. 

Most fracture problems the reader will encounter will be uniaxial or biaxial 

stress cases, but there are very important cases (projectile-impacted ceramic or 

glass armor plates, or localized Hertzian or sharp contact problems) where the 

stress state is triaxial; that is, there are stresses in all three directions, σ1, σ2, 

and σ3. Some of these are tension stresses that are denoted with a positive sign 

(e.g., +100 MPa), and some may be compressive stresses, denoted by a minus 

sign (e.g., -300 MPa). Mechanical engineers are very comfortable with 

multiaxial stress problems, and can convert the stresses in any arbitrary 

orientation at a given point to the specific orientation that has the principal 

stresses. (The X-Y-Z axes at any point can be rotated and are reoriented so that 

the small differential volume element has three normal stresses σ1, σ2, and σ3 

1-4 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Introduction 

and no shear stresses acting on its faces.) Mechanical engineers commonly use 

a Mohr’s circle analysis or finite element analysis to find this orientation and 

the maximum principal stresses. Maximum principal stresses usually control 

fracture in ceramics and glasses. Only in specialized cases do shear stresses 

cause fracture. This is an important difference in how ceramics and glasses 

break as compared to metals and polymers. 

It should also be borne in mind that most parts and laboratory specimens have 

stress gradients and may even have internal residual stresses. It is rare that a 

component has a uniform stress throughout. Simple loading configurations are 

typically covered in the first chapters of an elementary strength of materials 

engineering textbook, which may be consulted to gain a basic familiarity with 

stresses and strains. It is beyond the scope of this Guide to delve into 

multiaxial stress fracture criteria, but the reader should be aware that what 

causes fracture in a tensile-loaded ceramic part may be completely different 

than for a compressively-loaded ceramic part, or for a part with mixed tension 

and compression stresses. Engineers often use stress diagrams with fracture 

envelopes to illustrate fracture versus no-fracture boundaries in multiaxial 

stress problems. Mechanical engineers not versed in ceramic or glass behavior 

do not appreciate the nuances of multiaxial failure criteria for ceramics and 

glasses. For example, many finite element modelers who are new to the 

ceramic field erroneously use the Von Mises (or maximum distortion energy) 

criterion in their models. While this criterion works quite well for metals, it is 

completely wrong for ceramics and glasses. Many mechanical engineers are 

surprised at how different the Poisson’s ratios for glasses and ceramics are as 

compared to those for metals! 

There are two special stress conditions that warrant explanation. Mechanical 

engineers may say that parts are loaded in plane stress or plane strain 

conditions. Plane stress refers to a case where there are only two principal 

stresses, σ1 and σ2. The third stress σ3 = 0. In other words, all the stresses lie in 

a flat plane. This is commonly the case with thin parts, such as a thin-walled 

cylinders or pressurized bottles. The principal stresses are parallel to the 

surface. In contrast, with thick parts, there can be a σ3 internal stress. If the 

part is very thick, lateral thinning strains may be inhibited. The strain through 

the thickness, ε3 = 0. There are only two principal strains ε1 and ε2 and they are 

parallel to the surface. This is called the plane strain condition and it is 

especially important for fracture toughness test specimens. There are stringent 

requirements on the thickness of metallic fracture toughness specimens to 

ensure that they are in a plane strain condition. The matter is not important for 

ceramics and glasses, since the plastic yield zones in ceramics are tiny. 

1-5 
 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Part of the skill of a fractographer is knowing where to look and how to look. 

A specimen that breaks into only two pieces may not be too difficult to 

interpret. Medium to high strength specimens with multiple fractures can 

initially be confusing, but with a little experience and a few simple precautions, 

one can find the primary fracture. Component or service failures may be 

difficult since the loading conditions may be unknown. Key pieces may be 

missing or the evidence incomplete or contradictory. Multiple crack systems 

from different causes and events may intersect and confound an interpretation. 

The author suspects one reason that fractography is sometimes held with 

suspicion or is deemed subjective is that some incautious fractographers have 

jumped to conclusions and reached questionable interpretations. It is unwise to 

make conclusions based on limited or incomplete evidence. Novices often 

force fit a fracture scenario to their limited experience base. Some fracture 

patterns may be difficult to detect or be very subtle and may not be recognized. 

Good fractographers recognize their strengths and the limitations of their 

experience base and do not overreach or extrapolate. They should be alert to 

unusual or new fractographic markings or to failure modes with which they are 

not familiar and should be ready to search the literature, consult colleagues, or 

try to create comparable markings in the laboratory. No fractographer is born 

with a built-in data base of fractographic patterns in the brain, so step-by-step 

accumulation of experience is necessary. One may consult textbooks, reference 

articles, and this Guide to help acquire knowledge, but there is no substitute for 

hands-on direct eyeball and microscopy experience. 

There is a wealth of information in the technical literature about fractography 

of glasses and ceramics, but until the first edition of this book was written in 

2007, it was scattered in textbook chapters, in journals, and conference 

proceedings. Fréchette’s 1990 seminal book, Failure Analysis of Brittle 

Materials,

1 was a good starting point for the practical fractography of glasses. 

This Guide covers fractography of glasses and ceramics and presents practical 

information for both glasses and ceramics. It also includes an extensive 

bibliography. There now are formal standard practices in the American Society 

for Testing and Materials2

 and European Committee for Standardization4 for 

fractographic analyses of glasses and ceramics. 

Fractography is a tool for the broader topic of failure analysis. Figure 1.3 

suggests a simple but apt analogy. The fractographer is called upon to solve a 

fracture mystery. The detective uses his or her powers of observation to study 

the scene of the fracture and meticulously collects and preserves the available 

evidence. The clues are contemplated and weighed against the available 

background information furnished by witnesses. Comparable cases are 

1-6 
 

 

Introduction 

Figure 1.3 The fractographer as detective. 

considered to determine whether there are similarities and whether there is a 

pattern. Scientific analysis of the material may be done to verify its quality, 

composition, and authenticity. Pondering all the information, and keeping an 

open mind for all possible scenarios, the fractographer detective formulates a 

hypothesis and checks it against the known facts. The fractographer 

communicates the findings to the client, management, engineer or processor in 

a manner that is convincing and fathomable. Sometimes the findings must be 

presented in formal legal settings. In summary, fractographic analysis is not 

merely looking at fracture surfaces, but is the integration of knowledge from a 

variety of sources to solve the puzzle of how fracture occurred (Figure 1.4). 

Figure 1.4. Fractographic analysis is more than examining fracture 

surfaces.5 

1-7 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

One definition of failure analysis is: “A process that is performed in order to 

determine the causes or factors that have led to an undesired loss of function.”6 

Failure may be due to fracture, corrosion, excessive deformation, or wear. 

Hence, fracture is a subset of a larger class of failure causes. The readers 

interested in the broader topic of failure analysis in general, including concepts 

and philosophies about how to conduct an analysis, should refer to the ASM 

handbook on Failure Analysis and Prevention.6 That volume has multiple 

chapters and extensive discussions about failure analysis in general and also 

has such pragmatic recommendations as: “don’t jump to conclusions,” “keep 

an open mind,” and “avoid 5-minute best guess analyses.” 

This is a technical book. It is beyond the scope of this book to delve into issues 

of litigation, fault, or blame. Readers who are interested in these matters are 

referred to Chapter 11 “The Expert Witness” in Van Fréchette’s book,2

 an 

article by B. Adams in the second Alfred conference on fractography,

7

 an 

article by C. O. Smith in Volume 11 of the ASM handbook,8

 or a review by 

Bradt in the third Conference on Fractography of Advanced Ceramics in 

Slovakia.9 

A few words about nomenclature are appropriate here. Readers may be 

surprised that terms for ceramics and glasses are somewhat different than for 

other materials. For example, “hackle lines” and “lances” refer to steps or lines 

on the fracture surface running parallel to the direction of crack propagation. 

“Striae” or “striations” means completely different things to metallurgical 

fractographers, ceramic and glass fractographers, and glass manufacturers. In 

this Guide, I have adopted and added to the nomenclature system devised by 

Professor Van Fréchette since it is descriptive and logical. It was not without 

controversy when he introduced it formally in 1982 at a conference on 

fractography.10 Frechette’s concise paper was only three pages long, but was 

followed by a detailed three page commentary by R. Rice! I have resisted the 

temptation to devise new and colorful terms, since a proliferation of terms will 

only lead to confusion. Anne Roulin-Moloney, in the preface of her 

outstanding book on the fractography of polymers and composites,11 wrote 

about the “myriad of weird and wonderful terms to describe the features 

observed. Some examples are laceration, shish-kebab, cusps, welts, mist, and 

shards.” One cannot help but burst out laughing when one sees some of these 

terms. 

The history of fractography of brittle materials is a fascinating story of 

microscopy, material science, fracture mechanics, and people. Figure 1, 

adapted from my review articles,12,13 chronicles how the field evolved with 

1-8 
 

 

 

 

 

 

 

 

 

 

 

Introduction 

contributions from such luminaries as De Freminville, Preston, Griffith, 

Smekal, Schardin, Weibull, Wallner, Orr, Kerkhof, Shand, Ernsberger, Irwin, 

Fréchette, Kirchner, Rice, and others. 

1973 Kirchner 

2004 CEN 843-6 Ceramic 

Origins Morrell 

1992 MILHDBK 790 Origins 

Quinn, Swab, Slavin 

1993 ASTM C 1256 Glass 

Features DeMartino 

Digital image 

recording 

2008 ASTM C1678 

Fracture Mirrors Quinn 

1938 Schardin 

Varner 

1920 Griffith 

Scanning electron 

microscope 

1921 Preston 

1935 Smekal 

1970 Kerkhof’s 

book 

1943 Gölz 

1894 Brodmann 

1950 Kerkhof 

Optical microscopy Greenough 

stereoptical microscope 

1996 ASTM C1322 Ceramic 

Origins Quinn, Swab, Slavin 

1943 Murgatroyd 1939 Barstow and Edgarton 

1907 De Freminville 

1940 Orr 

1965 Fréchette 1960 Ernsberger 

Alfred Univ. 

1939 Weibull 

2016 Quinn’s book 

2nd edition 

1977 Alfred Univ. 

short course 

1958 Poncelet 

1990 Uematsu 

Danzer, Lube, Harrer 

2001 - Dusza Stara 

Lesna/Smolenice castle 

conference series 

Digital image 

processing, 

profilometry 

3-D imagery 

1974 Rice, Mecholsky, Freiman 

1954 Shand 

1939 Wallner 

1966 Johnson Holloway 

1984 ASTM STP 827 

Mecholsky, Powell 

2007 Quinn’s book 

1957 Irwin 

Transmission 

electron 

microscope 

1994 Bradt and 

Tressler’s book 

1990 

Fréchette’s 

book 

1999 Hull’s book 

1986-2011 Alfred 

conference series 

FRACTOGRAPHY MATERIALS SCIENCE MICROSCOPY STANDARDS 

1890 

1900 

1910 

1930 

1920 

1940 

1950 

1960 

1970 

1980 

1990 

2000 

2010 

Quinn 

1993 VAMAS Round Robin 

Ceramic Origins Swab Quinn 

Michalske 

Figure 1.5. The history of fractography of brittle materials 

1-9 
 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

At this point it is appropriate to point out a critical aspect of fracture analysis 

problems. Final fracture often occurs after a chain of events. One event may 

create a flaw, but then a second or even third event may be necessary to cause 

fracture. For example, drinking glasses commonly sustain internal impact 

damage from eating utensils that are carelessly dropped into the glass. 

Subsequent thermal stresses generated in the dishwasher may cause cracks to 

extend gradually around the base. If the large crack is not noticed, an unwary 

user who picks up a liquid filled glass may have an unpleasant surprise. What 

was the cause of failure in this case: The user pouring a liquid into the glass and 

lifting it to quench a thirst, the initial impact damage, or the thermal stresses 

that cause the crack to grow? The answer is all of the above. 

Chapter 1 References 

1.	 ASTM C 1322-15, “Standard Practice for Fractography and 

Characterization of Fracture Origins in Advanced Ceramics,” Annual Book 

of Standards, Vol. 15.01, ASTM Int., West Conshohocken, PA, 2015. (The 

original version was adopted in 1996.) 

2.	 V. D. Fréchette, “Failure Analysis of Brittle Materials,” Advances in 

Ceramics, Vol. 28, American Ceramic Society, Westerville, OH, 1990. 

3.	 R. W. Rice, “Microstructure Dependence of Mechanical Behavior of 

Ceramics,” pp. 199 – 381 in Treatise on Materials Science and Technology, 

Vol. 11, Academic Press, New York, 1977. 

4.	 CEN EN 843-6, “Advanced Technical Ceramics - Monolithic Ceramics – 

Mechanical Properties of Monolithic Ceramics at Room Temperature – Part 

6: Guidance for Fractographic Investigation,” European Committee for 

Standardization, Brussels, 2009. 

5.	 J. J. Swab and G. D. Quinn, “The VAMAS Fractography Round Robin: A 

Piece of the Fractography Puzzle,” pp. 55 -70 in Fractography of Glasses and 

Ceramics III, Ceramic Transactions, Vol. 64, eds., J. R. Varner, V. D. 

Fréchette, and G. D. Quinn, ACS, Westerville, OH, 1996. 

6.	 Failure Analysis and Prevention, ASM Handbook, Vol. 11, American 

Society for Materials, Materials Park, OH, 2002. 

7.	 P. B. Adams, “Fractography in the Courtroom,” pp. 219 – 226 in 

Fractography of Glasses and Ceramics, II, Ceramic Transactions Vol. 17, 

eds V. D. Fréchette and J. R. Varner, American Ceramic Society, 

Westerville, OH,1991. 

8.	 C. O. Smith, “Products Liability and Design,” pp. 71-78 in Ref. 6. 

9.	 R. C. Bradt, “Fractography, Some Legal Aspects and Challenges,” pp. 28 – 

42 in Fractography of Advanced Ceramics, III, eds., J. Dusza, et al., 

Transtech Publ., Zurich, 2009. 

1-10 
 

 

 

 

Introduction 

10. V. D. Fréchette, “Markings on Crack Surfaces of Brittle Materials: A 

Suggested Unified Nomenclature,” pp. 104-109 in Fractography of 

Ceramic and Metal Failures, ASTM STP 827, J.J. Mecholsky, Jr. and S.R. 

Powell, Jr. eds., American Society for Testing and Materails,1984. 

11. A. C. Roulin-Moloney, Fractography and Failure Mechanisms of 

Polymers and Composites, Elsevier, Essex, England, 1989. 

12. G. D. Quinn, “A History of the Fractography of Brittle Materials,” in 

Fractography of Advanced Ceramics, III, ed. J. Dusza, R. Danzer, R. 

Morrell and G. Quinn, TransTech Publ., Zurich, 2009. Key Engineering 

Materials Vol. 409 (2009) pp. 1-16. 

13. G. D. Quinn, “A History of the Fractography of Glasses and Ceramics,” 

pp. 1 – 56, in Fractography of Glasses and Ceramics, Vol. 6, Ceramic 

Transactions, Vol. 230, eds, J. R. Varner, M. Wightman, American 

Ceramic Society, Westerville, OH, 2011. 

1-11 

 

 

 

 

 

 

 

 

 

 

 

 

Resources 

2. Resources 

This Guide is designed to be a stand-alone resource. There is a wealth of 

additional information scattered in books and technical articles that delve into 

specific topics in more detail. 

Appendix A is a bibliography featuring many good publications of special 

interest to the glass and ceramic communities. The topics include books on 

fractography, compilations of conference proceedings, overview papers on 

fractography, reports on fractography round robins, compilations of fracture 

mechanics stress intensity factors, and papers on microscopy, fracture origin 

types, fracture mirrors, and fracture mechanics analyses of flaw sizes. 

There are two ongoing conference series devoted exclusively to the 

fractography of glasses and ceramics. One is in the United States and one is in 

Europe. The conferences have been held at four or five year intervals. The 

longest running and most famous was begun by Professors Van Derck 

Fréchette and James Varner in 1986 at Alfred University in New York State. 

Follow-on conferences were held in 1990, 1995, 2000, 2006, and 2011. The 

six conference proceedings have excellent information on a range of theoretical 

and applied problems by leaders in the field. Dr. Jan Dusza of the Institute of 

Materials Research of the Slovak Academy of Sciences, Košice, Slovakia 

began a conference series in 2001 that continued in 2004, 2008, and 2013. The 

next is scheduled for October 2016. Conference proceedings for the first three 

were published as bound books, but only a selection of papers from the 2013 

conference were published in a topical issue of the journal of the European 

Ceramic Society in 2014. 

Appendix B is a unique compilation of case studies that lists a series of 

publications dealing with specific case studies. These include fracture of 

windows, bottles, pressure vessels, and various ceramic parts. Some of the 

cases deal with the mundane and some with the exotic. It is included for the 

benefit of readers who wish to investigate specific problems. 

Appendix C is the most complete tabulation of fracture mirror constants ever 

compiled for ceramics and glasses. It is an expansion of a similar listing in 

ASTM C 1322-02a and the first 2007 edition of this Guide. 

Each chapter of this Guide also contains a list of references for topics covered 

in the chapter. 

2-1 
 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

There are now four formal documentary standards pertaining to fractographic 

analysis of glasses and ceramics: 

ASTM C 1256-13, “Standard Practice for Interpreting Glass Fracture Surface 

Features,” ASTM, International, West Conshohocken, PA, originally adopted 

1993, current version: 2013. 

ASTM C 1322-15, “Standard Practice for Fractography and Characterization of 

Fracture Origins in Advanced Ceramics,” ASTM, International, West 

Conshohocken, PA, originally adopted in 1996, current version: 2015. 

ASTM C 1678-15, “Standard Practice for Fractographic Analysis of Fracture 

Mirror Sizes in Ceramics and Glasses,” ASTM, International, West 

Conshohocken, PA, originally adopted in 2007, current version: 2015. 

EN 843-6, “Advanced Technical Ceramics, Monolithic Ceramics, Part 6: 

Guidelines for Fractographic Examination,” European Committee for 

Standardization, Brussels, originally adopted, 2004. 

2-2 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Tools and Equipment 

3. Tools and Equipment 

The most important tools of the fractographers craft are the binocular 

stereoptical microscope and the scanning electron microscope. The stereoptical 

microscope affords to the eyes a magnified, naturally-appearing, threedimensional view of the fracture surface that retains all color and reflectivity 

information. Fracture origins and flaws in ceramics and glasses can often be 

detected with the stereoptical microscope, but higher magnifications are often 

needed to see the flaw more clearly. The scanning electron microscope serves 

this need and has good depth of field, high magnification capability, and 

chemical analysis features. Other tools such as hand magnifiers, compound 

optical microscopes, or even atomic force microscopes play supportive roles. 

The “marriage” of the microscope with digital cameras and computers has 

revolutionized the field. The new technologies have rendered film photography 

obsolete and are opening up marvelous new capabilities not heretofore 

available. The shortcomings of conventional compound microscopes (e.g., 

limited depth of field) or of scanning electron microscopes (e.g., flat-appearing 

images) can be overcome. Virtual three-dimensional images can be 

constructed with automatically-rastering optical microscopes or by analysis of 

multiple scanning electron microscope images taken with slightly different 

specimen tilts. The pseudo three-dimensional images can be displayed, tilted, 

and rotated to afford different perspectives. Various quantitative numerical 

analyses of surface topography (e.g., surface roughness or even fractal 

dimension analysis) can be performed at almost no extra effort since the surface 

topography has been recorded digitally. 

Simple low-cost software programs are now available that can simply and 

conveniently overcome the depth of field limitations of virtually any 

microscope. A series of photos at the same magnification are taken while 

slightly readjusting the focus in sequential steps. The software interprets the 

regions that are in focus and stitches these together to create a single image 

with an infinite depth of field. 

In this chapter, traditional and new tools of the trade and are reviewed with an 

emphasis on their applicability to fractography. Digital cameras are discussed 

in four sections of this chapter: first in section 3.2 as simple tools for 

photographing an overall part, then in sections 3.4 and 3.5 as recording devices 

attached to conventional microscopes, and finally in section 3.7 as components 

of new digital microscopes which do not even have eyepieces for the viewer to 

look through. 

3-1 
 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

3.1 Low Power Optical Examination and Component Fracture 

Documentation 

The examination of fractured pieces begins with a visual examination using the 

unaided eye. Some fractographer prefer to use lint-free gloves when handling 

fragments, but gloves usually are not necessary. Ceramics and glasses are hard, 

durable materials and may be cleaned fairly easily. 

It may be tempting to rush to a microscopic examination of the fracture 

surfaces, but study of the general fracture and the overall condition of the 

component is time well spent. This examination and reconstruction of the 

specimen may be aided by a simple (1.2X to 1.4X) magnifying glass, whether 

hand-held or mounted on a stand as shown in Figure 3.1. A stand-mounted lens 

frees the fractographer’s hands for manipulation of the specimen. Larger 

versions with built-in ring lamps are ideal. A clean, tidy table or bench top is 

desirable, lest the fractographer inadvertently knock critical fragments onto the 

floor or against each other, or contaminate or damage the fracture surfaces. 

Figure 3.1 Simple magnifiers. 

Small jeweler’s loupes as shown in Figures 3.2 and 3.3 also are handy, 

particularly in the field. Their working distances range from 50 mm for the 

low power (5X) loupe, to about 25 mm for a 10X loupe. Some practitioners 

may be adept at holding a watchmaker’s loupe in one eye without using a hand, 

thereby freeing up both hands. Most fractographers will use one hand to hold 

3-2 
 

 

 

 

 

 

 

 

 

Tools and Equipment 

the loupe and the other to hold the examined fragment. Machinist loupes and 

pocket optical comparators with magnifications up to 20X (two are shown on 

the right of Figure 3.2) have built-in measuring scales with graduations as fine 

as 0.1 mm (0.005 in). Rapid measurements of mirror sizes or even fracture 

origin sizes may be made with these inexpensive tools. Remember, the more 

powerful the magnification, the smaller the field of view and the shallower the 

depth of field, so a good general duty loupe or comparator with a 7X or 10X 

magnification may be perfectly suitable and more versatile than one with a 20X 

magnification. At these low magnifications, illumination is usually not a 

problem and ambient light is adequate. 

Figure 3.2 Hand loupes and optical comparators. From the left, 5X and 7X 

watchmaker loupes; a 20X loupe, a 7X comparator with an internal 20 mm 

scale marked in 0.1 mm increments, and a 8X to 16X zoom comparator with an 

16 mm scale marked in 0.1 mm increments. 

Figure 3.3 Loupes (5X to 20X) may either be held by the eye or by hand. The 

part is a single crystal silicon rod. 

3-3 
 

 

 

 Fractography of Ceramics and Glasses 

Simple holders or clamps may be useful as shown in Figures 3.5 and 3.6, 

particularly if the fractographer is examining repetitive examples of a particular 

type specimen or component. It can be exasperating to hold a part in one hand 

at just the right angle, while holding a magnifier with the second, but then need 

an extra hand to apply some alcohol with a cotton swab to clean the part. 

Similarly, the fractographer may wish to jot down some important information 

as part of a written record. 

Figure 3.4 A jeweler’s or hobbyist’s head mounted magnifier (2X to 3.5X) can 

be helpful. 

Figure 3.5 Specimen holders. Clockwise from the lower left are: a bend bar 

holder made with an alligator clip on a bendable (10 gage electrical) wire with 

electrical tape wrapped around the teeth to cushion the specimen, two 

aluminum jigs with slots and grooves for bend bars or tension strength 

specimens, an aluminum block with hole, and three aluminum holders 

comparable in size to a glass microscope slide with groves to hold bend bars 

both on their flats and on edge for chamfer examination. 

3-4 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

Tools and Equipment 

Conventional clays should not be used for mounting specimens unless there is 

no alternative. Clay is a pernicious material that can easily contaminate the 

fracture surfaces and blend into the fracture features. Natural clays have colors 

and compositions similar to ceramic products. They are next to impossible to 

clean off. Their electrostatic surface charge and extremely fine plate-like 

structure cause them to cling to the fracture surface. Clays often have an oil to 

keep them plastic that makes things even worse, since the oils can smear over 

the fracture surfaces even on glasses. Clay often gets on one’s hands and then 

tends to get on everything. More information on clay contamination is in 

section 6.10. 

A completely satisfactory alternative is polymer clay that is available from arts 

and crafts stores.a Some polymer clays can be easily dissolved and cleaned 

with paint thinner, acetone, and even ethanol. The color in these polymer clays 

can easily be distinguished on the fracture surfaces. The best colors are deep 

blues or greens. White or gray should not be used since they cannot be 

distinguished from many ceramics. Garish colors should not be used, since 

they are a distraction in photographs. 

(a) (b) 

Figure 3.6 Component or specimen holders. (a) Clockwise from the lower 

left: a homemade plastic platform mounted on a swivel base; a pivoting 

wooden holder, a cork ring; a clamp on a stand, an alligator clip on a stand 

(with electrical tape on the teeth to cushion the specimen), an inverted 

miniature wood flower pot with a hole, and two wooden candleholders with 

dolls heads and polymer mounting clay. The doll’s head balls can be rotated 

to obtain optimum illumination angles onto the specimen’s surface. (b) shows 

a polymeric clay that is safe to use and is easily removable. 

a One example of a polymeric clay is Sculpey III Polymer, “Oven bake clay” made by 

Polyform products, Elk Grove, Ill. It is easily dissolved by paint thinner or acetone. 

3-5 
 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Simple cleaning tools and markers such as shown in Figure 3.7 should be 

conveniently located at the examination work station. Acetone, ethanol, or 

methanol with tissues and cotton swab applicators are handy for cleaning. 

Compressed air from air lines should be used with caution since there may be 

traces of oil mixed with the air. Water soluble ink pens (with fine tips) are 

handy for marking the specimen. Ceramic fracture surfaces should never be 

cleaned with metallic tools or probes. It may be tempting to try to scrape or 

poke off a contaminant, but metallic tools are abraded by the harder ceramic 

causing metallic traces to be smeared on the fracture surface. Fractured dental 

restorations often have plaque on them. The plaque can even be on a fracture 

surface if it has been exposed to the oral environment for any period of time. 

An example of what plaque looks like is shown in section 6.10 of this guide, 

where contaminants are discussed. Plaque can be removed by using an 

ultrasonic bath with a 16% solution of NaOCl (“bleach”) for 8 minutes, then 

rinsing well with water (courtesy S. Scherrer). 

There are two instances where specimens should be examined before cleaning. 

Firstly, staining or discoloration of a flaw may provide important information 

about the flaw itself. Secondly, fine debris of the material itself on the fracture 

surface may be a clue that cyclic fatigue has occurred. These are discussed at 

the very beginning of Chapter 5, Fracture Surface Examination. 

Figure 3.7 Convenient tools and cleaning agents 

3-6 
 

 

 

 

 

 

 

 

 

Tools and Equipment 

3.2 Photographing the Overall Component: Cameras, Cellphones, 

and Simple Digital Microscopes 

It is often prudent to photograph the entire specimen or component. An overall 

photo provides an essential context for subsequent close up photos. A variety 

of scales or rulers should be available, which may be set alongside the 

component in order to show the size or scale. It is prudent to get size markers 

on the recorded image as soon as possible, otherwise much time can be wasted 

afterwards deducing or guessing what magnification was used. It is very easy 

to make mistakes with the size markers added after the fact. Fractographers 

would do well to follow the example of the archeological community, wherein 

rulers and meter sticks are ubiquitous features of any formal record of a dig 

site. In lieu of scales, simple props such as coins or other common objects may 

be set alongside the component, but perhaps on the side or bottom where they 

subsequently can be cropped or replaced with marker bars in subsequent 

reports or publications. 

A variety of camera types may be used to photograph an entire component. 

Digital cameras have become the tool of choice for virtually all such work, but 

even they are being supplanted by cell phone cameras. Conventional filmbased photography (Figure 3.8) has been rendered obsolete for most 

applications. 

(a) (b) 

Figure 3.8 Obsolete conventional film photographic equipment. (a) is a 35 

mm single lens reflex camera tripod, and close-up lenses. (b) is a camera stand 

with instant film (Polaroid) camera, built-in flood lamps and a fiber optic light 

source for low-angle grazing illumination. 

3-7 
 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Modern digital single lens reflex (SLR) cameras and even simple consumer 

digital cameras with a macro option can be used at very close distances (~100 

mm or less) from the object. Digital cameras, either professional highresolution models or simple consumer versions such as shown in Figure 3.9, 

have revolutionized fractographic photography. They have auto focus and 

exposure controls and can capture good images in ambient lighting without a 

flash. Liquid crystal display (LCD) monitors on the rear of the cameras allow 

fractographer to see what the lens will capture when the image is snapped. 

The digital camera instantly displays the captured image. Remedial action can 

be taken immediately and a new image captured. This is very efficient and led 

to the rapid obsolescence of film-based photography. Another advantage is that 

images taken with digital cameras seem to be less sensitive to the illumination 

source color. Color film often produced odd-colored images if film balanced 

for outdoors was used indoors with tungsten filament bulbs or fluorescent 

lamps. Another great advantage of the digital camera is that it needs less light 

to make a good quality image Ambient illumination usually suffices. Flash 

lamps used with old film cameras often over- or underexposed an object and 

created harsh shadows. 

Figure 3.9 Digital camera with macro capability, auto exposure, and auto 

focus. 

Experience has shown that a digital camera chip with 2 million to 5 million 

pixels is adequate for most applications. Some publications require no less 

than 3 megapixel resolution and images must be presented at 120 dots per cm 

(300 dots per inch). The LCD monitor should be bright and large (at least 40 

mm). The camera should have auto focus and auto exposure options and, most 

importantly, have a macrophotography option for close ups at distances as short 

as 100 mm. File storage format should include a Joint Photographic Experts 

Group (JPEG) or Tagged Image File Format (TIFF) option. Sometimes Bitmap 

(BMP) format files are used with Microsoft Windows operating system 

computers. Some professionals prefer the raw image format (RAW), which is 

3-8 
 

 

 

 

 

 

 

 

 

 

 

 

Tools and Equipment 

image data with very little processing or compression. Image compression is 

important for the various formats. 

JPEG images have an option that they may be saved with variable amounts of 

compression, but users should record images with little or no compression. A 

copy of the image can be made at a later time with some compression to reduce 

the file size so that electronic documents or presentations with many images do 

not become unwieldy. JPEG files are as much as 1/5th to 1/10th the size of an 

uncompressed TIFF file, with very little or no loss of detail. JPEG is available 

only for true color or grayscale images. JPEG reduces the number of bits in the 

image by eliminating repetitive image data or image data that is hard to see. 

The software may allow various compression levels and the higher the level of 

compression, the lower the quality of the image. JPEG may compress more 

than other methods and the results may not be fully reversible. At higher 

compression ratios (> 20:1 or a 50% setting), JPEG can degrade the quality of 

the image. A JPEG file is automatically uncompressed when it is opened. 

TIFF is either an uncompressed format or a LZW (Lempel-Ziv-Welch) 

compressed format. LZW compression translates common patterns in an 

image into short codes and is best at preserving all the image data and 

achieving good compression. It doesn't achieve the high compression ratios 

that JPEG does. LZW is available for monochrome, grayscale, palette, and true 

color images. 

The RAW format is used by manufacturers to contain the raw (minimally 

processed) data from the image sensor in a digital camera. RAW file formats 

are proprietary and can differ greatly from one manufacturer to another and 

sometimes between cameras made by one manufacturer. RAW files are 

normally only used when additional computer processing is intended. The 

proprietary nature of RAW files poses severe problems that fractographers 

should keep in mind. Photographers will find their older images inaccessible, 

as future software versions lose support for older cameras and entire camera 

brands and RAW format subtypes may disappear! 

The author typically saves and presents his images in JPEG format with no 

compression. Virtually all the photographic images in this Guide were saved in 

the JPEG format. Line art and graphs were saved in LZW compressed TIF 

files. 

The digital camera should be capable of easily transferring images to a 

computer, either through a cable to a USB or FireWire port on the computer or 

via a docking port. Common consumer digital cameras that have these features 

3-9 
 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

now cost less than one hundred dollars. More elaborate single lens reflex 

(SLR) digital cameras have more options, more versatile lenses, and higher 

resolutions. Some may have extremely high resolutions, but it should be borne 

in mind that the stored image file sizes may be very large. SLR cameras should 

be ordered with close up lenses as an option. 

(a) (b) 

(c) (d) 

Figure 3.10 Cell phones can be used for fractographic analysis (a) shows an 

obsidian arrowhead as viewed at 2X magnification in a cell phone; (b) a street 

photo near Santo Stefano Basilica in Bologna, Italy in 2014; (c) a close-up of 

the tempered glass fracture pattern, and (d) the close-up at 5X of the internal 

inclusion origin (arrow). The author was on his hands and knees for (d), which 

was shot at 5X about 60 mm above and aiming through the glass at an angle. 

3-10 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

Tools and Equipment 

Built-in cell phone cameras have surprisingly high resolutions as shown in 

Figure 3.10. Some models can be adapted with free applications (“apps”) that 

enable the phone to be used in a magnifier mode. As this was being written in 

2015, the author uses a cell phone with 110 mm sized bright digital screen and 

autofocus with a 5X magnification capability. It can be used as close as 60 mm 

from an object. Some cell phone manufacturers have clip-on close-up lenses 

and simple holder-stand accessories that can reach even greater magnifications. 

Users can construct their own little stands to hold the camera steady when 

taking close-up images (e.g., some are shown on the Internet). The author has 

also been surprised in some of the courses on fractography that he teaches 

when students used cell phones to take images through one eyepiece of a 

stereoptical microscope. 

Digital camera photography has eliminated many of the color problems 

associated (e.g., sensitivity to incandescent bulb temperature) with older film 

technology. In some cases, the best results can be obtained by photographing 

the component outdoors with ambient lighting on a neutral-colored 

background, such as a white or gray poster board. The author often 

photographs specimens on the window sill of his office on a white poster board 

base. An additional white cardboard is placed vertically as a reflector on the 

inside of the room to reflect some external light to back-fill the lighting on the 

specimen. 

3.3 Stereo Binocular Microscope 

More powerful magnification will be needed in most cases to observe small 

fracture origins or markings on the fracture surfaces. The fractographer’s 

single most valuable tool is the stereoptical binocular microscope or 

stereomicroscope as shown in Figure 3.11. Basic binocular stereomicroscopes 

have magnification ranges to the eyes of about 8X to 32X or 10X to 40X, but 

these limit one’s view of a small fracture origin flaw and greater magnifications 

are needed. Upper magnifications of 100X or as high as 300X are available 

with many stereomicroscopes and are better for fractographic analysis. On the 

other hand, having a small magnification at the lowest limit (e.g., 5X) 

facilitates taking an overall picture of a small component. Hence, a stereoptical 

microscope with a broad zoom range (e.g., range of 10, 16, or even 20 power) 

is very advantageous. A 50/50 beam splitter (half the light is sent to the eyes 

and half is sent to a camera) in the stereo binocular microscope is very 

desirable since it allows one to look through the eyepieces at the same time an 

image is sent to the camera. The alternative, a lever which diverts light either 

to the eyes or to the camera, is cumbersome and less desirable. It can be 

exasperating to use. 

3-11 
 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Figure 3.11 Stereo binocular microscope with a fiber optical light source for 

directional illumination. This is an indispensable tool for the ceramic and 

glass fractographer. “Low-angle, grazing illumination” is illustrated in the 

picture. In other cases, a ring light mounted on the objective lens of the 

microscope may be suitable. This particular microscope has interchangeable 

0.63X, 1.0X and 2.0X objective lenses, a body zoom magnification range of 

.65X to 4X, and 25X eyepieces for an overall magnification range of 10X to 

200X . The moderately tall (300 mm) mounting post enables specimens up to 

150 mm tall to be placed under the microscope. 

The stereomicroscope is designed such that each eyepiece views the object 

from a slightly different angle, thereby creating two different images. The 

brain interprets these as a three-dimensional view. The stereomicroscope is in 

reality two separate microscopes. Older rudimentary Greenough design 

stereoscopes had two completely different optical paths, one for each eye. 

They had very limited capabilities. Most contemporary stereomicroscopes use 

a common objective lens and then two internal optical light paths in the body 

up to the eyepieces. Parfocality, whereby the object remains in focus while the 

magnification is changed, is a very advantageous feature of the microscope. 

Follow the manufacturer’s guidelines to set the eyepieces for Parfocality. 

Stereomicroscopes present an image that is right side up and is laterally correct. 

The images are therefore easy to correlate to a specimen held in the hands or on 

the stage. The long working distance from the objective lens and the good 

3-12 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Tools and Equipment 

depth of field make it possible to clearly see surface features despite the 

inherent roughness or curvatures of fracture surfaces. Some stereomicroscopes also have diaphragm apertures that can be stopped down to enhance the 

depth of field. Many stereomicroscopes are modular and easily modified or 

expanded. 

The total magnification of a stereomicroscope is the product of the magnifycations of the objective lens, the body, and the eyepieces. There often are 

different objective lenses (0.5X to 2.0X) that screw into the bottom of the 

microscope. Low power 0.5X or 0.63X lenses are useful for examining or 

photographing an entire test specimen or viewing large areas of the fracture 

surface. Higher power lenses (1.6X or 2.0X) may be screwed in for close-up 

examination. It may come as a surprise that the objective lenses have such low 

magnifying power, but a large diameter lens is needed to create the two optical 

views at slightly different viewing angles. The low power also means 

that the working distance from the objective lens to the object is very large, on 

the order of 20 mm to 200 mm. This is a very important and desirable feature 

for the examination of irregular components. One advantage of the long 

working distance is that illumination sources such as fiber optic gooseneck 

light guides can be brought into close proximity to the work piece and the angle 

can be adjusted with plenty of room as shown in Figure 3.11. A good 

fractographic examination set up should have both a low power objective lens 

(0.5X or 0.63X) and a high power lens (1.6X or 2.0X). An intermediate power 

lens (1.0X) may also be handy. Objective lenses can cost anywhere from $ 400 

to $ 4,000 each. The more expensive lenses have better sharpness at greater 

magnifications and have less chromatic aberration. 

Basic student-model stereoptical microscopes have a limited magnification 

range. They have little or no capability for expansion (e.g., addition of 

different objective lenses or camera ports) and have very limited working 

distances. They are not a wise investment for the serious fractographer. 

Nevertheless, when such a microscope is the only one available, it is possible to 

use a cell phone to take an image through one eyepiece. 

Much of the magnification occurs in the body of the stereomicroscope. Most 

bodies have multiple magnifications either through the use of discrete lenses 

that can be rotated into the optical path, or better, through stepless-zoom 

magnification. Modern stereomicroscopes have been vastly improved over 

earlier generation models that may have only had a 2X or 4X zoom range. 

Modern bodies can steplessly zoom over a 10X, 16X, or even 20X range while 

keeping the object in focus. A very desirable feature is discrete click stops on 

the zoom knob corresponding to specific magnifications. This feature is critical 

3-13 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

for size measurement purposes, either by the use of a graduated reticule in the 

eyepiece, or by a camera mounted on the microscope. 

Final magnification is provided by the wide field eyepieces that are usually 

10X, 16X, or 25X. Stereomicroscopes may be used either with or without 

eyeglasses. The distance between the eyepiece tubes usually can be adjusted 

for the interpupilary distance. Eyepieces may have rubber boots that attach to 

the eyepieces for screening out stray light coming into the eyes. These can be 

removed or slid out of the way for eyeglass wearers. Typical overall 

magnifications can range from as low as 5X to as high as 200X. Some modern 

stereomicroscopes can achieve as much as 300X to 400X magnification with 

16X or 20X body magnifications and 1.6X objectives. 

The stereoptical microscope should be mounted on a tall mounting post so that 

large objects can be placed on the stage. Often it is necessary to mount long 

specimens end on so that the fracture surface is viewed directly. Figure 3.11 

shows a model with about a 300 mm post that allows specimens as large as 100 

mm to 150 mm tall to be placed on the stage. 

It is also advisable to have one or both eyepieces adjustable, so that the 

operator can focus each eyepiece independently. One reticule should have a 

cross hair, which is valuable if one wishes to show a colleague a particular 

feature, or the stereomicroscope will be used for size measurement in 

conjunction with a traversing stage as will be described below. If no 

independent measurement system is incorporated, then a built-in graduated 

reticule should be mounted in one of the eyepieces. The reticule readings will 

have to be calibrated at various magnifications. 

A very valuable but expensive optional feature is a tilting eyepiece head that 

may be adjusted up or down so that the viewer does not have to extend or bend 

his or her neck to see through the eyepieces, or to have to move a chair or stool 

up or down. A tilt head can cost between $ 3,000 and $ 4,000. 

3.4 Stereoptical Microscope Camera Port and Digital Cameras for 

Microscopes 

Stereomicroscopes may be equipped with a camera port for a video or digital 

camera. The port will be in line with either the left or right optical paths. The 

camera cannot record or display both images that the eyes see and the brain 

interprets as a three-dimensional image. The camera commonly will be 

mounted on the camera port through a C-mount or an F-mount that must be 

chosen with care. It must fit the camera port’s diameter, but also have with an 

3-14 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Tools and Equipment 

appropriate internal lens to match the image size to the camera chip size. 

Some microscope bodies have a lever switch that directs the light to either the 

eyepiece or to the camera. This is extremely inconvenient. A much better 

scheme is a beam splitter that directs 50 % of the light to the eyepiece and 50 % 

to the camera so that the image can be viewed at the same time it is projected 

onto a monitor. 

Instant film cameras are now obsolete. They were supplanted first by video 

cameras, but now almost exclusively by digital cameras connected to a 

computer. Video cameras which had fast image capture capabilities had to be 

used with frame grabbers to snap an image. They had very low 640 x 480 or 

768 x 500 picture element resolutions so that many images per second could be 

taken. They have been rendered obsolete as digital cameras became more 

capable, less expensive, faster even with low light, and easier to use. 

Digital cameras designed to be attached to microscopes can show “live” images 

at up to 10 to 20 fames/sec with nominally reduced resolutions on the computer 

monitor. Once the object is framed and focused, a higher quality image can be 

acquired or “snapped.” A small piece of white paper is often handy, since it 

can be momentarily placed in the field of view and the camera color-balanced 

by pressing a camera software button. A fast “live” image mode makes 

focusing on a computer monitor very efficient. Fractographers should avoid 

systems with slow refresh rates since focusing is exasperating as one makes a 

slight focus adjustment, and then has to wait a few seconds to see the effect on 

a monitor. A digital camera chip with 2 million to 5 million pixels is adequate 

for most applications. The most common image formats in 2015 are JPEG, 

TIFF and sometimes BMP as discussed in section 3.2. Image compression 

should be minimized or not used at all when capturing and saving images. 

Sometimes excessive emphasis is placed on having large pixel counts in digital 

cameras. There is no harm in having digital cameras with larger pixel counts, 

but storing and handling very large files might become cumbersome, especially 

if the images are embedded in documents. Refresh rates on “live” images may 

also be slowed down. Nearly all the microscope images in this Guide were 

captured with a CCD (charge-coupled device) digital camera having a 11.8 mm 

x 8.9 mm RGB color mosaic chip having 2 megapixels (1600 x 1200) and 

saved as 24 bit color jpg files. The price of these cameras is cost effective and 

resolutions of 3 megapixels or more are common at prices less than $ 4,000 as 

this Guide was written. CCD chip cameras are commonly used but CMOS 

(complementary metal oxide semiconductor) chip cameras are also becoming 

more popular. (A microscope digital camera specialist or salesman can be 

consulted for the advantages and disadvantages of each.) 

3-15 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

The camera input can be fed to free software provided with the digital camera 

as part of a complete package, or to a separate image processing software 

package. In the past there inevitably were compatibility problems with 

software and cameras and all the usual annoying nuances of computer 

programs. Contemporary digital cameras and software packages are much 

more user friendly. Nevertheless, it probably is best to have digital cameras 

and their software installed by a camera or microscope company representative 

who has experience trouble shooting installation and compatibility problems. 

The simpler the software the better! The author has used digital software 

packages that are excessively complicated, in large part since they were written 

with lots of options and clever additions for experts. They are slow and 

cumbersome, and have so many menus and submenus that they are a nuisance 

to use. (Users should beware of salespersons who are eager to sell expensive 

software packages that the user does not really need.) A lot of time is lost in 

training students and new users. 

A very important attribute of a digital camera is the ease which a magnification 

marker bar can be added to an image. Some digital cameras come with free 

software that allows not only the usual image adjustments (brightness, contrast, 

sharpness, cropping, etc.) and file-saving options, but annotations and arrows as 

well. This is important, and the fractographer can mark a feature on an image 

immediately and not have to use a secondary software package back in his or 

her office. Even more importantly, it is crucial to be able to put a 

magnification – size marker bar onto an image. The fractographer should get 

into the habit of doing this immediately. It can be confusing afterwards to 

know just what magnification a photo was taken with, and an erroneous marker 

bar might be put on. One option is to save a file with an image name that 

includes the magnification, e.g., “specimen 6 side view at 32X.jpg”. If a 

microscope has a variety of lenses and body magnifications, then calibration 

factors for each combination will have to be established in the software. 

Usually this is not too difficult, and once these values are set, they can be 

retrieved automatically by the software afterwards. The author uses a scheme 

whereby the camera software has a list of files with all possible lens 

magnification combinations, with file names that are easy to understand, e.g., 

“1.6X obj 4X body” which is short for: “1.6 X objective lens, 4X body 

magnification.” The author routinely needs only 5 to 10 seconds to put an 

accurate magnification marker bar on his digital images. The author has gone 

to fractography laboratories around the world and has been pleasantly surprised 

that others have used simple schemes like this, and it is possible to start using 

the microscope with little or no confusion. 

3-16 
 

 

 

 

 

 

 

 

 

 

 

Tools and Equipment 

3.5 Digital Image Processing and Focus Stacking 

Digital image technology has revolutionized techniques for processing images. 

Brightness, contrast, sharpness, and gamma adjustments can easily be made to 

captured images. Complex discussions about the gamma adjustment and 

luminance of pixels can be found on the internet or in software packages, but 

can be confusing to newcomers. Simply put, gamma settings adjust the 

brightness of the pixels that are between the limits of the very bright and very 

dark regions in an image. In other words, a gamma adjustment does not alter 

the bright or dark regions, but everything in between. 

Inexpensive commercial software now overcomes the depth of field limitations 

of common microscopes. A series of ten to twenty JPEG or TIFF digital 

images of an object are taken by an ordinary digital camera with the focus 

adjusted slightly between each photo. Only a portion of the object is in focus in 

each image. The software evaluates each image, assesses which portions are in 

focus, and then constructs a composite final image with all elements in focus in 

just a few seconds. This is called “focus stacking” or “Z-axis stacking.” It 

does not matter whether the images are from a camera mounted on a tripod, an 

optical microscope, or a scanning electron microscope. Figures 3.12 and 3.13 

show examples. Even if the images shift slightly after refocusing in a 

stereoptical microscope, the software compensates for the shifts. The practical 

ramifications of this technology is that infinite depths of field are achievable 

with all microscopes, dramatically enhancing the versatility of all microscopes. 

(a) (b) (c) 

Figure 3.12 A silicon nitride 6 mm diameter rod broken in flexure (801 MPa). 

(a) and (b) show stereoptical microscope digital images. Even though the 

microscope has great depth of field, the cantilever curl and holder are out of 

focus in (a). (b) shows the holder in focus. (c) shows the digitally 

reconstructed composite image from 15 separate images. 

3-17 
 

 

 

 

 

 

 

 

The bottom is in focus. 

 Fractography of Ceramics and Glasses 

In focus. In focus. All in focus. 

(a) (b) (c) 

Figure 3.13 A 25 mm glass fragment. (a) and (b) show typical images made 

by a stereoptical microscope at low magnification. Although the microscope 

has excellent depth of field, portions of the fragment are out of focus in each 

view. (c) shows the digitally reconstructed image made from 16 separate 

images focused from the top to bottom of the fragment. The digital 

reconstruction took less than 1 minute. 

Figure 3.14 A discussion stereoptical microscope that allows two observers to 

simultaneously view an object. Note the very tall post (800 mm) on which the 

microscope body is mounted. The entire assembly can be moved up or down as 

required to facilitate examination of very large or very small specimens. 

Beam splitters in the optical path enable simultaneous viewing through the 

eyepieces at the same time the digital camera views the object. 

3-18 
 

 

 

 

 

 

 

 

Tools and Equipment 

3.6 Discussion Stereomicroscope 

Figure 3.14 shows a “discussion stereoptical microscope” that enables two 

observers to simultaneously see the image. A beam splitter diverts half of the 

light to the auxiliary viewing station. The configuration shown is a side-byside model which is preferred for most fractographic analysis. Discussion 

microscopes often have a built in moveable pointer that can be seen by both 

observers. Discussion microscopes are extremely valuable for training new 

fractographers or for showing features to clients or colleagues. The auxiliary 

viewing station should also have adjustable eyepieces so that the secondary 

observer can adjust the focus to match the focus of the primary observer. A 

camera port can be included with another beam splitter so that the image can be 

projected onto a monitor at the same time it is viewed through the eyepieces. 

3.7 Digital Microscopes 

Simple digital microscopes, also known as USB microscopes, connect directly 

to an ordinary computer. They do not have eyepieces to look through. 

Everything is projected onto a computer monitor. They are a new technology 

Figure 3.15 Digital microscopes are a new tool for the fractographer. This 

simple digital microscope (arrow) is mounted in a stand on the right. It is 

focused on the fracture surface of a broken glass rod. The monitor and 

computer (not shown) are separate components 

3-19 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

that is becoming increasingly common. They can range from inexpensive lowpower, hand-held models to elaborate, high-power, expensive models mounted 

on a rigid platform with full Z-axis control and digital image stitching 

capabilities. Figure 3.15 shows a mid-priced model (~$ 1,000 plus $130 for the 

stand) with variable magnifications from 20X to 250X and corresponding 

working distances of 48 mm to 2 mm. It may be difficult to obtain sharp, 

focused images with the simpler models that are hand held, so a stand is 

advisable. The digital microscope is not parfocal. As the magnification is 

changed, the image must be completely refocused and the working distance is 

changed significantly. This is a nuisance. Another limitation is illumination, 

which is usually provided by built-in light-emitting diodes surrounding the 

lens. This limits their ability to highlight or even discern critical fracture 

surface features. For example, shadowing, or low-angle grazing illumination, 

which is essential for ceramic examination and fracture mirror examination, 

cannot be done. A separate light source is recommended. 

At the other end of the spectrum are high-end digital microscopes mounted on 

very rigid fames and bases and which may have built-in vibration dampers. 

They may have fully-motorized positioning stages and automated Z-axis 

focusing and even have motorized tilting heads. Camera and stage controls are 

via joysticks, electronic knobs, or even touch-screen monitor manipulation. 

These systems can automatically conduct Z-axis focus stacking analysis and, 

combined with the stage control, can generate quantitative three dimensional 

profiles of the fracture surface. The virtual three-dimensional image can be 

displayed on the monitor and rotated and viewed from different perspectives. 

Despite this extraordinary versatility, the illumination systems with such 

microscopes (usually a ring of LEDs) are very limited. Many of the techniques 

described in this Guide require fully-independent illumination control. It is 

recommended that independent illumination sources such as shown in the next 

section be acquired for use with such high-end digital microscopes. The 

working distances between the lens and the piece should be large enough that 

independent light can be directed onto the specimen at variable angles. 

3.8 Illumination Sources 

Figure 3.16 shows some options for specimen illumination. Bright light 

sources are advisable to facilitate the examination of dark ceramics and the use 

of high magnifications. The most common illumination source for fractographic analysis is a ring light that clamps to the objective lens. The author 

rarely uses it. It provides uniform lighting, but may not be effective in 

3-20 
 

 

 

 

Tools and Equipment 

Figure 3.16 Illumination sources for use with a stereomicroscope. Clockwise 

from the middle left: a ring light intended to be clamped to the objective lens, a 

ring dark light illuminator on an adjustable stand, a line light source on a 

stand, and a transmitted light platform. 

Figure 3.17 Dual gooseneck fiber optic light guides allow flexibility in 

directional illumination. 

3-21 
 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) (b) 

Figure 3.18 Low-angle grazing illumination (a) is effective in highlighting 

fracture surface topography as shown in (b), which is the fracture surface of a 

ceramic rod broken in flexure. 

highlighting ridges or valleys. Directional illumination can be achieved with 

one or more fiber optic gooseneck light guides that direct an intense light beam 

onto the fracture surface at an angle. Dual gooseneck guides from a single light 

source as shown in Figure 3.17 are convenient so that one guide may be used 

for illuminating the key features while the second can be used as filler. 

Low-angle grazing illumination is very effective for accentuating peaks and 

valleys on fracture surfaces as shown in Figure 3.18. The ability to adjust the 

angle of the illumination source is critical. Easily-adjustable holders are a great 

time saver. 

This is one major revision in this 2nd edition of this Guide as compared to the 

first edition. Low-angle grazing illumination was erroneously termed “vicinal 

illumination” in the first edition. Vicinal illumination is different. 

“Vicinal illumination,” illustrated in Figures 3.19 and 3.20, is helpful in 

detecting cracks that may be difficult to see by other means in a translucent 

material. A bright light illuminates an area near to (“in the vicinity”), but not 

directly on the area to be studied. Light is scattered sideways in the translucent 

material. If the light encounters a crack, much of the light is reflected. Only 

some is transmitted across the crack. This creates a sharp delineation between 

light and dark areas that highlights the crack. 

3-22 
 

 

 

 

 

 

 

Tools and Equipment 

(a) (b) (c) 

Figure 3.19 Vicinal illumination of a PEG dental resin with amorphous 

calcium phosphate filler. The disk specimen (a) broke in a biaxial strength test, 

but there was insufficient stored elastic energy to fully break the specimen. (b) 

is a close-up of one crack tip, and (c) is a close-up of the tensile surface at the 

fracture origin. The flaw is a trapped hair or fiber. This was confirmed by 

examination of the fracture surfaces once the specimen was snapped in half by 

hand. 

(a) (b) 

Figure 3.20 Bright field reflected light (a) and vicinal illumination (b) of a 

multilayer ceramic capacitor. Both images are at the same magnification. 

Only one crack is apparent in (a). Vicinal illumination revealed the second 

crack on the right in (b). These are from a compound optical microscope. The 

aperture was stopped down for the image in (b). (Courtesy of S. Hull, NASA, 

Goddard Research Center.) 

3-23
 
 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Dark field illuminators such as shown in Figure 3.21 may also be handy. The 

dark field illuminator is similar to a normal ring light, but the light is directed 

sideways from the ring towards an object in the middle. It is in essence lowangle grazing illumination from all directions. This sideways illumination 

accentuates subtle surface features and scratches. 

Figure 3.21 Dark field illumination of the fracture surface of a fractured steel 

railroad coupler pin. The arrow points to a fatigue crack. 

Some stereomicroscopes have a light source in the base for transmitted 

illumination that is useful for glasses and transparent ceramics. Illuminated 

bases as shown in Figure 3.22 may also be used. 

Some translucent or transparent specimens may be illuminated by 

“transillumination.” The fiber optic light source is directed through the 

material from one side. Usually the side surface nearest the light guide will be 

washed out, but internal features and subtle fracture surface features may stand 

out elsewhere on the fracture surface. Figure 4.1 in the next chapter shows 

examples. 

3-24 
 

 

 

 

Tools and Equipment 

(a) (b) 

Figure 3.22 Transmitted light can be from an illuminator plate as shown in (a) 

or in (b), an improvised stage made of Plexiglas, and white paper or a mirror 

for transmitted illumination. A fiber optic light source can be directed onto 

the tiltable inclined plane so that light is reflected up through a specimen 

mounted on the top. (Thanks to Dr. D. Shechtman for suggesting this to the 

author.) 

Figure 3.23 Crossed polarizing filters or plastic sheets. Polarizers can either 

be dedicated accessories for the stereoscope, or can be improvised from 

ordinary single lens camera lens attachments or sheets of plastic polarizing 

film. 

3-25
 
 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Polarizers (Figure 3.23) may be valuable for examining transparent materials. 

They may reveal residual strains in glasses or twins in single crystals. One 

polarizing filter may be placed between the transmission light source and the 

specimen and an analyzer (another polarizing filter) attached to the objective 

lens. The analyzer can be rotated as required to achieve the desired color or 

brightness differences. Polariscopes are discussed in section 3.25 below. 

3.9 Coatings 

Coatings may be applied to specimens that have poor reflectivity or which are 

translucent and scatter light from beneath the surface, thereby washing out the 

surface detail. A thin layer of gold or platinum may be sputtered onto the 

specimen surface in a small vacuum chamber. Such equipment is commonly 

used to prepare a specimen for scanning electron microscope examination, but 

the coating may enhance optical examinations as well. Figure 3.24 shows an 

example. 

Gold and platinum coatings may be removed without harm to most ceramic 

specimens by soaking the pieces for about 1 minute in “aqua regia.” This is a 

combination of 1 part by volume of nitric acid and 3 to 4 parts hydrochloric 

acid. This is a hazardous combination and appropriate precautions must be 

taken (e.g., gloves and eye protection). Only mix the amount to be used in a 

cleaning session. Never store it, since it releases gasses that build up pressure 

in a sealed container. The specimens should be rinsed in water afterwards. 

When disposing of used aqua regia, water should not be added to the acid mix. 

Instead, the acid mix should be carefully poured into a large container of water 

or a sink with ample flowing water. 

Aqua regia can cause harm to a specimen if the specimen contains anything 

metallic, such as a metallic-inclusion fracture origin. The author once had a 

highly-filled resin matrix dental composite bend bar that had some trapped 

contamination covering the origin during the initial SEM examination. The 

original gold coating had to be removed, the specimen recleaned and recoated 

for a second SEM examination. The fracture origin was a calcium rich 

inclusion. The aqua regia not only removed the gold coating, but it also reacted 

with the fracture origin forming calcium chloride. The EDX analysis during 

the second SEM session showed chlorine at the origin, which was puzzling 

until the author remembered the aqua regia cleaning step and realized that the 

HCl had reacted with the flaw. 

3-26 
Tools and Equipment 

(a) 

(b) 

Figure 3.24 Two views of a fractured yttria-stabilized tetragonal zirconia disk 

with reflected light illumination. (a) shows the uncoated fracture surface and 

(b), the same region after a thin gold coating was applied. 

3-27 
 

 

 Fractography of Ceramics and Glasses 

Figure 3.25 A simple but effective coating technique for white ceramics is to 

stain the surface with a green felt tip pen. The staining can be watched in real 

time through the microscope. The green stain can be easily removed with 

ethanol alcohol on a cotton swab applicator and the process repeated. 

3-28 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

Tools and Equipment 

A simple expedient that works very well for many white or gray ceramics is a 

common felt tip pen as shown in Figures 3.25 and 3.26. Green felt tip pens 

work the best, partly because the eye is most sensitive to green wavelengths, 

but probably also due to the character of the green dye. b The dyes can easily 

be removed with ethanol within a few seconds or minutes and the process 

repeated in a matter of seconds. Most of the dye will come off, but some may 

penetrate into crack branches or undercuts in the fracture surface and may be 

more persistent. If the green dye is allowed to sit for longer periods of time, it 

can be removed by soaking the part for 1 to 2 minutes in a 30% solution of 

H2O2 (hydrogen peroxide). Note this is a much greater concentration than is 

use for topical cleaning of wounds. Gloves and eye protection should be used. 

Additional dye and staining techniques are discussed in section 3.24 later in 

this chapter. 

(a) (b) 

Figure 3.26 An example of green dye staining on a dental glass ceramic bend 

bar. (a) shows the entire fracture surface illustrating how the stain drains into 

valleys and accentuates hackle lines and ridges and how easily the fracture 

mirror region stands out (arrow). The bottom view shows the fracture origin: 

a surface crack (arrows) made by Knoop indentation for a fracture toughness 

measurement. The pictures were medium resolution video prints and the thin 

horizontal line is an artifact from a defect in the print head. 

b Papermate Flair, 1.1 mm, Sanford Corporation, Division of Newell, Rubbermaid. 

3-29 
 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

3.10 Size Measurements 

Sizes may be measured by means of a graduated reticule in a microscope 

eyepiece. The graduations must be calibrated by means of a stage micrometer 

such as shown in Figure 3.27. Users should be cautious about the accuracy of 

stage micrometers. Some are rudimentary and only meant to furnish an 

approximate magnification for photography and can be as much as several 

percent in error. Others may be very accurate. If a stereomicroscope has a 

zoom capability, it is best that the body have discrete click steps or detents so 

that precise magnifications can be calibrated and recorded at the discrete steps. 

(a)

 (b) 

Figure 3.27 Two stage micrometers (a). A close-up of the bottom one is in (b). 

The numbers show 1 mm spacings. The thin lines on the top right are 100 μm 

apart. The thin lines on the top left (arrow) are 10 μm apart. This model was 

accurate to an amazing 0.01 %. 

3-30 
 

 

 

 

 

 

 

 

Tools and Equipment 

A powerful accessory to a microscope is a traversing X-Y stage with 

micrometer positioning heads such as shown in Figure 3.28. This accessory is 

especially valuable for precisely aligning a specimen in the field of view, 

particularly when photographing an image. (A lot of time can be lost when 

nudging a specimen around with one’s hands.) The micrometer heads, in 

conjunction with a crosshair in one eyepiece, also facilitate accurate size 

measurements of features. One advantage of such a system is that if the 

fracture surface is viewed directly from above through both eyepieces, and the 

fracture plane is horizontal, then the micrometer-measured sizes need not be 

corrected for foreshortening, as long as the surface being examined is parallel 

to the X-Y stage. 

If size measurements are made on recorded images, care must be taken to 

ensure that the camera axis is near to perpendicular to the fracture surface 

plane. Cameras mounted on a stereomicroscope view the fracture surface at an 

small angle to the fracture surface, so its view may be foreshortened. The 

specimen may have to be tilted a few degrees to obtain a true view of the size 

of the fracture surface. So, for example, one can tilt a bend bar sideways a little 

bit until the camera does not “see” the sides of the bend bar. 

An additional advantage for the X-Y stage measurement is that the size 

readings are independent of the magnification used to view the specimen. In 

contrast, images captured by the camera must be calibrated for the exact 

magnification. 

Figure 3.28 A X-Y traversing stage with digital micrometer heads is very 

convenient for positioning a piece and making accurate size readings. 

3-31 
 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

One drawback to the traversing stage is that it takes up vertical space and hence 

the microscope body must be moved up higher and possibly may limit the 

working range under the objective if the microscope is mounted on a short post. 

As noted above, stereomicroscopes for fractographic examination should be 

mounted on tall posts for maximum versatility. 

3.11 Compound Optical Microscope 

Compound optical microscopes such as shown in Figures 3.29 and 3.30 may be 

used for fractographic analysis, but are less versatile than stereo-microscopes. 

Compound microscopes are best suited for small specimens with relatively flat 

surfaces. They magnify much more (up to 1600X) than a stereomicroscope, 

but three-dimensional topographical effects are lost. Virtually all work is now 

done with air lenses. Oil immersion lenses are no longer needed. 

Objective lenses from 5X to 100X may be used in the lens turret, but in practice 

the limited depth of field makes it difficult to see a fracture surface in focus at 

larger magnifications. Only portions of a fracture surface may be in focus 

even at lower magnifications. Working distances are also very small, and care 

must be taken not to contact the objective lens with a specimen. Long focal 

length objectives are ideal. These have greater working distances, but at the 

cost of some loss of light collecting power and resolution. The author uses a 

compound microscope with long focal length lenses. The 100X objective has a 

1.3 mm clearance and 0.8 numerical aperture. Some 100X lenses are available 

with working distances as long at 13 mm, but these have lower numerical 

apertures (0.5 to 0.6) and thus less resolving capability. The clearance between 

a stage and the objective lenses may also preclude examination of large 

fractured pieces. Low-angle grazing illumination is difficult to achieve with 

compound microscopes, especially at greater magnifications. 

Dark field illumination may be effective for examining cleavage planes in 

single crystals. Direct reflected light illumination may be too bright or wash 

out details due to the mirror-like reflection from a cleaved crystal surface. 

Dark field illumination blocks the central portion of the illumination light 

coming through an objective, allowing only oblique rays in a ring to strike the 

specimen. Only light that is scattered by features on the fracture surface is 

reflected back up into the objective for the viewer to see. Fracture surfaces 

appear as small bright lines and spots against a dark background. Dark field can 

highlight subtle cleavage steps and make the symmetry of a fracture mirror 

more evident. 

3-32 
 

 

Tools and Equipment 

Figure 3.29 Compound reflected light microscope. Although the microscope 

has twin eyepieces, they both view the same image collected by the objective 

lens. Overall magnifications can be as much as 1000X, but the working 

distances and the depths of field are very small. Only relatively flat specimens 

may be examined. 

3-33 
 

 

 

 Fractography of Ceramics and Glasses 

Figure 3.30 A research compound microscope with a digital camera and 

computer. Linear encoders with a 1 micrometer resolution are attached to the 

traversing X-Y stage and facilitate rapid size measurements. The instant film 

camera on the top was a common accessory in the 1970s to 1990s, but has been 

rendered obsolete by digital cameras such as the one marked by the arrow. A 

fractured glass rod mounted on blue polymeric clay is on the stage. The 

fracture mirror can be viewed through the eyepieces at the same time it is 

projected on the monitor. 

3-34 
 

 

 

 

 

 

 

 

 

 

 

 

 

Tools and Equipment 

Compound microscopes are useful for measuring very small features, such as 

tiny fracture mirrors in optical fibers or other strong test specimens. 

Compound microscopes can also be used in transmitted illumination mode. 

These microscopes are also optimal for measuring fracture mirror sizes in 

glasses. 

Compound microscopes may be very useful for examining fine detail in 

replicas. Thin polymer replicas of curved surfaces may be flattened by taping 

the ends of a replica on a glass microscope slide, or the replica may be placed 

between two glass slides so that the curved surface can be flattened and 

examined at high magnification in either transmitted or reflected light. 

Compound microscopes are essential for examining polished specimens for 

microstructural analysis, a useful adjunct to the fractographic analysis. 

3.12 Replication Equipment 

Replicas of fracture surfaces may be valuable in some instances. Some 

fractured parts may be so large or unwieldy that it may be difficult to position 

them under an optical microscope. In other instances, the fractured part may be 

a critical piece of evidence that cannot be removed to the fractography 

laboratory. Replicas can also be an important backup source of critical 

fracture surface information if the component is lost or damaged. They can be 

valuable teaching aids in a course on fractography since every student 

gets to examine the same object with the same features! 

In some cases, fracture surface features may actually be easier to observe on the 

replica than on the original surface. In many transparent or translucent 

ceramics such as porcelains or coarse grain ceramics, light scattering from 

below the surface can obscure the fracture surface details. For example, subtle 

wake hackle lines behind pores may be easier to see on replicas. Curved 

surfaces can be made more amenable to high magnification examination, since 

a replica can be flattened. Since some replicas are a negative of the fracture 

surface, depressed features such as pores or cracks are converted to raised 

features on the replica that sometimes are more easily imaged in the scanning 

electron microscope. 

There are a variety of materials that may be used to make a replica as shown in 

Figures 3.31 to 3.35. Two excellent summary papers with practical tips are by 

Varner.1,2

 The replica must be made without the introduction of artifacts. The 

replica material must also not damage the original material nor chemically react 

with it! 

3-35 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Figure 3.31 Thin cellulose acetate replica tape, acetone, applicator, and bend 

bar specimens. These are used to make reversed (negative) replicas. 

There are two types of replicas: “positive replicas” and “negative (or reversed) 

replicas.” Some replica materials are laid over a fracture surface then removed 

after they dry sufficiently to retain the critical surface features. The material 

preserves a mirror image of the fracture markings: what was left is now right, 

and so forth. This is a “reversed replica” or a “negative.” Other replicas may 

be made in multiple steps to create a “positive replica.” So for example, a 

silicone impression or mold may be made around an object, and then pulled off 

the fracture surface. An epoxy resin is then cast into the mold and allowed to 

dry. The mold is removed, and now the epoxy represents a positive replica 

(i.e., a clone) of the original part. The dental materials community has a good 

line of products suitable for making mold impressions of teeth and these 

materials are suitable for general fractographic applications as well, as 

discussed by Scherrer et al. 3 Fréchette also describes replication techniques in 

his book.4 

Cellulose acetate tapes or films (such as used in transmission electron 

microscopy) are convenient replica materials that may be made at room 

temperature as shown in Figure 3.31. The specimen surface should be clean. 

A short strip of 0.13 mm to 0.25 mm thick tape is cut from the sheet or roll and 

then moistened on one side with acetone. Thick tapes or films may be 

3-36 
 

 

 

 

 

 

Tools and Equipment 

(a) 

(b) 

(c) 

Glaze 

Porcelain 

Origin 

Figure 3.32 Cellulose acetate replica of a fracture surface of a porcelain 

electrical insulator. (a) shows the replica taped to a glass microscope slide. 

(b) and (c) show close-ups of the replica which reveals the fracture mirror and 

an inclusion flaw. 

3-37 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

immersed in acetone for 10 s to 15 s, whereas thin tape may be wetted by a few 

drops from an eye drop applicator. The specimen surface may also be wetted 

by acetone, although that step is not always necessary. The tacky film or tape 

is applied to one side of the fracture surface and then is gradually brought into 

contact with the remainder. The replica material should be applied so as not to 

entrap any air. Pressure is applied to the replica for 15 s up to two or three 

minutes as required, taking care to not slide or shift the replica material. A 

small pad (or if the specimen is small enough, a finger or thumb) may be used 

to apply the pressure. If a finger is used, a folded tissue or small pad between 

the finger and the tape should be used so that one’s fingerprints do not get 

superimposed on the back side of the replica tape! The pressure is carefully 

removed without disturbing the replica which remains in contact with the 

fracture surface for 10 min to 15 min more drying time. The replica is peeled 

from the fracture surface (or it may pop off) and may be trimmed with scissors 

or mounted on a glass microscope slide such as shown in Figure 3.32. 

Cellulose acetate replicas may be viewed with reflected or transmitted 

illumination in an optical microscope. An interesting advantage of a replica 

over an original surface is that even if the original surface is curved, the replica 

can be flattened out and examined in a compound optical microscope at greater 

magnifications (in either reflected or transmitted light mode) than if the replica 

were only examined with a stereoptical microscope. Gold or carbon coatings 

can be sputtered onto negative replicas to enhance viewing in reflected light or 

to prepare the replica for SEM examination. 

Since “reversed” replicas like the cellulose acetate tapes are flipped versions of 

the fracture surface, it is prudent to label them clearly as to their location on the 

component, lest the fractographer confuse top-bottom, or left-right orientations. 

Special care should be taken to avoid artifacts in the replicas such as trapped air 

pockets, tears, or handling artifacts. Replicas should be kept clean and 

carefully stored in containers. 

If it is possible to warm a specimen, then negative replicas made with the 

thermoplastic compound polyvinyl chloride (PVC) (180o

C) or, alternatively, 

silicon elastomers may also be used (100o

C).4 If PVC is used, the specimen is 

heated and the PVC applied with a Teflon rod. 

Positive replicas may be made with filled silicone rubbers, unfilled (transparent 

silicone rubber, and polyvinylsiloxane (PVS).1,2,3 The latter material is 

commonly used in dentistry to make impressions in patient’s mouths as shown 

in Figures 3.33 and 3.34. Extraction of a damaged dental crown from the 

mouth typically requires its destruction. Therefore it is prudent to make a 

3-38 
 

 

 

Tools and Equipment 

(a) 

(b) 

Figure 3.33 Dental replica equipment. (a) shows a typical PVS applicator 

gun which has interchangeable tips and compounds. (b) shows Dr. S. Scherrer 

demonstrating the technique in a dental materials fractography class. 

3-39 
 

 

 

 

 Fractography of Ceramics and Glasses 

replica before the crown is extracted as discussed by Scherrer et al 3 and shown 

in Figure 3.34. Sometimes just one PVS compound is adequate, but in other 

cases it is optimal to use two PVS compounds. The first is a very fluid PVS 

that can retain very fine detail (e.g., Ultra-Light Body) when applied to the 

fractured part. Once it sets, a second PVS that is a fast setter is applied over the 

first layer. It is used to build up a larger, more handleable mold. 

Once the mold has set, an epoxy resin may be poured into it. One interesting 

advantage of this procedure is that the mold can be reused to make multiple 

epoxy replicas. They are useful as teaching aids since every student examines 

the same object with the same features. Gold or carbon coatings can be 

sputtered onto positive replicas to enhance viewing in reflected light or to 

prepare the replica for SEM examination. 

Sections on replication techniques are in both the Visual Examination and 

Light Microscopy and the Transmission Electron Microscopy (TEM) chapters 

in the ASM Handbook on Fractography.5 

In general, some practice is needed to create successful replicas with good 

detail and without artifacts. Before a critical component is replicated, it is 

prudent to practice with the materials to optimize one’s procedure. 

3-40 
 

 

 

Tools and Equipment 

(a) Top view of the fractured all-ceramic 

crown (first upper left molar) in the 

patient’s mouth. The broken crown is 

cleaned with alcohol on a cotton swab. 

(b) The crown is water spray rinsed and air 

dried. 

(c) A light body precision impression material 

is injected around the fracture surface. 

(d) The whole crown is covered with more 

impression material. 

(e) The mold is removed. 

(f) A positive replica is cast with epoxy resin. 

It is then gold coated for SEM 

examination. 

Figure 3.34 Steps in the preparation of a dental replica that is a positive 

duplicate of a fractured crown. Photos a-d were taken in the patient’s mouth. 

Figure 6.80 has additional views of this fractured alumina crown. (Photos and 

technique courtesy of S. Scherrer, Univ. Geneva) 

3-41 
 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) 

1 2 3 

epoxy replica 

cellulose 

acetate 

replica 

(b) (c) 

 (d) 

Figure 3.35 Replicas of a broken whiteware plate. (a) shows both a cellulose 

acetate tape replica on the left, and a dental PVS mold and epoxy positive 

replica on the right. The PVS mold was cut into 3 pieces (labeled 1,2,3) and 

the middle one gold coated. (b) shows the cast epoxy resin positive from the 

plate rim area. (c) shows a close-up of the epoxy replica. Telltale small wake 

hackle lines (described in chapter 5) are seen in the glaze (arrows). These 

little “weathervanes” are more easily seen in the glaze than in the whiteware 

body and they help direct one’s attention back to an origin area shown in (d) . 

Coarse hackle lines (chapter 5) lead back to the origin site, a contact damage 

crack on the rim marked by the arrow. Coating the clear epoxy with a thin 

gold coat would have made the markings even easier to interpret. 

3-42 
 

 

 

 

 

 

 

 

Tools and Equipment 

Figure 3.36 A conventional scanning electron microscope. 

3.13 Scanning Electron Microscope 

Despite the versatility of optical microscopy, there inevitably will be cases 

where the fracture initiating flaw is too small to be clearly discerned and the 

scanning electron microscope (SEM) is required (Figure 3.36). The SEM is a 

complementary tool to the optical microscope. They “see things” differently. 

It is best to examine specimens with the optical microscope first, since much 

valuable information such as color, reflectivity, and internal flaws in translucent or transparent materials cannot be detected by the SEM. Nevertheless, 

the SEM is a versatile tool that can view large portions of a specimen at low 

magnifications and yet instantly zoom in for high-resolution close-ups of 

regions of interest. It can also provide compositional information. The 

practical limits of magnification for fractographic analysis are of the order of 

10X to 20,000X. 

Modern SEMs are not difficult to use and materials scientists and engineers can 

perform their own routine examinations without the need of an expert operator. 

The advent of commercial SEMs and transmission electron microscopes in the 

1960s led to a quantum jump in the quality of fractography of ceramics and 

glasses. Researchers and engineers could finally see and characterize the 

strength-limiting flaws that previously were too small to be clearly discerned 

with optical microscopy. It is fair to say that the SEM opened up a whole new 

3-43 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

world of possibilities for fractographers, processors, and mechanical testers and 

led to many major advances in materials science and engineering.c 

The SEM uses electrons as a signal source and, since the probe size can be 

made smaller than the wavelength of light, smaller features can be resolved and 

much greater magnifications achieved. A vacuum is necessary to remove gas 

that would interfere with the electron beam. The SEM has superb depth of 

field over a broad magnification range. For a working distance of 20 mm, the 

depth of field may be 0.4 mm at 100X or 40 μm at 1000X and 4 μm at 

10,000X.6 Sometimes the SEM is preferred for this reason at even low 

magnifications (100X to 500X) that are well within the range of optical 

microscopes. On the other hand, SEM images often look flat and tend to deemphasize large surface topography changes. The SEM image can also be 

deceptive and it may be confusing to interpret whether a feature projects above 

the surface or is a depression below the surface. Stereo SEM viewing, 

described below, can help solve this dilemma. A great advantage of the SEM is 

that compositional analysis can also be done using accessories that detect Xrays emitted from the fracture surface. 

Specimens must fit into the vacuum chamber and are usually cut and mounted 

on a small aluminum stub that will be placed into the SEM as shown in Figure 

3.37. The SEM should have a large opening to accommodate broken 

fragments and test specimens without having to cut or trim them to small sizes. 

The stage has external controls that permit the specimen to be positioned in X, 

Y, and Z directions, rotated, and tilted. The size of the chamber and the mount 

station, plus the clearance to the electron optics all put limits on the size of 

specimens than can be examined. A good SEM for fractography has a large 

chamber, a large entry port, and can pump down with a minimum of difficulty. 

Specimens on an aluminum stub are shown in Figure 3.37(b) and 3.38. It is 

convenient to cut the specimens so that they will sit square and facing upward. 

Irregular specimens can be propped up by small metallic pieces. It is best to 

cut multiple specimens to have similar lengths to minimize refocusing time 

when viewing in the SEM. A small laboratory cut off wheel with a thin 

diamond wafering saw is also handy for cutting long specimens into 

manageable lengths for the SEM. The specimens should be cut so that the 

fracture surface is approximately perpendicular to the incident beam in order to 

avoid distortions from foreshortening. Once in the SEM, the specimen 

The coincidental advent of modern fracture mechanics in the late 1950s, described 

in chapter 7, also led to major breakthroughs in the study of ceramics and glasses. 

3-44 

c 
 

 

Tools and Equipment 

(a) 

(b) 

Figure 3.37 (a) shows an open SEM vacuum chamber showing the specimen 

station on the left and the positioning controls on the right, outside of the 

chamber. (b) shows a close up of the two fracture specimens on their stub on 

the examination mount. Both fragments had to be cut from longer pieces to fit 

in the chamber. This model has a large opening and the stage can 

accommodate specimens as tall as 25 mm. 

3-45 
 

 

 

 

 

 Fractography of Ceramics and Glasses 

Figure 3.38 Two matching fracture halves of a broken round tension specimen 

on a stub in a protective plastic box. Silver paint was used to attach the 

specimens to the stub. Ideally, the halves should have been cut to similar 

lengths to save time refocusing in the SEM. 

Figure 3.39 SEM specimen preparation tools. Clockwise from the top left are 

two SEM length calibration standards, two silver paint bottles, a double-sided 

conductive adhesive tape roll, two sheets of double-sided conductive adhesive 

disks, and an assortment of aluminum stubs. 

3-46 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Tools and Equipment 

can be tilted to enhance the emissions and signal responses if necessary. 

Figure 3.39 shows some common accessories and tools for SEM examination. 

The specimens must be cleaned carefully of greases, dust, or debris. It can be 

exasperating to find a tiny piece of dust right on top of the origin site when the 

specimen is in the chamber. Soaking the specimen in a small beaker of acetone 

in an ultrasonic bath for a minute or so is often effective. After removal, a 

gentle blast of compressed air can eliminate any acetone residue. Sometimes it 

may be sufficient to swipe the fracture surface with a cotton swab soaked in 

acetone. Small trays or plastic boxes are convenient for holding the specimens 

and protecting them from contamination prior to examination. Special 

containers are available for mounting stubs. The best procedure is to keep the 

specimens as clean as possible from the moment the fragments are recovered. 

There must be a good electrical contact between the specimen and the stub. 

The specimen can be affixed to the stub by silver paint or by conductive carbon 

double-sided sticky tape. If multiple specimens are placed on a stub, it is a 

good idea to make a hand drawn sketch or map to show which specimen is 

where on the stub, lest the viewer become confused when viewing the array in 

the SEM. Furthermore all the specimens should be oriented in a similar way. 

For example, if a series of bend bar fracture surfaces are being examined, they 

all should have their tensile surfaces facing the same way. Otherwise, much 

time can be lost getting oriented or searching the wrong areas during the SEM 

session. A set of optical photos of the fracture surface can be handy when the 

specimens are examined in the SEM. 

The next step for most SEM examinations is to apply a thin 5 nm to 40 nm 

conductive gold, gold-palladium, or carbon coating by means of a sputter 

coater shown in Figure 3.40 or a thermal evaporator. The former is preferred 

for rough surfaces. The coatings are needed for ceramics and glasses in order 

to eliminate surface charge build up from the electron bombardment. For most 

applications, the coating does not interfere with the image. A coating is not 

needed for conductive ceramics or for examinations with low accelerating 

voltages (1 keV to 5 keV). Low accelerating voltage examinations may have a 

diminished resolution, however. Carbon coatings are used if X-ray information 

of lighter atomic weight elements is desired. Coating with gold or carbon is a 

simple procedure that only takes a few minutes. Sometimes a gold coating may 

be applied to facilitate some optical microscopy that is scattered back up 

through a fracture surface in translucent or transparent specimens (see Section 

3.9, page 3-26). The sputter coater will preferentially coat from one direction, 

and sometimes it is advantageous to position the specimen and the stub at an 

angle so that shadow coating will occur on the surface. Slight charging on 

3-47 
 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Figure 3.40 A gold sputter coater used to apply conductive coating. It is 

helpful for both optical and scanning electron microscopy. 

some edges of the fracture surface may enhance the image by promoting 

contrast. Coating along the specimen sides may be incomplete, so sometimes it 

is wise to add a thin silver paint line along the specimen sides from the fracture 

surface to the stub. Of course, the silver paint should be kept well away from 

the fracture origin or any features of interest on the fracture surface. The 

conductive coating step may be omitted for some modern SEMs that operate at 

low accelerating voltages and low vacuums as described in a later section on 

environmental SEMs. Additional details and tips on coatings for SEM 

examination may be found in the paper by Healy and Mecholsky.7 

The SEM can be operated in several modes. A focused beam of high-energy 

electrons (typically 1 keV to 30 keV) is rastered over a rectangular portion of 

the specimen. Some electrons interact with the specimen surface and some 

penetrate into the volume. Figure 3.41 shows that the interaction volume is 

larger than the size of the incident beam, which may be as small as 10 nm 

diameter.

8 The three most important signals are secondary electrons, 

backscattered electrons, and X-rays. Secondary electrons reveal the surface 

topography. Backscattered electrons give a combination of the topography and 

atomic number and crystallographic information. X-rays reveal compositional 

information. 

3-48 
 

 

 

 

 

 

Tools and Equipment 

Figure 3.41 The electron beam interacts with the specimen surface and volume 

creating emissions that may be analyzed. Signals are generated in all 

directions, but they come from different depths of the material. The beam is 

rastered over the surface and an image is sequentially formed. 

The initial working distance can be adjusted to enhance the field of view and 

the depth of field. Long working distances allow a lower minimum 

magnification which is sometimes helpful in initial orientation or in 

photographing the whole part or as much as possible. A longer working 

distance also allows the specimen to be tilted more. On the other hand, longer 

working distances reduce resolution and may reduce X-ray signal acquisition. 

Image clarity is also affected by specimen tilt. As noted above, having the 

surface perpendicular to the incident beam minimizes foreshortening, but 

secondary electron and X-ray collection can be maximized by tilting the 

specimen towards the detectors. 

Initial inspections on older SEMs with a cathode ray tube (CRT) display often 

are typically with 500 line scans, but when a image is recorded on film, a 

higher resolution CRT with 2000 or more line scans is used. The raster of the 

beam on the specimen is synchronized with that of the electron beam that 

rasters across the display CRT. Modern SEMs with digital monitors also have 

differences in resolution between “live image” displays and captured images. 

3-49 
 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

The secondary electron mode of operation provides the best spatial resolution 

and produces a surface topographical image. Secondary electrons are low 

energy electrons (10 eV to 50 eV) primarily produced from the material close 

to the surface. Since they are low energy, a positively charged detector off to 

one side can capture the emitted electrons for the image. These electrons may 

even follow a curved trajectory from the surface to the detector. Edge charging 

can be a problem with very rough surfaces or those with sharp corners and 

cracks, in which case gamma adjustments to the image may help. Gamma 

adjustments can suppress bright areas and enhance dark areas. For some edge 

charging cases, backscattered imaging may be preferred. 

Backscattered electrons are higher energy electrons (5 keV to 40 keV), up to 

the energy of the incident beam and are emitted from a larger volume than the 

secondary electrons. The backscattered electrons travel in straight lines and are 

registered by a different detector that may have to be slid into proximity with 

the fracture surface. They often have a central hole for the incident electron 

beam. Backscattered electrons include both topographical and compositional 

responses. Detectors can be split ring or four quadrant (or more) versions and, 

by turning different elements on or off in the detector, it is possible to 

emphasize topographical or compositional information. Backscattered electron 

imaging is effective in detecting phase distributions in heterogeneous or 

composite microstructures and also aberrant features such as inclusions or 

second phase concentrations. Such information is often completely missed by 

secondary electron imaging. 

Figures 3.42 through 3.45 show matching secondary electron and backscattered 

electron images of fracture origins. Figure 3.42 shows how the backscatter 

mode can even help highlight a controlled semi-elliptical flaw made by a 

Knoop indenter. 

X-ray emissions provide valuable information about the elemental composition. 

This energy-dispersive spectroscopy by X-ray analysis (EDS or EDX) is 

especially useful for detecting and characterizing inclusions, second phase 

variations, or compositional inhomogeneities, which all may act as strengthlimiting flaws. The analysis is over a region of the order of size of a few 

micrometers, but penetrates into the interior 1 μm to 2 μm. With most SEMs 

and with coated surfaces, elements from sodium (Z = 11) or heavier can be 

detected simultaneously. Lighter elements such as boron or carbon are not 

readily detected. A better resolution of the lighter elements is obtained with 

very thin gold or carbon coatings, or even uncoated specimens with low 

accelerating voltage. Newer detectors and vacuum-tight polymer windows 

3-50 
 

 

 

 

 

 

 

 

 

Tools and Equipment 

200 μm 

200 μm 

(a) (b) 

Figure 3.42 Comparative secondary electron (a) and backscattered electron 

(b) images of the fracture surface of a Knoop semi-elliptical precrack in a 

silicon nitride fracture toughness test specimen. The edge charging in (a) 

emphasizes certain features, but (b) shows that the flattening effect of the 

backscattering mode tends to better reveal the overall crack (arrows). 

(a) 

(b) (c) 

Figure 3.43 Backscattered mode electron scanning electron microscopy is 

effective in showing compositional difference of fracture origins flaws relative 

to the background. (a) is an optical image of a silicon nitride with 30 mass % 

TiN bend bar illustrating the origin has a gold color (722 MPa). (b) is an 

SEM secondary electron image of the origin. (c) is a backscattered electron 

image of the flaw. This origin may be termed either a TiN agglomerate or a 

compositional inhomogeneity. (Courtesy J. Kübler, EMPA). 

3-51 
 

 

 Fractography of Ceramics and Glasses 

(a) 

(b) 

(c) 

Figure 3.44 A cluster of iron particles acting as an origin in a silicon nitride 

rod broken in flexure. The origin is below the surface. The stress at the origin 

was 682 MPa. (a) and (b) show secondary electron images and (c) is the 

backscattered image at the same magnification as (b). 

3-52 
 

 

Tools and Equipment 

(a) (b)

 (c) 

Figure 3.45 An iron inclusion in a silicon nitride rod (684 MPa). Matching 

secondary electron (a) and backscattered electron images (b). (c) is the EDX 

pattern of the inclusion revealing it is iron. 

3-53 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

allow lighter elements including boron, carbon, nitrogen and fluorine to be 

detected with modern equipment.9 

An X-ray scan may require several minutes as the electron beam illuminates the 

surface and the detector accumulates a sufficient signal of a range of emitted Xrays. A typical scan appears as shown in Figure 3.45c where signal intensity is 

plotted versus X-ray energy or wavelength. The specimen response is 

compared to known characteristic wavelengths of the various elements to 

identify which elements are present. A common procedure is to scan the 

critical feature, then rescan a different area away from the feature and compare 

the patterns. The accelerating voltages should be high enough to activate 

emissions from the elements. This should be 1.5 times the threshold energy of 

the element. So for example, to detect copper which has Kα X-rays of 8.0 keV, 

it is best to use an accelerating voltage of 12 keV or greater. Otherwise if the 

voltage is too low, then the X-rays will not be emitted or detected, even if 

copper is present.10 Similarly, for zirconium, which has Kα X-rays of 15.7 

keV, it is best to use an accelerating voltage of 25 keV or greater. 

X-ray compositional maps, tailored to a particular element, may be made of the 

region being examined. The areas that contain the element in question emit 

X-rays and create bright spots on the map, whereas regions without the element 

appear dark. Digital image maps may be constructed with false colors 

representing areas of similar composition. 

Figures 3.45a,b shows matching secondary electron and backscattered electron 

images of an inclusion in a silicon nitride tension specimen. Figure 3.45(c) 

shows the EDX analysis from a close-up view of the inclusion. In this case the 

iron peak stands out prominently. At other times, the inclusion peaks may 

appear only as a small fraction of the intensity of the primary constituent (Si in 

this case) due to the averaging effect of the area being imaged. Several peaks 

at different energies and intensities characterize each element and sometimes 

these will overlap with peaks from other elements, which can confuse the 

interpretation. Gold or palladium peaks from the coating may also be present. 

EDX analysis is usually semi quantitative, and even under the most ideal 

conditions, the mass fraction concentration is only precise and accurate to 

± 1 %.11 The minimum concentrations that can be detected are 0.5 to 1.0 mass 

fraction % under ideal conditions, but more typically is several mass fraction % 

or more. 

Electron backscatter diffraction (EBSD) is a valuable tool for studying 

microstructures and can discern the variation of crystallographic orientation of 

individual grains, but it requires a flat specimen surface. An example is shown 

3-54 
 

 

 

 

 

 

 

 

 

 

Tools and Equipment 

Figure 3.46 Electron backscatter diffraction map of the flat exterior surface of 

a tungsten carbide/11 mass % cobalt fracture toughness test specimen. The 

black crack has propagated through the microstructure. Individual grains with 

different crystallographic orientations are shown in color. Crack propagation 

was under static loading and was mixed inter- and transgranular, with the 

latter through the grains on [101�0] slip planes. (Image courtesy of R. Morrell 

and K. Mingard, NPL) 

in Figure 3.4612 which images the exterior surface of a fracture toughness test 

specimen. A focused electron beam is scanned across a crystalline material and 

at each point an electron diffraction pattern is created and analyzed. At the 

present time EBSD is not amenable to analysis of rough fracture surfaces. 

Magnifications displayed on the SEM monitor (e.g., 1000X) may not be 

accurate. Brooks and McGill state that accuracies of 5 % are typical.11 The 

superimposed micrometer bars on the photos or images are usually better than a 

nominal magnification, but are not necessarily exact either. If quantitative 

measurements are made with the SEM, then care should be taken to ensure the 

fracture surface is not foreshortened. An SEM length calibration standard such 

as shown in Figure 3.47 should be used to check the magnifications. Some 

older SEMs have magnifications that vary across the field of view and which 

are not the same vertically and horizontally. 

A paper by Healey and Mecholsky on SEM techniques and their application to 

brittle materials is very instructive.7 An excellent, detailed article on the use of 

the SEM for fractography is by Brooks and McGill.11 

3-55 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Figure 3.47 A NIST Standard Reference Material SEM length calibration 

standard. Finely spaced reference lines are located at intervals along the sides 

of the box on the right close up view. 

3.14 Stereo SEM Imaging 

The SEM takes pictures that make the surface look flat and do not show the 

topography clearly. A much better sense of the full topography can be revealed 

by the simple process of taking matching stereo photographs. Two photos of 

the same area are taken, but with the specimen tilted 5o to 15o in the SEM 

chamber between shots. A minor adjustment to the brightness and contrast 

may be necessary for the second photo, but is not essential. Some people are 

able to hold the two images in front of their eyes and focus the left eye on the 

left photo and the right eye on the right photo. The distance between the photos 

is adjusted and the eyes are “relaxed.” The brain merges the images and 

interprets them as a single three-dimensional view. Most people find it easier 

to put the images under a stereoscope, which is a stereo viewing optical device 

as shown in Figure 3.48. The fractographer looks into two eyepieces and the 

photos are spaced apart and adjusted slightly back and forth until the threedimensional effect is achieved. It may take a little practice to get the two 

images to register, but the usual trick is to focus each eye on an identical 

feature or spot in the two photos. Initially the two eyes will discern the spot as 

two separate images, but by carefully sliding one or the other photos around, 

the spots can be made to merge. At that instant the three-dimensional effect 

should be detected since each eye will be viewing the same surface, but from 

different perspectives. This may take a little practice, but the viewer helps 

make it easy for most people. One curious effect is that the perception of what 

3-56 
 

 

 

 

 

Tools and Equipment 

features are up or down may change depending upon which photo is on the left 

and which is on the right. It may be necessary to swap the photos to get the 

correct rendition. The orientation of the photos may have to be turned 90o as 

well in order to match the tilt angle used in the SEM. 

Figure 3.48 Stereoscope. The two photos are placed under the viewer and the 

viewer looks into the two eyepieces. The photos are moved around slightly 

until they appear as one image in three dimensions. 

Some modern SEMs can digitally capture the pair or of tilted images, or even 

three images at different tilts. Image analysis software can then interpret the 

multiple images and constructs a pseudo three-dimensional rendition that is 

displayed on a monitor. The computer image can be rotated or displayed from 

different vantage points and height profiles through various sections can be 

displayed. Quantitative analysis may be done by a procedure known as 

stereophotogrammetry. A few applications to ceramics so far were hampered 

by contrast limitations13 or lack of computing power,14 but the latter problem 

probably is less of a concern now. 

3-57 
 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

3.15 Field Emission SEM 

Field emission scanning electron microscopes (FESEM) offer dramatic 

increases in magnifications (600,000X) compared to the SEM and improved 

resolutions down to 1 nm (Figure 3.49). They rival the resolution and contrast 

of many transmission electron microscopes. The field emission cathode 

replaces conventional tungsten cathodes on the common SEM. A smaller 

probe beam is used. Accelerating voltages are lower and conductive coatings 

are not needed in many cases. Energy dispersive X-ray spectroscopy (EDX) 

can be applied to smaller regions than in a conventional SEM. The FESEM is 

used much the same way that a conventional SEM is used for fractographic 

analyses. One problem with such microscopes is that the working distances 

and the pre-evacuation chambers are very small and it may not be possible to 

get large pieces into the microscope. The author used such a FESEM to take 

Figures 6.69 and 8.21 which show origin flaws in single crystal silicon as small 

as 200 nm, possibly the smallest fracture origin flaws ever detected.15 

Figure 3.49 A field emission scanning electron microscope 

3-58 
 

 

 

 

 

 

 

 

 

Tools and Equipment 

3.16 Environmental SEM 

The environmental scanning electron microscope (ESEM) is designed to 

operate either in a high or low vacuum unlike conventional scanning electron 

microscopes, which require a high vacuum. The environmental SEM can 

operate with environments having water vapor or other gases. A specimen 

coating is not needed, as static charges do not develop on the specimen since 

the gasses ionize and can neutralize excess electron charge build up on the 

specimen. Magnifications up to 100,000 are possible, but resolutions are less 

than that from a conventional SEM. One advantage of the ESEM is that 

dynamic processes such as corrosion or chemical reactions can be studied by 

time-lapse imaging. The ESEM is primarily used for biological materials or 

with materials when it is undesirable to apply a coating. 

3.17 Transmission Electron Microscope 

The transmission electron microscope (TEM) is capable of extraordinarily high 

magnifications and high resolutions (< 0.1 nm). It is limited to specimens 

about 0.5 μm or less in thickness, however. The TEM can reveal critical 

microstructural details. Prior to the advent of modern scanning electron 

microscopy, TEM was applied to replicas of fracture surfaces in order to study 

the fracture morphology. Today, the TEM is not used for routine fractographic 

analysis, but is a critical tool for research projects studying the effect of 

microstructure on properties. It can reveal crucial information about defect 

densities in grains, grain boundary sliding and cavitation, twin formation, 

damage accumulation, phase changes, fracture mode (intergranular or transgranular), and grain boundary phases and their thickness. TEM specimens are 

cut from the bulk and then electrochemically or ion milled to very thin foils that 

are transparent to electrons. Alternatively, carbon replicas may be made for 

examination. Since the TEM examines only tiny portions of the material, it is 

almost never used to find and characterize fracture origins or classical fracture 

surface markings. 

3-59 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

3.18 Atomic Force Microscope (AFM) 

The AFM is a powerful microscope that uses a tiny tip (radius about 2 nm to 20 

nm) at the end of a cantilever to probe the surface. The tip is brought into 

close proximity to the sample surface and either gently contacts the surface or 

taps it while the tip is oscillated close to the cantilever’s resonant frequency. 

The probe is held at a fixed location in the X-Y plane and the surface to be 

examined is rastered in the X-Y plane. The Z-deflection of the cantilever is 

magnified by a laser beam that reflects off the cantilever into a position 

sensitive photo detector or interferometer. A computer records the X-Y and Z 

displacements and a three-dimensional image of the surface can be constructed. 

The AFM is capable of extremely fine resolution (10 pm) and does not require 

any special specimen preparation or testing environment, but the specimen 

must be very nearly flat and has to be mounted flat on a holder. The AFM is 

most frequently used in microelectronic and biological studies. It has been 

applied to the study of fatigue crack markings in metals and fracture mirror 

roughness measurements in glasses. There are limitations to the AFM. The 

scanned area is usually quite small, depending upon the resolution it can be as 

small as 1 μm square to as large as 150 μm square. The vertical surface 

features can only be a few micrometers in size. Acquisition times may be long. 

Color and reflectivity information are lost. Therefore the AFM is almost never 

used to search for fracture origins or other telltale fracture markings, but 

instead is used for very close up, high-resolution examination of surface 

roughness detail or coating evaluations. Hull used it to study the nature of mist 

and hackle lines in brittle epoxies.6 

3.19 Optical Profilometer 

Optical profilometers are designed to make non-contact three-dimensional 

surface roughness measurements from 0.1 nm to several millimeters. The sub 

nanometer resolution for an optical system with white light illumination source 

(halogen bulb) may seem surprising, but interferometry allows the small height 

variations to be resolved. Optical profilometers either translate the specimen 

vertically or, if the specimen is tilted slightly, laterally in front of an 

interferometer. Reflected light from the specimen and a reference surface are 

combined to form fringes that are projected onto a digital camera. Height 

variations are denoted by color and contrast differences in a computer 

generated image as shown in Figure 3.50. One advantage of the optical 

interferometer is that precise height and roughness variations are quantitatively 

recorded and can be interpreted by the program software. 

3-60 
 

 

 

Tools and Equipment 

(a) 

(b) 

Figure 3.50 Matching (a) optical microscope and (b) lateral scanning optical 

profilometer images of a fracture mirror in a ground glass rod. (b) shows an 

area 1.0 mm tall x 1.3 mm wide. The colors correspond to the scale range 

shown on the right (+7.9 μm to -16.3 μm). Wallner lines and mirror mist are 

readily detected and displayed. The origin is a surface crack from a scratch. 

(b is courtesy of R. Gates) 

3-61 
 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

3.20 Confocal Scanning Light Microscope 

This is a variant of an optical microscope. Rather than broadly illuminate a 

fracture surface, a point source of light (often from a laser) is scanned over the 

surface and reflected into the microscope through a pinhole aperture. A lens 

with a large numerical aperture is used and on uneven surfaces only a small 

portion of the surface will be in focus. A photo-multiplier tube records the 

strong return signal from the in-focus regions. Optical slices of in-focus 

surface are registered and recorded by a computer as the focus is finely 

adjusted. The final product is a computer generated three-dimensional 

topographic map of the specimen surface, often in false color corresponding to 

the surface heights. Figure 3.51 shows an example for a glass fracture mirror. 

Resolutions are of the order of 0.25 μm or greater. The primary virtue of the 

confocal microscope is that in-focus images can be collected that would not 

otherwise be possible due to the limited depth of field of the light microscope. 

Quantitative information about height differentials can also be recorded. 

Confocal microscopes usually have longer image acquisition times than 

common optical microscopes with cameras, although some contemporary 

models can acquire and display images in as fast as 1/30th s which is suitable 

for video rate imaging. A full topographical characterization can require 10 s 

or more. 

Some fracture surfaces are difficult to view with a confocal microscope. 

Carbon and graphite surfaces reflect very little light back up to the lens. Single 

crystals can act as specular mirrors and reflect light away from the lens. 

Diffuse reflection may be enhanced by application of a coating in such cases. 

3-62 
 

 

 

 

 

 

Tools and Equipment 

2 

1 

Origin 

(a) (b) 

(c) 

(d) 

Trace 2, Mist 

5 μm 

40 μm 

Peaks from Wallner lines 

50 μm 

Trace 1, Mirror 

1 μm 

Peaks from Wallner lines 

Figure 3.51 Matching stereoptical (a) and confocal microscope (b) images of 

a glass rod fracture mirror. Height profiles for tracks 1 and 2 are shown in (c) 

and (d). The major peaks are from Wallner lines. The minor peaks show the 

inherent roughness in the mirror and mist regions. (b) and traces are courtesy 

of J. López-Cepero Borrego and A. Ramirez de Arellano Lopez, Univ. Seville.) 

3-63 
 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

3.21 Stress Wave and Ultrasonic Fractography 

Ultrasonic or stress wave fractography is an effective fractographic tool in the 

laboratory that was pioneered by Kerkhof.16,17 A transducer attached to the 

specimen creates continuous elastic waves at a specific frequency. The elastic 

waves create ripples on the fracture surface as shown in Figure 3.52. Local 

crack velocity can be simply calculated on the basis of the spacing between the 

lines and the frequency of the pulses. Richter and Kerhof 17 used 5 MHz lead 

zirconate titanate transducers to measure cracks running at terminal velocities 

in glass. Their review article has a wealth of additional details about this 

method and many fascinating examples of its use. Additional examples of their 

work are in section 5.4.4 on tertiary Wallner lines. 

Tsirk improvised18 with a simple hand-held woodworking palm sander that 

vibrated in a circular motion at 11,000 rpm (183 Hz) and was able to create 

makings on the fracture surfaces of obsidian pieces that he was knapping. 

Figures 5.48 d - f in chapter 5 show his images. 

Figure 3.52 Fracture in a glass plate. Fracture originated at the left and 

accelerated to the right as attested to by the increased spacing between the 

timing ripples. (Courtesy H. Richter) 

3.22 High Speed Photography 

Studies of impact fractures have benefited from high speed photographic 

techniques. Field’s review is good starting point.19 The first extensive use of 

high-speed photography to study fracture in glasses was by Schardin and 

colleagues.20 In 1937, Schardin and Struth used a multiple spark camera to 

observe that the terminal velocity of cracks in soda lime glass was about 1500 

m/s. 

3-64 
 

 

 

 

 

 

 

 

 

 

 

 

 

Tools and Equipment 

3.23 X-Ray Topography 

X-ray topography is a specialized technique for single crystals.21,22 An X-ray 

beam illuminates the test specimen and images formed using the diffracted 

beams are recorded. The image from an ideal crystal is completely 

homogeneous, but imperfections that distort or “strain” the crystal lattice 

produce image contrast. Often the defects themselves are not visible. It is the 

disturbances to the lattices that are revealed, and these are often magnified in 

size relative to the defect size. X-ray topographs will reveal dislocations, 

stacking faults, inclusions, and surface damage. Since diffracted X-rays form 

the image, interpretation is not always straightforward. Figure 3.53 shows 

examples of X-ray topographs. They reveal handling damage on and below the 

polished surfaces of a sapphire hemispherical domes and which were not 

visible with the optical or scanning electron microscopes. 

Polished outer surface 

Fracture surface 

Polished outer surface 

TH 

Polished outer surface 

(a) (b) 

Figure 3.53 X-ray topographs of a fractured sapphire hemispherical dome. 

(a) shows the outer polished surfaces of the matching fracture halves. The 

broken halves were carefully realigned for this image. It reveals handling 

damage (HD) cracks not readily visible to the optical or scanning electron 

microscopes. (b) shows a close up of one piece tilted to show both the fracture 

surface and outer polished surface. Twist hackle – cleavage step lines (TH), 

dislocations (D), and fabrication damage (FD) are apparent. (Courtesy D. 

Black). 

3-65 
 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

3.24 Dye Penetration and Staining 

Dye penetration is a staple of nondestructive testing, but it also has value for 

fractographic analysis. It is commonly used to search for hidden cracks in a 

ceramic part prior to usage or after exposure to some loading or environmental 

conditions. Simple staining with common dyes such as the felt tip pens shown 

in section 3.9 can aid post-fracture interpretation in porous or coarse-grained 

materials that are otherwise very difficult to interpret. The dye can be applied 

to the fractured parts. For example, Figure 3.54 shows a portion of a fractured 

porcelain toilet tank. Simple staining revealed tiny incomplete crack branches 

that gave a clear indication of the direction of crack propagation. Figure 4-30 

shows thermal shock cracks highlighted by red dye. 

Figure 3.54 Fractured porcelain part with dye penetration by a common felttip pen. The dye reveals crack branching patterns (small arrows) and thus the 

direction of crack propagation (large arrow). 

Dye penetration is sometimes used with ceramics to try to penetrate grinding 

induced cracks, but these are often very small and tight, making penetration 

problematic. Examinations of the fracture surfaces are done after fracture to 

ascertain whether the dye has penetrated beneath the ground surface. Dye 

“bleeding” can be a problem if the dye migrates into the fracture surface either 

during the fracture event or afterwards. 

3-66 
 

 

 

 

 

 

 

 

 

 

 

Tools and Equipment 

Simple colored staining dyes or felt tip pens may be used with white ceramics. 

Level 4 ultrahigh sensitivity fluorescent dyes are used with dark ceramics. 

Ultrahigh penetrating fluorescent dyes are needed to get into the tiny cracks. 

Special precleaning and soaking steps sometimes are necessary to facilitate dye 

penetration. Sometimes it may even be necessary to heat or vacuum treat a part 

to eliminate adsorbed gasses on the crack faces prior to dye penetration. After 

penetration, ultraviolet or “black lamp” illumination (Figure 3.55) is used to 

excite the dye such that it fluoresces and can be seen by the eye. Parts may be 

examined under a stereoscope, but the images are dim since the amount of light 

emitted is very slight. It is difficult or impossible to photograph features under 

these conditions. Compound optical microscopes, on the other hand, may be 

used with intense xenon light sources and fluorescent filters so that images may 

be captured as shown in Figure 3.56. The photo shows an image made on an 

older generation instant developing film that was exposed for several minutes. 

Digital cameras would have made a better recorded image. Figure 3.57 shows 

a C contact damage crack in a large notched silicon nitride ball prior to and 

after a controlled fracture.23 Fischer et al.24 have used fluorescent dyes on 

flaws in dental ceramics. 

Figure 3.55 Fluorescent dye penetration equipment. Two double torsion 

specimens are shown. The white one is an alumina, the black one, a silicon 

nitride. 

3-67 
 

 

 

 

 

 Fractography of Ceramics and Glasses 

Figure 3.56 Fracture surface of a silicon nitride specimen with a Knoop 

semielliptical surface crack that was penetrated by a fluorescent dye prior to 

fracture. The dashed line is the specimen bottom surface. This photo was 

made with a compound microscope with xenon bulb illumination and a 

fluorescence filter at a nominal 200 X magnification. The dim blurred image is 

typical for the weakly-emitting dye penetrant. This photo was made with old 

style instant-developing film that was exposed for several minutes. 

a b 

c 

Figure 3.57 C contact damage crack in a silicon nitride ball. (a) shows the 

outer surface and the crack highlighted by fluorescent dye penetrant. (b) shows 

one half of the notched ball after it was broken. (c) shows the fracture surface. 

(Courtesy T. Lube and R. Danzer, Mountain University Leoben). 

3-68 
 

 

 

 

 

 

 

 

 

 

Tools and Equipment 

Machining cracks are small and tight, and difficult to dye penetrate. Part of the 

problem with them is that there usually is a thin zone (of the order of only a 

few micrometers) on the surface that is in residual compression. It clamps the 

cracks tight at the surface. Mizuno et al.25 devised a clever procedure to use a 

metallic solution with Pd(NO3)2 impregnation in a vacuum. The solution 

normally would not penetrate the as-ground surfaces, but this was overcome by 

two methods. One was to polish off the surface compression layer. The 

alternative method was to apply a small amount of bending to open up the 

cracks for penetration on the tensile surface. The palladium-penetrated cracks 

were evident on the fracture surfaces when viewed by SEM, presumably in the 

backscattered mode. Kanematsu used this methodology but with an electron 

micro probe to study machining cracks by using cold-isostatic pressing to 

infiltrate the dye. 26,27,28, He further used plasma etching to highlight the 

surface and subsurface damage around grinding damage and indentations.27,29 

Figure 3.58 from reference 28 shows matching SEM images and palladium 

maps for three different specimens ground with a 100 grit, 75% concentration, 

resin-bonded diamond wheel. 

(a) 

(b) 

(c) 

Figure 3.58 Grinding cracks in silicon nitride revealed by palladium nitrate 

solution impregnation. These are paired views of fracture surfaces for three 

different specimens. The first image in a pair is a palladium map. The second 

is a conventional backscattered SEM image at the same magnification. 

(Courtesy of W. Kanematsu). 

3-69 
 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

The Nagaoka University of Technology in Japan has done fine work with novel 

dye penetration methods to investigate three-dimensional structures and flaws 

in powdered compacts and fired ceramic structures. For example, Professors 

Uematsu and Tanaka and students have used potassium iodide and iodine in 

diluted water missed with a saturated boric acid solutions to study the 

uniformity and lack of uniformity in pressed compacts.30,31 

3.25 Polariscope 

As discussed in section 3.7 on lighting, transparent objects may be viewed 

through crossed polarizers. A dedicated polariscope such as shown in Figure 

3.59 may be used to examine large objects such as bottles or plates held 

between the rear polarizer and illumination source and the front polarizer (or 

eyepieces). 

In photoelastic materials the existence of residual stresses may be revealed by 

light and dark patterns or color bands. Glasses and plastics are photoelastic. 

The method is based on the property of birefringence, whereby a light ray 

passing through the material has two refractive indices. Many crystals are 

birefringent, but plastics and glasses become birefringent when stressed. The 

strains in the body cause the horizontally- and vertically-polarized light waves 

(a) (b) 

Figure 3.59 (a) shows an old but still very useful polariscope. (b) shows a twin 

band (arrows) in a complex single crystal sapphire part. The photo was taken 

by a simple digital consumer camera held up to the eyepiece. 

3-70 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

Tools and Equipment 

to travel at different velocities through the stressed material in the directions of 

the maximum principal stresses. Photoelastic analysis is usually applied to 

two-dimensional X-Y problems, such as plates or bottles and vials. The 

stresses that are revealed are those that are in the X-Y plane of the structure and 

not those that may exist through the thickness. If the stress optical coefficient 

for the material is known, then quantitative stress estimates may be made. 

Highly stressed parts will have a sequence of alternating light and dark bands, 

or “fringes,” that can be counted and used with the stress-optic law to estimate 

the stresses. One thing a fractographer should keep in mind is that once a part 

breaks, the internal strains can be relieved to some extent, and the stress 

distribution will be different than that in the unbroken part. 

Polariscopes are also valuable with single crystals. Twins may be revealed as 

shown in Figure 3.59b. 

3.26 Other Equipment 

A common scanner such as shown in Figure 3.60a is handy for converting print 

images to digital JPEG or TIFF files. 

Figure 3.60b shows a simple ultrasonic cleaner that is helpful for cleaning 

small specimens prior to SEM examination. The specimens may be placed into 

small beakers with a solvent such as acetone and then placed into the sonic 

cleaner bath which is partially filled with water. Methyl alcohol is effective in 

removing finger prints. A few minutes vibration is usually sufficient to remove 

common contaminants and cutting debris. 

(a) (b) 

Figure 3.60 Other useful equipment for a fractography laboratory. Digital 

scanner (a) and ultrasonic cleaning bath (b). 

3-71 
 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Simple cutoff or wafering machines (Figure 3.61) are handy for cutting large 

specimens to sizes that can be mounted in optical or scanning electron 

microscopes. Cutting should be done wet to flush debris away and to minimize 

heating of the specimen. 

Figure 3.61 Wafering and cut off wheel machines. 

3.27 Other Resources 

The ASM Handbooks Volume 11 on Failure Analysis and Prevention and 

Volume 12 on Fractography have articles on fractographic equipment.5,32 

Appendix A of this Guide lists additional sources of information. 

3.28 The Future 

The merger of the microscope with digital cameras and computers will 

continue and virtual imaging with computer-generated topographical images 

will become commonplace. Despite all the advantages of the new tools and 

advanced technologies, they will never replace direct optical imaging. The 

stereoptical microscope will always be a primary tool since it is easy to use and 

retains color and reflectivity information and can capture internal reflections. 

Most of the new digital technologies emphasize topographic detail to the 

exclusion of all else. Very often a key feature will be fleetingly observed 

through the microscope eyepieces during the rotation of a specimen or the 

angle of lighting. 

New computer software and digital camera technologies will enhance the 

capabilities of the conventional compound optical microscope so that it 

3-72 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Tools and Equipment 

partially replaces the scanning electron microscope. The latter is most 

commonly used at 200X to 1000X for fractographic work, since this is all the 

magnification that is ordinarily needed to find and record 5 μm to 100 μm size 

flaws. These magnifications are well within the reach of enhanced compound 

optical microscopes with digital cameras and Z-axis image stacking. On the 

other hand, the chemical (energy dispersive x-ray analysis) and structural 

analysis (electron diffraction) capabilities of the SEM will ensure it is a staple 

of the fractographers tool box for many years to come. 

Microanalytical techniques in the SEM will continue to improve. Spatial 

variations in microstructure can have a strong influence on crack propagation. 

We tend to envision a microstructure as uniform with all the grains being 

randomly oriented. Variations in composition, grain orientations, and second 

phase distributions can have a profound influence on local properties and local 

crack propagation. Electron backscatter diffraction is limited at the present to 

very flat surfaces, but imagine how our eyes will be opened if it or some other 

technology can be applied to fracture surfaces. 

Perhaps one of the greatest current limitations of the art of image storage and 

display is that it is all two-dimensional. Fracture surfaces are by their very 

nature three-dimensional, so something is lost when an image of a fracture 

surface is captured and stored in a computer’s memory. This is partially solved 

by new image analysis computer programs used in conjunction with multiple 

SEM images (taken with different tilts) or optical microscopes with Z-axis 

rastering such that a computer generated model or depiction of the surface can 

be displayed. Such virtual images can be color enhanced and can be tilted, 

rotated, or viewed from different angles. Yet, they are all displayed on twodimensional computer monitors. 

One can imagine the future fractographer donning special eyepieces or a visor 

that enable him or her to traverse over or “fly over” a virtual fracture surface. 

These will certainly have considerable novelty value, but one wonders whether 

they will lead to better diagnoses than those made by an experienced observer 

looking through a simple stereoptical microscope. It is very common for the 

fractographer, while wiggling and tilting a fracture surface under a stereoptical 

microscope, to detect a fleeting reflection or stray subtle fracture line, which 

upon further examination unlocks the secret of a particular problem. Will 

future computer software and image analysis programs ever match this skill? 

Perhaps the great breakthrough will occur when technology is invented that can 

capture genuine three-dimensional images that retain full color, reflectivity, and 

internal reflection information and display them with full fidelity in three 

3-73 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

dimensions. The author can personally attest to the hours of desperate struggle 

to acquire the best two-dimensional image to capture a critical detail, only to 

feel despair when showing the image later and realizing that it “just doesn’t 

show what the eyes saw.” 

Chapter 3 References 

1. J. R. Varner, “Using Replicas in Fractography of Glasses and Ceramics – 

Advantages and Pitfalls,” pp. 299- 307 in Fractography of Glasses and Ceramics, 

VI, eds. J.R. Varner and M. Wightman, Ceramic Transactions, Vol. 230, Wiley, 

Hoboken, NJ, 2012. 

2. J. R. Varner, “Replicas as a Technique for Examining Fracture Surfaces of 

Ceramics,” pp. 107-112 in Fractography of Advanced Ceramics, III, eds. J. Dusza 

et al., Transtech Publ. Zurich, 2009. 

3. S. S. Scherrer, J. B. Quinn, G. D. Quinn, H. W. Anselm Wiskott, “Fractographic 

Ceramic Failure Analysis Using the Replica Technique,” Dental Materials, 23 [11] 

(2007) 1397-1404. 

4. V. D. Fréchette, Failure Analysis of Brittle Materials, Advances in Ceramics, Vol. 

28, American Ceramic Society, Westerville, OH, 1990. 

5. Fractography, Vol. 12, ASM Handbook, ASM Int., Materials Park, OH, 1987. 

6. D. Hull, Fractography, Observing, Measuring and Interpreting Fracture Surface 

Topography, Cambridge University Press, Cambridge, 1999. 

7. J. T. Healy, and J. J. Mecholsky, Jr., “Scanning Electron Microscopy Techniques 

and Their Application to Failure Analysis of Brittle Materials,” pp. 157–181 in 

Fractography of Ceramic and Metal Failures, ASTM STP 827, ASTM, West 

Conshohocken, PA,1984. 

8. B. L. Gabriel, “Scanning Electron Microscopy,” pp. 166 – 178 in Fractography, 

Vol. 12, ASM Handbook, ASM Int., Materials Park, OH, 1987. 

9.	 S. Burgess, X. Li, and J. Holland, “High Resolution Energy Dispersive X-ray 

Spectrometry in the SEM and the Detection of Light Elements Including Lithium,” 

Microscopy and Analysis, May 2013, pp. S8-S13. 

10. L. S. Chumbley and L. D. Hanke, “Scanning Electron Microscopy,” pp. 516 – 526 

in Failure Analysis and Prevention, Vol. 11, ASM Handbook, ASM Int., Materials 

Park, OH, 2002. 

11. C. R. Brooks and B. L. McGill, “The Application of Scanning Electron 

Microscopy to Fractography,” Mater. Charact., 33 (1994) 195-243. 

12.	 K.P. Mingard, H. G. Jones, M. G. Gee, B. Roebuck, J. W. Nunn, “In situ 

Observation of Crack Growth in a WC-Co Hardmetal and Characterization of 

Crack Growth Morphologies by EBSD,” Int. J. Refr. Met. Hard Mater., 36 (2013) 

136-142. 

13. H. U. Marschall, R. Danzer, and R. Pippan, “Three-Dimensional Analysis of 

Decorated Ceramic Fracture Surfaces by Automatic Stereophoto-grammetry,” J. 

Am. Ceram. Soc., 83 [1] (2000) 223 – 225. 

14. T. D. Fletcher, J. J. Petrovic, and J. E. Hack, “Quantitative Fractography of SiC 

Whisker-Si3N4 Matrix Composites,” J. Mat. Sci., 26 (1991) 4491 – 4498. 

3-74 
 

 

 

 

 

 

 

 

 

 

 

 

Tools and Equipment 

15. G. D. Quinn, “Fractographic Analysis of Miniature Theta Specimens,” Ceram. 

Eng. Sci. Proc., 29 [3] (2008) 189 – 199. 

16. F. Kerkhof, “Bestimmung der Maximalen Bruchgeschwindigkeit Verschiedener 

Gläser Nach Der Ultraschalmethode,” Glasstechn. Bericht, 35 (1955) 267 - 272. 

17. H. G. Richter and F. Kerkhof, “Stress Wave Fractography,” pp. 75 – 109 in 

Fractography of Glass, eds., R. C. Bradt and R. E. Tressler, Plenum, NY, 1994. 

18.	 A. Tsirk, Fractures in Knapping, Archaeopress, Oxford, UK, 2014. 

19. J. E. Field, “High-Speed Photography,” Contemp. Phys., 24 [5] (1983) 439 – 459. 

20. H. Schardin and W. Struth, “Neuere Ergibnisse der Funken-kinemographie,” 

Zeitschrift Tech. Phys., 18 (1937) 474 - 477. 

21. D. R. Black and G. G. Long, X-Ray Topography, NIST Recommended Practice 

Guide, NIST Special Publication 960-10, April 2004. 

22. D. Black and G. D. Quinn, “X-ray Topography for Fractography of Ceramic Single 

Crystals,” J. Fail. Anal. Prev., 6 [3] (2006) 69 -76. 

23. T. Lube, S. Witschnig, P. Supancic, R. Danzer, and O. Schöppl, “The Notched Ball 

Test – Characterization of Surface Defects and Their Influence on Strength,” pp. 

225 – 234 in Fractography of Glasses and Ceramics, VI, eds. J.R. Varner and M. 

Wightman, Ceramic Transactions, Vol. 230, Wiley, Hoboken, NJ, 2012. 

24. H. Fischer, F. Karaca, and R. Marx, “Detection of Microscopic Cracks in Dental 

Ceramic Materials by Fluorescent Penetrant Method,” J. Biomed. Mater. Res., 61 

[1] (2002) 153-158. 

25.	 S. Mizuno, N. Kobayashi, and K. Miyata, “Analysis of Sintered Silicon Nitride 

Grinding Damage,” SAE Technical Paper 930163, March 1-5, 1993, SAE Int., 

Warrendale, PA, 1993. 

26.	 W. Kanematsu, M. Sando, L. K. Ives, R. Marinenko, and G. D. Quinn, 

“Determination of Machining Crack Geometry by Dye Penetration Technique,” J. 

Am. Ceram. Soc., 84 [4] (2001) 795-800. 

27.	 W. Kanematsu, “Visualization of Subsurface Damage in Silicon Nitride from 

Grinding by a Plasma Etching and Dye Impregnation Method,” J. Am. Ceram. 

Soc., 89 [8] (2006) 2564-2570. 

28.	 W. Kanematsu and L. K. Ives, “Propagation Behavior of Machining Cracks in 

Delayed Fracture,” J. Am. Ceram. Soc., 87 [3] (2004) 500-503. 

29.	 W. Kanematsu, T. Miyajima, and M. Sando, “Visualization of Knoop Indentation 

Damage of Silicon Nitride by Plasma Etching,” J. Am. Ceram. Soc., 84 [10] (2001) 

2427-2429. 

30. Z. Kato, S. Tanaka, N. Uchida, and K. Uematsu, “A Novel Method to Investigate 

the Non-uniformity of the Internal Structure of Granule Compacts,” in Key 

Engineering Materials, Vols. 206-213, Proc. 7th Eur. Cer. Soc., (2002) pp. 341 – 

344, Trans Tech Publ., Zurich, Switzerland. 

31.	 Y. Saito, S. Tanaka, N. Uchida, and K. Uematsu, “Direct Observation of ThreeDimensional Structure in Powder Compact,” in Key Engineering Materials, Vols. 

206-213, Proceedings of the 7th Eur. Cer. Soc., (2002) pp. 657 – 660, Trans Tech 

Publ., Zurich, Switzerland. 

32. Failure Analysis and Prevention, Vol. 11, ASM Handbook, ASM Int., Materials 

Park, OH, 2002. 

3-75 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

General Examination 

4. GENERAL EXAMINATION AND FRACTURE PATTERNS 

Analysis begins with a simple visual examination of the broken pieces. The 

general patterns of crack extension and branching not only point the way back 

to a fracture origin, but also provide information about the cause of fracture, the 

energy of fracture, and the stress state. The majority of brittle fractures start at 

one origin and radiate outward. Severe thermal stresses or impact loadings can 

create multiple origin sites. Sometimes once one primary fracture occurs, large 

fragments remain that are still under load, and they can have secondary 

fractures. 

Fréchette defined the fracture origin as: “that flaw (discontinuity) from which 

cracking begins.”1 A slightly broader definition considers that the fracture 

origin is both a location and a flaw: 

fracture origin: the source from which brittle fracture begins. 

Even if the primary objective is to find the fracture origin, the overall breakage 

pattern should be examined first. Most of this chapter is concerned with 

breakage patterns in general. Examination of the fracture surfaces is covered in 

Chapter 5. An experienced fractographer may simultaneously examine the 

general breakage pattern and the fracture surfaces. Chapter 6 covers the types 

of fracture origins that one might find. 

Breakage may be sudden, unstable, and catastrophic as in a laboratory strength 

test or a dropped glass. Alternatively, it may be in stages as in a fractured 

dental crown restoration, or a crack growing in an automobile windshield. 

Most but not all origins in glass are from surface- or edge-located origins. 

Ceramics fracture from volume, surface, or edge sources. The fracture origin 

is the site that had the worst combination of tensile stress and flaw severity. 

The latter is determined by the flaw size, shape, and thermoelastic or chemical 

interactions with the matrix. A large flaw may not necessarily trigger fracture, 

especially if it is located in a portion of a part that has low stress or it is very 

blunt such as a spherical pore. The severity of a flaw and whether local stresses 

are sufficient to trigger fracture are covered by principles of fracture mechanics 

presented in Chapter 7. This chapter begins with some generic breakage 

patterns and shows some common examples, both for laboratory tested 

specimens and component fractures. The end of the chapter has some general 

observations about fractures caused by mechanical overload, impact or contact, 

thermal, corrosion or oxidation, and some observations on how residual stresses 

may affect fracture patterns. 

4-1 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

4.1 Specimen Reconstruction 

The general examination often begins with specimen reassembly. Ideally all 

fragments are available. Even if the pieces with the origin are lost, valuable 

information can be gleaned from the fragments. With some experience and 

with a familiar breakage pattern, irrelevant shards or ordinary secondary 

fracture fragments may be ignored. At this point the fractographer should take 

care not to introduce contamination to the fragments and especially the fracture 

surfaces. 

Sometimes an experienced fractographer is able to look at fragments of a 

familiar component or test specimen type and be able to immediately find the 

piece with the fracture origin. For example, experienced fractographers are 

able to sift through a bucket of tempered glass fragments and find the origin 

pieces on the basis of their telltale morphology. 

A common concern is whether it is safe to put the pieces together. This step is 

discouraged when conducting analysis of polymer or metal fractures. In 

contrast, ceramic and glasses are hard and durable materials. The pieces 

usually fit perfectly together. It is acceptable to fit pieces together gently, 

taking care to avoid abrasion or chipping. Single crystals and glasses are more 

sensitive to chipping than polycrystalline ceramics and must be handled more 

carefully. 

Simple adhesive tapes may be used to temporarily hold pieces together. Apply 

the tape sparingly and only on an external surfaces and not between fracture 

surfaces. Small strips of masking tape may be used judiciously, but it should 

be borne in mind there is a risk of contamination. Whatever tape is used, it 

should be a type that is easily removable. Keep in mind that a reassembled 

body held together by tape strips may be unwieldy. 

In complex cases with multiple fragments, it may even be necessary to glue 

fragments together. Glues should be used sparingly and preferably in a tackbonding mode away from candidate fracture origin regions. Avoid getting 

glue on the fracture surfaces. The overall assembly will fit better if there is no 

build up or misfits from glue joints between fracture planes. Simple cheap 

nitrocellulose air-drying cementsa are suitable and set fast. They can easily be 

removed by peeling off with a fingernail or razor knife or by dissolving with 

common solvents. In special cases it may be necessary to use a strong cement 

to glue the pieces together on their fracture surfaces. Cyanoacrylate ester 

a Duco cement or Scotch “All purpose adhesive, Super strength”. 

4-2 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

General Examination 

(“super”) glues may be used since they can set with very thin bonds such that 

cumulative misfits are held to a minimum. Such glues can be removed by 

soaking in acetone or nail polish remover. In any case, glues should be used 

sparingly and every precaution taken to prevent them from coming into contact 

with fracture origin areas. 

During reconstruction it may be helpful to mark the pieces with a number or 

letter scheme in order to expedite the examination and to minimize specimen– 

to-specimen refitting exercises. Judicious marking with a pencil or fine tip felt 

tip marker pens are effective, but care should be taken to avoid contaminating 

the fracture surfaces or the fracture origin. Felt tip pen markings can usually 

be removed with alcohol or acetone. 

During the initial examination it may be advantageous to try transillumination 

on white and translucent ceramics. Transillumination, whereby the part is 

illuminated from behind or from the side, may reveal internal cracks, flaws, and 

microstructural irregularities not detected by examination of the outer surface. 

Figure 4.1 shows three examples. 

4.2 Crack Branching Patterns 

Crack branching and radiating patterns may be interpreted before the fracture 

surfaces are examined. Much of the fractographer’s craft is in the observation 

and interpretation of these patterns. Crack branch patterns can lead back to an 

origin and also tell much about the stress state and stress magnitude. A simple 

example is a glass biaxial strength test specimen shown in Figure 4.2. The 

radiating pattern is intuitive and leads the observer back the origin which is in 

the disk middle. Sometimes multiple fractures can be simultaneously triggered. 

Examples are severe thermal shock fractures, or mid-body fractures on each 

side of an impacted bottle. Secondary origins are often present in components 

that initially break from one origin, but then the fragments themselves fracture. 

The first general rule is the: 

Law of Normal Crack Propagation: a crack propagates normal to the 

direction of the maximum local principal tension stress. 

The word “normal” means “perpendicular.” As the crack propagates, minor 

perturbations in the direction of local principal tension can modify the plane of 

cracking. These minor perturbations can create telltale markings such as 

Wallner lines or hackle on the fracture surface as described in Chapter 5. 

Internal stresses or inhomogeneities can also cause crack local deviations. 

4-3 
 

 

 

 

 Fractography of Ceramics and Glasses 

(a) (b) 

(c) 

(d) (e) (f) 

Figure 4.1 Transillumination. (a) and (b) are an extracted human tooth from 

an approximately 80 year old patient. Transillumination from the side shows 

internal enamel cracks. (courtesy S. Scherrer) (c) shows transmitted (through) 

illumination of a chemical vapor deposited SiC micro tensile specimen. Very 

coarse grains acted as light pipes through the body. Relatively pure SiCs are 

green. (d) – (f) show an aluminum-rich inclusion origin, in a highly-filled 

resin-matrix dental composite. (d) is a trans-illumination view with light from 

below. Although the fracture surface features are blurred by the material’s 

translucency, the origin flaw stands out at the chamfered corner. The SEM 

images (e) and (f) show the inclusion. 

4-4 
 

 

 

 

 

 

 

General Examination 

(a) 

(b) (c) 

Figure 4.2 Fracture patterns in a moderately-strong glass disk tested in ringon-ring biaxial strength test. (a) shows the test arrangement, (b) shows a 

schematic of a fracture pattern. The arrow marks the origin. (c) shows an 

example in a BK-7 borosilicate crown glass disk. 

Once fracture has commenced, the propagating crack can fork or branch into 

multiple propagating cracks. These may be at non-normal angles to the farfield stress direction. Nevertheless, the initial fracture starts out as a single 

radiating crack that separates the part into two pieces. Often the search for the 

fracture origin entails studying the branching patterns and tracing them back to 

the region of first fracture where the two pieces with the origin join across a flat 

surface as shown in the center of Figure 4.2b. 

Branching may not occur at all if the failure stresses are low, or if they drop off 

in the direction of crack propagation, or if the opening up of the crack makes a 

specimen more compliant such that the stresses are relieved. 

4-5 
 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

θ θ

origin 

(a) (b) 

Figure 4.3 Crack branching indicates the direction of crack propagation. 

4.3 Crack Branching Angles 

Cracks which have reached their terminal velocity (normally about 50 % to 60 

% of the shear wave velocity) may split into two cracks with an acute angle 

between them as shown in Figure 4.3a. Fréchette1 terms this velocity 

branching, velocity forking, or velocity bifurcation. Branching patterns are a 

valuable aid in diagnosing the direction of crack propagation. For example, a 

pair of branches in opposite directions typically brackets the origin as shown in 

Figure 4.3b. 

Another very useful pattern is that the angle of forking varies with the stress 

state as shown in Figure 4.4. The trend shown by the solid line was reported by 

Preston in 1935 for the fracture of glass articles.2 Uniaxially-stressed parts, 

such as direct tension strength or bend bar specimens, fork at about 45o

. 

Equibiaxially-stressed parts, such as a uniformly-loaded thick window that has 

an origin near the center, fork at as much as 180o

. Pressurized bottles, which 

have circumferential hoop stresses that are double the axial stresses, fork at 

about 90o

. Torsion fractures, wherein the ratio of the principal stresses is 

negative (one stress is tensile, the other is compressive as shown in Fig. 1.2), 

have very small forking angles. Preston said that the angles were very 

consistent and varied by only a few degrees, but other investigators have 

observed much greater variability. 

It now appears that systematic variations can occur due to component shape, 

the exact loading configuration, the stress level, and possibly the material 

properties. For example, the 180o

 branching angle for equibiaxial tension 

(“drum skin tension”) reported by Preston occurs for pressure loaded, thick 

windows that fracture at high loads. Shinkai3 shows such a fracture (his Figure 

22a) in a 1.7 m x 0.9 m x 6 mm building window. On the other hand, most 

laboratory scale disk specimens tested with ring-on-ring equibiaxial 

4-6 
 

 

 

 

 

 

 

 

 

General Examination 

Figure 4.4 Branching angle versus stress state. The solid line is the trend for 

glass shown by Preston.2

 The shaded bands show Rice’s data4 ranges for 

several ceramics and Bullock and Kaae’s range for a glassy carbon.5 The 

author’s data for silicon nitride ring-on-ring disks are also shown.6 

loading (Figure 4.2) have smaller angles from 120o to 155o

 and only approach 

180o

 if the specimen is very strong and has much stored elastic energy at 

fracture. Laboratory scale specimen data by Rice4 on a range of ceramic 

materials, Bullock and Kaae5

 on glassy carbon, and the author on silicon 

nitride,

6

 tend to confirm Preston’s general trend, but with smaller angles. Most 

of Bullock and Kaae’s uniaxial angles were 30o

. The exact biaxial stress ratio 

for Rice’s data for his thermally-shocked and ring-on-ring loaded disks is 

unclear and therefore is shown as a band in Figure 4.4. Some of the fractures 

were outside the inner loading ring where the hoop and radial stresses were not 

the same and thus σ2/σ1 < 1. Some of his data was for ball-on-ring loaded 

specimens which were also not equibiaxial unless fracture began at the exact 

center. It is also unclear how he defined his angles. They may have been the 

angle between any two crack segments or the included angle between the two 

crack segments that were the furthest apart. Based on Figure 12 of his paper, it 

may have been the former. Nevertheless, Rice made a number of useful 

4-7 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

observations about the details of branching angle. A number of measurements 

were made and the standard deviations indicated there was some variability in 

branch angles even in uniaxially-loaded specimens. In some instances, the 

branching angle depended somewhat on whether the crack split into two 

branches or into three, including an extension of the original crack. He also 

noted that even along a given branch the angle did not necessarily stay 

constant. If a branch started out at an unusually low angle, it tended to curve to 

increase the angle. The converse was also observed. 

Rice suggested that branching patterns tend to split the broken component into 

regions of approximately similar area.4 This implies that the cracks partition 

the component into pieces with comparable areas and strain energies, a 

reasonable supposition. Bullock and Kaae also noted that the branching angles 

diminished with progressive branching.5 

Unfortunately there is no consensus on how to measure branching angles. Both 

Preston and Fréchette recommended that the angle be measured close to the 

point of forking, but it may be more appropriate to measure the angle once it 

has stabilized. Close inspection shows that a branch starts at a shallow angle 

and then increases to its stable configuration as shown in Figures 4.5 and 4.6. 

Preston’s original sketch (Ref. 2, his Figure 2) for the equibiaxial disk pattern 

in fact shows such a gradual curvature at the origin and the angle reached 180o 

only after a small extension. J. Quinn used magnification to study branch 

angles in biaxially stressed glass disks and also observed angles that started 

small, but then stabilized at a larger angle.7

 She observed that there was a 

stress magnitude dependence of the final angle: highly stressed disks had 

angles that approached the 180 degree limit, whereas lower stressed disks had 

smaller final angles. Rice evidently measured the angles after they had been 

fully developed. Branching angles often vary within a broken part, usually due 

to spatial variations of the stresses. So for example, in a ring-on-ring loaded 

disk the stresses are equal in the middle circle, but the radial and hoop stresses 

drop off at different rates outside the inner circle out to the rim. Furthermore, 

as cracks progressively branch they may begin to interact with other branching 

cracks. Hence, it is common for an initial branch angle to be large in the 

middle near an origin in a biaxially stressed disk, but the subsequent branch 

angles diminish out towards the rim. 

So although there are discrepancies in the reported data, and the optimum 

manner of measuring the angles is unresolved, branching angles provide at least 

a qualitative indication of the stress state. There may be material property or 

component shape effects. This may be a good topic for a graduate thesis. 

4-8 
 

 

 

 

General Examination 

θ

Figure 4.5 Close up of the branching point. 

(a) (b) 

Figure 4.6 Fracture pattern in a 44.5 mm silicon nitride ring-on-ring (10 mm 

x 40 mm) equibiaxial strength test disk. The stress at fracture in the inner ring 

was 447 MPa. The branching is curved near the origin. Successive secondary 

branching occurs over short distances. (b) is a close up of the origin area 

showing that the initial branch angles are 140o

 and 145o. Additional branching 

is minimized as the crack propagated into lower stress areas away from the 

disk center. 

4.4 Crack Branching Distances 

The distance a propagating crack travels before branching is directly related to 

the stresses and stored energy in the component. The greater the stored energy, 

the shorter the distance to branching. If the crack travels through a region of 

constant stress, the empirical relationship is: 

σ Rb = Ab (4.1) 

4-9 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

where σ is the stress, Rb is the branching distance (radius), and Ab is a material 

constant called the fracture branching constant. A table of branching constants 

is in Appendix C. Additional details about equation 4.1 and Ab, including the 

latter’s similarity to fracture toughness, are discussed in chapter 7 on 

Quantitative Analysis. There is some evidence that Ab may not be a material 

constant8,9

 and may depend upon geometry and stress state. 

4.5 Fragmentation Patterns 

A general qualitative assessment about the stresses in a part can be made from 

the number of fragments. Quantitative analysis is discussed in Chapter 7. 

Low stress, low energy fractures create minimal branching and hence few 

fragments. High energy fractures cause extensive fragmentation. The extent 

of fragmentation depends upon the stress state throughout the body and the 

total energy available for fracture. A small pebble or BB gun shot creates only 

localized damage in a window, but a hurled brick will cause window bending in 

addition to localized impact damage. Sometimes a general observation about 

fragmentation patterns is sufficient for a diagnosis. For example, if a fracture 

occurs at an unexpectedly low stress the fractographer might suspect and search 

for a gross material, handling, or impact flaw that has weakened a part. Many 

thermal breaks are low energy fractures. High local or transient stresses may 

precipitate initial crack propagation, but the stress levels may be low in most 

other portions of a part. Low energy fractures often have few or no telltale 

markings on the fracture surfaces, as discussed in the next chapter. 

Processing cracks that occur on firing also generate minimal fragmentation. 

That is because the cracks are so large that they significantly weaken the part. 

Differential shrinkage or out-gassing often create local stresses that can initiate 

a crack in the green state or during final firing, but the stresses may relax as the 

crack opens. Such cracks are often rough, twisted, and winding. Their internal 

fracture surfaces may have telltale thermal rounding or discoloration from 

being exposed during processing. Section 6.8.1 in the chapter on Origins has 

examples of large processing and firing cracks. 

Moderate-stress failures create more fragments that can be reassembled so that 

branching patterns can be interpreted. There may be more fracture surface 

markings. 

On the other end of the spectrum, high-energy fracture may be so virulent and 

create so many fragments that reconstruction may be impractical. Critical 

pieces may be lost. That is not to say fractographic analysis is hopeless. 

Sometimes telltale pieces that contain the origin can be culled out of a mass of 

4-10 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

General Examination 

fragments and the critical information gleaned while the majority of pieces are 

ignored. Analysis of the fracture surfaces can reveal whether fracture was 

caused by tension or bending, as will be discussed in the next chapter. 

Uneven stress states may be manifested in many ways. For example, Fréchette1 

showed an example of a tempered, high-powered lamp that cracked 

spontaneously as a result of heating in-service. In the fracture origin area, the 

fragments were large, attesting to a low to moderate stress level. Elsewhere in 

the lamp, extensive fragmentation and dicing typical of tempered glass was 

observed. Fréchette concluded that the intense heating in service caused the 

middle of the lamp to lose its temper and weaken such that it was more 

susceptible to fracture from thermal stresses. 

Although similar techniques are applied to analyze laboratory-generated 

fractures and component service fractures, it is convenient to treat them 

separately in the following paragraphs. Even if one is primarily interested in 

component analysis, the experience gained with laboratory fractures can help. 

4.6 Laboratory Test Fracture Patterns 

In the laboratory the cause of fracture is known. Fractographic analysis may be 

done to identify strength-limiting flaws or to study crack microstructure 

interactions. It also may be done to verify that the test was done correctly and 

that the fixtures and specimen were aligned properly. It also may be done to 

gain valuable experience in examining fracture patterns in a particular material 

that can help with component failure analysis. 

A simple precaution should be taken before the testing commences: Shielding 

or buffering material should be placed around the specimens to catch all the 

fragments and to minimize secondary fractures and impact damage. Cotton or 

tissue buffering and paper shields can be placed around the test specimens. In 

some instances the specimens may be partially taped before the test so that the 

pieces remain attached to the tape after fracture. This must be done 

judiciously and in no instance should tape interfere with the load application. 

Tape should never be on the tensile stressed surfaces. Biaxial ring-on-ring 

loaded disks may be taped on the compression side of the specimen with a tape 

circle inside the loading ring and a tape annulus outside the inner loading ring. 

The portion of the surface in direct contact with the loading ring should not be 

taped. 

Specimens should be premarked with orientation or location marks. This 

preliminary step can dramatically aid post fracture analysis. For example, the 

4-11 
 

 

 

 

 

 Fractography of Ceramics and Glasses 

orientation of round specimens tested in tension should be marked to ascertain 

whether fracture occurs preferentially from one side or another. Similarly, 

flexural strength specimens should always have the loading points marked on 

the side once the specimen has been inserted and preloaded in the bend fixtures. 

The marks will help ascertain whether breakages occur at the load pins. 

Sometimes a grid may be marked on a specimen to help with reconstruction. 

This may help with biaxial disks loaded in flexure, but again, any such 

markings should be on the compression side of the specimen. 

Figure 4.7 Tension specimen and fiber fractures. 

Low- Moderate 

Strength 

Moderate- High

Strength 

End of 

Gage

Section 

Break Fiber /

Filament 

Break 

High

Strength 

Fracture 

4-12 
 

 

 

 

 

 

 

 

 

 

 

General Examination 

4.6.1 Tension strength 

This category includes machined specimens with grip holes or shoulders, glass 

optical fibers pulled in tension, or long tubes or rods that are pulled apart. 

Direct tension is conceptually one of the simplest loading configurations, but in 

practice it can be tricky to achieve. Slight loading misalignments or test 

specimen irregularities can create superimposed bending stresses. Breakage 

patterns in properly aligned specimens are usually quite simple as shown in 

Figure 4.7. Fractures initiate and run perpendicularly to the loading axis until 

and if they branch. Moderate-to-high strength specimens may branch with the 

telltale 30o to 45o angle as discussed in section 4.3. Fractures at a 

nonperpendicular angle to the loading direction (or with compression curls as 

discussed in the next section) should be cause for concern and suggest 

misalignments in the load train. It is especially important to ascertain whether 

fracture origins are located on the specimen surface, in the interior, or at an 

edge as will be discussed in the next chapter. Breakages in a group of test 

specimens should be randomly distributed within the gage section. An 

occasional break near the end of the gage section is probably not too serious, 

but if all specimens break from the end, then the tester should review the test 

specimen design. The stress concentration at the end of the gage section, the 

fixture and grip alignments, and the machining of the blend junction at the end 

of the gage section should be checked. 

High strength fibers have the added complication that secondary fractures are 

common. After initial fracture, the intense stored elastic energy in the fiber 

releases, sending stress waves through each broken portion. Reverberations 

and fiber whipping cause additional breakages. Fragmentation may be so 

thorough that it may be impractical to find a primary fracture. Testing in oil or 

coating the fibers with grease may minimize the whipping and secondary 

breakage, at the risk of contaminating the primary fracture and altering the 

environmental sensitivity of fiber strengths. 

4-13 
 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

4.6.2 Flexural strength 

Flexural strength testing with rod or bar specimens is often done in lieu of 

direct tension testing. Flexure is much easier to do. Specimens and testing 

equipment cost considerably less. A drawback is that much less volume and 

area are exposed to the full tensile stresses. Hence, flexural strengths are 

typically greater than direct tension test strengths. Reference 10 is a review of 

flexural strength testing of ceramics. For our purpose here, it is convenient to 

briefly review the stress state as shown in Figure 4.8. The bending loading 

creates a stress distribution such that the maximum tensile stress is on the 

bottom as shown in (b) and (c). The stresses diminish into the interior until 

they are zero in the middle, at the “neutral axis.” The tensile stresses are 

balanced by compression stresses on the opposite side. The maximum stress 

exists only directly opposite the middle loading point in three-point loading and 

diminishes linearly with distance from the middle loading point to the outer 

support points. In four-point flexure, the maximum stress occurs over a larger 

region between the two inner loading points. 

(a) (b) 

(c) (c) 

tension 

stresses 

Figure 4.8 Three- and four-point flexure test configurations are shown in (a). 

The stress distribution on a cross section is illustrated in (b) and (c). 

4-14
 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

General Examination 

A very telltale feature of bend fractures is the compression curl, also known as 

the cantilever curl, shown in Figures 4.9 and 4.10. Once the crack has 

propagated from the tensile half of the specimen into the compression side, it 

slows down and changes direction. In stronger specimens the crack can branch, 

creating a double curl. The existence of a compression curl is an important 

sign that the specimen either was loaded primarily in bending or had a strong 

(a) 

force 

Four-point silicon nitride 

rectangular bend bar Moderate strength glass 

rod that branched early 

(b) (c) 

Round glass rod 

Bovine dentin 

rectangular bend bar 

(d) (e) 

Figure 4.9 Compression or cantilever curls are a telltale feature of flexural 

fractures. (b) – (e) show side views of four-point flexure specimens. In each 

case shown, the tension surface is on the bottom and the compression surface is 

on top. 

4-15 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

bending component. Compression curls are not normally present in direct 

tension specimens, unless they were badly misaligned such that some bending 

was superimposed on the gage section. The origin of fracture on a fracture 

surface is directly opposite the compression curl. The presence or absence of 

compression curls may be important for interpreting not only test specimens, 

but component fractures as well. Thermal fractures usually do not have 

compression curls. Tsirk argued in his book11 that in simply supported beams 

as shown here in Figure 4.9a, the curl tends to run back towards the support. 

He argued this is due to shear stresses and that the maximum principle stresses 

steers the crack in that direction. To verify this, I tested fourteen glass rods 

supported in a vice with cushioned grips, by tapping with a small mallet 

hammer on the rod ends as shown in Figure 4.9. The rods had a small score 

mark on the tensile surface near the support to weaken them. If the score mark 

was within the contact area of the grips, the curl came out toward the applied 

force. If the score was right at the end of the cushioned support, the curl went 

back into the grips. If the score was just outside the gripped area, the curl went 

back to the grips, but sometimes came outward towards the force as shown in 

Figure 4.9. The angle of the force was a critical parameter, and if the force 

were angled slightly away from the vice, the curl came outward toward the 

force. I concluded that the direction of the curl varied with the location of the 

origin and the precise angle of the force. 

Kolsky,12 using very high speed photography in glass, showed that the 

compression curl in beams in bending is formed late in a fracture sequence, 

often due to elastic wave reverberations interacting with the slowly moving 

crack in its final stages of breakthrough. This may seem counter-intuitive since 

the fracture surface shows that fracture very quickly reaches terminal velocity 

(if a fracture mirror and hackle form as discussed in the next chapter), but 

cracks may in fact slow down very rapidly once they get to the compressivelyloaded half of the beam. Furthermore, Kolsky showed that the local stress state 

around the crack is dramatically altered during the fracture event. It is very 

much a dynamic situation. Most of this is of secondary importance, however. 

The key point here is quite simple: a compression curl indicates a part was 

loaded in bending. 

Figure 4.10 shows a variety of fracture patterns in bend bars. Low strength 

fractures are often the easiest to interpret since the specimen breaks into only 

two or three pieces. Often a small “T” shaped piece will be opposite the origin, 

but this fragment can be ignored since it never has the origin on it. The origin 

is below the “T” fragment on the two matching pieces nearer to the tensile 

stressed surface. Sometimes it is handy to keep the “T” fragments to confirm 

the other fragments do fit together. 

4-16 
 

 

 

General Examination 

Cushioning material 

Figure 4.10 Fracture patterns in four-point flexure specimens. Cushioning 

material under the specimens minimizes damage to the fracture surface edge. 

Higher strength bars have two or more fractures and it must be deduced which 

was primary. Secondary fractures are caused by reverberations and stress 

reflections after the first fracture has occurred.12,13 Secondary breaks 

frequently occur at the loading points. If there are two fractures and one is 

4-17 
 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

located at an inner loading point and the other is in the inner gage section, the 

latter is probably the primary fracture with the first origin. If the two breaks are 

close to the two inner loading points, then the exact locations on the tensile 

surface should be examined closely. Often one fracture will be right under a 

loading point, whereas the other may be close, but not right at the loading 

point. The latter is usually the primary fracture. If both are exactly under the 

two inner points, it may be difficult to deduce which was first. There is also the 

possibility the specimen might have been loaded unevenly causing stress 

concentrations at both sites. If this occurs for more than a few specimens in a 

set, then the fixture should be checked as well as the specimen’s parallelism. 

Properly aligned specimens and fixtures almost always cause a perpendicular 

primary fracture in the middle gage section. Secondary fractures are often but 

not always at a slight non-perpendicular angle to the specimen axis. 

Misaligned specimens or fixtures can cause twisted crack planes as shown in 

Figure 4.11 and described in Ref. 14. 

An unusually large flaw outside the gage section can trigger a fracture with 

crack plane tilted towards an inner loading pin. Specimens with such low 

strength fractures usually do not have secondary breaks. Fractures outside the 

inner gage section in properly aligned specimens and fixtures are usually 

completely valid and a warning that unusually large flaws are present. They 

should not be cavalierly dismissed as invalid tests. 

σ 

(a) (b) 

Figure 4.11 A twisted fracture surface in a misaligned alumina bend bar in a 

fixture that did not articulate properly. (a) shows the fracture surface, and (b) 

the tensile surface of both halves. The principal stress direction (solid black 

arrow) was 14 degrees off the expected direction (dotted arrows). Twisted or 

angled fracture planes should prompt the tester to check the fixture articulation 

and specimen alignment procedures. 

4-18 
t

 

 

 

 

 

 

 

General Examination 

4.6.3 Biaxial flexural strength 

Plates or disks may be tested in biaxial tension by a variety of methods 

including: pressure on-ring, ring-on-ring, piston-on-ring, ball-on-ring, and ballon-three balls. 

Pressure-on-ring loading (whereby a specimen is supported on a ring and is 

loaded by a uniform pressure on the opposite surface) has the virtue that a large 

area is stressed biaxially. There are no loading point or stress concentration 

problems with the inner loaded area. Ring-on-ring loading, which is more 

commonly done since it is simpler to do, creates an Equibiaxially- stressed 

region in the inner circle. Pressure-on-ring and ring-on-ring are best used with 

specimens that have flat and parallel surfaces for even contact with the loading 

and support fixture rings. Figure 4.12 shows fracture patterns in ring-on-ring 

specimens. 

low 

strength 

medium 

strength 

high 

strength 

Figure 4.12 Fragmentation patterns in ring-on-ring loaded specimens. The 

small arrows mark likely origin areas. In low strength parts, after the first 

break has occurred, the larger remnant may still bear load from the inner 

loading ring, and it in turn will break in bending. In moderate to high strength 

disks, secondary circumferential cracking occurs near the inner loading ring. 

d 

4-19 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

The single ball loading schemes have the drawback that only a tiny spot 

directly under the loading point experiences the maximum stress. The 

concentrated forces of the ball distort the stress field in the vicinity. Fracture 

almost always starts directly below the loading ball. 

The ball-on-three balls or piston-on-three balls schemes are better suited for 

uneven specimens, but the stress state is not very symmetric. Cracks often 

align themselves to run between the supporting balls. Excessively thick test 

specimens should not be used, since breaking forces are so high that contact 

cone cracks can be initiated by the middle loading ball. These disturb the 

stress state and, in the worst case, the cone cracks will propagate through the 

specimen thickness and reach the tensile side and cause unintended fracture. 

This is an invalid test. 

A common misperception about ring-on-ring testing is that the edge condition 

is unimportant, since the maximum stresses are within the inner loading circle. 

This notion is only partly correct. Figure 4.13 shows one of several fractured 

BK-7 borosilicate crown glass specimens that were part of a large set of 

specimens intended to generate a comprehensive design and reliability 

database. The overall fracture pattern showed the origin was on the rim. 

Fracture surface examination confirmed that outer rim grinding cracks caused 

fracture. Stress analyses for ring-on-ring specimens (e.g., Fessler and Fricker15 

or Salem and Powers16) show that the maximum hoop and radial stresses are in 

the inner circle. The radial stresses gradually diminish to zero out at the rim. 

Hoop stresses also diminish, but are finite at the rim and depend upon the 

specimen dimensions, the fixture sizes, and Poisson’s ratio. In the case of the 

disk shown in the Figure, the hoop stresses were as large as 48 % of the 

maximum stresses in the middle. It was not surprising that with a highlypolished tensile surface, the disk fractured from grinding flaws on the rim. In 

this testing episode, the problem was detected early after only six disks had 

been tested. The remaining disks were reground and HF acid etched to 

diminish the severity of the grinding cracks. Subsequent fracture testing was 

successful and almost always initiated fracture from the central disk region. If 

this precaution had not been taken and the remaining disks tested to fracture, 

the database would have been nearly worthless. The point here is that 

fractography is a valuable aid to mechanical property testing. 

4-20 
 

 

 

 

 

General Examination 

(a) 

(b) 

Tensile surface 

Outer 

rim 

σ

Figure 4.13 A fractured glass ring-on-ring tested disk. (a) shows the overall 

breakage pattern. Fracture started on the rim (large arrow) in response to the 

uniaxial hoop stress σ (small arrows), propagated towards the middle, and 

branched in response due to the greater biaxial stresses in the middle. (b) 

shows the fracture surface. The origin is a grinding crack on the outer rim 

(large arrow). 

4-21 
 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

4.6.4 Laboratory test specimen analysis: additional tips 

It is often helpful to look at contact marks or scuff marks where the specimen 

was loaded, whether it was in tension, uniaxial or biaxial bending. Such 

surface marks are one type of witness marks: 

Witness marks: marks on a specimen surface that attest to contact with a 

foreign body. 

In this manner it may be possible to ascertain whether loading was in fact 

uniform or not. Contact marks in a bend specimen may also help with the 

interpretation of whether a break occurred directly opposite a loading pin. 

Test specimens are often broken in batches of ten, thirty, or more. Which 

specimens should be examined first? The pragmatic answer is that the weakest 

ones should be examined first, since they are the easiest to interpret. The 

experience gained from these easy specimens facilitates examination of the 

more complex fractured specimens, since the fractographer will be better 

prepared to recognize primary fractures and distinguish them from secondary 

breaks. 

4.7 Component Fracture Patterns – General 

Component fractures have the complication that the state of stress or the cause 

of failure may be unknown. The general fragmentation pattern can provide 

powerful clues. The fractographer first should obtain whatever ancillary 

information is possible about the component. What exactly was the material? 

What were the circumstances of the fracture? Was it an isolated case or part of 

a series? Had some change been made to the material or the loading 

conditions? Was there a noise (e.g., shattering sounds or water hammer)? 

Were there other parts or evidence found with the fractured material (impact 

debris)? Was there a witness? Is the witness honest? Does the witness have a 

good memory? Background information such as this may be crucial for the 

broader failure analysis. Nevertheless, the fractographer must keep an open 

mind, and let the pieces tell their story. 

A wise precaution is to retrieve as many pieces as possible. In some instances 

where there is an obvious pattern, the fractographer can cull out important 

fragments. On the other hand, if other people have the fragments, you cannot 

expect them to know which are the important pieces. They should retrieve all 

pieces. If remnants are in a mount or assembly, then they should not be 

4-22 
 

 

 

 

 

 

 

 

 

 

 

General Examination 

retrieved or extracted until after the fractographer has had an opportunity to 

examine them in situ. Of course, this may not always be practical, but the 

fractographer should make an effort to see the fracture scene and fragments in 

as close to in situ state as possible. At the minimum, photographs of the overall 

fracture pattern and close-ups as necessary should be taken. Figures 4.14 

illustrate some of these recommendations for a practical failure problem. This 

window fracture also illustrates how witness marks or traces of the impactor 

material aided the interpretation. 

a b

c d 

Figure 4.14 Dr. J. Quinn with a window fracture problem at NIST. Evidence 

collected at the scene included foreign material attached to the impact site as 

shown in (c). The crack patterns indicated impact from the exterior. That, plus 

the height of the impact above the ground, the tuft of fur, the time of the year 

(November, the rutting season) and similar occurrences elsewhere suggested 

the cause of failure shown in (d). The deer saw its reflection in the window and 

attacked it! 

4-23 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

The impulse is to jump directly to an examination of the fracture surfaces to 

find an origin. While there may be no harm in this (and it is human nature for 

most fractographers to take this step), there is one extra step that is prudent to 

take immediately, especially if a moderate or great amount of time will be spent 

in specimen reconstruction or the part will be involved in serious failure 

analysis or litigation:

 The entire component or part should be photographed or sketched. 

There are several reasons why it is wise to capture an image of a whole 

structure or sketch it at an early stage. There are obvious instances, for 

technical or legal purposes, why a permanent record is required. It is best to 

make the record as soon as possible, lest pieces be lost, misplaced, or mangled 

in handling. An early record is essential if pieces need be cut up for 

microscopic examination. If the pattern of fracture is complex and an origin 

location not obvious, the fractographer may have to examine many pieces in 

order to track the cracking pattern back to the origin. A sketch or photo that 

can be marked up aids this process. Figure 4.15 shows an example of a broken 

component and an accompanying sketch with work notes. 

An overall sketch or photo will help the analyst show other fractographers, 

engineers, managers, or clients the overall context of the fracture origin or 

fracture pattern. Is the fracture from the same location as other parts, is the 

pattern of breakage the same, or has a new failure mode been triggered? Much 

time and accuracy may be lost in orally explaining or drawing sketches from 

memory the fracture patterns days or weeks later. A handy overall photo of the 

part can save time and reassure others that the correct fracture origin photos 

indeed have been taken. A ruler, magnification marker, or a common object 

such as a coin should be placed next to the component when a photo is taken. 

Many close-up images of fracture features in the technical literature are 

unconvincing or may in fact have nothing to do with a bona fide fracture origin. 

If many examples of a particular part are to be studied, then a sketch or photo 

of an unbroken part may be useful as a blank onto which fracture patterns may 

be sketched or drawn. 

There is another very important and overlooked reason why a fractographer 

ought to make sketches. In the very process of drawing a line, the 

fractographer must think: “where exactly is that line, exactly what does it mean, 

and which way does it go?” Why did it start here and not there? The 

fractographer must contemplate and interpret what has happened as the sketch 

4-24 
 

 

 

General Examination 

(b) 

 (a) 

Figure 4.15 A fractured single crystal sapphire hemispherical dome that 

broke in a high temperature, high Mach wind tunnel. The crack progression 

along individual pieces was used to back track to the origin (arrow). The 

sketch (b) helped integrate and keep track of the information from examination 

of the individual pieces. The origin is marked by the arrow in (a) and is 

labeled “origin” in (b). Figure 8.19 shows the actual origin flaw which was a 

tiny particle impact site. Thermal stresses then caused the dome fracture. 

4-25
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

is being made. This is an important step in an interpretation and diagnosis of 

fracture. My colleague Professor James Varner at Alfred University put the 

matter rather succinctly: “One can take a photo without thinking, but one 

cannot make a drawing without thinking.” One study noted that students who 

took a museum tour had more difficulty remembering the objects if they only 

photographed them, rather than if they looked at the objects without taking a 

photo. Indeed, the popular author David Macauley, who has written a series of 

well-illustrated books such as Cathedral,17 or City,18 has said that when he 

creates his drawings he must contemplate how the structures were designed, 

what materials were used, and how they were assembled. Leonardo da Vinci’s 

drawings underscore the depth of his knowledge about objects both natural and 

manmade and how they worked. 

Some common fracture patterns are shown in the following examples. The 

reader is referred to the bibliography for more information. Fractographers 

who are dealing with specific component failures are strongly advised to search 

the literature for prior examples and documentation of similar fracture patterns. 

Appendix B of this guide is a compilation of fractographic case studies. 

4.8 T Intersections and Crossing Cracks 

Intersecting cracks may pose interpretation problems in component fracture 

analysis. A fractographer may initially be confused by the multitude of 

intersecting cracks. Which came first? Which way was the crack going? With 

a little experience and applying the process of deduction, it usually is not too 

difficult to answer these questions. Figure 4.16 offers some guidance. 

Crack branching creates patterns shown in (a). The branch angle depends upon 

the stress state as shown previously in section 4.3. Even with biaxial loading 

and branch angles near 180o

, a close examination of the branching point will 

show which way the crack was moving as it split. Hence the determination of 

the crack propagation direction is easy. Intersecting cracks as shown in (b) are 

also easy to interpret. The first crack passes through the plate undisturbed. The 

second crack approaches and is stopped at the intersection since it is unable to 

traverse the previously cleaved material. Intersections are commonly at 90o 

since the second crack moves at right angles to tensile stresses and tensile 

stresses cannot be carried across the previously cleaved crack. Examination of 

the fracture surfaces (to be discussed in the next chapter) will also corroborate 

the interpretation, since the fracture surface of the first crack is planar and 

undisturbed though the intersection. Wallner lines and hackle lines for the first 

crack plane are continuous and connected on each side of the intersection. 

4-26 
 

 

 

 

 

 

 

 

 

 

(a) Branching crack (b) T intersection 

General Examination 

(c) Crossing cracks
	

Figure 4.16 Intersecting cracks. (a) shows common crack branching, which 

can be distinguished by the gradual forking. (b) shows a T intersection. The 

numerals indicate the first and second cracks. (c) shows a crossing crack, 

which may occur if the first crack does not completely cleave the part. 

The scenario shown in Figure 4.16c is less common, but occurs in cases where 

an initial crack does not completely cleave the plate. This can easily occur with 

bending fractures wherein the crack leads on one side that is in tension, but 

does not necessarily go all the way through on the compression side. The 

window in Figure 4.14b has examples. In such cases, the remnant ligament of 

unbroken material provides a path for a crossing crack to traverse the first 

crack. The key feature to look for is a disturbance in the fracture surface 

features. The fracture surface markings (Wallner lines and hackle) on the first 

plane are continuous on either side of the intersection. On the other hand, the 

second crack has a pronounced change in the markings on its crack plane, often 

a jog, or a hook around and up to complete the fracture just on the other side of 

the first crack. Sometimes two cracks seem to cross exactly over each other at 

exactly 90o

, but a very close inspection of the exact intersection will show that 

one (the second crack) has a slight offset or step where it intersects the other. 

4-27 
 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

4.9 Invisible Cracks 

Incomplete fractures sometimes create cracks that are difficult or impossible to 

see. The interfacial separations may be less than the wavelengths of light and 

the cracks are not visible even under a microscope. The cracks may even be 

partially healed. Interfacial adhesion may be due to hydrogen bonding from 

water molecules adsorbed onto the fresh surfaces following the initial crack 

opening.19,20,21,22 The healing can occur even in inert environments and 

vacuum, and indeed, the literature shows that dangling ruptured silicate bonds 

do link back up. When a crack is created but closes quickly, as in the case of a 

window impact, pockets of air may be trapped between the crack faces. These 

pockets may be visible if light is reflected off the glass-air interface. Sometimes the light is refracted creating elusive but colorful reflections that are 

visible only at certain angles. (This is one reason why it is wise to rotate and 

shift a piece while inspecting it, or to move around it if the piece is stationary.) 

“Chill check” cracks, described in section 6.8.3 and shown in Figure 6.57, are 

tiny cracks formed by transient, very-localized thermal shock. They can be 

very difficult to see. Sometimes crack segments will appear to be isolated and 

not connected to other cracks or to an origin, but they are in fact connected by 

invisible segments. Fréchette1 described a case wherein an invisible long crack 

in an intravenous bottle went unnoticed in a hospital, leading to a fatality. 

4.10 Plates and Windows 

Plates and windows can exhibit a wide variety of fracture patterns. Some 

common modes are shown in Figure 4.17. An excellent article on the fracture 

of flat glass is by Shinkai, reference 3. His article shows subtle variations in 

the fracture patterns depending upon whether the plate edges are freely 

supported or are built in. Variations also occur with plate thickness. The extent 

of residual stress in tempered glass plates may also be estimated from the size 

of the diced fragments as will be discussed in 7.2.2. Heat strengthened glasses 

also have compressive residual stresses such that the glass is twice as strong as 

an annealed plate, but the pieces do not dice into small fragments. Heat 

strengthened pieces break into triangular shards like annealed plates. 

Blunt objects may create a cone crack that penetrates partially or completely 

through the plate. Such a flaw may or may not necessarily constitute a failure, 

since the plate is essentially intact. At higher velocities, radial cracks may be 

generated from the impact site. Radial cracks also may be triggered by impact 

of sharp objects. 

4-28 
 

 

 

 

 

 

 

 

 

 

` Uniaxial bending Pressure loaded Diagonal compression 

Supported on the sides Supported all around (after Ref. 3) 

Pressure loaded Pressure loaded Pressure loaded 

Freely supported Freely supported Freely supported 

(After Ref. 3) (After Ref. 3) (After Ref. 3) 

Ball impact Center impact Center impact 

Freely supported Secondary ring cracks Secondary ring cracks 

General Examination 

Sharp impact Blunt impact 

Tempered glass (BB hit) 

Figure 4.17 Plate and window fracture patterns. Figure 4.14b shows an 

example of a center impact plate with secondary ring cracks 

4-29 
 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

In other cases, the bending forces from the impact may cause the crack (once it 

is away from the immediate impact site) to run on the opposite side of the plate 

from the direction of impact. At even higher velocities, many radial cracks fan 

outward from the impact site. The continued loading of the plate causes the 

radially-fractured segments to bend inward, causing them to break in bending 

leading to circumferential secondary cracking. In these secondary fractures, the 

maximum tension is on the impacted side of the plate. These rings around the 

impact often have offsets at the radial cracks, confirming that the radial cracks 

occurred first. Figures 4.14b and 4.17 show examples. Sometimes a secondary 

ring crack may step across a radial crack if the latter has not completely severed 

the plate. Very blunt objects may not necessarily create cracking at the impact 

side. Bending forces can create tension opposite the impact side and can 

trigger crack growth from flaws on the opposite side. In some instances, the 

very blunt object can generate bending forces in the plate that triggers a crack 

from a flaw at the plate edge. A crack runs to the impact site, and then radiates 

and branches repeatedly outward. In analyzing such fracture sequences, it is 

often very useful to combine the overall crack pattern analysis with an 

examination of the fracture surfaces, as described in the next chapter. The 

fracture surface examination reveals whether the crack was running uniformly 

through the thickness or was leading on the inside or outside plate surfaces. 

The latter information can establish whether the plate was in bending, and if so, 

which side was in tension. 

4.11 Tempered Windows 

Tempered glasses fracture into many small fragments. Occasionally such 

fractures are triggered by internal flaws such as nickel sulfide inclusions 

(discussed in Chapter 6). More commonly, the fracture is triggered by impact 

or sharp contact loading on the surface. If the loading is sufficient, it drives a 

surface crack through the surface compression temper zone and into the interior 

tensile stresses. Once this occurs, the plate will spontaneously fracture into 

many small fragments, since the internal stresses are more than sufficient to 

cause propagation and branching. Often the diced glass remains in place and 

does not fall apart. Figure 3-10 showed an example of a street skylight. The 

origins can often be deduced from the branch crack patterns as shown in 

Figures 3-10 and 4.18. Close examination of the origin area often reveals that 

in moderate impact or loading cases, the crack branching creates matching 

symmetrical hexagonal shaped pieces at the origin as shown in Figures 3-10 

and 4.18c. Indeed, as will be discussed in section 7.3.2 and Figure 7.6, the first 

branch distances may be used to estimate the temper stress. An excellent 

article on the fracture of tempered flat glass is by Shinkai, reference 3. Shinkai 

states that some higher energy impacts may not have the two telltale fragments 

4-30 
 

 

 

 

 

 

 

 

General Examination 

at the origins, and may instead have cone cracks that penetrate from the 

surface. 

Finding the origin fragments may seem hopeless for extensively fractured 

tempered plates if the fragments have fallen out of a frame, but the hexagonal 

(or pentagonal or heptagonal) fragments are distinctive and can be found by 

simply sifting through the rubble. With persistence and a little luck, the patient 

fractographer may be rewarded. Professor J. Varner has singled out origin 

pieces from a bucket full of swept up fragments. In some instances, tempered 

plates do not break up into uniform fragments. The existence of atypical 

fragments is often a sign that the temper was non-uniform either by design 

(e.g., Shinkai, Ref. 3, p. 288) or not (e.g., Fréchette, Ref. 1, case 10.18), or 

relieved by high temperature exposure (e.g., Fréchette, Ref. 1, case 10.22). 

a b 

c d

Figure 4.18 Tempered glass fractures in the author’s son’s car. Flying debris 

from a tornado in College Park, Maryland in September, 2001 caused the 

impact fractures (a, b, d). The arrows show the actual origin in one window 

and (c) shows some variations of shape of typical matching origin pieces. The 

fracture surface of the origin piece is shown in Figure 5.19b 

4-31 
 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

4.12 Thermally Induced Plate and Window Fractures 

A center-heated plate is one which is heated in the middle while the rim is 

cooler. Differential strains put the middle of the plate into compression and the 

rim into tension. The plates fracture as shown in Figures 4.19 and 4.20. The 

crack starts from an edge origin in tension and initially propagates at 90o to the 

edge, but then changes to a meandering wavy pattern as it approaches the 

initially biaxial compression stressed middle portions of the plate. The waves 

are often periodic as shown in (b). The waviness is an example of Hopf 

(a) (b) (c) 

Figure 4.19 Fractures in center-heated plates. (a) and (b) are low stress 

fractures (less than 10 MPa {1,500 psi} in glass), and (c) is a higher stress 

fracture such that branching occurred. 

(a) (b) 

Figure 4.20 Thermal fractures with origin marked by arrows. (a) shows an 

alumina furnace baffle plate. Fracture started from grinding cracks associated 

with a hole. (b) shows glass slides broken by simply heating them on a hot 

plate. 

4-32 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

General Examination 

bifurcation problem that can occur in nonlinear systems.23,24 Figure 4.20 

shows a thermal failure in an alumina plate. Case 6 in Chapter 10 shows 

another example in a SiC furnace plate. Fréchette1 noted that meandering 

cracks are not proof of thermal stresses and they can also form in impact cases 

in the final stages of breakup. 

Thermal stress fractures often have localized tensile stresses that start fracture, 

but diminish away from the origin site. Branching may be minimal and the 

fracture surfaces may be relatively featureless and flat in regions away from the 

origin. Wallner lines or hackle lines (discussed in the next chapter) are rare. On 

the other hand, there may be multiple arrest lines attesting to stepwise 

propagation of the crack. Scarps may be present on the initial parts of the 

fracture surface if the part was thermally shocked by a cooler fluid. 

The relatively featureless fracture surfaces mean that the crack is moving 

slowly when the wavy pattern forms. The crack alters the stress distribution and 

it in turn alters the crack path. This nonlinear interaction causes the crack to 

curve back and forth. 

4.13 Bottles and Pressure Vessels 

Bottles and pressure vessels are susceptible to a variety of failure scenarios. 

Figure 4.21 shows some common modes, but this illustration is by no means 

complete. The reader is referred to the excellent review article by Kepple and 

Wasylyk25 for a more detailed presentation with more illustrations. Preston’s 

article26 in 1939 was an outstanding early contribution. Kepple and Wasylyk 

cite earlier work by F. W. Preston from the 1930s and 1940s in their review. 

Fréchette1 also illustrated a number of interesting glass bottle fractures. 

Fractures from internally-pressurized bottles (e.g., carbonated beverages, Fig. 

4.21 a-c, have vertical initial cracks since hoop tensile stresses are double the 

axial stresses (as discussed in detail below). Fractures usually initiate on the 

outside wall and branch in a symmetrical pattern about the vertical axis. The 

number of branches is proportional to the stress in the glass. Fractures start on 

the outside wall since larger flaws are more apt to be present there, but as the 

crack propagates it will extend to the interior and then lead slightly on the 

internal surface, since the stress is slightly higher there. Secondary breaks from 

bending have crack fronts that strongly lead on the inside surface. Pressure 

fractures from the base will also have extensive forking as shown in Figure 

4.21d. 

4-33 
 

 

 

 

 Fractography of Ceramics and Glasses 

a b c d e 

f g h i j 

Figure 4.21 Bottle fracture patterns. (a-c) sidewall fractures due to internal 

pressure at progressively greater pressures; (d) internal pressure fracture 

starting from the base; (e) is a thermal fracture from sudden cooling of the 

base; (f) is an impact fracture on the right side (labeled 1) with a hinge 

fracture on the side (labeled 2); (g-i) show water hammer fractures and (j) 

shows a fracture from diametral rim clamping. 

4-34 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

General Examination 

Thermal fractures have very little or no branching as shown in Figure 4.21e. 

Sudden temperature differentials can occur in a variety of different ways. 

Tensile stresses are generated when a portion of the body is suddenly cooled 

and attempts to shrink. The warm portions resist the contraction and put the 

cooler portions into tension. The temperature gradients can be either through 

the thickness or from one part of the body to another. Exposure of a hot vessel 

to a cold liquid often leads to thermal shock fracture. Fracture often starts and 

runs around the base and may extend up into the sidewall. Fréchette1

 and 

Kepple and Wasylyk25 cover more thermal stress fractures of glass containers. 

Additional information about thermal fractures is in section 4.19.2 later in this 

chapter. 

Sometimes impact damage sites may be found around the base, either on the 

inside or the outside. Impact damage sites on the inside of drinking glasses or 

beer mugs are often due to careless dropping of eating utensils into the glasses. 

External impact fractures on the side walls often have a starburst pattern at the 

impact site as shown in Figure 4.21f and 4.22. The principal crack system does 

not necessarily start at the impact site, however. The force of impact can cause 

the sidewall of the vessel to flex outwards creating a bending stress in the side 

of the vessel. This can trigger so-called “hinge fractures” 45o to the side of the 

impact site. Wallner lines indicate that fracture starts on the outside surface of 

the hinge fracture sites. A network of cracks from the impact site can start on 

either the inside or outside of the impact site, depending upon the sharpness of 

the impactor. The impact site may create the initial fractures, but in many cases 

the hinge origins fracture first. They then send a leader crack over to the 

impact site that triggers or exacerbates the fragmentation at the original impact 

site. Figure 4.22 shows an example of this in a glass tube. 

Filled glass containers dropped on their side trigger fracture from origins on the 

inside surface. 

Water hammer is a phenomenon that can cause vessel fractures as shown in 

Figures 4.21g-i. The sudden arrest of a liquid in motion or a surge of liquid 

that collapses a void in a partially-filled vessel can create transient intense 

internal localized pressures. Figure 4.21g shows a base fracture initiated by 

water hammer when a filled bottle was dropped. There may be a high 

concentration of cracks at the origin, but little branching afterwards since the 

stresses were focused at the base. Collapse of a void can occur in vacuumsealed or incompletely-filled vessels. Base fractures shown in Figures 4.21h,i 

can result. Vessel necks are vulnerable to many forms of damage from chipping 

4-35 
 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

due to capping-clamping stresses. Figure 4.21j shows a crack triggered by 

diametral clamping around the rim. 

(a) (b) 

(c) 

A B A C B

B C 

Figure 4.22 Small glass tube loaded diametrally at opposing points A (only 

the front is shown). Fracture started at the side hinge point B which generated 

an array of radiating cracks. One of these is the “leader crack” C (arrows in b 

and c) which went to the loading site A where it branched repeatedly and 

fanned out yet again. (Tube courtesy of W. Haller.) 

4-36 
 

 

 

 

 

 

 

 

General Examination 

(a) 

(b) (c) (d) 

(e) (f) 

Figure 4.23 Pressurized thin-walled vessels. (a) shows that the side walls are 

stressed in biaxial tension. Hoop stress is double the axial stress and therefore 

the first fracture is usually vertical as shown in the medicinal vial shown in 

(b)–(f). (e) shows the matching fracture surfaces at the origin. There is a 

distinct fracture mirror centered on the origin. (f) shows the origin was an arc 

shaped crack in the outer wall, a clear indication of blunt contact damage. 

Inner wall Outer Inner wall 

walls 

Origin 

Origin 

Outer 

surface 

Fracture 

surface 

Label 

removed 

and vial 

reoriented 

Origin 

4-37 
 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

As noted above, fractures from internally-pressurized bottles (e.g., carbonated 

beverages, Fig. 4.21 a-c, or medicinal vials Fig. 4.23) usually have vertical 

initial cracks since hoop tensile stresses are double the axial stresses, 

irrespective of the pressure. This is a good opportunity to show a stress 

analysis of a simple structure. 

Imagine making a vertical cut through the middle of the bottle with height L, 

radius r, and wall thickness t, under pressure P. This is a common practice for 

mechanical engineers to do when they construct “free body diagrams” to 

analyze stress problems. A part is cut in an imaginary sense. The external 

forces and moments acting on one half of the component are balanced with 

internal stresses in the walls. For the bottle, the forces from the sideways 

expanding pressures that act on the arc of the inside surface are balanced 

against the constraining forces in the side walls. These are the sum of the hoop 

stresses in the two side walls that have cross-sectional area 2Lt: 

� PPPPPPPPPPrrPPPPPP = 2rrPPPP = 2PPππσσhoooooo

ππ�2

−ππ�2

(4.2) 

σσhoooooo = PPrr

ππ (4.3) 

For the axial stress (vertical direction), imagine splitting the vessel in a 

horizontal plane midway along its height. The force created by the sum of the 

pressures on the bottom circular surface of area πr2 is balanced by the sum of 

the vertical axial stresses in the vessel side walls which have a total cross 

sectional area of 2πrt: 

PPPPrr2 = 2PPrrππσσaaaaaaaaaa (4.4) 

PPrr σσaaaaaaaaaa = (4.5) 2ππ

Hence: (4.6) σσhPPPPoo = 2σσaaaaaaaaaa

This solution is for thin-walled vessels where it may be assumed there is 

negligible stress gradient through the wall thickness. Engineering textbooks 

should be consulted for thick–walled vessels since there will be a stress 

gradient through the wall thickness. 

4-38 
 

 

 

 

 

 

 

 

 

 

 

General Examination 

4.14 Torsional Fractures 

Torsional loadings such as shown in Figure 4.24 produce twisted fracture 

surfaces at an angle to the part length. Torsional stresses usually do not induce 

a special failure mode in ceramics or glasses. Fracture starts at a flaw and the 

crack propagates normal to the plane of maximum tensile stress, which is 

aligned at 45o to the shear stress direction. In other words, the shear stress (τ) 

state is equivalent to one in which a material element is pulled in tension (σ1) in 

one direction and compressed laterally (σ2) as shown in Figure 4.24a. Fracture 

primarily occurs due to σ1 tensile stress and fracture propagates perpendicular 

to σ1. Many parts that fail in torsional loading also have some bending as 

well. Fracture surfaces often have ample twist hackle markings as discussed 

later in section 5.3.3. Figure 4.11 shows an example of a fracture in a bend bar 

that had superimposed bending and torsional stresses. Figure 5.31 shows the 

fracture surface of a glass rod with mixed tension and torsional loading. 

τ

σ1 

σ2 

(a) (b) 

Figure 4.24 Torsional loading causes fractures at an angle to the longitudinal 

axis of a part. The shear stresses are τ. These generate equivalent principal 

stresses σ1 in tension and σ2 in compression as shown in (a). Fracture occurs 

perpendicular to σ1. (b) shows a schematic of a glass rod fractured in torsion 

(after Michalske, ref. 27). The origin is at the arrow. Note how branching 

occurs to either side of the origin, but that only one branch in each case 

continues, giving rise to the curved fracture surface. 

4-39 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

4.15 Chipping 

Concentrated loads near an edge can chip off a portion of the body. Figure 

4.25 shows some examples. Chipping can either be a primary mode of fracture, 

or it can be a secondary nuisance. For example, it is one of the leading causes 

of failure of all-ceramic dental restorations such as crowns or bridges, and 

hardmetal (e.g., tungsten carbide) or ceramic cutting tools. On the other hand, 

it may be a nuisance if there is so much chipping that it interferes with 

interpretation of a fracture from some other cause. Thus, it is important to be 

careful when fitting pieces together during an initial examination. 

The force to cause fracture depends upon the shape of the object that applies 

the contacting force, the distance from the edge, the angle of the applied force 

P, the angle of the edge, and the material’s fracture toughness. In general, the 

greater the distance away from an edge, the greater the force that is necessary 

to create a chip. Modern use of this test for engineering materials was started 

in 1986 at the National Physical Laboratory in United Kingdom by Almond and 

McCormick.28 A remarkable finding was that the shape of the chip was 

invariant with material. Edge chipping is a field of growing attention, and there 

is a growing literature on it. Quite interestingly, Charles de Freminville’s 

seminal 1907 treatise on fracture of brittle materials included excellent 

illustrations (Fig. 5) of edge chips in glass.29 Several papers on edge chipping 

were given in a session at the 2000 Alfred Conference on Fractography of 

Ceramics and Glasses.30 Edge chipping also has been a topic in the lithic 

fracture literature and flint knappers have a good practical knowledge on the 

topic (e.g., Refs. 11, 31). 

The exact force-distance trend depends upon the type of indenter (sharp or 

blunt). In a laboratory setting, sharp diamond indenters (Rockwell, Vickers, 

Knoop) or sharp conical tools are used, but in practice any contact type loading 

can create a chip at an edge. Lab-scale indenters initiate cracks around the 

contact site which then grow to form a chip. On the other hand, some loadings 

with blunt contactors near an edge cause a pre-existing crack to propagate 

parallel to the specimen side and to form a long flake. The top of the flake, 

where the force is applied, is called the “platform” in the lithics literature. The 

fracture mechanics of chip formation and how long slender flakes may be 

formed by stable crack extension are fascinating.31,32 The long slender flakes 

can be used as cutting tools. Cotterell makes the case that the ability to form 

such long slender stone tools was a critical step in the evolution of mankind.31 

4-40 
 

 

 

 

 

 

 

General Examination 

(a) (b)

 (c) (d) 

(e) (f) 

Figure 4.25 Examples of chip fractures. (a) is a schematic that shows how a 

chip is formed and the telltale shape. (b) shows a top view of a chip that is just 

about to pop off a lithium disilicate foundation dental glass-ceramic. 

The loading point is obvious. (c) shows a row of edge chips created in 

porcelain under controlled conditions in the laboratory. (d) is a close-up. (e) 

and (f) show edge chips in an ancient Roman (circa 300 AD) 22 cm square x 

3.4 cm thick hypocaust stone from Trier, Germany. (b,c,d, courtesy J. Quinn) 

4-41 
 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

The shape of chips varies with the angle of the applied load as shown in Figure 

4.26.33 This has practical significance, since the shape of the chip can tell a 

fractographer the direction of the force that caused the chip. Since the first 

edition of this book was written, I have followed up on the edge chipping work 

of my departed spouse Dr. Janet Quinn. We have published a number of papers 

on edge chipping of dental ceramic and composite restorative materials and 

structural engineering ceramics and we have a new quantitative model.34,35,36 

Early work emphasized a linear relationship between force and distance the 

(a) 

(b) 

Figure 4.26 The angle of the applied load affects the shape of edge chips. (a) 

is a schematic that shows the patterns and (b) shows front views of two chips in 

glass (at different loads and distances from the edge) made by a conical 

indenter tool angled away from the edge at 105o

. (after J. Quinn and R. 

Mohan, Ref. 33). 

4-42 
 

 

 

 

 

 

 

 

 

 

General Examination 

force is applied away from the edge, but new results indicate nonlinear trends 

can result from the partitioning of indenter energy into competing deformation 

and fracture modes in the material. Figures 4.27 and 4.28 show some of the 

new chips. 

(a) 

(b) (c) 

Figure 4.27 Edge chipping. (a) shows the test machine with a 120o sharp 

diamond scriber making a chip in a dental stone material. (b) and (c) show a 

ceramic ball indenter chipping brittle denture material tooth specimens. 

(a) (b) 

(c) (d) 

Figure 4.28 Edge chips formed in (a) dental lithium disilicate glass ceramic; 

(b) dental 3Y-TZP zirconia; (c) dense alumina; and (d) laminated 

alumina/alumina-zirconia ceramic. The chips were stained with a green felt tip 

pen to make the fracture surfaces easier to see in these translucent materials. 

4-43 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Edge chips are very common as secondary fractures on broken ceramic or glass 

fragments. As secondary fractures, they may have had nothing to do with the 

original cause of fracture. They easily occur if fragments bump into each other 

or impact other objects during breakage or subsequent handling. If there is any 

doubt as to whether an edge chip is a secondary or a primary fracture, compare 

the matching fracture halves. If the chip is only on one half, then it is a 

secondary fracture. It is quite common, especially with glasses, to have one 

fracture half rub against the edge of the matching fracture half, causing chips in 

the latter. Undesirable secondary fracture chips can occur in bend bar strength 

specimens and other test coupons if the pieces fly around and impact the 

fixtures after the first break. Often the tensile edge of a bend bar will hit first, 

and the genuine fracture origin may be chipped off! Soft cushioning material 

should be placed in the fixtures under the bars as discussed in section 4.6. 

Some additional means to prevent the fragments from flying about should also 

be used (e.g., cotton or rolled paper or tape pieces). 

4.16 Laboratory Test Fractures and Component Fractures 

Examination of laboratory test coupon fracture surfaces can often aid 

interpretation of component failures. It is not unusual to be confronted with 

only one or two component fractures. The fracture surface markings may be 

unfamiliar and common fracture markings such as hackle lines or cantilever 

curls may be hard to recognize. A few broken bend bars, tension specimens, 

or fracture toughness specimens of the same material may be invaluable in such 

cases. The fractographer can study the crack propagation behavior and 

markings for these familiar configurations. 

The author used this approach to aid the interpreting of broken large (≈1 m) 

sintered porous fused-silica missile radomes in the mid-1980s. This is Case 7 

discussed in Chapter 10. Only a few broken components were available and 

an initial examination of the weak material with its rough fracture surfaces was 

unsuccessful. Examination of bend bar and double torsion fracture toughness 

specimens gave the author more experience in interpreting crack propagation 

directions. Optimum examination conditions for detecting hackle and Wallner 

lines were discovered. This know-how was then applied to the full size 

components, with the result that fracture surfaces which initially confounded 

the author became tractable on the reexamination. The origins were found. 

4.17 Controlled Component Fractures 

Sometimes one must resort to controlled laboratory fractures of components 

themselves in order to simulate expected or actual fracture modes. Morrell et 

4-44 
 

 

 

 

 

 

 

 

 

 

 

 

 

General Examination 

al.37 and Richter38 have documented two superb examples of this approach for 

ceramic hip joint balls. Morrell et al.

37 supplemented their component testing 

with biaxial disk and bend bar fracture testing. 

An excellent recent analysis of 27 in vivo (in the patient’s mouth) all-ceramic 

dental crown fractures was recently described by Øilo et al.39 Crowns are 

shell-like ceramic restorations which are cemented onto a tooth that is too 

damaged for repair by amalgam or composite fillings. The crowns were a 

dense alumina, primarily from one source. Most of the clinical fractures split 

the crowns vertically into two or more pieces such as shown in Figure 4.29. 

Laboratory testing (in vitro) of 30 model crowns of the same material system 

was done using three different loading schemes. One of the loading schemes, 

occlusal (top surface) loading of the crowns with a large 30 mm diameter ball 

pressing on a 3 mm thick rubber disk (to distribute the load and prevent 

localized contact damage) generated nearly identical fracture patterns as in in 

vivo fractures. Fractographic analysis of both the model and clinical 

(a) 

(c) 

(b) 

A 

B 

Zinc phosphate 

cement 

Rubber 

insert 

Alumina 

crown 

Occlusal surface 

margin origin 

Piece A 

Figure 4.29 Alumina molar crown (clinical 33) which fractured after 6 years 

in the patient’s mouth. (a) shows an occlusal surface (top) view. (b) shows a 

side view and the fracture surface of piece A. Careful examination showed the 

origin was on the margin on the left, and the fracture ran up through the crown 

and to the other side, splitting it in half. (c) shows the laboratory test set up 

which produced nearly identical fracture patterns. The epoxy base simulated 

the dentin of the reduced human tooth. (courtesy M. Øilo) 

4-45 
 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

fractures confirmed that the occlusal loading generated internal expansion of 

the shell-like crown, creating high hoop stresses along the margins (the open 

end of the crown that is in contact with the gum line). Subsequent work 

verified that the clinically-fractured pieces had excessive faults (chips, cracks, 

irregularities) at much too sharp (~0.1 mm) margins. Figure 6.70 is one 

example, and several more are shown in Case 12 in Chapter 10. The 

combination of the faults with the hoop stresses led to premature splitting 

fractures. 

Another dental example, but on an actual human tooth, was by Dr. J. Quinn.40 

She mimicked a fracture of an upper right canine tooth that broke at the root 

several days after the patient had heart surgery. It was thought that blunt force 

damage during intubation by an anesthesiologist had caused damage that led to 

the fracture. Fractographic analysis showed that fracture started from the 

lingual (tongue) side and caused a curved fracture plane from cantilever 

loading. Dr. Quinn obtained a comparable tooth from an extraction (from 

another patient) and loaded it in cantilever loading in a laboratory setting. It 

broke in an almost identical manner. This, plus tooth decay in the original 

tooth, evidence of step wise crack growth, and evidence of bruxism, showed 

the tooth was damaged before the surgery. The anesthesiologist was not at 

fault. Nevertheless, it was recommended that mouth guards (commonly used 

by bruxers when sleeping) ought to be used during intubation for surgery as 

well. 

4.18 Finite Element Analysis 

Stress solutions may be straightforward for simple shapes and loadings, but 

many components have complex shapes and asymmetrical loadings. Closedform solutions for the stress state may not be available. Finite element analysis 

(FEA) is a valuable tool for such problems. FEA models partition a body into 

small elements and uses computer analysis to obtain the stresses and strains in 

each using the theory of elasticity. In principle, the entire stress distribution in 

a body can be analyzed, provided that the loading conditions are accurately 

simulated. Software for FEA is becoming increasingly easy to use, but some 

skill is needed in setting up the problem, applying a good mesh, and identifying 

the correct boundary conditions. Richter utilized FEA results to help interpret 

the ceramic hip ball fractures.38 

FEA analysis can be a case of “garbage in and garbage out.” Modelers may be 

so enamored of the models with their assumed loading conditions and anticipated fracture modes that they are unreceptive to obvious evidence that a 

component is not fracturing “the way it is supposed to break.” A good 

4-46 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

General Examination 

modeler should be flexible and not afraid to modify or adapt the simulation 

based on the fractographic findings. Reference 41 reviews several case studies 

and compared model predictions to actual fracture modes. 

4.19 Characteristics of Some Common Fracture Modes 

4.19.1 Mechanical overload fracture 

This chapter has illustrated only a sampling of mechanical fractures. The 

general breakage patterns can sometimes lead one directly back to an origin 

site, but often it is necessary to examine the fracture surfaces to interpret the 

directions of crack propagation as shown in Chapter 5. Once a fracture pattern 

is recognized and an origin site identified, the next step is to assess whether the 

flaw is intrinsic to the manufacture of the part (as in a pore or inclusion), or 

from surface finishing, or from handling, contact or impact damage, or 

environmental attack (e.g., corrosion, oxidation, or erosion). Origins are 

covered in chapter 6. 

4.19.2 Thermal fracture and thermal shock 

Thermal stresses are created by differential strains created when portions of the 

body try to stretch or contract, but are constrained by other portions of the body 

at different temperatures. Thermal stresses may either be steady state (invariant 

with time) or time dependent. Gradual or linear spatial temperature gradients 

usually do not generate thermal stresses. Two types of thermal stress cases 

occur: (a) components with stresses arising from overall temperature gradients 

between portions of the body, and (b), surface and interior stresses caused by 

through-thickness temperature gradients. 

The center-heated plate problem discussed in section 4.12 is a classic case of 

(a). The differential expansion of the middle regions of the structure relative to 

the rim creates uniaxial hoop type tensile stresses at the rim that trigger fracture 

from an edge origin. The crack initially propagates perpendicularly to the edge, 

but then slows down and meanders in the interior biaxial compression stress 

field. The interior stresses are of insufficient magnitude to cause crack 

branching. The fracture surface markings (discussed in the next chapter) may 

be sharp and clear near the origin but then fade away to a very smooth 

featureless appearance as the crack slows down in the interior. Case 6 in 

Chapter 10 is an example of a center heated plate fracture in a silicon carbide 

furnace part. Center-cooled plates have the opposite stress state: tension in the 

interior and compression at the rim. Fracture is less likely to occur since the 

edge, which is likely to have the more serious flaws, is in compression. 

4-47 
 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Thermal stresses and strains occur during transient cool down or heat up 

events. Sudden temperature changes are called “thermal shock.” Thermal 

shock can cause strains between portions of a body or through-thickness 

stresses. Ceramics and glasses are most susceptible to sudden cool down 

thermal stresses, since tensile stresses are created at the surface. The surface 

layer of the part cools quickly and tries to contract, but the hotter interior 

portion prevents the contraction thereby creating tensile strains and stresses in 

the outer layer. Compressive stresses are generated in the interior. Suddenly 

cooling the surface with a fluid such as water can generate substantial tensile 

stresses and fracture since the fluid promotes rapid heat transfer from the 

surface. Some procedures, such as laboratory quenching of hot bend bars into 

water, create sharp stress gradients with a biaxial stress state at the surface. 

This leads to the formation of a two-dimensional network of surface cracks that 

may not necessarily penetrate completely into the interior. Examples are 

shown in Figure 4-30. Shao et al. 42 showed some images showing how 

cracking patterns in alumina plates changed with increasing temperature shock. 

At low shock, a few looping cracks formed. With greater shock, a dense 

cellular structure formed. With the maximum shock, a cellular structure with 

enveloping larger cracks appeared. 

(a) (b) 

Figure 4.30 Thermal shock cooldown cracks in alumina bend bars (a), and a 

disk (b). The cracks are highlighted by red dye penetrant. (courtesy of R. 

Danzer and T. Lube, Mountain Univ. Leoben) 

Very-localized thermal fractures from contact of a component surface with a 

cooler contacting object can create “chill checks” that are very localized tiny 

cracks (Figure 6.57). These do not penetrate far or very deep since thermal 

strains are very localized. The cracks are susceptible to extension leading to 

fracture by subsequent events or loadings. 

4-48 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

General Examination 

4.19.3 Impact or contact fractures 

Impacts may create localized damage sites that can become the origin of 

fracture if the impactor is small or sharp, but can also cause fracture by 

imposing bending stresses on a structure such that fracture initiates on the side 

opposite the impact site or elsewhere in the body. Section 4.13 shows both 

scenarios. Sharp and blunt impacting objects create different types of localized 

damage origins as shown in sections 6.7.4 and 6.7.5. When an impact or 

contact fracture is suspected, the surface should be carefully examined for 

“witness marks” as discussed in section 4.6.4. 

4.19.4 Corrosion or oxidation 

Oxidation and corrosion can weaken a part and leave it susceptible to fracture. 

Telltale surface reactions are often readily apparent on the surface in the form 

of numerous pits, discoloration, and surface scale. Sometimes the damage is 

confined to small localized pits, bubbles, or blisters as shown in Figure 6.26. In 

some difficult to interpret cases, the reactions are localized along grain 

boundaries and can only be discerned with the SEM. Chemical analysis can 

show cation migration to or from a reaction site. It can also indicate 

contamination or reaction with the environment or furnace linings or insulation. 

Richerson43 shows several examples of corrosion-oxidation damaged silicon 

nitrides. 

4.19.5 Residual stresses 

Residual stresses can arise from many sources. They can be intentional as in 

tempered or heat-strengthened glass or unintentional as in ceramics sintered in 

furnaces with thermal gradients. Residual stresses can exist at different scales 

within a structure. It is convenient to differentiate these as: bulk, through the 

thickness, localized, or microstructural residual stresses. 

There may be bulk residual stresses in a part whereby one portion or side is 

straining against other portions due to differential contractions during cooling. 

They can even cause spontaneous fractures in the absence of externally applied 

stresses. For example, there are reported cases of improperly-sintered porcelain 

toilets that fractured spontaneously leading to catastrophic house flooding. 

Residual stresses may also exist through the thickness of a part, as in heatstrengthened or tempered glass plates. Residual stresses may be very localized, 

such as around a sharp impact or contact damage site (e.g., a sharp particle 

impact site or Vickers indentation). Shallow residual stresses are also common 

4-49 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

on machined and ground surfaces. They typically are compressive within a few 

micrometers of the outer surface and change to tensile further beneath the 

surface. As such they can affect origins located on the surface. Finally, 

residual stresses may exist at a microstructural level, such as at the grain 

boundaries between grains or between two phases in polycrystalline materials, 

or between an inclusion and the matrix. The inclusion may have very different 

thermal contraction or elastic properties than the matrix. Thus, it is simplistic 

to say: “the piece had residual stresses.” One must be more specific and 

recognize that there may be multiple types of residual stress in a part. Furthermore, the existence of a residual stress may not necessarily be harmful. Some 

are compressive and some are tensile. Where an origin is located and the type 

of residual stress it is exposed to are important factors to consider. Crack 

propagation and branching may be dramatically affected as a crack moves from 

one type of stress field into another (e.g., compressive to tensile). Residual 

stresses may be highly variable within a part. Reference 44 is an interesting 

new paper on the variability of residual stresses in air-jet tempered glass plates. 

Reference 45 has other examples of nonuniformly-tempered glass including 

one illustration of a large automobile windshield. Uneven fragmentation and 

crack patterns are telltale signs of nonuniform tempering. 

The fractographer may be tipped off to the existence of residual stresses by the 

overall fracture patterns as shown in this chapter. A sudden change in crack 

propagation direction may be another sign. They can also be indicated by 

abrupt, seemingly spontaneous fractures when a component is being cut or 

machined. They may also be signaled by Wallner line markings on the fracture 

surfaces as shown in the next chapter. Distortions in the size and shape of a 

fracture mirror can be evidence of residual stresses. Analysis of flaw sizes can 

also be used to estimate residual stresses as shown in section 7.14. 

Fréchette described a case of boron carbide blanks fracturing during machining.46 Fractographic analysis showed that the fracture started at the root of 

the machine cut and then spread rapidly along the centerline of the slabs and 

then outward to the free surfaces. This indicated the interior was in tension and 

the outer portions were in compression. The firing schedule was reviewed and 

it was learned that the part was allowed to free cool between 1950 o

C to 1800 o

C followed by very gradual cooling thereafter. The free cooling occurred in a 

range where creep was possible which could have set up a final stress state not 

unlike tempered glass. Afterwards, no matter how carefully the slicing was 

done, once it reached the tensile interior stresses the part fractured. The 

problem was simply remedied by changing the cool down rate through the 

creep regime to minimize differential strains during the cool down, thereby 

eliminating the source of the residual stresses. 

4-50 
 

 

 

 

 

 

 

 

 

 

 

 

 

General Examination 

4.19.6 Time-dependent fracture 

Delayed fracture can occur due to slow crack growth from preexistent flaws. 

The loading conditions and stresses may initially be insufficient to cause 

breakage, but a flaw may slowly grow when it is under load until it reaches a 

critical condition and then triggers breakage. The rate of crack growth is very 

sensitive to the stresses and stress intensities acting on a flaw, and growth can 

be very slow or rather fast. For constant stress, cracks accelerate with time in 

accordance with fracture mechanics analyses shown in section 7.10. There 

may be little or no warning of the imminent failure since the extent the crack 

grows to become critical may be very small. Sometimes the slow crack 

growth leaves telltale markings on a fracture surface as shown in section 5.9. 

4.19.7 Compression fracture 

Genuine compression fractures are rare. The compression strength of ceramics 

and glasses is ten or more times greater than the tensile strength. Unless 

special precautions are taken in testing many specimens actually break from 

tensile stresses and not from genuine compression. Tensile stresses can arise 

from even slight misalignments or elastic property mismatches, and the tensile 

stresses will cause fracture.

47,48,49,50 Specimen design is critical. End splitting 

can occur if the compliance of the loading anvils creates sideways stresses at 

the ends of the specimen. Right circular cylindrical specimens are especially 

(a) (b) (c) 

Figure 4.31 Compression strength specimens. Cylindrical specimens (a), are 

very susceptible to end splitting shown by the red cracks. Dumbbell specimens 

(b) are also susceptible to end splitting if the loading anvils are not carefully 

chosen or are misaligned. (c) shows a valid compression fracture whereby the 

gage section is pulverized but the end caps are intact. (c is courtesy C. Tracy). 

4-51 
 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) (b) 

(c) 

σ

σ

σt 

σt 

σ

σ

Figure 4.32 Mechanisms of compression fracture. (a) shows microcracks 

(“wing cracks”) that form from preexisting flaws such as pores. (b) is a closeup of the flow lines of stress around an inclined semielliptical pore. Localized 

tensile stress occurs at the flaw ends. (c) shows a simulated microstructure 

with hexagonal grains. Microcracks can be generated at grain boundaries 

(left), or from elastic property mismatch at grain boundaries (middle), or from 

twinning or dislocation activity within grains (middle and right side). 

4-52 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

General Examination 

prone to these problems, and often fractures occur due to end splitting as shown 

in Figure 4.31a. Such results are invalid. Dumbbell tapered specimens are 

preferred, as shown in Figure 4.31b,c. A properly-tested specimen fracture is 

shown in Fig 4.31c. The extraordinary fragmentation of the gage section is 

notable. There is no point in looking for a single fracture origin. The goal of a 

fractographic analysis is more to verify that a test was done correctly, and to 

identify the mechanism of compression fracture as described in the next 

paragraphs. 

The primary mechanisms of compression fracture are shown in Figure 4.32. 

One mechanism shown in (a) and (b) involves preexisting flaws.51,52 Although 

the overall stress state is compression, the flow lines of stress cannot pass 

through an open flaw such as an inclined pore. They bend around the obstacle 

and create localized, sideways tensile stresses at the pore ends. (Shear stresses 

exist across the flaw and at its ends, but they can be resolved to localized 

tensile stresses as shown in Figure 4-32b.) Microcracks initiate, then grow 

stably and alter the flow lines of stress such that the local tensile stresses 

gradually rotate to a horizontal orientation. The microcracks propagate and 

align themselves parallel to the compression loading axis and eventually stop as 

shown in Figure 4.32a. These cracks are often called “wing cracks.” New 

microcracks initiate at other pores with different orientations and sizes. As 

loading continues, the structure becomes saturated with a high density of 

microcracks. Final fracture occurs with a sudden dramatic collapse and 

“explosion” of the specimen. Enormous stored elastic energies are released in 

the form of pulverization, kinetic energy of the fragments, sound waves and 

reverberations, and even triboluminescent light emission. The key point here is 

that compression strength is not controlled by a single large flaw. It is 

controlled by the flaw density and size distribution. Pores are used in the 

explanation above, but other flaws (e.g., inclusions) can induce the same 

response if the elastic properties of the flaw do not match the matrix or there 

are localized residual stresses around the flaws. 

Figure 4-32c shows several alternative modes of fracture that do not necessarily 

require initial flaws. Microcracks can be generated in a microstructure by 

several mechanisms.

53,54 Grain boundary microcracks (on the left in Figure 4-

32c) can be initiated due to elastic property mismatches between the adjoining 

grains that have different orientations. The elastic moduli and crystal constants 

can be dramatically different on different orientations, especially in non-cubic 

materials. Furthermore, strong residual stresses can exist along grain boundaries due to differential shrinkage between adjoining grains with different 

orientations. The middle of Figure 4-32c shows two grains with different 

orientations and moduli and, as vertical force is applied, the lateral strains can 

4-53 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

be very different, causing the nucleation of wedging cracks in one or the other 

grain. On the right of the Figure 4-32c, twins or dislocations within grains can 

initiate microcracks. Usually there are not enough independent slip systems 

active at low temperatures to lead to generalized plasticity, but the very large 

compression stresses can create sufficient dislocation activity and buildups to 

form microcracks within grains or at grain boundaries. 

Chapter 4 References 

1. V. D. Fréchette, Failure Analysis of Brittle Materials, Advances in Ceramics, Vol. 

28, American Ceramic Society, Westerville, OH 1990. 

2. F. W. Preston, “The Angle of Forking of Glass Cracks as an Indicator of the Stress 

System,” J. Am. Ceram. Soc., 18 [6] (1935) 175 - 176. 

3. N. Shinkai, “The Fracture and Fractography of Flat Glass,” pp. 253 – 297 in 

Fractography of Glass, eds. R. C. Bradt and R. E. Tressler, Plenum Press, NY, 

1994. 

4. R. W. Rice, “Ceramic Fracture Features, Observations, Mechanism and Uses,” pp. 

5 – 103 in Fractography of Ceramic and Metal Failures, Special Technical 

Publication 827, American Society for Testing and Materials, West Conshohocken, 

PA, 1984. 

5. R. E. Bullock and J. L. Kaae, “Size Effect on the Strength of Glassy Carbon,” J. 

Mat. Sci., 14 (1979) 920 – 930. 

6. G. D. Quinn and G. Wirth, “Multiaxial Strength and Stress Rupture of Hot Pressed 

Silicon Nitride,” J. Eur. Ceram. Soc., 6 (1990) 169 - 177. 

7. J. B. Quinn, “Extrapolation of Fracture Mirror and Crack-Branch Sizes to Large 

Dimensions in Biaxial Strength Tests of Glass,” J. Am. Ceram. Soc., 82 [8] (1999) 

2126 – 2132. 

8. D. Shetty, A Rosenfield, and W. Duckworth, “Crack Branching in Ceramic Disks 

Subjected to Biaxial Flexure,” Comm. Amer. Ceram. Soc., Jan. 1983, C10-C12. 

9. D. Hull, “Influence of Stress Intensity and Crack Speed on the Fracture Surface 

Topography: Mirror to Mist to Macroscopic Bifurcation,” J. Mat. Sci., 31 (1996) 

4483 – 4492. 

10.	 G. D. Quinn and R. Morrell, “Design Data for Engineering Ceramics: A Review of 

the Flexure Test,” J. Am. Ceram. Soc., 74 [9] (1991) 2037 - 2066. 

11.	 A. Tsirk, Fractures in Knapping, Archeopress, Oxford, United Kingdom, 2014, 

page 171, Figure 11.3. 

12. H. Kolsky, “The Waves Generated by Brittle Fracture in Glass,” Trans. Soc. 

Rheology, 20 [3] (1976) 441 – 454. 

13. J. Miklowitz, “Elastic Waves Created During Tensile Fracture, The Phenomenon 

of a Second Fracture,” J. Appl. Mech., Mar. 1953, 122 – 130. 

14.	 G. D. Quinn, “Twisting and Friction Errors in Flexure Testing,” Ceram. Eng. Sci. 

Proc., 13 [7-8] (1992) 319-330. 

15. H. Fessler and D. C. Fricker, “A Theoretical Analysis of the Ring-on-Ring 

Loading Disk Test,” J. Am. Ceram. Soc., 67 [9] (1984) 582 – 588. 

4-54 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

General Examination 

16.	 J. A. Salem and L. Powers, “Guidelines for the Testing of Plates,” Ceram. Eng. 

Sci. Proc., 24 [4] (2003) 357-364. 

17.	 D. Maccauley, Cathedral: The Story of Its Construction, Houghton Mifflin, New 

York, 1981. 

18. D. Maccauley, City: A Story of Roman Planning and Construction, Houghton 

Mifflin, New York, 1974. 

19. D. G. Holloway, “The Fracture Behavior of Glass,” Glass Technol., 27 [4] (1986) 

120 -133. 

20.	 T. A. Michalske and E. R. Fuller, Jr., “Closure and Repropagation of Healed 

Cracks in Silicate Glasses,” J. Am. Ceram. Soc., 68 [11] (1985) 586-590. 

21.	 S. M. Wiederhorn and P. R. Townsend, “Crack Healing in Glass,” J. Am. Ceram. 

Soc., 53 [9] (1970) 486 – 490. 

22.	 B. Stavrinidis and D. G. Holloway, “Crack Healing in Glass,” Phys. Chem. 

Glasses, 24 [1] (1983) 19-25. 

23.	 A. Yuse and M. Sano, “Instabilities of Quasi-static Crack Patterns in Quenched 

Glass Plates,” Physica D, 108 (1997) 365-378. 

24.	 A. Yuse and M. Sano, “Transition Between Crack Patterns in Quenched Glass 

Plates,” Nature, 362, 25 March 1993, 329-331. 

25. J. B. Kepple and J. S. Wasylak, “Fracture of Glass Containers,” pp. 207 – 252 in 

Fractography of Glass, ed. R. C. Bradt and R. E. Tressler, Plenum, NY, 1994. 

26. F. W. Preston, “Bottle Breakage – Causes and Types of Fractures,” Bull. Am. 

Ceram. Soc., 18 [2] (1939) 35 -60. 

27. T. A. Michalske, “Quantitative Fracture Surface Analysis,” pp. 652 – 662 in 

Engineered Materials Handbook Ceramics and Glasses, Volume 4, ed. S. 

Schneider, ASM Int., Metals Park, OH 1991. 

28.	 E. A. Almond and N. J. McCormick, “Constant-Geometry Edge-Flaking of Brittle 

Materials,” Nature, 321 [6065] (1986) 53-55. 

29.	 Ch De Freminville, “Caractères Des Vibrations Accompagnant Le Choc Dé 

L’Examen Des Cassures,” Review de Mettalugie, 4 (1907) 833-884. 

30.	 Fractography of Glasses and Ceramics IV, eds. J. R. Varner and G. D. Quinn, 

Ceramic Transactions, Vol. 122, Amer. Ceram. Society, Westerville, OH, 2001. 

31.	 B. Cotterall, Fracture and Life, Imperial College Press, London, 2010. 

32.	 M. D. Thouless, A. G. Evans, M. F. Ashby, and J. W. Hutchinson, “The Edge 

Cracking and Spalling of Brittle Plates,” Acta Metall., 35 [6] (1987) 1333-1341. 

33. J. B. Quinn and R. Mohan, “Geometry of Edge Chips Formed at Different 

Angles,” Ceram. Eng. Sci. Proc., 26 [2] (2005) 85 – 92. 

34. G. D. Quinn and J. B. Quinn, “A New Analysis of the Edge Chipping Resistance 

of Brittle Materials,” Ceram. Eng. Sci. Proc., 35 [1] (2015) 209 -224. 

35. G. D. Quinn, A. A. Giuseppetti, and K. H. Hoffman, “Chipping Fracture 

Resistance of Dental CAD/CAM Restorative Materials: Part 1, Procedures and 

Results,” Dent. Mater., 30 (2014) e99 – e111. 

36. G. D. Quinn, A. A. Giuseppetti, and K. H. Hoffman, “Chipping Fracture 

Resistance of Dental CAD/CAM Restorative Materials: Part 2, Phenomenological 

Model and the Effect of Indenter Type,” Dent. Mater., 30 (2014) e112 - e123. 

4-55 
 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

37. R. Morrell, L. Byrne, and M. Murray, “Fractography of Ceramic Femoral Heads,” 

pp. 253 – 266 in Fractography of Glasses and Ceramics IV, eds. J. Varner and G. 

D. Quinn, Cer. Trans., Vol. 122, Amer. Ceram. Society, Westerville, OH, 2001. 

38. H. Richter, “Fractography of Bioceramics,” pp. 157 - 180 in Fractography of 

Advanced Ceramics, ed. J. Dusza, Key Engineering Materials, Vol. 223 (2003), 

Trans Tech Publ., Switzerland. 

39. M. Øilo, A. K. Kvam, J. E. Tibballs, and N. Gjerdet, “Clinically Relevant Fracture 

Testing of All-Ceramic Crowns,” Dent. Mater., 29 (2013) 815-823. 

40. J. B. Quinn, L. W. Schultheis, and G. E. Schumacher, “Failure Analysis of a 

Broken Tooth,” J. Fail. Anal. Prev., 4 [1] (2004) 41-46 

41. G. D. Quinn, “Design and Reliability of Ceramics, Do Modelers, Designers and 

Fractographers See the Same World?” Ceram. Eng. Sci. Proc., 26 [2] (2005) 239 – 

252. 

42.	 Y. Shao, X. Xu, S. Meng, G. Bai, C. Jiang, and F. Song, “Crack Patterns in 

Ceramic Plates after Quenching,” J. Am. Ceram. Soc., 93 [10] (2010) 3006 – 3008. 

43. D. W. Richerson, Chapter 16 in Modern Ceramic Engineering, 2nd edition, Marcel 

Decker, NY, 1992. 

44.	 J. Anton, A. Errapart, M. Paemurru, D. Lochegnies, S. Hödemann, and H. Aben, 

“On the Inhomogeneity of Residual Stresses in Tempered Glass Panels,” Estonian 

J. Engineering, 18 [1] (2012) 3-11. 

45.	 K. Blank, “Some Practical Aspects of Thermally Strengthened Glass,” pp. 485 – 

499 in Strength of Inorganic Glass, ed. C. Kurkjian, Plenum, NY, 1985. 

46. V. D. Fréchette, “Fractography and Quality Assurance of Glass and Ceramics,” 

pp. 227 - 236 in Quality Assurance in Ceramic Industries, eds. V. D. Fréchette, L. 

D. Pye, and D. E. Rase, Plenum, NY, 1978. 

47.	 A. Rudnick, C. W. Marschall, W. H. Duckworth, and B. R. Enrick. “The 

Evaluation and Interpretation of Mechanical Properties of Brittle Materials,” U.S. 

Air Force Materials Laboratory Technical Report, AFML TR 67-316, April 1968. 

48.	 C. A. Tracy, “A Compression Test for High Strength Ceramics,” J. Test. and Eval., 

15 [1] (1987) 14-19. 

49.	 W. Dunlay, C. Tracy, P. Perrone, “A Proposed Uniaxial Compression Test for 

High Strength Ceramics,” MTL TR 89-89, US Army Materials Technology 

Laboratory, Watertown, Mass., 1989. 

50.	 G. Sines and M. Adams, “Compression Testing of Ceramics,” pp. 403-434 in 

Fracture Mechanics of Ceramics, Vol. 3, eds. R.C. Bradt et al., Plenum, 1978. 

51.	 M. Adams and G. Sines, “Crack Extension from Flaws in a Brittle Material 

Subjected to Compression,” Tectonophysics, 49 (1978) 97-118. 

52.	 F.M. Ernsberger, “Tensile and Compressive Strength of Pristine Glasses by an 

Oblate Bubble Technique,” Phys. Chem. Glasses, 10 (6) 1969 pp. 240-245. 

53.	 J. Lankford, “Compressive Strength and Microplasticity in Polycrystalline 

Alumina,” J. Mat. Sci., 12 (1977) 791 – 796. 

54.	 R. W. Rice, “The Compressive Strength of Ceramics,” pp 195 – 229 in Ceramics 

in Severe Environments, eds. W. W. Kriegel and H. Palmour III, Plenum, NY, 

1971. 

4-56 

 

 

 

 

 

 

 

 

 

Fracture Surface Examination 

5. Fracture Surface Examination 

5.1 Introduction 

Fracture surfaces have a wealth of information that may help determine the 

cause of fracture and even the fracture origin itself. The fracture surfaces 

should have an initial examination to look for any signs of contamination, and 

they should be cleaned as discussed in section 3.1 if necessary. There are two 

important cases where it may be wise to examine fracture surfaces before 

cleaning. First, surface contaminants may provide clues about the fracture 

event and may not be merely contamination picked up afterwards from 

handling. Figure 10.26e in Chapter 10 is a good example where discoloration 

at the origin was an important clue to understanding the initial flaw in a broken 

dental crown. Second, fine debris on a fracture surface may be a clue that 

cyclic fatigue may have occurred, if the debris is crushed material that was 

dislodged from the microstructure. This is covered in section 5.9 on page 5-84. 

The origin may range from a type and location anticipated by the fractographer 

(as in a broken bend bar) to an unexpected type from an unexpected location or 

cause (as in a broken component). With familiar laboratory strength test 

specimens, the fractographer may proceed directly to the primary fracture 

surface and quickly locate the origin area. Usually the fractographer seeks a 

fracture mirror, a key feature centered on the origin. Figure 5.1 shows two 

examples. The origins are often very small and magnification by a microscope 

is necessary. Stereoptical microscopes are ideal for initial fracture surface 

examinations. Sometimes this examination may be all that is needed for a full 

characterization of the origin. In other instances, the flaw at the origin is so 

small that scanning electron microscopy is needed. 

Unexpected failure modes or component fractures require more time to study. 

The crack propagation markings on fracture surfaces allow one to evaluate the 

state of stress, the crack velocities, and interpret crack propagation paths back 

to an origin. In instances where an origin is not obvious, it may be necessary to 

study many fragments and carefully determine the local direction of crack 

propagation. Stereoptical microscopes are ideal for this initial fracture surface 

examination. Wallner lines and hackle markings are helpful indicators of the 

direction of crack propagation (dcp) in each piece. It is helpful to mark which 

way the crack was running in a fragment. This can be repeated for other 

fragments and then the pieces reassembled so that the lines of crack 

propagation can be tracked back to an origin site. Even very complex fractures 

may be interpreted if the fractographer is patient and systematic. 

5-1 
 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

The greater the stress in the part at fracture, the more the stored energy, and the 

richer the fracture markings. Weak parts with low stored energy are often 

difficult to interpret. The part breaks into two pieces with relatively featureless 

fracture surfaces. Very porous or coarse-grained materials also can mask 

fracture surface fractographic markings. On the other hand, very strong 

(a)

 (b) 

Fracture 

mirror 

Compression curl 

Origin 

Mirror and Origin 

Compression curl 

Figure 5.1 Classic fracture patterns. (a) shows a low-power stereoptical 

microscope image of a fused silica rod fractured in bending at 88.5 MPa (12.8 

ksi) (b) shows an SEM image of a chalcogenide fiber fractured in bending at 

368 MPa (53.4 ksi). (b is courtesy of J. Quinn and V. Nguyen) 

5-2 
 

 

 

 

 

 

 

 

 

 

 

 

 

Fracture Surface Examination 

components such as some glass optical fibers may have so much stored elastic 

energy at fracture that excessive fragmentation occurs and the origin cannot be 

found. 

The nomenclature of fracture markings is varied. For example, “twist hackle” 

is also referred to as “river delta” lines. The Guide adopts the nomenclature 

system devised by Fréchette.1 His system is practical, logical, and designed to 

be easily remembered. Some of the definitions in the next sections are from 

Fréchette and many are new. 

“Conchoidal fracture” and “cleavage fracture” are terms commonly used in the 

geological community. Cleavage fractures are smooth and planar fractures 

such as those on preferred breakage planes in single crystals. Conchoidal 

fractures do not follow natural planes of separation and often are curved, in a 

manner reminiscent of curved seashells. Conchoidal is derived from the Greek 

word for a mussel shell. Glass and many ceramic fractures are conchoidal. 

“Joint” is a common term in the geological community and means a fracture 

surface. Joints are defined as a fracture or crack in a rock mass along which no 

appreciable movement has occurred. “Parting” is another geological term 

describing flat fractures in a mineral, but is not necessarily along preferred 

cleavage planes. Parting is sometimes associated with twinning. 

5.2 Fracture Mirrors, A Fractographer’s Best Friend 

5.2.1 Fracture mirrors in glass 

Fracture mirrors: Relatively smooth regions surrounding and centered on the 

fracture origin. (ASTM C 1322 and C 1145) 

Once a fractographer finds a mirror, he is reasonably sure that the fracture 

origin has been found. Figures 5.1 to 5.3 show mirrors in glass. They are 

telltale features that can be recognized by even novice fractographers. In 

strong specimens they are approximately circular or semicircular in shape, but 

significant deviations may occur as discussed later. Usually there is only one 

mirror corresponding to a single fracture origin, but there are exceptions that 

are discussed later in this chapter. There won’t always be a fracture mirror 

and, if the fractographer cannot find one, he or she should not worry. More 

about this will be covered later in this chapter, but for now let us take a closer 

look at the fractographer’s best friend. 

The fracture mirror is the region where a crack radiates outwards from a flaw at 

5-3 
 

 

 

 

 Fractography of Ceramics and Glasses 

(a) 

(b) 

(c) 

Figure 5.2 Three views of a fracture mirror in fused silica glass rod broken in 

bending at a high stress level (96.2 MPa, 13.9 ksi). (a) shows the entire 

fracture surface. (b) is a close-up of the mirror, and (c) is the same region but 

with the illumination adjusted. Maximum tension was at bottom center in view 

(a). The mirror is asymmetric due to the bending stress field. 

5-4 
 

 

Fracture Surface Examination 

(a) 

(b) 

mirror 

hackle mist 

Figure 5.3 Two views of a fracture mirror in a moderate strength (76.2 MPa, 

11.0 ksi) fused silica glass rod broken in bending. The mirror is incomplete 

since it was large relative to the cross section size. The advancing crack 

propagated into a decreasing stress field in the interior and did not have 

sufficient energy to form the mist and hackle in a complete circle. If this rod 

had been loaded in direct tension instead of bending, the mirror would have 

been fully formed and semicircular. The incomplete mirror in the direction of 

the interior is a telltale sign of bending. 

5-5 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

the fracture origin. The crack accelerates from near zero velocity to terminal 

velocity within the mirror region. Using ultrasonic stress wave fractography on 

soda lime silica specimens, Richter and Kerkhof2 and Field3 showed that a 

crack accelerates to between one-half to the full terminal velocity (≈ 1500 m/s) 

by the time the mirror-mist markings formed. This remarkable acceleration 

over a very short distance occurs in microseconds. Smekal4 arrived at the same 

conclusion from his analysis of the Wallner lines (described later in this 

chapter) inside fracture mirrors. 

The terminal velocity in glasses ranges from 700 m/s to 2500 m/s depending 

upon composition and elastic properties. Terminal velocities are greater for 

ceramics. Table 5.1 lists some values. The terminal velocity is about 0.5 to 0.6 

times the Raleigh surface wave velocity in an elastic body, or about 0.38 times 

the longitudinal (dilatational) stress wave velocity. The Raleigh wave velocity 

is less than the longitudinal wave velocity, vl = √E/ρ, where E is the elastic 

modulus and ρ is the density. The terminal velocity of the crack is limited by 

the speed that elastic energy is transferred to the crack tip. In principle, the 

crack can travel as fast as the Raleigh wave speed, but in practice the local 

deviations of the crack front that create mist and hackle limit the crack speed. 

Some atypical shock loading cases can drive a crack faster than the normal 

terminal velocity provided that the crack does not develop mist, hackle, or 

branching. Maximum crack velocities in single crystals whereby mist and 

hackle formation are suppressed can reach a greater fraction of the Raleigh 

wave speed than glass and polycrystalline ceramics. In some ballistic impact 

conditions, a damage wave of cracking in glass can proceed at velocities 80 % 

to 90 % of the longitudinal wave speed.5,7 High speed photography revealed 

that individual crack segments traveled no more than the conventional terminal 

velocity of 1480 m/s, but that secondary cracks nucleated right behind the 

shock wave. 

If the stresses acting on the advancing crack decrease, the crack can decelerate 

and even stop. This is not unusual in thermal shock cases where stress 

gradients can be large. 

The term mirror came in to being as a result of early optical microscopy of 

glasses wherein the mirror region was so smooth that it reflected light like a 

mirror. From my research into the history of fractography of brittle materials, 

it appears that Brodmann6 first used the word “mirror” in 1894. Figure 5.2c 

shows a bright mirror region. (The illumination in this figure also accentuates 

the faint curved Wallner lines within the mirror, about which more shall be said 

later in this chapter.) In polycrystalline ceramics or composites, the qualifier 

5-6 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Fracture Surface Examination 

Material Maximum Crack 

Velocity (m/s) Ref. 

Soda lime glass 1510 8 

Soda lime glass 1580 3 

Soda lime glass 1540 12 

Soda lime glass 1600 9 

Soda lime glass 1500 11 

Soda lime glass 1460 7 

Various flint glasses 750 - 1400 8 

Heavy flint glasses 700 - 1095 9 

Mirror glass 1520 9 

Borosilicate crown glass (BK-7) 1677 9 

Borosilicate crown glass (BK-7) 1800 4 

Fused silica 2100 4 

Fused silica 2500 11 

Fused silica 2155 8 

Reaction bonded Si3N4 (AME) 2950 10 

Hot pressed Si3N4 (HS 130) 3500 10 

KCl, single crystal 700 3 

LiF, single crystal 1700 3 

MgO, polycrystalline (85% dense, 15 μm) 1800 9 

MgO, single crystal, cleavage {100} 5100 3 

Diamond, cleavage {111} 7200 3 

Al2O3, single crystal sapphire, cleavage on {10Ī0} 4500 3 

Al2O3, polycrystalline (Lucalox 15 μm) 2000 10 

Tungsten 2200 3 

Tool Steel 1700 10 

Resin Plastic 420 8 

PMMA 770 10 

Table 5.1 Terminal crack velocities from Field (3), Smekal (4), Winkler et al. 

(7), Schardin et al. (8), Kerkhof (9), Congleton and Denton (10), Chaudhri 

(11), and Barstow and Edgarton (12) 

“relatively,” as in “relatively smooth,” must be used, since there is an inherent 

microstructural roughness even in the area immediately around the origin flaw. 

The underlying principal that accounts for the micro and macro branching has 

been attributed to the crack reaching a critical velocity,13 a critical energy 

level,14 a critical stress intensity,15,16,17,18 or a critical strain intensity.19 A 

velocity criterion for crack branching was discounted by data shown by 

Congleton and Petch.16 Richter and Kerkhof,

2 Field,3 and Döll20 used 

ultrasonic fractography to show that cracks approached terminal velocity before 

the formation of the mist boundary and there was no pronounced change in 

velocity when the mirror formed. Kirchner et al.

17,18 showed convincing 

evidence that a stress intensity criterion accounts for the mirror sizes and 

5-7 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

shapes. Congleton and Petch16 demonstrated that a stress intensity criterion 

controlled branching. They also showed that the Johnson and Holloway energy 

criterion14 was related to the stress intensity criterion. Tsai and Mecholsky21 

have pointed out that for isotropic materials it is difficult to distinguish between 

the energy, stress intensity, or strain intensity criteria since they are all related. 

This topic will be discussed again in Chapter 7 on Quantitative Analysis and in 

Appendix D on Guidelines for How to Measure Fracture Mirrors. 

Careful electron microscopy and atomic force microscopy now have shown that 

the formation of the mirror is a gradual progression of very localized crack path 

deviations from the main plane. Figures 5.4 a-e show superb transmission 

electron microscope images of various regions in a glass fracture mirror 

collected by Beauchamp.22,23 Poncelet24 showed comparable images in 1958, 

but Gölz may have been the earliest with extraordinary electron microscope 

photos published in 1943.25 Rounded ridges that are elongated in the direction 

of crack propagation gradually coarsen in amplitude. At some point, they 

develop slight hackle steps where over running and under running fingerlike 

crack segments link. This point, where nano-micro scale hackle lines begin to 

form, could be an important transition point. Close examination of 

Beauchamp’s Figure 5.4 b shows they are not present inside the mirror, but 

they are just visible in his Figure 5.4c near the mirror-mist boundary. 

Attempts to correlate the mirror boundaries with surface roughness parameters 

have had only limited success, as discussed in more detail in Appendix D. 

Hull26 pointed out that different surface roughness characterization devices 

such as atomic force microscopes (AFMs), mechanical profilometers, and laser 

optical profilometers all have different advantages, disadvantages, sensitivities, 

and scanning zone sizes. AFMs can measure tiny regions with very high 

sensitivities, but may miss large hackle steps in a mist or hackle zone. These 

latter features can dramatically alter the average or root mean square roughness. 

Mist and hackle regions may have different roughness at different scales. Hull 

discusses these various scales of roughness in some detail in his book.26 

The mirror markings can be explained with reference to Figures 5.2, 5.3b, and 

5.4. The explanation below closely follows Fréchette’s1 general discussion as 

well as Beauchamp’s22 convincing analysis. As the crack accelerates away 

from the origin, micro portions of the crack front begin to twist slightly or tilt 

up and down out of the main fracture plane. These local deviations occur as 

consequence of the stress field in front of a fast crack having maxima that are 

out of plane, unlike the case for a static or slowly moving crack.13 The 

momentary tilting or twisting does not persist for very long since crack plane 

5-8 
 

 

 

 

 

 

 

 

Fracture Surface Examination 

(a) 

(b) 

Mirror near origin 

Mirror, 1⁄3 of the way to the mist 

Figure 5.4 Transmission electron microscope pictures from parts of the 

fracture mirror in a lightly tempered soda lime glass sheet. (courtesy E. 

Beauchamp, refs. 22,23) (a) is in the smooth mirror area near the critical flaw. 

The tiny black dots are an artifact of the specimen preparation, but the larger 

(50 nm) round bumps are genuine. (b) is in the mirror region about 1/3rd of the 

way to the mirror/mist boundary. The arrow shows the direction of crack 

propagation (dcp). The ridges and valleys are elongated in the direction of 

crack propagation, but they are too small to be detected optically. 

5-9 
 

 

 

 Fractography of Ceramics and Glasses 

(c) 

(d) 

Mist 

Mirror near the mist 

Figure 5.4 continued (c) is a region still inside mirror but close to the mist 

boundary. The ridges are longer and wider. (d) is in the mist, just beyond the 

mirror/mist boundary. The features are similar to the prior figure, except that 

they are now large enough to be detected optically. Close examination of (c) 

and (d) reveals faint but distinct twist hackle lines in some of the large features, 

attesting to some localized lateral crack growth. 

5-10 
 

 

 

 

 

 

 

Fracture Surface Examination 

(e) 

Hackle 

Figure 5.4 continued (e) shows features in the hackle zone. The features are 

now larger than 10 μm, but are similar in character to those in the mist and 

mirror zones. 

deviations are restricted by the energetic cost of creating additional crack 

surface. The slight tilt or twist variations in crack plane quickly rejoin the main 

propagating crack plane. These tiny local crack perturbations exist well within 

the mirror region, but are too small to be optically discernable. As the crack 

advances, they eventually become large enough to be just discernable with the 

optical microscope as the “mist” zone surrounding the origin (Figure 5.4d). 

The mist has a slight frosty appearance such as when water condenses on a 

reflecting mirror. As the crack continues to advance, the local perturbations 

increase and begin to oscillate and form larger tongue-like segments that may 

deviate from the average fracture plane to the degree that micro steps are 

generated running parallel to the direction of crack propagation. (Some have 

described the out-of-plane perturbations as “fingers.”) The perturbations 

gradually coarsen such that the tongue-like elements can overcut other portions 

of the crack front thereby generating large “velocity hackle” lines that run 

parallel to the direction of crack propagation. Beauchamp22,23 pointed out that 

many of these features are similar in character from within the mirror out to the 

hackle zone. They differ only in scale. The transitions between the regions are 

gradual and are usually described as not abrupt. One intriguing observation by 

Beauchamp22 was that tilted or twisted hackle segments that relink with the 

main crack plane may generate elastic pulses and Wallner lines (described later 

in this chapter) that could trigger additional hackle along the crack front. This 

5-11 
 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

could be the impetus to form a band of hackle along the crack front and hence 

the mist-hackle boundary. 

In summary, although the process of roughening through the boundaries has 

often been described as gradual, there are two potential transition points. The 

mirror–mist boundary may be a transition where nano-micro steps form 

between ridges that are tens of nanometers tall. The mist-hackle transition 

boundary may be where a band of microhackle forms that is triggered by selfgenerated or external elastic pulses (Wallner lines). 

Fréchette1 defined hackle and mist hackle. This Guide adopts his definition for 

the former, but makes a slight modification to the latter: 

Hackle: A line on the surface running in the local direction of cracking, 

separating parallel, but noncoplanar portions of the crack surface. 

(Ref.1, Fréchette) 

Mist hackle: Markings on the surface of an accelerating crack close to its 

effective terminal velocity, observable first as a misty appearance and with 

increasing velocity revealing a fibrous texture, elongated in the direction of 

cracking. 

It is implied in these definitions that hackle lines are physical steps on the 

fracture surface, as opposed to undulations. Mist is fairly obvious in glasses, 

but is difficult or impossible to discern in ceramics due to the microstructure. 

Fréchette’s definition of mist hackle originally also included the words “and 

coarsening up to the stage at which the crack bifurcates” which suggests it 

encompassed all mirror markings including hackle right up to the point of crack 

branching. It seems reasonable to distinguish the mist region from hackle 

region for practical reasons that will become evident when ceramic mirrors are 

discussed below. Therefore this Guide adds the following term: 

Velocity hackle: Hackle markings formed on the surface of the crack close to 

its terminal velocity, observable as discrete elongated steps aligned in the 

direction of cracking. 

Mirrors are very important since they not only draw attention to the fracture 

origin, but their size may be used to estimate the stress in the part at the 

moment of fracture. For a long time it was common practice to interpret the 

mirror as having distinct boundaries between the mirror and mist and between 

the mist and hackle. Extensive empirical work showed that the relationship 

between the mirror size and the stress at the instant of fracture is: 

5-12 
 

 

 

 

 

 

 

 

 

 

Fracture Surface Examination 

σ R = A (5-1) 

where σ is the tensile stress at the origin at the instant of fracture, R is the 

mirror radius, and A is a constant, known as the “mirror constant” and is 

considered by many to be a material property constant. It has units of 

stress√length, which are the same units as stress intensity which is discussed in 

detail in section 7.5.1. The mirror radius may either be the mirror-mist 

boundary or the mist-hackle boundary for glasses. There are different A’s for 

each boundary. For polycrystalline ceramics there usually is simply a single 

mirror-hackle boundary. Equation 5-1 is discussed in more detail in Chapter 7, 

Quantitative Analysis and in Appendix D (where guidance is given on how to 

actually make the measurement), but for the moment we need only consider 

some general ramifications. 

This simple relationship has tremendous practical value. It can be used to 

estimate the fracture stress. It is not necessary to have a-priori information 

about the how the part was loaded. The smaller the mirror, the larger is the 

stress at the origin site. A small mirror is proof that the part was strong and had 

a small strength-limiting flaw. Conversely, large mirrors mean the failure 

stress was lower and implies a large defect. In some instances, a part may be so 

weak that the mirror size is larger than the part cross-section and hence the 

mirror markings are not visible. That, in and of itself, is valuable information. 

The existence of a mirror boundary implies that the part was stressed to a 

moderate or high level. Alternately, very strong parts (e.g., pristine optical 

fibers) may have such high strengths that mirrors are so small as to be not 

practically measurable with optical microscopes or the mirror may not even be 

found due to excessive fragmentation. 

Mist and hackle are sometimes triggered earlier than they would otherwise 

form by material flaws or microstructural features. Figures 5.5a-c shows 

examples. Strong Wallner lines, described later in this chapter, may also 

trigger mist formation in fracture mirrors such as shown in Figures 5.6. Figures 

5.2 and 5.3 shown earlier also hint at this possibility. Beauchamp has described 

cases like these where perturbations of the crack front can trigger early mist 

and hackle27 and long, persistent “hackle streaks.” 

Surface irregularities often cause a fracture mirror centered on a surface origin 

to depart from a semicircular shape and turn slightly inward towards an origin 

as shown in Figures 5.2 and 5.3. 

5-13 
 

 

 

 

 

 Fractography of Ceramics and Glasses 

dcp 

SC 

A 

(a) 

(b) 

(c) 

Figure 5.5 Early mist and hackle streaks associated with defects. (a) shows 

mist and hackle triggered by a surface crack at site “A” in a soda lime glass 

specimen at 65 MPa. The origin was well to the right, beyond the photo edge. 

(b) shows early mist and hackle triggered by a large bubble “B” in a 70 MPa 

strength specimen. The origin is a surface crack “SC.” (c) shows mist behind 

inclusions in obsidian. Gull wings (described later in this chapter) are also 

evident. (a and b courtesy R. Rice, c is courtesy A. Tsirk.) 

5-14 
 

 

 

 

 

 

Fracture Surface Examination 

(a) 

(b) 

Wallner lines 

origin 

Early mist 

Figure 5.6 Early mist and hackle. (a) and the enlargement (b) show an 

example of early formation of mist. The Wallner lines in (a) suggest the origin 

is off to the left in the image. The crack radiating outwards encounters other 

cracks in the surface shown by the arrows in (b). A strong primary Wallner 

line is generated and it triggers mist earlier than normal during mirror 

formation. 

5-15 
 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

(c) 

(a) (b) 

Figure 5.7 Mixed-mode loading creates fracture mirrors without mist but with 

a systematic array of “lance hackle.” (a) shows a normal mirror from pure 

opening mode I loading. (b) and (c) show mirrors from mixed opening/shear 

modes I/II, (c is courtesy J. Mecholsky). 

Gopalakrishnan and Mecholsky et al.28,29 showed that mist does not form 

around glass fracture mirrors if the loading has some in-plane shear (Mode II 

loading) acting on the crack at the same time it is under opening Mode I 

loading (see chapter 7). Hackle markings are very distinct in such cases as 

shown in Figure 5.7. The boundary is comprised of “lances” or “lance hackle” 

that are very similar to twist hackle. These steps on the fracture surface 

probably all go up (or go down) systematically as one goes around the crack 

front periphery, unlike the case of Mode I loading, where hackle steps are 

random or alternating up and down out of the fracture surface. The size of the 

mirror is not altered by the mixed-mode loading. The mirror/mist boundary 

for pure Mode I loading and the onset of lances in mixed Mode I/II loadings are 

comparable in size. The practical significance of this difference in appearance 

is that if a glass fracture mirror is comprised of lances with no mist, then the 

loading was not pure tension, Mode I, but mixed mode with some shear. 

Sommer30 has shown similar fracture surfaces with lance hackle and no mist in 

glass rods loaded in combined tension and torsion (i.e., mixed Mode I and 

Mode III loading). 

5.2.2 Fracture mirrors in ceramics 

The microstructure and polycrystalline nature of most ceramics obscures most 

of the fine details of fracture mirrors. Figures 5.8 - 5.10 show some variants of 

mirrors in ceramics. Mist is usually not recognizable. Sometimes coarse 

microstructural elements or flaws in a mirror may trigger early hackle lines 

within the mirror, prior to the onset of the generalized velocity hackle that 

5-16 
 

 

 

 

 

 

 

 

 

 

Fracture Surface Examination 

forms the mirror boundary. The mode of crack propagation, whether it is transor intergranular, also affects the mirror markings. Transgranular fracture is a 

mode whereby the crack propagates directly through grains on cleavage planes, 

whereas intergranular fracture is a mode whereby the crack takes paths along 

grain boundaries between grains. These modes will be discussed later in 

section 5.8. Mirrors may be easier to see in materials that start out with 

transgranular fracture, but then develop increasing fractions of intergranular 

fracture. On the other hand, it may be more difficult to define mirrors when the 

fracture mode is intergranular from the origin right through the entire mirror. 

For these reasons it is appropriate to add the qualifier “relatively” as in 

relatively smooth when describing the mirror region since there is an inherent 

roughness from the microstructure even in the area immediately surrounding 

the origin. Rice observed that mirrors are more detectable in strong specimens 

and are less clear in weak specimens.31 Mirrors may not be detectable in coarse 

grained materials or those with more than 10 % porosity. Mist can only be 

detected in fine-grained ceramics or those that fracture transgranularly.32 

(a) (b) 

(c) 

Figure 5.8 Fracture mirrors in ceramics. (a) is a SEM image of a strong 

bend bar (1024 MPa) of Y-tetragonal zirconia polycrystal (Y-TZP) for dental 

crowns. (courtesy J. Quinn) (b) is a stereoptical microscope of another Y-TZP 

bend bar (486 MPa) for an engine application. (c) is a stereoptical image of a 

silicon nitride bend bar (487 MPa). 

5-17 
 

 

 

 

 Fractography of Ceramics and Glasses 

(a) (b) 

(c) 

Figure 5.9 Fracture mirrors in ceramics. (a) is a hot-pressed silicon carbide 

tensile specimen (371 MPa). (b) is a silicon carbide tension specimen and (c) 

is a silicon nitride rod broken in flexure. Many beams and rods broken in 

flexure have internal flaw origins if the specimens are ground carefully. 

5-18 
 

 

 

Fracture Surface Examination 

(a) 

(b) 

Figure 5.10 Examples of incomplete ceramic mirrors due to bending stress 

fields. (a) shows an optical image of a gold-coated biaxial Y-TZP zirconia 

disk. The arrows mark the origin and the mirror. (b) is an alumina bend bar 

with low-angle grazing illumination to reveal the mirror in this stereoptical 

microscope image of this translucent material. The small white arrows mark 

the origin (a grinding crack) and the large black arrows mark the mirror. The 

small black arrows mark one of many coarse, arced Wallner lines. 

5-19 
 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

5.2.3 Fracture mirrors, special cases 

Multiple fracture mirrors 

In many cases, only one origin is activated, and there is only one fracture 

mirror. The crack propagates and reaches terminal velocity after only a short 

extension and in a period of microseconds. In rare instances, two virtually 

identical flaws may be triggered simultaneously in a laboratory strength test 

such as shown in Figure 5.11. 

Sometimes there will be multiple simultaneously-activated origins in a 

component or test specimen due to severe loadings such as laser heating or 

thermal shock by water quenching.33 Another example is a side-impacted 

bottle as shown in Figure 4.21f whereby the impact causes the sidewalls of the 

vessels to flex outwards. The bending stresses activate “hinge fractures” to 

either side of the impact site. When there are multiple origins, each generates a 

crack system and the fractographer must deduce the sequence of events that 

causes them. Sometimes, once one origin starts a crack system, the component 

stresses are redistributed and a second origin is then activated. The burst gas 

turbine rotor Case 1 in chapter 10 is an example. 

Secondary fractures can occur due to a variety of causes. Elastic wave 

reverberations after an initial fracture can cause secondary fractures. These 

complicate the interpretation of high strength four-point bend specimens and it 

can be difficult to determine which one was the primary fracture as shown in 

Figure 4.10. An observation that a fracture mirror is on what once was the 

compressively-loaded surface in a bend bar is proof the break was secondary. 

If there are two fracture planes with mirrors on the tensile side in the inner gage 

section, the larger mirror may be the primary break and the smaller mirror may 

be secondary due to stress reverberations that momentarily created a greater 

local stress. 

Concentric mirrors have been reported by Ramulu et al.34 and correspond to a 

crack front that is alternately decelerating and accelerating. Such markings 

also occur in high-strength tensile specimens when a crack branches repeatedly. 

At each branch, the crack slows down somewhat, then accelerates back to near 

terminal velocity and branches again. 

Steps or jogs in a mirror 

In some cases there may be a pronounced step or jog starting at the origin. This 

usually occurs when the critical flaw is irregular, three-dimensional, or inclined 

at an angle to the principal stress. Crack extension from the irregular flaw may 

occur on slightly-offset, parallel planes and thereby create a step or jog which 

5-20 
 

 

 

Fracture Surface Examination 

(a) 

(b) 

Figure 5.11 Multiple mirrors from grinding damage in a silicon nitride bend 

bar. (b) is a close-up of (a). Grinding damage caused fracture in these 

examples and the origins in each mirror were nearly identical. (Ref. 35) 

5-21 
 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) 

(b) 

Mirror 

Polished surface 

Figure 5.12 Bent mirror in a BK-7 borosilicate crown glass ring-on-ring disk 

strength specimen. The origin was a Vickers indentation flaw (large white 

arrows) with multiple radial and median cracks. (a) and the close-up (b) show 

the fracture surface with the specimen tilted back to show the outer polished 

surface where the indentation (white arrows) started fracture. Fracture started 

from Vickers indentation radial cracks oriented in different directions. A 

hackle line extends from the origin straight up through the mirror. Notice how 

main crack branching occurs (small white arrows) soon after the mirror 

boundaries formed in this equibiaxial-loaded specimen. Notice also that the 

branching angle initially is small then flares out to its final configuration. 

5-22 
 

 

 

 

 

 

Fracture Surface Examination 

may persist up into the mirror or even further. These steps are one type of 

“hackle” as described later in section 5.3. They usually point back to the origin 

and are an additional aid to locating the origin flaw. 

Curved or bent mirrors 

Cracks propagate perpendicular to the maximum principal tension stress, but 

some parts may have maximum stresses that are invariant with orientation. 

Equibiaxial stressed ring-on-ring test coupons (Figure 4.2) are an example. 

Curved or wavy fracture mirrors may occur in equibiaxial-loading cases. 

Figure 5.12 shows an example of a bent mirror. 

Non circular mirrors 

Mirror shape distortions may reveal important information about flaws or the 

stress distributions in a part. Mirrors that are not circular or semicircular 

indicate stress gradients, residual stresses, or elongated flaws at the origin. 

Parts with mirrors centered on an interior origin in part are usually circular. 

Oblong or elliptical mirrors in the interior are less common. Figure 5.9 shows 

three internal fracture origins and mirrors in ceramic test specimens. In one 

case (Figure 5.9c), the part was tested in flexure, but the mirror was relatively 

small and showed no obvious distortion or elongation. 

On the other hand, parts with surface origins often have flared or elongated 

mirrors, especially if the part was stressed in flexure. The bending loading 

mode usually creates a stress gradient with a maximum tension at the surface, 

zero stress at the middle (“the neutral axis”), and compression on the opposite 

surface as shown in Figure 4.18. As the crack accelerates from the surface 

origin, it traverses through a diminishing stress field towards the interior. This 

slows down the crack and elongates the mirror shape towards the interior as 

shown in Figures 5.10, 5.13, and 5.14. It can even eliminate the mirror 

markings altogether if the mirror is large relative to the specimen cross-section 

size. The stress gradient does not affect the markings along the tensile surface. 

There often is a small cusp on either side of the mirror pointing back to the 

origin in glasses. The cusp is accounted for by fracture mechanics 

considerations. (A crack that connects to a free surface experiences a slightly 

greater stress intensity factor right at the surface.) The small cusp is 

inconsequential for most analysis, except when fracture mirror sizes are 

measured as discussed in chapter 7 and Appendix D. 

5-23 
 

 

 

 

 Fractography of Ceramics and Glasses 

(a) (b) (c) 

Figure 5.13 Mirror shapes around surface origins. (a) is common for tension 

specimens, or flexure specimens or plates in bending if the mirror is small 

relative to the thickness. (b) and (c) show mirrors that are incomplete into the 

interior. This occurs when the mirror is large in a part loaded in bending. 

Figure 5.14 Fracture mirrors in four glass rods broken in flexure. The 

maximum tensile stress was on the bottom in each case. Fracture mirrors on 

the rod sides had pronounced flair or were incomplete due to the stress 

gradient. 

5-24 
 

 

 

 

 

 

 

 

 

 

 

 

Fracture Surface Examination 

The elongations in Figures 5.13 and 5.14 are common. Flexural stress 

gradients always cause elongations towards the interior. On the other hand, 

there are several important cases whereby elongations along the outer surface 

of the part occur as shown in Figure 5.15. Elongated flaws may generate 

elongated mirrors and sometimes even mirrors with side protuberances. Figure 

5.16 shows two possible scenarios for the formation of such elongated mirrors. 

Machining/grinding cracks and polishing scratches cause such markings.35,36 

Figures 5.17 and 5.18 show examples. 

(a) (b) (c) 

Figure 5.15 Mirror elongations along the surface can occur if the origin is a 

scratch or a long parallel grinding crack. (a) shows a single elongated crack 

at the origin, (b) shows a grinding zipper crack (see section 6.7.6), and (c) 

shows a zipper crack that generated lobes on the sides of the mirror. 

(a) (b) 

Figure 5.16 Two scenarios that could account for mirror elongations along 

the surface from long flaws. (a) shows how a crack could unzip or propagate 

quickly along the surface. (b) shows that criticality could be reached at several 

points on the flaw periphery and fracture commences accordingly. 

5-25 
 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Figure 5.17 Elongated mirror in a fully-annealed BK-7 borosilicate crown 

glass ring-on-ring disk strength specimen tested in inert conditions (dry 

nitrogen). (146 MPa, 21.2 ksi). The origin was a 7.3 μm deep long crack 

presumably from a scratch in the polished surface although all surface traces 

of the offending scratch had been eliminated by the final polishing. An SEM 

image of the origin flaw is shown in Figures 6.28 a,b and 7.22 a,b,c. In this 

optical photo the initial flaw appears as a thin white band on the bottom. 

Notice the slightly deeper penetration of the flaw in the middle of the mirror 

suggesting that it grew to criticality here first. The crack then fanned out to 

form the elongated non-circular mirror, but also unzipped along the surface to 

the left and right. The shadow bands extending up from the flaw are Wallner 

lines generated by minor flaw perturbations extending up into the mirror. 

Differential interference contrast image from a reflected-light compound 

microscope. 

5-26 
 

 

Fracture Surface Examination 

(a) 

(b) 

Figure 5.18 Sideways-elongated mirrors from grinding flaws in ceramic bend 

bars. (a) “zipper machining crack” (section 6.7.6) in sintered reaction bonded 

silicon nitride. (542 MPa) (b) zipper crack in sintered silicon nitride. (432 

MPa) 

5-27 
 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Residual surface stresses may also alter mirror shapes. The shape distortion 

depends upon the magnitude, sign, and depth of the residual stresses. Figure 

5.19 shows a typical parabolic residual stress distribution in glass from thermal 

tempering.37 Stresses are strongly compressive at the surface, decrease to zero 

about 21% of the thickness in from each surface (assuming the two surfaces are 

cooled at the same rates), and change to tensile in the interior. The 

compressive stress at the surface is twice the magnitude of the interior tensile 

stress. Figure 5.20 shows mirrors in a heat-strengthened glass consumerwares. 

For a surface origin, the compressive surface residual stresses suppress mirror 

boundary markings along the outer surface. Compensating residual internal 

tensile stresses shorten the mirror dimensions into the interior. A more detailed 

discussion of discussion of mirror shapes in tempered plates is in Appendix D 

and Figure D.8 in particular. Diced pieces from fully-tempered glass show mist 

and velocity hackle in the interior in Figure 5.21. Stress 

0 

σt, max 

σc, max Compression Tension 

side middle side 

-0.5 -0.4 -0.3 -0.2 -0.1 0 0.1 0.2 0.3 0.4 0.5 

Distance Through the Plate 

Figure 5.19 Residual stresses in thermally-tempered glass pieces. 

5-28 
 

 

 

Fracture Surface Examination 

(a) 

(b) 

Figure 5.20 Residual stresses alter the mirror shape. These images show 

thermally-strengthened glass consumerwares. The piece in (a) shows a 

complete suppression of mirror boundary markings along the lower right 

surface, but they are present to the left of the origin. This indicates the piece 

was thermally-strengthened with emphasis to protect the rim. The mirror is 

well-developed into the depth since residual tensile stresses there help form the 

mist and hackle. (b) shows a piece with mirror markings into the interior, but 

not along the bottom surface. The crack branched from the mirror sides in the 

interior of the piece then hooked back around to sever the fragment as shown 

by the arrows. The origin is one or a series of chatter cracks (black arrows) 

from contact damage on the outer surface. 

5-29 
 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) 

(b) 

Figure 5.21 Mist and velocity hackle in tempered glass. These photos are of 

the same auto window pieces in Figures 4.18b and d. (a) shows the crack 

reached terminal velocity in the interior where the tensile stresses were 

greatest. The black arrow shows the direction of crack propagation and the 

onset of mist. Each diced piece usually has this pattern since cracks slow at a 

branch then accelerate to terminal velocity again. Crack propagation was 

retarded at the outer surfaces due to the compression. Final breakage through 

to the surface has vertical twist hackle (white dashed ring). The arc shaped 

lines are secondary Wallner lines described later in this chapter. (b) shows the 

exact single impact origin (arrow) and mirror for the window. 

5-30 
 

 

 

 

 

 

 

 

 

Fracture Surface Examination 

5.3 Hackle 

In addition to the mist hackle and velocity hackle (described on page 5-12 in 

the mirror section), that are self-initiated by the crack at terminal velocity, there 

are other types of hackle that are very useful for ascertaining the direction of 

crack propagation. These markings enable the fractographer to trace crack 

propagation back to an origin. Hackle lines are sometimes called “lances” and 

less commonly “striations,” but this latter term is not recommended. Hackle 

can be generated from irregularities in the origin flaw. A good example is 

“grinding crack hackle” shown in section 6.7.6 in the chapter on origins. 

5.3.1 Coarse microstructural hackle 

Coarse microstructural hackle: Large broad hackle lines that form from nonspecific sources, but most likely variations in the microstructure. 

Ceramics often develop coarse hackle even in the absence of mirrors or other 

well-defined fracture surface markings. They occur in low-strength or porous 

ceramics and often are the only marking on a fracture surface that indicate the 

direction of crack propagation. The microstructure or some geometric 

irregularity causes portions of the crack to advance on non-coplanar regions 

separated by rounded ridges as opposed to sharp steps. Figures 5.22 and 5.23 

show examples. “Fine microstructural hackle” is described on page 5-36. 

Figure 5.22 Coarse microstructural hackle. Matching fracture surface halves 

of a split cylinder flexural strength specimen of a coarse-grained (50 μm to 80 

μm) Mg partially-stabilized zirconia. It has a very rough clumpy fracture 

surface. Low-angle grazing illumination shows the coarse hackle lines (black 

arrows) radiating from the origin region (white arrow). 

5-31 
 

 

 Fractography of Ceramics and Glasses 

(a) 

(b) 

Figure 5.23 Coarse microstructural hackle. (a) is a broken O-ring specimen 

of a ceramic membrane material. (b) is a silicon nitride bend bar with a Knoop 

indentation flaw on the bottom middle (black arrow). Coarse microstructural 

hackle lines (white arrows) guide the observer back to the origins. The only 

other fractographic feature observable is the compression curl in (b). Both 

specimens are so weak that fracture mirrors are larger than the cross sections. 

5-32 
 

 

 

 

 

 

 

 

Fracture Surface Examination 

5.3.2 Wake hackle 

Wake hackle: A hackle mark extending from a singularity at the crack front in 

the direction of cracking. 

When an advancing crack encounters an elastic singularity such as an inclusion 

or a pore, the crack front may split at the object and sweep past it on both sides. 

The two fronts often pass the obstacle on slightly different planes and create a 

step or “tail” between them. The tail may fade away quickly or persist for long 

distances. These markings are very recognizable and may either be large or 

very small. They serve as miniature “weather vanes” which tell the direction 

of local crack propagation. Figures 5.24 – 5.26 show examples. 

Wake hackle in glazes or veneers often is the only recognizable fracture surface 

feature in kitchenware, electrical insulators, and some dental ceramics. The 

core ceramic materials in these instances are often coarse-grained, porous, or 

weak, and few fracture surface markings are evident. Sometimes in even the 

most difficult materials, tiny porosity will aid the fractographer by serving as a 

source of wake hackle. A superb example was shown by Ko, reference 38. 

dcp 

Figure 5.24 Wake hackle, the fractographers “weather vane” for crack 

propagation. This shows wake hackle from bubbles in glass. The direction of 

crack propagation (dcp) is shown by the arrow. Many bubbles are beneath the 

surface in this view. 

5-33 
 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) 

(b) 

dcp 

Wake hackle 

Figure 5.25 Wake hackle. (a) shows an array of wake hackle from pores in 

an Empress II dental crown veneer. The direction of crack propagation is 

marked by the arrow. Notice that some hackle is triggered even though the 

crack did not cleave the pore (small white arrow). (b) shows an example of 

transgranular fracture from a large grain origin in polycrystalline AlON. 

Wake hackle was triggered by grain boundaries and tiny pores (small arrows). 

5-34 
 

 

 

 

 

Fracture Surface Examination 

(a) 

(b) (c) 

dcp 

dcp 

Figure 5.26 Wake hackle (a) shows wake hackle from inclusions in a lead 

zirconium titanate bend bar. They point back to the origin on the bottom. The 

tensile surface is on the bottom and a compression curl is on the top. The 

origin is a semielliptical surface crack (arrow) generated by a Knoop 

indentation. (b) and (c) show parabolic double wake hackle tails triggered by 

inclusions in obsidian glass. (b,c courtesy A. Tsirk) 

5-35 
 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Irregularities such as pores or inclusions are not the only sources of wake 

hackle in ceramics. Grain boundaries and local crack path redirection as a 

crack propagates from one grain to another with a different crystallographic 

orientation, or local-density or other microstructural variations all can trigger 

“fine microstructural hackle:” 

Fine microstructural hackle: Small hackle lines of the order of size of the 

microstructure that occur as a crack interacts with the local microstructure. 

Sometimes meticulous observations of within-grain cleavage, micro-wake 

hackle, and micro-twist hackle lines can allow one to backtrack to an origin. 

Kerkhof and Sommer used ultrasonic fractography to show that a crack 

accelerates in the immediate vicinity of a hole in glass, but then is retarded as 

the crack runs past the hole.39 Figure 5.45d later in this chapter shows an 

example. 

5.3.3 Twist hackle 

Twist hackle: Hackle that separates portions of the crack surface, each of 

which has rotated from the original crack plane in response to a lateral 

rotation or twist in the axis of principal tension. (Fréchette, Ref 1.) 

Twist hackle are very telltale markings as illustrated in Figures 5.27 and 5.28. 

They are also known as delta patterns, river deltas, river patterns, lances, and 

sometimes even in the geological community as striations. (Striations in this 

instance are not to be confused with fatigue striations, a well-known metalurgical fracture surface marking). The roughly parallel segments point in the 

direction of local crack propagation. Twist hackle can be generated by the 

primary crack as it travels directly through the body, especially as it goes 

around corners or geometric irregularities, or when stress conditions changed. 

Twist hackle can also occur from final breakage between local crack segments, 

or through to the side of a specimen that had had strong thermal or stress 

gradients. Twist hackle shows the local direction of crack propagation. Plates 

in bending often have curved cracks that run quickly on the tension side, but do 

not quite break through to the opposite surface. The crack later snaps through 

to the opposite surface leaving twist hackle markings as shown in Figure 5.29a. 

Double torsion fracture mechanics specimens often have such markings. 

Figures 5.21a and 28b show a similar circumstance whereby final breakthrough 

of tempered glass has twist hackle near both outer surfaces. 

5-36 
 

 

 

 

 

 

 

Fracture Surface Examination 

(a) 

Crack 

propagation 

direction 

End on view 

(b) ( c) 

Figure 5.27 Twist hackle. (a) shows how twist hackle is formed. A crack 

propagating from left to right runs perpendicular to the initial stress field as 

shown on the left. The axis of principal stress then tilts as shown on the top 

middle. The crack is unable to rotate all at once in response to the new stress 

direction, so it breaks into small, unconnected segments. The steps are “hackle 

lines” or “lances.” Lateral breakthroughs occur between the micro segments 

as shown in (b) and (c). Note in (c) how the micro segments can link either by a 

top segment breaking though to a bottom segment or vice versa. On any given 

lance, the link up can alternate between either segment creating a barbed 

appearance. Sometimes an incomplete segment can breakthrough long after the 

main fracture has occurred, generating faint tinkling sounds and creating very 

sharp needle like fragments that fall free from the fracture surface. The small 

hackle lines gradually merge into coarser hackle lines, and so on, until the new 

crack plane is aligned perpendicular to the new tensile stress axis. Usually 

crack propagation is in the direction of fine to coarse hackle. 

5-37 
 

 

 Fractography of Ceramics and Glasses 

(a) 

(b) 

dcp 

Figure 5.28 Twist hackle. (a) and (b) show close-ups of twist hackle in glass. 

The crack was running in the direction of the arrow in (a). (b courtesy J. 

Varner) 

5-38 
 

 

 

 

 

Fracture Surface Examination 

(a) 

(b) 

dcp 

Figure 5.29 Twist hackle. (a) shows a glass microscope slide that was broken 

in bending with a little superimposed torsion. The main crack fracture from 

left to right, leading on the bottom surface as attested to by the profusion of 

mist and secondary Wallner lines and an occasional primary Wallner line 

(described in the next section). Final break-through to the top surface (white 

arrow) created the twist hackle along the top. (b) shows a similar pattern but 

with a tempered glass fragment. . 

5-39 
 

 

dcp

 Fractography of Ceramics and Glasses 

(a) 

(b) 

side 

groove 

side 

groove 

side 

groove 

Figure 5.30 Twist hackle in grains. (a) and (b) show fracture surfaces of 

chemical vapor deposition (CVD) silicon carbide micro-tensile specimens. 

CVD materials often have coarse, columnar grains. Such grains run from top 

to bottom in this view. Twist hackle in these grains in both images show that 

fracture propagated from left to right as shown by the arrows. Fracture 

occurred from cracks located at the bottom of the grooves on the left side in 

both cases. 

5-40 
 

 

 

 

 

 

 

Fracture Surface Examination 

A 

Figure 5.31 Twist hackle around the circumference of a 15 mm diameter 

glass rod broken in tension (induced by lateral fluid pressure compression 

around the periphery while the rod ends were free) with some superimposed 

torsion. The wavy lines are tertiary Wallner lines (described in the next 

section) created by ultrasonic fractography. Fracture started at the bottom at 

point A and progressed into the rod interior. Final breakthrough to the outer 

surface generated the twist hackle. (Sommer, ref. 30, Kerkhof and Sommer, ref. 

39, and Richter and Kerkhof, ref. 2.) 

In polycrystalline ceramics, twist hackle markings within coarse grains can 

serve as very helpful markers. They are local indications of crack propagation 

and it is often possible to track a propagating crack back to an origin by reading 

the local “weather vanes.” Figure 5.30 shows examples. The fractographer 

may be surprised to see that twist hackle line within separate grains may not all 

line up the same way. It is quite common for cracks to take detours through the 

microstructure in response to local conditions. 

Sommer30 analyzed the formation conditions of twist hackle and concluded that 

in glasses a stress redirection of 3.3o is needed to initiate the twist hackle. 

Beauchamp27 observed hackle with only 1o of twist. Torsional stresses 

superimposed on tension stresses can create twist hackle markings as shown in 

Figure 5.31 which shows a rod loaded in mixed mode I, opening mode, and 

mode III, out-of-plane shear, as described in section 7.1. An interesting paper40 

reported on how a continuum phase-field approach could be used to model 

twist hackle formation and propagation. Very small initial perturbations along 

the initial crack front would form parallel “daughter crack segments” that 

gradually coalesce, as the main crack plane rotated. 

5-41 
 

 

 

 

 

 Fractography of Ceramics and Glasses 

5.3.4 Shear hackle 

Shear hackle: A particular form of twist hackle that occurs in the lazy loops 

generated in the latter stages of fracture of a hollow specimen.
 

(Fréchette, Ref 1.)
 

This definition may seem peculiar, but once the fractographer has seen an 

example such as Figure 5.32, it becomes clear what Fréchette had in mind.1 He 

said: “Far from the fracture origin, it is common to see the crack turn through 

90o and develop a cupped surface, inclined at 45o to the free surface, from 

whose centerline a spray of twist hackle emerges. It has little significance in 

reconstructing a failure event in industry, but it is useful in problems of rock 

fractures.” 

dcp 

Figure 5.32 Shear hackle in a glass plate. The crack was perpendicular to 

the glass surface and was slowing as it moved in the direction of the white 

arrow. It then turned 90o parallel to the surface and finally broke through to 

the surface at about 45o

, producing the spectacular fan-shaped shear hackle. 

Shear hackle is a variant of twist hackle. 

5-42 
 

 

 

 

 

 

 

 

 

 

 

Fracture Surface Examination 

5.3.5 Corner hackle 

Corner hackle: A fan like array of hackle lines created when a crack curves in 

a plate or goes around a corner of a component. 

Figure 5.33 shows two examples. The fan like array spreads outwards in the 

direction of crack propagation. 

(a) (b) 

Figure 5.33 Two types of corner hackle. (a) shows a fan like pattern that 

occurs when a crack that is leading on the bottom surface turns to another 

direction. (b) shows a fan-like pattern that occurs when a crack goes around a 

corner of a container or vessel. 

5.3.6 Step hackle 

Step hackle: A type of twist hackle in the form of a single line that is an arc 

shaped line on a fracture surface of a plate broken in bending and twisting. 

Step hackle is a line that is a step on the fracture surface such as shown in 

Figure 5.34. The line is parallel to the local direction of crack propagation. 

Unlike most twist hackle, it is an isolated line. The arc curvature of a step 

hackle line is opposite to the curvature of arrest lines or Wallner lines that may 

exist on the same fracture surface as shown in Figures 5.34b and c. The 

fracture surface is usually tilted or curved and not perpendicular to the plate 

outer surfaces (end-on view Figure 5.34a.). The step can initiate by itself if the 

tilt is strong enough, or may be triggered by other features such as bubbles, 

inclusions or mist. Step hackle is shown here for glasses, but can occur as 

“cleavage step hackle” as described in Chapter 8 on single crystals. 

5-43 
 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) 

fracture surface view end on view 

twist hackle 

tension side 

compression side 

step hackle line dcp 

dcp 

(b) 

(c) 

Figure 5.34 Step hackle schematic (a). (b) and (c) show a piece of a large 

(1.5 m by 0.75 m) glass mirror that fell off a wall in the author’s home. The 

mirror broke into over 100 pieces. Many fragments had evidence of bending 

and twisting. Notice the opposite arcs of the step hackle line and the Wallner 

lines (thin dashed lines in a). 

step hackle line Wallner 

lines 

mist and velocity hackle 

twist hackle 

twist hackle on the compression side 

Step hackle line 

dcp 

Wallner 

lines 

5-44 
 

 

 

 

 

 

 

 

 

 

 

 

 

Fracture Surface Examination 

5.4 Wallner Lines 

5.4.1 Introduction 

Wallner line: A rib shaped mark with a wavelike contour caused by a 

temporary excursion of the crack front out of plane in response to a tilt in the 

axis of principal tension. It may also form from passage of the crack front 

through a region with a locally shifted stress field, as at an inclusion, pore, or 

surface discontinuity. (adapted from Fréchette, Ref. 1) 

Also known as “ripples” or “rib marks,” Wallner lines were named for Helmut 

Wallner who first explained how they formed in 1939.41 He showed beautiful 

images (similar to Figures 5.2 and 5.3 of this Guide) of fracture mirrors in 

flexurally tested glass rods with the lines within the mirrors. He presented a 

splendid sketch showing how an expanding crack front was overtaken by an 

expanding elastic wave. The Wallner line is the locus of the intersection of 

the elastic wave and the crack front. The elastic wave momentarily causes the 

crack to ripple out of plane like a wave on a pond surface. The shape of the 

line depends upon the respective shapes and speeds of the elastic wave and the 

crack front, as well as the direction the wave approaches the crack front. 

Wallner lines are invaluable in determining the direction of crack propagation. 

As will be shown, they usually do not have the exact same shape as the crack 

front, but they are usually curved (bowed) in the direction of crack propagation. 

They are best understood by the examples that follow. 

Wallner lines are very shallow hillocks on a fracture surface. They look like a 

thin shadow band and they tend to shift slightly as illumination is adjusted. 

This is in marked contrast to crack arrest lines that are sharp and do not appear 

to move as light is adjusted. Wallner lines are very easy to see on fracture 

surfaces once the lighting is adjusted, but they are very difficult to see on the 

scanning electron microscope, since they produce almost no contrast and are 

very shallow. 

There are several types of Wallner lines depending upon the source of the 

elastic wave. Fréchette1 thoughtfully categorized these as primary, secondary, 

and tertiary. This Guide follows his scheme and adopts his definitions. His 

figures showing the evolution of the loci are instructive, but have been revised 

and expanded into separate illustrations in this Guide to make them easier for 

the beginner to understand. Keep in mind that a crack front through a plate 

may be straight or have different shapes according the stress state as shown in 

Figure 5.35. 

5-45 
 

 

 Fractography of Ceramics and Glasses 

(a) 

(b) 

(c) 

(d) 

Figure 5.35 Crack front variations due to stress distributions through the 

thickness. (a) shows a crack front in a uniform tensile loaded plate. (b) a 

plate with greater tensile stress on the bottom. (c) a plate in bending such that 

the bottom is in tension and the top in compression. (d) a tempered plate with 

tensile stresses in the interior and compressive stresses at the surface. 

5-46 
 

 

 

 

 

 

Fracture Surface Examination 

5.4.2 Primary Wallner lines 

Primary Wallner line: A Wallner line formed by an elastic pulse generated by 

some portion of the crack front with a singularity in the specimen such as a 

discontinuity at the free surface or within the specimen, or with any localized 

stress field or elastic discontinuity. (Adapted from Fréchette, ref. 1) 

Primary Wallner lines commonly are generated by surface blemishes and 

irregularities such as scratches, pits, or edge chips. Figure 5.36 shows an 

example and Figure 5.37 illustrates how they are formed. If a Wallner line is 

seen to emerge from such a surface blemish, it is evidence that the surface fault 

was present at the instant the crack ran by. Alternatively, if a large blemish or 

chip is on a glass fracture surface, but there is no Wallner line from it, the 

blemish occurred afterwards and was not initially on the piece. 

Unless the crack is moving very slowly (and the elastic wave traverses the 

entire crack front very quickly), primary Wallner lines do not show the exact 

crack front profile. Nevertheless, they are curved in the same direction and can 

help the fractographer deduce which way a curved crack was running. It is 

fairly easy to show that straight crack fronts generate straight Wallner lines, but 

they will tilt at a different angle than the original crack front depending upon 

the crack velocity. 

dcp 

Figure 5.36 Primary Wallner line in a glass slide broken in bending. A 

simple scratch on the surface (arrow) triggered the strong Wallner line as the 

crack ran from left to right. Other tiny faults triggered fainter primary Wallner 

lines. 

5-47 
 

 

 

 Fractography of Ceramics and Glasses 

front 

radiating 

elastic pulse is 

generated here 

Figure 5.37 Primary Wallner lines. This schematic series shows the 

progressive advance of the crack at successive time intervals. The elastic wave 

that forms the Wallner line (thick line) is generated when the advancing crack 

hits a surface irregularity “O”. The crack leads on the bottom surface, but is 

moving at constant velocity in this example. 

5-48 
 

 

 

 

 

 

 

Fracture Surface Examination 

dcp 

Figure 5.38 Wallner lines are more difficult to see in polycrystalline 

ceramics. Low-angle grazing illumination is effective. This is an optical image 

of a silicon nitride double torsion fracture toughness specimen. The crack was 

running in the direction of the white arrow. The double torsion configuration 

bends a specimen so that the bottom is in tension and the top in compression. 

The cracks lead on the bottom surface. Primary Wallner lines formed with arcs 

shown by the white dashed curve. See also Figure 5.10b. 

dcp 

(a) (b) 

Figure 5.39 Gull wings are a variant of primary Wallner lines. (b) is a closeup of (a). Most of these pores also have wake hackle as well. These images 

are from obsidian with crack propagation from bottom to top as shown by the 

arrows. The wing angles can be used to measure crack velocity (section 7.9). 

(specimen courtesy A. Tsirk) 

5-49 
 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Wallner lines are less easy to see in ceramics. Figures 5.10b and 5.38 show 

examples. In porous, granular, or coarse-grained ceramics, large broad ridges 

that are Wallner lines may be the only recognizable marking on the fracture 

surface. If so, the ridge may have been caused by microstructural irregularities 

or surface flaws on one side of the origin or to a geometric irregularity (e.g., an 

inside corner in a formed-ware, or a ridge on a plate). 

Figure 5.39 shows “gull wings,” a variant of primary Wallner lines. They are 

usually formed by a stress pulse created as the crack encounters pores, 

inclusions, or other irregularities. There even are instances where they may be 

triggered by secondary fracturing behind the primary crack front, and hence 

will be detached from obvious irregularities.42 

5.4.3 Secondary Wallner lines 

Secondary Wallner line: A Wallner line generated by an elastic pulse released 

by a discontinuity in the progress of the crack front, typically one of the rough 

details which arise as the crack approaches its effective terminal velocity. 

(After Fréchette, Ref 1.) 

Unlike the primary Wallner lines that form as a result of a crack encountering 

an external feature, secondary Wallner lines occur from features the crack 

generates itself. Mist and hackle are disturbances created by crack running at 

or near terminal velocity. The mist and hackle generate copious secondary 

Wallner lines. Figure 5.40 shows their formation sequence and Figure 5.41 

shows examples. It is very common for both primary and secondary Wallner 

lines to be present. They can be distinguished by their different curvatures. 

Secondary Wallner lines are more hooked (like a fish hook) than primary 

Wallner lines. As was the case with primary Wallner lines, secondary Wallner 

lines do not show the exact crack front profile. Even so, it is helpful to deduce 

from their curvature which way a crack was running. 

5-50 
 

 

 

 

 

 

Fracture Surface Examination 

crack front 

mist and velocity hackle 

Wallner line elastic wave 

t = 0 

t = 1 

t = 2 

t = 3 

Final 

Wallner line 

dcp 

radiating 

Figure 5.40 Secondary Wallner lines are formed by the crack itself. This 

schematic series shows the progressive advance of the crack and the elastic 

pulse that forms the secondary Wallner line (dark line) in a plate with a stress 

gradient. Notice the hook final shape, a consequence of the stress gradient 

caused by the variation in the crack velocity from a maximum on the bottom to 

close to zero at the top. 

5-51 
 

 

 

 

 Fractography of Ceramics and Glasses 

(a) 

(b) 

(c) 

dcp 

Figure 5.41 Examples of secondary Wallner lines in glass microscope slides. 

(a) and (b) show Wallner lines from the sides of a mirror. (b) is a close-up of 

the right side of (a). Mist and hackle triggered the secondary Wallner lines. 

There also are fainter primary Wallner lines inside the mirror. Notice their 

different arcs. (c) is a slide broken in bending. 

5-52 
 

 

 

 

Fracture Surface Examination 

5.4.4 Tertiary Wallner lines 

Tertiary Wallner lines: Wallner lines caused by elastic pulses generated from 

outside the crack front. (Fréchette, Ref. 1) 

Vibration or shock from a source other than the crack itself can create waves 

that intersect with and cause momentary deviations in the crack front. 

Depending upon the source of the elastic pulse relative to the crack plane, the 

interaction may generate Wallner lines that may be very close to the shape of 

the crack front, or it may produce an arc formed by a progression of 

intersections as in the examples shown above. For example, if the crack is 

moving relatively slowly, the fast elastic wave will overtake all portions of the 

crack front nearly simultaneously, producing a Wallner line “snapshot” of the 

crack front. 

If the pulse source is located at a site normal to the fracture plane, the pulse 

may reach all portions of the crack front simultaneously also producing a 

Figure 5.42 Tertiary Wallner lines in a large chunk of glass that was 

hammered into pieces. There are multiple impact sites marked by arrows, the 

most prominent of which is marked by the larger white arrow. Notice how the 

tertiary Wallner lines are concentric about the impact sites. This was one of 

Professor Fréchette’s demonstration pieces. 

5-53 
 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

“snapshot” of the crack front. The impacted glass shown in Figure 5.42 is a 

good example. Notice how the tertiary Wallner lines are concentric about the 

impact sites. The impact site not only generated the cracks, but the elastic 

pulses as well. (In contrast note how the primary and secondary Wallner lines 

in Figures 5.3 and 5.41a cross over each other on top of the origin.) Impact 

typically results in a ringing within the specimen that may be short lived, but 

will be long enough to endure for the period a crack traverses through a body. 

On the other hand, a sudden stress relief may create a brief vibration when a 

crack pops in or propagates. This may be accompanied by an audible snap, but 

only generates faint elastic pulses so that only a few tertiary Wallner lines form 

near the origin. Fréchette1 noted that window panes that are cracked by thermal 

stresses may have only a few tertiary Wallner lines near the origin and are 

featureless elsewhere. 

A schematic case of a pulse overtaking a crack from the rear is shown in Figure 

5.43. The reader is reminded that Wallner lines do not always create a 

“snapshot” of the crack front shape. Tertiary Wallner lines are not always 

formed by pulses arriving from a source so conveniently located as in test 

pieces studied with ultrasonic fractography. They may come from any 

direction including from sources that are out of the crack plane. 

5-54 
 

 

 

Fracture Surface Examination 

overtaking 

Figure 5.43 Schematic showing the stages of formation of a tertiary Wallner 

line for a curved crack front overtaken by an elastic wave from behind the 

crack. The Wallner line has a similar concavity as the crack front, but the loci 

are different. 

5-55 
 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Ultrasonic or stress wave fractography,

2,6 previously mentioned in section 3.21, 

deliberately creates elastic waves to make tertiary Wallner lines on the running 

crack. Figures 5.44 and 5.45 show some examples from Richter and Kerkhof’s 

work.

2 Local crack velocity can be simply calculated on the basis of the 

spacing between the lines and the frequency of the pulses. They were able to 

measure cracks running at terminal velocities in glass with 5 MHz lead 

zirconium titanate transducers. Their review article2 has many fascinating 

examples of the application of this method to study crack front profiles and 

velocities. 

(a) 

(b) 

Figure 5.44 Tertiary Wallner lines from ultrasonic fractography. A 

transducer created elastic pulses that made “snapshots” of the crack fronts as 

they propagated through glass test specimens. (a) is a notched tension 

specimen loaded to fracture. The crack started at “SC” and ran to the right 

forming mist at M, and branching at B. The crack reached a terminal velocity 

of 1540 m/s, two waves to the right of the letters SC, well before the mist 

formed. (b) shows an accelerating crack that grew from an indentation flaw in 

a bend bar. The ellipse shapes are in perfect accordance with fracture 

mechanics predictions as discussed in section 7.5.2. These two images show 

how these Wallner lines are wave-like undulations on the fracture surface, 

unlike the sharp arrest lines in Fig. 5.47c. (a + b courtesy, H. Richter) 

5-56 
 

 

 

Fracture Surface Examination 

(a) 

(b) 

100 μm 

dcp 

dcp 

500 μm 

b 

a 

dcp 

dcp 

Figure 5.45 Tertiary Wallner lines examples from ultrasonic fractography. 

(a) shows a surface crack in a bend bar that went unstable from one side when 

growing by slow crack growth in water. The bending stress field caused the 

crack to grow faster along the surface (arrow a) than into the depth (arrow b). 

The crack went unstable on the right side (yellow arrows) and then rapidly 

fanned out and around the still slowly-advancing main crack. See also Figure 

5.56. (d) shows a crack front approaching a capillary tube void at 60 m/s. The 

arrow shows the crack propagation direction. The crack locally accelerated to 

60 m/s as it approached the void, but then slowed down to 20 m/s as it passed 

around it, and then jumped at a faster velocity after passing around the void. 

(Ref. 2, photos courtesy of H. Richter) 

5-57 
 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

5.4.5 Wallner lambda lines 

Elastic waves can reflect off free surfaces creating a lambda shaped line on the 

fracture surface as shown in Figure 5.46. These are typically formed by 

primary or secondary Wallner lines. Crack velocity can be calculated from the 

spacing of the Wallner lines as discussed in section 7.9 and Figure 7.31. 

Propagating 

crack front 

a 

t 

Figure 5.46 Wallner lambda line mark. The schematic shows the advancing 

crack (thinner line) encountering a discontinuity at point O. An elastic pulse 

creates a radiating elastic wave. The intersection of the wave and the 

advancing crack front creates a primary Wallner line that runs upward to the 

top surface. When the elastic wave reflects from the opposite surface, the 

Wallner line continues back downward. 

5.5 Arrest Lines 

Arrest Line: A sharp line on the fracture surface defining the crack front 

shape of an arrested, or momentarily-hesitated crack. Resumed crack 

propagation occurs under a more or less altered stress configuration. 

This is an adaptation of Fréchette’s original definition.1 His definition included 

the words “rib mark” as in “rib shaped,” but this author has seen straight arrest 

lines in thermally-stressed plates, so a more general definition seems in order. 

The words “momentarily-hesitated” are also added to Fréchette’s definition 

since it is not necessary for the crack to come to a complete halt. Fréchette also 

explained this in his book. The arrest line occurs because the crack 

repropagation is on a slightly different plane, usually in response to a slight 

change in the axis of principle tension. 

5-58 
 

 

 

 

 

 

Fracture Surface Examination 

Arrest lines are different than Wallner lines in two key respects: arrest lines are 

sharp, and, unlike most Wallner profiles, arrest lines show the crack front 

profile at an instant in time. Sometimes, if the crack repropagation occurs in a 

slightly different direction, twist hackle lines will form that lead away from the 

arrest line. 

Michalske showed that cracks that arrested after slow crack growth sometimes 

require extra loading to repropagate after aging.43 The crack tip blunts slightly 

during the aging, and the crack has to resharpen before repropagation. This 

process occurs with slight variations along the crack front so that when the 

crack repropagates there are slight humps and valleys along the crack front and 

tiny hackle lines leading away from the arrest line. 

Concentric arrest lines in ceramics and glasses sometimes bear a resemblance 

to classical mechanical-fatigue markings in metals, but they are not the same. 

Classical mechanical fatigue in metals is iterative crack extension by cyclic 

loading. Repeated loading on a crack causes damage nucleation, 

accumulation, and stepwise growth in metals. This leaves periodic, spaced 

striation lines on the fracture surface. This mechanism is not operative in most 

fine-grained equiaxed-microstructure ceramics and glasses. Section 5.9.5 

below has more on mechanical fatigue. 

5-59 
 

 

 

 

 Fractography of Ceramics and Glasses 

(a) 

(b) 

(c) 

arrest line 

Wallner lines 

Figure 5.47 Crack arrest line examples. (a) shows a glass microscope slide 

broken in bending and the arrest line is on the right (arrow). The lines to the 

left are primary Wallner lines. Wallner lines are not as sharp as arrest lines. 

(b) shows an arrest line in a glass slide that broke from thermal stresses. (c) 

shows concentric arrest lines from repeated cyclic loading of a glass plate in 

bending. Compare the sharp lines in this photo to the wavy tertiary Wallner 

lines in Figure 5.44b. (c is courtesy D. Green and is the same as Fig. 5.66). 

5-60 
 

 

 

 

 

 

 

 

 

 

 

 

Fracture Surface Examination 

5.6 Scarps 

Scarps shown in Figure 5.48 are subtle curved lines on a fracture surface 

caused by interaction of a propagating crack and a liquid or a reactive 

environment. The liquid can act to retard a crack front or can cause it to 

accelerate, sometimes in a discontinuous fashion. The scarps mark sudden 

transitions in crack velocities that can occur due to interactions of the crack 

with liquid. They usually correspond to transitions in the crack tip 

environment, including changing from wet to dry, dry to wet, humid air to dry 

air, etc. They even can occur in totally wet environments if the crack 

accelerates so fast that the viscous liquid cannot keep pace and instead cavitates 

behind the crack tip. Scarps are usually only detected in glasses and single 

crystals. Scarps could exist in polycrystalline ceramics, but the inherent 

roughness from the microstructure masks their presence. 

The practical significance of scarps is that they are evidence that fracture 

occurred in the presence of a fluid, most commonly water. They also can help 

confirm the local direction of crack propagation. 

Fréchette has identified several types of scarps and accounted for their 

formation.1 Scarps can vary from as simple as a single faint line across the 

fracture surface perpendicular to the direction of crack propagation (cavitation 

scarp), to a complex network of curves such as shown in Figure 5.48. Figure 

5.48b of a scarp in single crystal sapphire is, to the best of my knowledge, the 

only documented example of a scarp in a ceramic material. 

Fréchette called some of these “Sierra scarps” if they are pointed in the 

direction of crack propagation as shown here as Figure 5.48c (and also in 

Frechette’s book, Figures 2.28 and 2.29) due to their similarity to his favorite 

mountain range. They form when water (in the form of spaced fingerlets as in 

a Taylor instability) overtakes a decelerating dry crack, causing local advances 

forward by the crack front. 

“Cavitation scarps” occur when an accelerating crack causes void formation or 

cavitation in the water between the fracture surfaces that is trying to keep up 

with the crack tip. Once cavitation occurs the crack outruns the fluid water and 

a scarp line is left behind. The faint vertical lines within the scarps in Figure 

5.48d have been described as “subcavitation hackle” by Michalske and 

Fréchette, who describe it as a microhackle precursor to complete cavitation.1,44 

The faint lines have alternatively been described as “streaks” by Beauchamp, 

since they may not in fact be steps, but slight perturbations in the crack plane. 

5-61 
 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

(c) (d) 

dry 

wet 

(e) (f) 

dcp 

(a) (b) 

Figure 5.48 Examples of scarps. The arrows show the direction of crack 

propagation. (a) is a fractured glass medicinal vial. The liquid was the 

medicine within the vial. (b) is from a single crystal sapphire hemispherical 

dome broken by water quench thermal shock. This differential interference 

contrast optical photo shows the scarps have variability in height. (c) shows 

Sierra scarps in glass caused by water “fingerlets” overtaking the advancing 

crack (courtesy E. Beauchamp) (d) - (f) are obsidian flakes formed by 

knapping. (courtesy A. Tsirk) The faint vertical lines within the scarps in (d) 

are either “streaks” or “subcavitation hackle,” a precursor to complete 

cavitation. The parallel bands in (c), (e) and (f) are tertiary Wallner lines 

intentionally introduced by vibrations from an engraver (c), or a hand palm 

sander (e,f) to study the crack velocities. In (f) water that was available from 

the right side facilitated crack advance causing the scarp. 

5-62 
 

 

 

 

 

 

 

 

 

 

 

 

Fracture Surface Examination 

Scarps are very subtle and can be very difficult to see. They are easily 

overlooked. They are best seen in an optical microscope with the specimen 

tilted to create a mirror-like reflection from the fracture surface. The specimen 

or the lighting must be tilted back and forth until one catches just the right 

reflection. They usually exist in only a portion of the specimen and usually 

where the crack is moving at a slow to moderate speed. 

Tsirk45,46 has studied a variety of scarps in knapped obsidian arrowheads and 

knives formed by pressure flaking. Some of the scarps were generated in 

antiquity when they chipped during usage. Tsirk has postulated that the shape 

of scarps may depend upon the nature of the fluid and possibly the viscosity. 

Scarps that have been observed associated with blood are different from those 

formed with water or saliva.45,46 The reader is referred to Fréchette’s book1 and 

References 43-46,47,48 for details on scarps. 

5.7 Glue Chips 

Glue chipping is a process used by artists to create surface decorations on glass. 

Figure 5.49 illustrates the geometries. Figures 5.50 and 5.51 show examples. 

Glue is coated onto a lightly-abraded glass surface. The plate is warmed 

causing the glue to contract. Shallow cracks enter the glass at the edge of the 

glue at about 45 degrees then run parallel to the surface. The glue is peeled 

away removing scalloped chips in the glass. Fréchette noted that the chips 

formed by this artistic process closely resemble flaws generated in glasses in 

architectural and structural applications.49 

glue metal 

σ

Ceramic or glass Ceramic or glass 

(a) (b) 

Figure 5.49 Glue chipping cracks. (a) show a side view with glue applied to 

the surface of the glass of ceramic. Contraction of the glue causes a crack to 

enter the ceramic or glass at about 45o to the surface. (b) shows a comparable 

geometry with a metal bonded to a ceramic or glass. If the metal contracts 

more than the ceramic, edge cracks can be created. 

5-63 
 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) (b) 

(c)

 (d) 

glue chips on bottom surface 

fracture surface 

dcp 

twist hackle 

dcp 

bevel on edge of the glass disk 

flat outer disk surface 

Glue chip origin 

fracture surface 

Figure 5.50 Examples of glass glue chips. (a) shows an Alfred University 

demo plate covered with glue chips. (b) is a close-up of isolated glue chips on 

the surface. (c) shows a side view of a broken plate that shows glue chips on 

the bottom that affect crack propagation. (d) shows large glass disk that 

fractured from a glue chip caused by an epoxy bond to a metal part. 

5-64 
 

 

 

 

 

 

 

 

100 μm

Fracture Surface Examination 

Glass damage 

(a) (b) 

glass glass alumina 

substrate 

100 μm 100 μm 

(c) (d) 

Figure 5.51 Glue chipping damage in a thick-film resistor that cracked during 

thermal cycling. (a) shows the component in its epoxy encapsulant. (b) shows 

the exposed electrical circuitry. Scallop cracks (white-colored reflective 

patches) formed in an inner glass encapsulating layer. (c) and (d) show two 

glue chips in the glass. The arrows show the direction of crack propagation. 

(all photos and interpretation courtesy R. Tandon and S. Monroe) 

The conventional wisdom in many joining problems is to choose a metal or 

other material that contracts more than the ceramic or glass. During cool down 

from the elevated temperature joining process, the ceramic or glass is put into 

residual compression. This indeed is what happens, except right at the edge of 

the joint where a localized tensile stress can introduce cracks in the ceramic or 

glass as shown in Figure 5.49. If the cracks remain shallow and come back to 

the surface joint, they are harmless. Alternately, if they remain in the ceramic 

or glass interior, where subsequent tensile loadings activate them, or they tilt 

down away from the joint, they can cause severe strength degradation. Glue 

chips, when exposed so as to reveal their fracture surfaces, have ample twist 

hackle, arrest lines, and tertiary Wallner lines. 

5-65 
 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Glue chips were the origins of fracture in spectacular failures of very large 

windows in the John Hancock skyscraper in Boston in the early to mid-1970s.

49 

The windows were 25 mm – 32 mm (1 inch to 1.25 inch) double-glazed glass 

panels with two 6.35 mm (0.25 inch) thick individual panels bonded with a lead 

spacer between them. A copper coating was flame sprayed to the inside rims of 

the glass panels. The glass panels were then soldered to the square lead spacer 

around the rim. Differential contraction of the lead versus the glass created 

incomplete glue chips in cold weather. (A silver reflective coating on the 

inside of the outer glass panel may also have been a factor.) Subsequent 

mechanical and thermal loadings caused the cracks to propagate, and the 

windows eventually fractured. The author and his future spouse, who were 

undergraduate students at Northeastern University, looked at the building in the 

1970s as we came out of engineering classes. At any time, hundreds of glass 

panels were replaced temporarily by plywood sheets. Eventually, all the 

windows in the skyscraper were replaced. The above interpretation of the 

cause of the John Hancock window failures is based on my private 

communications with Prof. Fréchette, Mr. Leighton Orr, and information in 

Reference 49. 

Frechette49 reported a second instance where glue chipping cracks caused 

tempered float glass spandrels to fracture in a high-rise building. In this 

instance, the actual cause was spots of copper metal from weld spatter. The 

high shrinkage of the metal droplets on cooling caused glue chip type cracks. 

5.8 Transgranular and Intergranular Fracture 

In addition to finding the fracture origin, in polycrystalline ceramics it is often 

important to observe whether crack propagation is transgranular and cleaves 

through grains or is intergranular whereby the crack runs around the grains 

along grain boundaries. It is usually necessary to have scanning electron 

micrographs of the fracture surface for this characterization. Figures 5.52 and 

5.53 show examples. Sometimes the word intragranular is used interchangeably with transgranular, but the former word is not recommended, since 

it is so similar to intergranular that it often creates confusion. 

Material scientists are concerned with the mode of propagation since it 

provides information about how cracks interact with the microstructure. 

Intergranular fracture may indicate weak boundary strength. It may be a 

manifestation of between-grain stresses due to elastic or thermal anisotropies, 

especially in non-cubic ceramics. Mixed inter- and transgranular fracture 

modes are common. It is customary to make approximations or estimates of 

5-66 
 

 

 

 

 

Fracture Surface Examination 

the fraction of inter- and transgranular fracture and to show an illustration. 

There is clearly some interpretation involved and what may be a 50 % to 50 % 

mix for one observer might be reported as a 40 % to 60 % mix by another. 

(a) 

(b) 

Figure 5.52 Transgranular fracture examples around fracture origins. (a) is 

aluminum-oxynitride. Notice the wake hackle triggered by the grain 

boundaries and some small pores. (b) is a sintered alpha silicon carbide. 

5-67 
 

 

 Fractography of Ceramics and Glasses 

(a) 

(b) 

Figure 5.53 Examples of intergranular fracture surfaces. (a) is an elongated 

grain silicon nitride. (b) is a 99.9 % sintered alumina. 

5-68 
 

 

 

 

 

 

 

 

 

 

 

 

 

Fracture Surface Examination 

Fracture mirrors in polycrystalline ceramics may initially have transgranular 

fracture but change to mixed trans- and intergranular fracture as the crack 

radiates outwards.32 

Slow crack growth or environmentally-assisted crack growth usually occurs 

intergranularly. Grain boundaries are weaker than matrix grains and secondary 

phases or glasses often are distributed along the grain boundaries. There are 

exceptions, however. 

Transgranular fracture though grains often generates twist hackle and/or 

cleavage steps within the grains as the crack twists and turns to follow 

preferred cleavage planes within the microcrystals. These markings can be 

interpreted to evaluate the local direction of crack propagation. 

5.9 Stable Crack Growth 

5.9.1 Slow crack growth (SCG) at ambient temperature 

Many glasses, oxide ceramics, and nonoxide ceramics with an oxide sintering 

aid are susceptible to environmentally-assisted slow crack growth (SCG). The 

expression “slow crack growth” usually is used in the context of stable crack 

extension due to combined effects of stress and environment. (It is also 

commonly used to describe high temperature stable crack extension as 

described later in section 5.9.3.) Water, either in liquid or as a vapor, is the 

most common reactive species that causes slow crack growth, although other 

small polar molecules such as hydrazine (N2H4), ammonia (NH3), and methanol 

(CH3OH) can enhance crack growth in silicates.

50,51 Ceramics such as alumina 

and magnesium fluoride are also susceptible to environmentally-enhanced 

crack growth. Acetonitrile (CH3CN) can attack magnesium fluoridate, but not 

alumina or silica.

50,51 Slow crack growth requires both a stress to open the 

crack faces and a reactive species to chemically react with the strained atomic 

bonds at the crack tip. Hence, it is sometime called “stress corrosion cracking.” 

Figure 5.54 illustrates the leading theory for how environmentally-assisted slow 

crack growth occurs in glasses. Slow crack growth can lead to weakening 

and/or time dependent fracture. Slow crack growth is also sometimes called 

“static fatigue” if the loading is steady or constant. If the stable crack growth 

occurs while loading increases at a constant rate, as in common strength tests, it 

is sometimes referred to as “dynamic fatigue.” 

All glasses are susceptible to slow crack growth. Many other ceramics are 

too, particularly if they are bonded by a glassy boundary phase, no matter how 

thin. Covalently-bonded ceramics are not susceptible to SCG at room temper5-69 
 

 

 Fractography of Ceramics and Glasses 

ature, unless there is an oxide boundary phase. Section 7.10 has more qualitative information about the mechanics of slow crack growth and Table 7.2 lists 

materials that are or are not susceptible. 

(a) 

(b) 

H 

Si 

O 

O 

Si 

H H O

H 

Si 

Si 

O 

Si 

Si 

O 

H 

H 

O 

H2O 

stress 

Figure 5.54 Stress corrosion cracking can cause slow crack growth in glasses. 

The polar water molecule, either in liquid or gaseous form, attacks strained 

silicate bonds. (after Michalske and Freiman50 and Michalske, Bunker, and 

Freiman51) 

5-70 
 

 

 

 

 

 

 

 

 

 

Fracture Surface Examination 

Sometimes it is possible to detect evidence of the stable crack extension. Slow 

crack growth is usually intergranular in polycrystalline ceramics. If the fast, 

unstable fracture mode is transgranular or mixed, then the difference in 

propagation morphology may be detectable as a SCG “halo” on a fracture 

surface as shown in Figure 5.55. 

Slow crack growth halo: A region or band around a flaw in a ceramic caused 

by stable crack extension. The band is caused by changes in the mode of crack 

propagation. 

SCG halos can often be detected with both optical and SEM microscopy as 

shown in Figure 5.55 for the case of a high-purity alumina.

52 Usually the 

intergranular zone appears darker when viewed with an optical microscope 

since the rougher fracture surface scatters light in different directions. In 

contrast, the slow crack growth zone looks lighter in the electron microscope 

possibly due to slight charging effects on edges of the intergranularly-fractured 

grains. Mecholsky53 showed superb pictures of intergranular slow crack 

growth (from constant stress-to-fracture tests, “static fatigue”) versus 

transgranular fast fracture in magnesium fluoride. Morrell presented a case of a 

fractured alumina hip ball with an unusual circular intergranular zone of slow 

crack growth that appeared shiny in the optical microscope.54 

mixed 

transgranular / intergranular 

mixed 

mixed mixed 

intergranular 

intergranular 

(a) (b) 

Figure 5.55 Knoop semicircular flaws in two specimens with SCG “halos” in 

a sintered 99.9 % alumina. (a) is an optical image of a 29 N indentation flaw 

(Courtesy J. Swab), and (b) is a SEM image of a 49 N flaw. The precracks are 

mixed transgranular-intergranular. The halo is mostly intergranular. The fast 

fracture is mixed transgranular-intergranular. The crack growth occurred in 

only a few seconds during an ordinary strength test. Notice that the rougher 

intergranular surface appears dark in the optical photo since it scatters light, 

whereas it appears light-colored in the SEM. 

5-71 
 

 

 

 Fractography of Ceramics and Glasses 

(a) 

(b) 

Figure 5.56 Slow crack growth in glasses sometimes leaves a partial 

semielliptical line that approximates the critical crack size. (a) shows an 

example for a glass disk tested in ring-on-ring biaxial flexure in water. The 

original flaw was an 8 μm to 10 μm deep crack from a polishing scratch. It 

grew approximately 20 times deeper in the bending stress field before it went 

unstable along the right side. Only a very faint cavitation scarp exists at that 

position, but a more distinct line forms around the crack front periphery as the 

unstable crack wraps around it. (b) shows a schematic of this sequence. See 

also Figure 5.45a. 

The usual remedies to eliminate SCG in laboratory strength testing are to test 

either at a very fast loading rate or in an inert atmosphere. The SCG halos 

disappear when testing is done in inert atmosphere.52 

5-72 
 

 

 

 

 

 

 

 

Fracture Surface Examination 

In glasses a subtle partial elliptical line that sometimes appears on the fracture 

surface provides a strong clue that SCG has enlarged a flaw as shown in 

Figures 5.45a and 5.56.55,56 The part elliptical line is a slight step that forms as 

a result of the intersection of the unstable crack that wraps around the portion 

of the crack that has not gone unstable. 

5.9.2 Stable extension from local residual stress or R-curve effects 

Stable crack extension can occur, even in the absence of environmental effects, 

if a flaw has a strong local residual tensile stresses. Indentation crack flaws 

made by either Vickers or Knoop indentation are good examples. Damage 

zones underneath the indentation can provide the tensile stresses that are the 

driving force for crack extension. As the crack extends away from damage 

zone and the crack periphery gets larger, the effect of the tensile stresses 

diminish and the crack slows down and arrests in the absence of externallyapplied stresses. Alternatively, in fast fracture tests where the specimen is 

loaded to fracture, there can be momentary stable crack extension of the flaw 

prior to rupture. Figure 5.57 shows an example of stable crack extension at 

room temperature due to local residual stresses in a hot-pressed silicon nitride. 

Rising R-curves, which will be discussed in more detail in chapter 7, can also 

lead to stable crack extension at room temperature. There is only a small 

fractographic literature on this effect, in part due to the fact that the amount of 

stable crack extension prior to a flaw going unstable is often very small. 

(Stable crack extension from local residual stresses or due to R-curve effects is 

usually not termed “slow crack growth” in the literature.) 

Figure 5.57 Knoop flaw (dark semiellipse) in silicon nitride showing stable 

crack extension (white arrows). The local residual stresses underneath the 

indentation provided the impetus for the extension. Traces of the very shallow, 

elongated Knoop indentation can be seen on the bottom surface. 

5-73 
 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

5.9.3 High-temperature slow crack growth in ceramics 

At elevated temperatures, increased atom mobility and decreases in boundary 

phase viscosity can lead to intergranular crack growth. This also can cause 

weakening or time dependent fracture. The environment may or may not play 

a role in slow crack growth. Figures 5.58 and 5.59 show flaws that have been 

altered as a result of slow crack growth. The slow crack growth zones are 

usually rougher than the normal fast fracture surface and are usually easy to 

detect due to the topography difference in the SEM or by a topography or 

reflectivity difference in an optical microscope. Oxidation of the material in a 

slow crack zone often causes a color or reflectivity difference that is also easy 

to detect in an optical microscope. The slow crack growth zones can be very 

small and may enlarge a flaw just enough to cause it to reach criticality. 

Alternatively, at lower stresses they may have extensive crack growth zones 

that can cover as much as one half of a fracture surface as show in Figure 5.60. 

In such cases, it often is impossible to locate the initial flaw from which the 

crack growth occurred. 

Figure 5.58 Intergranular slow crack growth in a pure AlON at elevated 

temperature. The slow crack growth zone (arrows) formed in the short time 

that was needed to break the specimen in a strength test. (1200o C, 155 MPa, 

22.5 ksi) 

5-74 
 

 

 

 

 

Fracture Surface Examination 

(a) (b) 

(c) 

Figure 5.59 Elevated-temperature slow crack growth in sintered alpha SiC 

tested in stress rupture at 1400 o

C in air (a). The close-up (b) shows the 

boundary between the intergranular SCG and the transgranular fast fracture 

zones. (c) shows another alpha SiC specimen wherein the slow crack growth 

appears to have originated near a pore. 

5-75 
 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) (b) 

(c) (d) 

Figure 5.60 Elevated-temperature slow crack growth zones in hot-pressed 

silicon nitride. These four bend specimens were tested at 1200 o

C in air. The 

black arrows mark the intergranular crack growth zones, which appear lighter 

since they have oxidized somewhat, unlike the final fracture regions. Notice the 

roughness of mirror-like hackle on the left side of (a) and the right of (b) as 

marked by the white arrows. At low magnification, this roughness could be 

confused with a slow crack growth zone, but closer examination shows it has 

directional lines typical of hackle. 

5-76 
 

 

 

 

 

 

 

 

 

Fracture Surface Examination 

5.9.4 High-temperature creep fracture 

Creep can cause fracture at higher temperatures and lower stresses. There is so 

much boundary phase ductility that sharp flaws are blunted. Fracture is due to 

the nucleation and growth of cavities, microcrack formation, and their 

coalescence. Creep fracture surfaces often look like the slow crack growth 

zones shown above, but with growth zones that extend across a major fraction 

of the fracture surface. Final rupture frequently occurs from creep crack zones 

connected to specimen corners. Tensile surfaces often reveal extensive 

tortuous microcrack patterns. A key difference between slow crack growth and 

creep fractures is that the former can occur with negligible bulk deformation, 

whereas the latter inevitably have significant permanent deformations with 

strains of 1% or more. Figure 5.61 shows several examples. 

stress 

stress 

(a) (b) 

(c) 

Figure 5.61 Bend specimens tested under creep fracture conditions. (a) shows 

the tensile surface of a hot-pressed silicon nitride specimen. Notice the 

extraordinarily rough fracture surface on the right side of (a). (b) shows the 

tensile surface of a siliconized silicon carbide specimen. Note that the 

microcracking is perpendicular to the axis of tensile stresses (double arrow). 

(c) shows a side view of a bend bar tested in creep fracture conditions. 

5-77 
 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Figure 5.62 illustrates the differences between slow crack growth and creep 

fracture. Glassy boundary phases such as shown in Figure 5.63 usually 

account for both processes. 

Slow crack growth occurs from preexistent flaws that grow until they reach a 

critical condition. Fracture toughness at elevated temperature is often greater 

than at room temperature due to the onset of limited plasticity. Hence, critical 

flaw sizes can actually be much larger at elevated temperature than at room 

temperature. Specimens that fail from slow crack growth have negligible (< 

0.2 %) permanent strain. 

Creep fracture on the other hand usually does not involve the initial flaws. It 

blunts them. Voids and microcracks are nucleated, then coalesce, and then 

grow leading to rupture. Creep fracture specimens have considerable 

permanent strain at fracture. Failure is more a matter of accumulation of 

damage. The intergranular crack growth zone on creep fractured bend bars 

often extends over 33 % or more of the fracture surface. 

(a) (b) 

Figure 5.62 Creep fracture and slow crack growth usually are intergranular 

at high temperature. In creep fracture (a), cavitation and microcrack 

nucleation occur throughout the body. In slow crack growth (b), the processes 

are confined to the crack tip region. 

5-78 
 

 

 

Fracture Surface Examination 

(a) (b) 

(c) 

Figure 5.63 Transmission electron microscopy images show the glassy 

boundary phase that often is the source of intergranular slow crack growth and 

creep crack growth. (a) shows the boundary phase which has partially 

devitrified in a silicon nitride with a yttrium aluminosilicate glass sintering aid. 

There still is a thin glassy layer that led to SCG and creep fracture. (b) shows a 

creep cavity opening up at a triple-grain junction. (c) shows strain sworls 

(black arrows) from grain boundary sliding. (Ref. 57, Photos courtesy W. 

Braue). 

5-79 
 

 

 

 Fractography of Ceramics and Glasses 

Figure 5.64 Fracture mechanism map for hot-pressed silicon nitride from 

Reference 58. This map is for flexurally-loaded specimens. 

5-80 
 

 

 

 

 

 

 

 

 

 

 

 

 

Fracture Surface Examination 

Figure 5.64 is a fracture mechanism map that puts these pheneomena into 

context.58 It is for a commercial silicon nitride hot-pressed with about 1 mass 

% magnesium oxide. The sintering aid reacted with silica on the surface of the 

starting powders and formed a very thin silicate boundary phase glass. 

Depending upon the stress and temperature, flexural specimens fractured from 

any one of several mechanisms. At low temperature, material and grinding 

flaws controlled strength. At moderate temperatures in an oxidizing 

environment, flaws either remained unchanged and controled strength, or they 

healed or blunted causing some strengthening. On the other hand, oxidation 

sometimes created pit flaws that weakened the body. At 1000o

C and greater, 

integranular slow crack growth caused flaw growth and eventual fracture in 

times shown in hours on the dashed lines. At low stresses and high 

temperature, creep fracture controlled lifetime. 

5.9.5 Mechanical fatigue from cyclic loading 

Fine-grained ceramics and glasses are usually not susceptible to classical 

mechanical fatigue, in the sense of stepwise stable crack growth from 

accumulated mechanical damage from repeated loading and unloading. Figure 

5.65 shows classical metal fatigue markings that are called “fatigue crack 

striations.” It is exceedingly rare to see genuine fatigue bands like this on 

ceramic fracture surfaces. In metals, each band-arc line may correspond to a 

number of loading cycles during which a damage process may occur, leading to 

one crack jump forward and then arrest, with the process repeating with 

additional cycling. 

With ceramics and glasses, if lines like these are observed, they are more likely 

to be simple arrest lines from iterative loading such as shown in Figure 5.66 

which is from Reference 59. The lines are visible due to slight variations in the 

crack plane between loading cycles. Incremental crack growth during cyclic 

loadings might be simply due to the accumulated effect of the slow crack 

growth phenomena described in section 5.9.1. Evans showed another good 

example in glass in Reference 60. 

The fracture surfaces of laboratory coupons (e.g., tension or bend bar 

specimens), even if they do fracture from fatigue cycling, will often not show 

such markings. The initial flaws grow a small amount before they go unstable 

because fatigue often occurs only for specimens that are loaded near the fast 

fracture strength. 

5-81 
 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) (b) (d) 

(c) 

(e) 

Figure 5.65 Classic metal fatigue striations in (a-d,) a steel railroad coupler 

pin (see also Fig. 3.21), and (e) an aluminum bicycle bracket. An arc line does 

not necessarily correspond to one loading cycle. 

Figure 5.66 Concentric arrest lines in a plate of soda lime silica that was 

cyclically loaded in flexure. The lines are due to slight variations in the crack 

growth planes on each loading cycle. (courtesy of D. Green, Penn State 

University, same as Figure 5.47c) 

5-82 
 

 

 

 

 

Fracture Surface Examination 

(a) (b) 

(c) 

Figure 5.67 Bend bar that was tested to fracture in cyclic fatigue loading. 

The material was a yttria and alumina additive sintered silicon nitride with an 

elongated grain structure for enhanced fracture toughness. (a) shows the bend 

bar fracture surfaces with the origin marked by the arrow. (b) shows a typical 

fast fracture portion of the fracture surface. (c) shows the origin region and 

fine debris (arrows) created by fatigue cycling. (courtesy T. Lube and R. 

Danzer) 

5-83 
 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

On the other hand, polycrystalline ceramics having coarse-grained or 

elongated-grain microstructures and anisotropic properties may be susceptible 

to fatigue crack growth. This is a curious tradeoff: the enhanced fracture 

toughness of an elongated-microstructure ceramic may be offset to some degree 

by sensitivity to mechanical fatigue. Unfortunately, the microstructures creates 

rough fracture surfaces such that it is difficult or impossible to discern evidence 

of progressive crack extension. 

One subtle, but effective sign of genuine fatigue damage cyclic, is small debris 

particles on the fracture surfaces of polycrystalline ceramics such as shown in 

Figure 5.67, which are from Reference 61. The debris is small and detectable 

only in the SEM. The backscatter mode may help detect such debris. The fine 

debris can occur if the loading entails cyclic tension and compression, (e.g., the 

R factor, which is the ratio of the maximum stress to the minimum stress, is 

negative.) Dislodged or fractured grains near the crack tip may be crushed 

between the crack faces if the crack faces go into compression. The debris may 

or may not be present if the crack is loaded in cyclic tension-tension (e.g., the R 

factor is zero or positive). One suspects that too zealous a cleaning of a 

fracture surfaces beforehand might be harmful! On the other hand, it is 

possible that the crushing may help the debris adhere to the fracture surface and 

a light cleaning will not necessarily remove it. 

5.10 Organizing all the Information – Fractographic Montages 

Some component fracture analyses are complex, and an origin site may not be 

immediately obvious. In such cases, the entire fracture surface may be 

examined with the goal of observing the local directions of crack propagation. 

Wake hackle, gull wings, Wallner lines or arrest lines as described in this 

chapter provide such clues. A map of the overall fracture surface can be 

constructed in the form of a large montage, around which key local closeup 

photos of the crack propagation directions can be marked. Little arrows labeled 

“dcp” can then be put on the overall image. The localized direction of crack 

propagation can be “backtracked” to an origin site that might not have been 

immediately obvious. A good example is Figure 5.68 for an alumina dental 

crown which split into two pieces in a patient’s mouth. It was a relatively low 

energy fracture that created only two pieces. The core ceramic microstructure 

was not easy to interpret, but there were enough markings in it and the outer 

veneer porcelain layer that it became obvious that fracture started at a margin. 

Ko also showed marvelous montage examples for an alumina that broke from 

fatigue loading.37 There were many tiny wake hackle lines from microporosity 

in an otherwise rough fracture surface that he used to back track to the origin. 

5-84 
 

 

 

 

 

 

 

Fracture Surface Examination 

A fracture map like this also has great value in communicating a failure 

analysis. It illustrates how much work has gone into an anlysis, and it can also 

provide context. It can be much more effective in convincing a client or reader 

than merely showing a handful of closeup photos and relying on the client to 

trust the fractograpgher’s judgement. 

Montages may also be used to integrate fractographic information with 

mechanical property test data such as a Weibull graph as shown in Section 6.19 

and Figure 6.78. 

Figure 5.68 A fractographic map - montage that integrates a lot of local 

information into a global interpretation. The photos are too small here to see 

the important details. The montage shows the concept. The local photos are 

organized in a systematic manner around the overall image that provides 

context and location information. The white arrow marks the fracture origin. 

Notice the sketch in the top left. Closeup photos and the actual fracture origin 

are presented as Figures 10.26 and 10.27 at the end of this book as the 

Procera alumina molar crown, one of the examples in Chapter 10. (SEM 

photos and the montage courtesy of S. Scherrer) 

5-85 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Chapter 5 References 

1.	 V. D. Fréchette, Failure Analysis of Brittle Materials, Advances in Ceramics, Vol. 

28, American Ceramic Society, Westerville, OH, 1990. 

2.	 H. G. Richter and F. Kerkhof, “Stress Wave Fractography,” pp. 75 – 109 in 

Fractography of Glass, eds., R. C. Bradt and R. E. Tressler, Plenum, NY, 1994. 

3.	 J. E. Field, “Brittle Fracture: Its Study and Application,” Contemporary Physics, 

12 [1] (1971) 1 - 31. 

4.	 A. Smekal, “Procedures for the Measurement of Fracture Propagation Velocities 

on Fracture Surfaces,” Glastechn. Ber., 23 [3] (1950) 57 – 67. 

5.	 U. Hornemann, H. Rothenhäuser, H. Senf, J. F. Kalthoff, and S. Winkler, 

“Experimental Investigation of Wave and Fracture Propagation in Glass Slabs 

Loaded by Steel Cylinders at High Impact Velocities, pp. 291 – 298, in Int. Phys. 

Conf. Series No. 70, 3rd Conference on Mechanical Properties at High Rates of 

Strain, Oxford, The Institute of Physics, London, 1984. 

6.	 C. Brodmann, Nachrichten von der Gessellschaft der Wissenschaften zu Göttingen, 

Mathematisch-Physikalische Klasse, 1 (1894) 44-58. 

7.	 S. Winkler, H. Senf, and H. Rothenhäusler, “High Velocity Fracture Investigation 

in Alumina,” pp. 165 – 183 in Fractography of Glasses and Ceramics, Advances in 

Ceramics, Vol. 22, eds., J. Varner, and V. Fréchette, American Ceramic Society, 

Westerville, OH, 1988. 

8.	 H. Schardin, L. Mücke, and W. Struth, “Bruchgescwindigkeit von Gläsern,” 

Glastechn. Ber., 141 [5] May 1954, 19 -147. Also available translated into 

English in Glass Industry, 36 [3] (1955) 133 – 168. 

9.	 F. Kerkhof, “Bestimmung der Maximalen Bruchgeschwindigkeit Verschiedener 

Gläser nach der Ultraschall-Methode,” Glastechn. Ber., 35 [6](1962) 267 – 272. 

10.	 J. Congleton and B. K. Denton, “Dynamic Fracture Toughness Experiment,” Final 

Technical Report, University of Newcastle on Tyne to U. S. Army European 

Research Office, July, 1975. 

11.	 M. M. Chaudhri, “High-Speed Photographic Investigations of the Dynamic 

Localized Loading of Some Oxide Glasses,” pp. 87 – 113 in Strength of Inorganic 

Glass, Plenum, NY, 1985. 

12.	 F. E. Barstow and H. E. Edgarton, “Glass Fracture Velocity,” J. Am. Ceram. Soc., 

22 [9] (1939) 302 – 307. 

13.	 E. H. Yoffe, “The Moving Griffith Crack,” Phil. Mag., 42 (1951) 739 – 750. 

14.	 J. W. Johnson and D. G. Holloway, “On the Shape and Size of the Fracture Zones 

on Glass Fracture Surfaces,” Phil Mag., 14 (1966) 731 - 743. 

15.	 A. B. J. Clark and G. R. Irwin, “Crack-Propagation Behaviors,” Exp. Mech., 6 

(1966) 321 – 330. 

16.	 J. Congleton and N. J. Petch, “Crack-Branching,” Phil. Mag., 16 (1967) 749 – 760. 

17.	 H. P. Kirchner and J. W. Kirchner, “Fracture Mechanics of Fracture Mirrors,” J. 

Am. Ceram. Soc., 62 [3-4] (1979) 198 – 202. 

18	 H. P. Kirchner and J. C. Conway, “Criteria for Crack Branching in Cylindrical 

Rods: I, Tension, and II, Flexure, ” J. Amer. Ceram. Soc., 70 [6] (1987) 413 – 425. 

5-86 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Fracture Surface Examination 

19	 H. P. Kirchner, “Brittleness Dependence of Crack Branching in Ceramics,” J. Am. 

Ceram. Soc., 69 [4] (1986) 339 – 342. 

20	 W. Doll, “Investigations of the Crack Branching Energy,” Int. J. Fract., 11 (1975) 

184 – 186. 

21.	 Y. L. Tsai and J. J. Mecholsky, Jr., “Fracture Mechanics Description of Fracture 

Mirrors Formation in Single Crystals,” Int. J. Fract., 57 (1992) 167 – 182. 

22.	 E. K. Beauchamp, “Mechanisms of Hackle Formation and Crack Branching,” pp. 

409 – 446 in Fractography of Glasses and Ceramics III, eds., J. Varner V. D. 

Fréchette, and G. D. Quinn, Ceramic Transactions, Vol. 64, American Ceramic 

Society, Westerville, OH, 1996. 

23.	 E. K. Beauchamp, “Fracture Branching and Dicing in Stressed Glass,” Sandia 

Laboratories Research Report, SC-RR-70-766, Jan., 1971. 

24.	 E. F. Poncelet, “The Markings on Fracture Surfaces,” J. Soc. Glass Technol., 42 

(1958) 279T – 288T. 

25.	 E. Gölz, “Übermikroskopische Feinstructuren an Glassbruchflächen,” 

(Ultramicroscopic Fine Structures on Glass Fracture Surfaces), Zeitschrift Physik, 

120 (1943) 773 – 777. 

26.	 D. Hull, Chapter 5, in Fractography, Observing, Measuring and Interpreting 

Fracture Surface Topography, Cambridge Univ. Press., Cambridge, 1999. 

27.	 E. K. Beauchamp, “Crack Front Stability and Hackle Formation in High Velocity 

Glass Fracture,” J. Amer. Ceram. Soc., 78 [3] (1995) 689-697. 

28.	 K. Gopalakrishnan and J. J. Mecholsky, Jr., “Quantitative Fractography of MixedMode Fracture in Soda Lime Silica Glass,” J. Am. Ceram. Soc., 95 [11] (2012) 

3622-3627. 

29.	 K. Gopalakrishnan and J. J. Mecholsky, Jr., “Quantitative Fractography of MixedMode Fracture in an R-curve Materials,” J. Mat. Sci., 48 (2013) 7081 – 7087. 

30.	 E. Sommer, “Formation of Fracture ‘Lances’ in Glass,” Eng. Fract. Mech., 1 

(1969) 539 - 546. 

31.	 R. W. Rice, “Fracture Topography of Ceramics,” pp. 439 – 472 in Surface 

Interfaces of Glasses and Ceramics, eds., V. Fréchette, W. LaCourse, V. L. 

Burlick, Plenum, NY, 1973. 

32.	 R. W. Rice, “The Difference in Mirror-to-Flaw Size Ratios Between Dense 

Glasses and Polycrystals,” J. Am. Ceram. Soc., 62 [9-10] (1979) 533 – 535. 

33.	 D. Lewis, III and J. R. Spann, “Fracture Features at Internal Fracture Origins in a 

Commercial Glass Ceramic,” Comm. Am. Ceram. Soc., Oct. 1982, C-173 – C-174. 

34.	 M. Ramulu, R. C. Bradt, A. S. Kobayashi, and K. H. Yang, “A Dynamic Fracture 

Mechanics Interpretation of Multiple Mist Regions on Soda-Lime-Silicate Glass 

Fracture Features,” pp. 215 – 227 in Fractography of Glasses and Ceramics, 

Advances in Ceramics, Vol. 22, J. Varner, and V. Fréchette, eds., American 

Ceramic Society, Westerville, OH, 1988. 

35.	 G. D. Quinn, L. K, Ives, and S. Jahanmir, “On the Fractographic Analysis of 

Machining Cracks in Ground Ceramics: A Case Study on Silicon Nitride,” Special 

Publication 996, NIST, Gaithersburg, MD, May, 2003. 

5-87 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

36.	 G. D. Quinn, L. K, Ives, and S. Jahanmir, “On the Nature of Machining Cracks in 

Ground Ceramics: Part I: SRBSN Strengths and Fractographic Analysis,” Mach. 

Sci. Technol., 9 (2005) 169 - 210. 

37.	 J. Barsom, “Fracture of Tempered Glass,” J. Am. Ceram. Soc. 51 [2] (1968) 75 – 

78. 

38.	 H. Ko, “Fracture Behavior of Sintered Alumina Under Rotary Bending and Static 

Fatigue,” pp. 15 – 29 in Fracture Mechanics of Ceramics 12, eds. R. C. Bradt, D. 

P. H. Hasselman, D. Munz, M. Sakai, V. Ya Shevchenko, Plenum, NY, 1996. 

39.	 F. Kerkhoff and E. Sommer, “Das Bild der Bruchfläche (von Gläsern),” in 

Handbuch der Mikroscopie der Technik, Vol. IV, Mikroskopie der Silicate, 

Umschau Verlag, Frankfurt am Main, Germany, 1964. 

40.	 A. J. Pons and A. Karma, “Helical Crack-Front Instability in Mixed-mode 

Fracture,” Nature, 464 (4 March 2010) 85 - 89 

41.	 H. Wallner, “Line Structures on Fracture Surfaces,” Zeitschrift fur Physik, 114 

(1939) 368 – 378. 

42.	 A. Tsirk, “Formation and Utility of a Class of Anomalous Wallner Lines on 

Obsidian,” pp. 57 -69 in Fractography of Glasses and Ceramics, Advances in 

Ceramics, Vol. 22, eds., J. Varner, and V. Fréchette, American Ceramic Society, 

Westerville, OH, 1988. 

43.	 T. A. Michalske, “Fractography of Stress Corrosion Cracking in Glass,” pp. 111 – 

142 in Fractography of Glass, eds., R. C. Bradt and R. E. Tressler, Plenum, NY, 

1994. 

44.	 T. A. Michalske and V. D. Fréchette, “Dynamic Effects of Liquids on Crack 

Growth Leading to Catastrophic Failure in Glass,” J. Am. Ceram. Soc., 63 [11-12] 

(1980) 603 – 609. 

45.	 A. Tsirk, “An Exploration of Liquid-Induced Fracture Markings,” pp. 87 – 101 in 

Fractography of Glasses and Ceramics, IV, Ceramic Transactions, Vol. 122, eds., 

J. Varner, and G. D. Quinn, American Ceramic Society, Westerville, OH, 2001. 

46.	 A. Tsirk, Fractures in Knapping, Archaeopress, Oxford, UK, 2014. 

47.	 V. D. Fréchette, “Markings Associated with the Presence of H2O During 

Cracking,” pp. 71 – 76 in Fractography of Glasses and Ceramics, Advances in 

Ceramics, Vol. 22, eds., J. Varner, and V. Fréchette, American Ceramic Society, 

Westerville, OH, 1988. 

48.	 T. A. Michalske, J. R. Varner, and V. D. Fréchette, “Growth of Cracks Partly 

Filled With Water,” pp. 639 – 649 in Fracture Mechanics of Ceramics, Vol. 4, eds. 

R. C. Bradt, D. P. H. Hasselman, and F. F. Lange, Plenum, NY, 1978. 

49.	 V. D. Fréchette and M. Donovan, “Some Effects of The “Glue-Chipping” Process 

on Strength,” pp. 407 –411 in Fractography of Glasses and Ceramics, II, Ceramic 

Transactions, Vol. 17, eds., V. Fréchette, and J. Varner, American Ceramic 

Society, Westerville, OH, 1988. 

50.	 T. A. Michalske and S. W. Freiman, “A Molecular Mechanism for Stress 

Corrosion in Vitreous Silica,” J. Amer. Ceram. Soc., 66 [4] (1983) 284-288. 

51. T. A Michalske, B. C. Bunker, and S. W. Freiman, “Stress Corrosion of Ionic and 

Mixed Ionic/Covalent Solids,” J. Amer. Ceram. Soc., 69 [10] (1986) 721 – 724. 

5-88 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Fracture Surface Examination 

52.	 J. J. Swab and G. D. Quinn, “Effect of Precrack ‘Halos’ on KIc Determined by the 

Surface Crack in Flexure Method,” J. Am. Ceram. Soc., 81 [9] (1998) 2261 - 2268. 

53.	 J. J. Mecholsky, Jr., “Intergranular Slow Crack Growth in MgF2,” J. Am. Ceram. 

Soc., 64 [9] (1981) 563-566. 

54.	 R. Morrell, “Dealing with Component Failures,” pp. 353 -368 in Fractography of 

Glasses and Ceramics V, eds., J. R. Varner, G.D. Quinn, M. Wightman, Ceramic 

Transactions, Vol. 199, American Ceramic Society, Westerville, OH, 2007. 

55.	 T. A. Michalske, “Fractography of Slow Fracture in Glass,” pp. 121 – 136 in 

Fractography of Ceramic and Metal Failures, ASTM Special Technical 

Publication 827, eds., J. J. Mecholsky, Jr. and S. R. Powell, Jr., ASTM, West 

Conshohocken, PA, 1984. 

56.	 T. A. Michalske, V. D. Fréchette, and R. Hudson, “Dynamic Effects of Liquids on 

Surface Crack Extension in Glass,” pp. 1091 – 1097 in Advances in Fracture 

Research, Vol. 2, ed., D. Francis, Pergamon, NY, 1981. 

57.	 G. D. Quinn and W. R. Braue, “Fracture Mechanism Maps for Advanced Structural 

Ceramics, Part 2, Sintered Silicon Nitride,” J. Mat. Sci., 25 (1990) 4377 – 4392. 

58.	 G. D. Quinn, “Fracture Mechanism Maps for Advanced Structural Ceramics, Part I, 

Methodology and Hot-Pressed Silicon Nitride,” J. Mat. Sci, 25 (1990) 4361 - 4376. 

59.	 P. J. Dwivedi and D. J. Green, “Determination of Subcritical Crack Growth 

Parameters by In Situ Observation of Indentation Cracks,” J. Amer. Ceram. Soc., 

78 [8] (1995) 2122-2128. 

60.	 A. G. Evans, “Engineering Property Requirements for High Performance 

Ceramics,” Mat. Sci. and Eng., 71 (1985) 3 – 21. 

61.	 T. Lube, I Khader , A. Kailer, U. Degenhardt, and K. Berroth, “Ermüdungsbruch in 

Siliziumnitrid,” Prakt. Met. Sonderband, 46 (2014) 353-358. 

5-89 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Origins of Fracture 

6. Origins of Fracture 

6.1 Origins, Flaws, and Defects 

The origin is not only a location, but also an object. In the parlance of the 

engineer or scientist, the irregularities that initiate fracture are termed flaws or 

defects. The fractographer should use the terms “defect” and “flaw” with 

discretion. Scientists and engineers understand that virtually all brittle 

materials are imperfect and contain irregularities that can behave as flaws, but 

nontechnical people may misunderstand these terms. That the material 

contains flaws or defects does not necessarily mean that the material has been 

prepared improperly or is somehow faulty. Conversely, there are instances 

where defects or flaws do indicate defective material. The author usually avoids 

using the word “defect” because of its similarity to the word “defective.” 

“Flaw” seems a little less severe and sounds to many as a more technical term. 

If there is any doubt, then the fractographer may fall back upon the most 

innocuous term: “fracture origin.” Ceramists, engineers, and researchers 

occasionally use the term “Griffith flaw:” 

Griffith flaw: A hypothetical sharp crack, usually envisioned as a tiny, 

slender, elliptically-shaped flaw, that intensifies stress at sharp tips. 

Griffith flaws and fracture mechanics are discussed in chapter 7. 

There are hundreds if not thousands of papers on flaws in ceramics and glasses. 

Only a handful are cited in this chapter, but many other key papers and books 

are listed in Appendix A, in the sections: “Origins in Ceramics” and “Origins 

in Glass.” The contributions of Mr. R. Rice, Prof. K. Uematsu, and Dr. F. 

Lange are particularly valuable. Starting in the 1990, Profs. Keizo Uematsu 

and Satoshi Tanaka at Nagaoka University did fine work to characterize flaws 

in ceramics, starting from the earliest processing stages and monitoring flaw 

evolution during subsequent processing. They confirmed the link between the 

flaw size and the strength distributions. This topic is discussed further in 

Weibull Analysis section 7.15 of this Guide. 

6.2 The Spatial Distribution of Flaws 

Flaws are either surface- or volume-distributed. For example, inclusions are 

almost always volume-distributed. Grinding cracks are inherently surfacedistributed. Some flaws may even be edge-distributed. The distinction is 

important for understanding how flaws originated and whether they are from 

processing or from subsequent handling, finishing, or service damage. The 

distinction is also crucial for brittle material design engineers who often must 

6-1 
 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Fracture 

surface 

Volume located Edge located
	

Surface located Near-surface located 

Figure 6.1 Schematics that show the four possible locations of a volumedistributed flaw. 

scale the strengths of laboratory test data to predict component strengths via the 

Weibull theory.1,2,3,4,5

 Strength scaling and reliability predictions may be 

dramatically different for area and volume scaling. Glasses usually, but not 

always, break from surface-distributed flaws. Ceramics can fracture from 

surface or internal flaws. 

One aspect about the spatial distribution of flaws needs clarification as shown 

in Figure 6.1. A single volume-distributed flaw may be located in the volume 

(in the bulk), at the surface, near the surface, or at an edge. For example, an 

inclusion originally may have been in the interior of a ceramic plate, but when 

a test bar was cut out, the inclusion ended up on the surface of the test bar. 

Surface-distributed flaws can only exist at the surface or at an edge. While the 

above statements may seem obvious, there are two reasons to keep the 

distinction clear. Firstly, if a flaw is found at a surface location, it does not 

necessarily mean it is a surface type flaw. It could be a volume type flaw 

located at the surface. Some causal analysts simply look at the locations of the 

fracture origins. If they are located at the surface they simply interpret the 

origin as a surface type flaw. This is often wrong, particularly if the test data 

set is from bend bars or biaxially-loaded plates. It is a common misunderstanding that these two configurations always break from surface flaws. 

6-2 
 

 

 

 

 

 

 

 

 

 

 

 

Origins of Fracture 

Origins indeed are often located at the surface of these bending configurations, 

but if reasonable precautions are taken to control the grinding damage, then 

they are quite likely to break from volume-distributed flaws. The following 

pages show many such examples. Secondly, even if the data are not intended 

for design, a common mistake is to assume that all surface-located flaws are 

machining damage. 

Thus, the fractographer should be careful when answering questions like: 

“What were the origins?” A cavalier answer such as “surface flaws” does not 

convey enough information. The answer should be more specific such as: 

“volume-distributed inclusions located at the surface,” or “volume–type flaws 

located at or near the surface,” or “surface-distributed grinding cracks.” 

Near surface if 0 < l < d Near surface if 0 < l < d1 

l 

d1 

d d2

l 

Fracture surface 

Specimen surface 

Figure 6.2 Near surface (NS) origins are no deeper (l) than the flaw diameter 

or major axis length. 

One might ask at what point is a volume type flaw “near surface.” A review of 

the fracture mechanics stress intensity shape factors for flaws located at various 

distances in proximity to a surface suggests the criterion shown in Figure 6.2. 

At distances closer than one diameter, the stress intensity shape factor for such 

flaws increase by more than a few percent. Sometimes the characterization of 

location is important for service-performance issues or fracture mechanics 

analyses of the flaws. For example, some near-surface located origins may be 

more susceptible to time-dependent crack growth than equivalent volumelocated origins. A classic example of this for time-dependent fractures in 

sintered silicon carbide is Case Study 3 in Chapter 10 of this Guide. Pores 

located at or near the surface behaved differently than identical pores located in 

the bulk. Near-surface-located origins may also be likely to link up with 

surface machining or impact damage, or to extend subcritically to the surface 

prior to catastrophic fracture. 

6-3 
 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

6.3 Are Flaws Intrinsic or Extrinsic? 

The irregularities that act as fracture origins in advanced ceramics can develop 

during or after fabrication of the material. Large irregularities (relative to the 

average size of the base microstructure) such as pores, agglomerates, and 

inclusions are typically introduced during processing. Some might deem 

inclusions and other irregularities as extrinsic flaws in a material, but to the 

extent that they occur naturally during fabrication and cannot be avoided, they 

could be considered intrinsic. For example, if a material customarily has a ball 

milling step in the process, and mill fragments flake off and become inclusions 

in every batch, then these may be considered intrinsic. Other origins can be 

introduced after processing as a result of machining, handling, impact, wear, 

oxidation, and corrosion. These usually can be considered extrinsic origins. 

However, machining damage may be considered either extrinsic or intrinsic. It 

is intrinsic to the manufacture to the extent that machining is a natural 

consequence of producing a finished specimen or component. In any case, the 

differentiation between intrinsic and extrinsic is often of no practical concern 

and it is pointless to argue the distinction. 

Wherever possible, the origin should be characterized by what it is rather than a 

description of how it appears, since the latter may depend upon the mode of 

viewing. Descriptions such as the origin was a “white spot” should be avoided. 

It may be white in an optical microscope, but it most likely will not in a SEM 

(Figure 6.3). Sometimes descriptive terms may be used, but as qualifiers to the 

true flaw identity, e.g., “pores that appear white when viewed optically.” 

(a) (b) 

Figure 6.3 Wherever possible, describe flaws by what they are, not how they 

appear. (a) shows a white spot at the origin in a reaction-bonded silicon 

nitride bend bar. The SEM image of the origin in (b) shows the origin is a 

pore. FS denotes fracture surface and T denotes tensile surface. 

6-4 
 

 

 

 

 

 

Origins of Fracture 

(a) (b) 

(c) 

Figure 6.4 Fracture origin in a sintered alpha silicon carbide flexure 

specimen. (a) suggests the origin is a pore, but (b), which shows the mating 

half, shows it is an agglomerate. (c) shows a polished microstructural section 

for this material that reveals the agglomerates (arrows). 

6.4 Matching Fracture Halves 

Both halves of the fracture surface should be examined whenever possible, 

since each contains information about the fracture origin. For example, Figure 

6.4 illustrates a case wherein an agglomerate was strength limiting. If only the 

(a) half was examined, the fractographer may have erroneously interpreted the 

flaw as a pore. 

6-5 
 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

6.5 External Surfaces 

If an origin is located at the surface, it is wise to examine the exterior surface 

by tilting the specimen back or by directly looking at the exterior surfaces for 

clues as to the flaw type as shown in Figure 6-5. A common question is: “Was 

it a scratch or a grinding flaw?” If the surface-trace lines up with the striations 

that show the grinding direction as in Figure 6.5a, it is likely a grinding fault. 

If the fault has a different axis, as in Figure 6.5b, then it is a scratch. Section 

6.7.3 later in this chapter has more on scratches. 

(a) 

(b) 

Fracture surface 

External surface 

Figure 6.5 The exterior surface should be checked if the flaw is located on the 

surface. (a) shows an irregular, deep grinding flaw with a distinctly different 

ground-surface trace at the origin in a reaction-bonded silicon nitride 

component. It is very likely that a large “renegade” abrasive grit in the 

grinding wheel caused this damage. (b) shows the fracture origin of a fused 

silica rod broken in flexure. This tilted-back view shows a scratch on the 

surface (arrow) coincident with the origin. 

6-6 
 

 

 

 

 

 

 

 

 

 

 

 

Origins of Fracture 

6.6 Volume-Distributed Flaws 

6.6.1 Pores and bubbles 

Pores are one of the easiest flaws to find and identify. In ceramics they are 

often equiaxed, but can easily assume a myriad of odd shapes as shown in 

Figure 6.6. Pores are volume-distributed flaws that are discrete cavities. For 

more information about pores as origins, see Rice’s article, Reference 6. 

Pores in glass or porcelain are often very smooth and round, and may contain 

trapped gas. As such they may be called “bubbles.” Their effect upon strength 

is very different. Since they are very smooth, they usually act as simple stress 

concentrators and not as sharp flaws. This distinction is important as discussed 

in detail in Chapter 7. A material may have large internal pores that are 

relatively harmless, and they may not necessarily act as fracture origins. On 

the other hand, they can be extremely deleterious if they contact an exterior 

surface or each other, as illustrated in Figure 6.7. Bubbles that are close to an 

outer surface are also prone to interact with grinding or machining damage that 

can link with the bubble, turning an otherwise harmless bubble into an 

extremely serious flaw. A good example of this was shown as Figure 5.5 in the 

previous chapter. Figures 6.8 and 6.9 show other examples of bubble origins. 

Reference 7 has more information about the feldspathic porcelain strengths and 

bubble/pore fracture origins. This dichotomy in behavior poses a dilemma to 

design engineers working with flaw sizes and Weibull statistics. A volumedistributed bubble can be either harmless or a dominant flaw depending upon 

where it is located in a piece. 

Pore-bubbles in ceramics or glass ceramics, can act as strength-limiting flaws 

even if they are located in the interior. It is often suggested (but rarely shown) 

that rim cracks a few grain diameters long extend from the bubble surface into 

the ceramic. These are a “Rings of Saturn” type flaw (e.g., Figure 7.26c) and 

fracture mechanics solutions for the stress intensity factors may be found in the 

literature. It is not likely that the actual flaws around the periphery of a real 

pore will be so idealized, however. Figures 6.10a,b show a fascinating example 

by Lewis8

 of a large (100 μm diameter) internal bubble fracture origin (Figure 

6.10) in a cordierite glass ceramic bend bar. The spherical bubble was 

probably smooth in the parent glass, but roughened during crystallization. 

Bernthaler9 used polymeric spheres in alumina to create spherical fracture 

origin flaws in an alumina such as shown in Figure 6.10c. 

Bubbles in glasses and porcelains can be very helpful to fractographers, since 

they are a good source of wake hackle lines on fracture surfaces that show the 

direction of crack propagation. 

6-7 
 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) (b) 

(c) (d) 

(e) (f) 

Figure 6.6 Examples of pore fracture origins in ceramics. (a) shows a SEM 

image of a pore in a 99.9% alumina bend bar. Electron charging created the 

bright spot. (b) shows a pore in an aluminum oxynitride bend bar that is within 

5 μm of the surface. This is a “near surface” location example. The 

transgranular fracture near the pore helps the pore stand out very clearly. (c) is 

an irregular pore in a sintered alpha silicon carbide bend bar. (d) is a pore in a 

porcelain electrical insulator (courtesy J. Taylor and R. Rice). (e) is a 

cylindrical pore in a silicon nitride probably from a fiber or hair burn out. (f) is 

a pore in a yttria-stabilized tetragonal zirconia polycrystal (3Y-TZP) (f is 

courtesy J. Quinn). 

6-8 
 

 

 

 

 

 

 

Origins of Fracture 

(a) (d) 

(b) (e) 

(c) (f) 

Exterior surface 

Exterior surface 

Exterior surface 

(g) 

Figure 6.7 Bubbles are smooth pores and may be relatively harmless or 

extremely deleterious depending upon location. (a) shows a bubble just 

“kissing” the exterior surface, creating very sharp cusps from which cracks 

will likely pop in and grow. (b) shows a near-surface bubble which can easily 

link up with grinding or other surface damage. (c) shows an interior bubble 

that is relatively harmless. Similar behavior can occur with two bubbles in 

proximity to each other. (d) and (e) are likely to form a strength-controlling 

flaw, whereas configuration (f) is relatively harmless. (g) is an example of (a) 

in alumina. (courtesy T. Bernthaler) 

6-9 
 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) 

(c) 

(b) 

Figure 6.8 Bubble origin in a mock-up of a dental crown. Dental stone is cast 

into a cup shape that is loaded in a vise with an inner rubber stopper. When 

the vise is tightened, the stopper expands and the “crown” splits in half. Blue 

tape holds the fragments together. Hoop tensile stresses in the wall triggered 

fracture from a bubble touching the inner surface (c). Grazing illumination 

accentuates hackle lines. Fracture started from the cusp marked with a black 

arrow. These mock-up crowns simulate real crown fractures. 

(a) (b) 

(c) 

Figure 6.9 Strength-limiting flaws in Vita Mark 2 feldspathic dental porcelain 

bend bars. The strengths were: (a)120 MPa, (b)119 MPa, (c) 115 MPa. 

6-10 
 

 

 

 

 

 

 

 

 

 

 

Origins of Fracture 

20 μm 200 μm 

(a) (b) 

(c) 

20 μm 

Figure 6.10 Internal pore-bubbles in ceramics and glass ceramics. (a) and 

(b) show an internal spherical pore fracture origin in a cordierite glass 

ceramic bend bar (315 MPa). The pore was from a bubble in the original 

glass. The material crystalized around the bubble during the ceraming step. 

The arrows mark the hackle boundary of the fracture mirror. (courtesy D. 

Lewis, III) (c) shows a large internal seeded-pore origin in alumina. 

(courtesy T. Bernthaler) 

Pores may often be very irregular and jagged. One way these may occur is by 

the formation of pressing defects when preparing the green body. Spray-dry 

powder agglomerates may not be eliminated during final pressureless sintering. 

Figure 6.11 shows examples. Hot pressing or hot isopressing would likely 

eliminate such flaws. These flaws are sometimes called: “pressing flaws” or 

“blank pressing flaws.” 

6-11 
 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) (b) 

(c) (d) 

(e) (f) 

Figure 6.11 Irregular pore origins. (a) through (d) were formed during cold 

pressing of alumina spray-dried powders. (courtesy T. Bernthaler). 

(e) and (f) are in a sintered 3Y TZP zirconia, (courtesy S. Scherrer). 

6-12 
 

 

 

 

 

Origins of Fracture 

(a) (b) 

(c) (d) 

(e) (f) 

Figure 6.12 Examples of porous region fracture origins in ceramics. (a) is in 

sintered aluminum oxynitride. (b) is in a reaction-bonded silicon nitride. (c ) 

and (d) are both from one sintered alpha silicon carbide bend bar, 

(e) is also from a sintered silicon carbide. (f) is a common type found in 

tetragonal zirconia polycrystals (3Y-TZP). 

6-13 
 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

6.6.2 Porous regions 

Porous regions are volume-distributed flaws that are three-dimensional zones 

of porosity or microporosity regions. Examples are shown in Figure 6.12. 

These can be obvious or very subtle. SEM microscopy is usually needed to 

identify these flaws. Optical microscopy may not be effective since the color 

and contrast at the origin match the surrounding material. 

6.6.3 Porous seams 

Porous seams are similar to porous regions, but are more planar or twodimensional. Examples are shown in Figure 6.13. If the material separated 

completely between these seams, the flaw might be more aptly described as a 

processing crack, as described below in section 6.8.1. 

(a) (b) 

(c ) (d) 

Figure 6.13 Examples of porous seam fracture origins. (a) and (b) show the 

origin in a fine-grained sintered 99.9% alumina. (c) and (d) show the origin in 

matching halves of another fine-grained sintered 99.9 % alumina bend bar. 

6-14 
 

 

 

 

 

 

 

 

Origins of Fracture 

6.6.4 Agglomerates 

Agglomerates are volume-distributed flaws that are a cluster of grains, 

particles, platelets, or whiskers or a combination thereof. They are a common 

flaw in ceramics made with powders that are prepared by spray drying. Spray 

dry agglomerates are often hollow. Agglomerates often sinter away from the 

matrix, creating a shell like void around the agglomerate. Since the 

composition is identical to the matrix, the color and reflectivity of this flaw are 

identical to the matrix. If agglomerates are suspected, examine more of the 

fracture surface for similar spherical or ring like features, since if there is one, 

there usually are many. Examples are shown in Figures 6.4 and 6.14. 

Lange et al.

10,11,12 wrote an excellent set of articles about agglomerates as 

fracture origins in alumina/zirconia composites. 

(a) (b) 

(c) 

Figure 6.14 Examples of agglomerate fracture origins. (a – c) show spray 

dry agglomerate fracture origins in sintered alpha silicon carbide. 

Agglomerates often sinter away from the matrix leaving a void-like crack. 

6-15 
 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Up to this point, agglomerates have been considered as discrete, isolated flaws. 

On the other hand, they may be pervasive and may make up the entire 

microstructure of a body that has not been fully-densified. Green bodies made 

up of dry-pressed, spray-dried particles can be sintered or hot-pressed to full 

density. If not, then the microstructure may comprise particles with incomplete 

bonding and inter-particle porosity such as shown in Figure 6-15. Polished 

microstructural sections may reveal the incomplete bonding, which may be 

subtle. It may be very difficult to discern this on a fracture surface, but it may 

be seen with optical microscopy if the lighting is carefully adjusted as shown in 

Figure 6-15c. Gee and Morrell19 showed similar features in dry-pressed, spraydried β-alumina. Hangl13 showed good images for an industrial grade zinc 

oxide for a varistor application.

 (a) 

(b) (c) 

Figure 6.15 Incomplete densification can lead to interparticle porosity 

throughout a body. This example shows fracture surface of a reaction-bonded 

silicon nitride bend bar. The fracture origin was grinding cracking and is 

marked by the large white arrows. Although this material had low fracture 

toughness, the grinding cracks penetrated only one granule deep, since the 

intergranular porosity impeded further extension. The spherical granules are 

revealed by careful lighting in (c). The small yellow circle shows their size. 

6-16 
 

 

 

 

 

 

 

Origins of Fracture 

6.6.5 Inclusions 

Inclusions are also volume-distributed flaws that are a foreign body with a 

composition different than the normal composition of the glass or ceramic. 

They are often easy to detect due to color or reflectivity differences compared 

to the matrix. Sometimes they are round and equiaxed, other times they are 

odd shaped or fragmented. Figures 6.16 and 6.17 shows some examples. 

(a) (b) 

(c) (d) 

(e) (f) 

Figure 6.16 Examples of inclusion origins. (a) is a red iron inclusion in 

sintered silicon nitride bend bar at an edge chamfer. (b) is an inclusion in a 

silicon nitride tension specimen. (c) and (d) are an iron inclusion in a silicon 

nitride. (e) is an alumina inclusion in a 3Y-TZP zirconia biaxial disk specimen 

(courtesy J. Eichler), and (f) is a vanadium inclusion in a binderless tungsten 

carbide bend bar. (courtesy J. Swab) 

6-17 
 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

The nickel sulfide inclusion shown in Figure 6.17a is a particularly deleterious 

flaw in tempered glasses, since it can undergo a spontaneously phase 

transformation at room temperature with an attendant volume expansion of 

several percent. If the inclusion is located in the interior, where it is exposed to 

tensile residual stresses, it can cause spontaneous catastrophic breakage at any 

time with no warning. Several cases are listed in Appendix B. 

The tungsten inclusion shown in Figure 6.17c and d is also unusually 

deleterious. The tungsten impurity, from tungsten carbide ball milling media, 

can form very brittle tungsten silicides that lower the fracture toughness around 

and in the flaw.14 

(a) (b) 

(c) (d) 

Figure 6.17 Examples of inclusion origins. (a) is nickel sulfide in tempered 

glass plate that spontaneously fractured. (b) is an irregular iron and chrome 

inclusion in a sintered SiAlON bend bar. (c) and (d) are views of a tungsten 

inclusion in a hot-pressed silicon nitride bend bar. 

6-18 
 

 

 

 

Origins of Fracture 

Figure 6.18 The mismatch in properties between an inclusion and a matrix 

determine how severe an inclusion will be. From Evans, ref. 15. 

Not all inclusions behave the same. Their effect depends upon whether their 

elastic and thermal properties match or mismatch those of the matrix.15 

Inclusions can cause cracking in the matrix, may crack themselves, or may 

detach and pull away from the matrix creating a void like flaw. Figure 6.18 

from a paper by Evans15 illustrates the possibilities. 

6-19 
 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Inclusions can also control strength in highly-filled resin-matrix dental 

restorative materials such as shown in Figures 6.19 and 4.1d-f. Transillumination can be effective in finding them in such translucent materials. 

(a) (b) 

(c) (d) 

(e) 

Figure 6.19 Examples of inclusion origins in a highly-filled resin-matrix 

dental composite (Paradigm MZ 100) (a) and (b) show a calcium rich 

inclusion. (c) – (e) show the same aluminum inclusion with: transillumination 

(c), conventional reflected light (d), and in the SEM (e). 

6-20 
 

 

 

 

 

 

 

Origins of Fracture 

6.6.6 Compositional inhomogeneities 

These are microstructural irregularities related to the nonuniform distribution of 

the primary constituents, an additive, or a second phase. They may have a 

color or reflectivity difference compared to the matrix. Sometimes they are 

pockets of glass in a ceramic. Examples are shown in Figure 6.20 and 6.21. 

(a) (b) 

(c) (d) 

Figure 6.20 Examples of compositional inhomogeneity fracture origins. (a) 

and (b) show an example of an aluminum rich zone in a hot-pressed silicon 

carbide. Aluminum is an element in the sintering aid. (c) shows an unreacted 

silicon globule that pulled away from the matrix in a reaction-bonded silicon 

nitride. (d) shows a silicon lake in a siliconized silicon carbide. 

6-21 
 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) (b) 

(c) 

Figure 6.21 Examples of compositional inhomogeneity fracture origins. (a) 

and (b) are an example of a local excess of sintering aid in a sintered reactionbonded silicon nitride. (c) is a yttria-alumina agglomerate in a silicon nitride 

bend bar. (c is courtesy T. Lube and R. Danzer) 

6.6.7 Large grains 

Large grains can occur due to local exaggerated grain growth. They may form 

due to a slight compositional or density variations during sintering. The 

compositional variability can be from an impurity or a sintering aid. Large 

grains can be a single or a cluster of grains having a size significantly greater 

than the range of the normal grain size distribution. They can often be seen 

with the optical microscope due to a difference in reflectivity compared to the 

matrix. Examples are in Figure 6.22. Rice’s book16 on grain size and 

mechanical properties of ceramics has a good discussion of large grains as 

fracture origins on pages 13 – 16. 

6-22 
 

 

 

 

Origins of Fracture 

(a) (b) 

(c) (d) 

(e) (f) 

Figure 6.22 Examples of large grain fracture origins. (a,b) show a cluster of 

large grains at a chamfer in a 99.9% alumina bend bar. (c,d) show an oversized grain in a silicon nitride. The material had small interlocking elongated 

grains intended to enhance fracture toughness. Notice how these are legitimate 

material flaws, and the chamfer preparation was completely satisfactory. 

(e, f) show examples of acicular large grains in sintered alpha silicon carbide. 

6-23 
 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Large quartz grains can act as strength-limiting flaws in porcelains. Quartz has 

several polymorphs and displacive transformations during cooldown from 

processing temperatures can cause cracking around and within the large quartz 

grains as shown in Figure 6.23. This is referred to as the β - α quartz inversion. 

If the quartz grains are much larger than the normal microstructure, they may 

stand out clearly on a fracture surface. If the quartz grains are similar in size to 

the normal microstructure, then the origin flaw may blend into the rough 

fracture surface and be difficult to discern. Good examples have been shown 

by Gee and Morrell, 17 Carlstrom and Carlsson,18 Southan19 (who had superb 

high-quality glossy color-prints), and Banda and Messer.20 

(a) 

(b)

 (c) 

(d) (e) (f) 

Figure 6.23 Cracking around and through quartz grains. (a)-(c) are from a 

porcelain mortar that was broken and are courtesy R. Morrell. (a) shows a 

polished and etched microstructural view. (b) shows an SEM image of one 

fracture origin (circled). (c) is a close-up SEM image showing microcracking 

around quartz particles. (d) – (f) are from porcelain dental crowns, courtesy of 

D. Southan.19 The thin-section polarized-light color image (d) reveals a 

reaction zone and microcracking around the quartz grains. 

6-24 
 

 

 

 

 

 

 

 

Origins of Fracture 

6.6.8 Grain boundaries 

Grain boundaries may be vulnerable areas and, if the grain size is large enough, 

the boundaries can act as strength limiting flaws. Such flaws are very difficult 

to unequivocally identify as origins since the fracture path may be intergranular 

and follow other grain boundaries thereby creating a very rough fracture 

surface. There are many papers on the effect of grain boundary cracking upon 

strength, but few actually show a grain boundary origin. Davidge and Tappin21 

showed grain boundary origins in alumina. 

Sometimes the grain boundary cracking occurs from internal strains set up from 

differential contractions in anisotropic, noncubic ceramic grains. In other 

cases, impurities or sintering aids can gravitate to grain boundaries and weaken 

them. For example, Salem22 found tiny alumina particles at grain boundaries of 

coarse-grained MgAl2O4 cubic spinel transparent plates. These so weakened 

the material that jagged crack paths formed and debris separated from the 

material upon fast fracture. Fracture origins were either large grains or the 

grain boundaries in disk strength specimen. Swab et al.23 also studied a coarsegrained spinel which had irregular crack surfaces and even had granular debris 

generated during fracture. Lithium and fluoride from the sintering aid 

concentrated at the grain boundaries and dramatically weakened the material.23 

I examined one large (40 cm square) plate made of this material and there was 

so much grain boundary cracking that green dye used to stain the fracture 

surface bled profusely into the bulk. Grains literately fell out of the fracture 

surface while I was doing this. 

Coarse-grained (50 μm to 70 μm) magnesia partially-stabilized zirconia 

typically fails from grain boundary cracks as shown in Figure 6.24. Microstructural analysis of polished specimens is essential for the diagnosis in such 

cases. Gee and Morrell17 showed another good example in a partiallystabilized zirconia. 

6-25 
 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) (b) 

(c) (d) 

(e) 

Figure 6.24 Grain boundaries may be flaws in coarse-grained materials such 

as Mg partially stabilized zirconia (50 μm to 70 μm grain size). (a) and (b) 

show optical images of the rough fracture surface of a split rod flexural 

strength specimen. The rough surface is due in part to weak grain boundaries. 

(c) and (d) show SEM images of a polished surface revealing the gaps and 

cracks along the grain boundaries. (e) shows a grain boundary fracture origin. 

Tiny wake hackle from several small pores radiate away from the origin. 

6-26 
 

 

 

 

Origins of Fracture 

6.7 Surface-Distributed Flaws 

6.7.1 Surface voids or pits from processing 

Surface voids are cavities on the surface of a component created by a reaction 

with the processing environment. They are similar to pores, except that they 

only arise on the surface. Surface voids can be in a reaction layer from hotisostatic pressing with a glass encapsulation cladding or from trapped 

outgassing bubbles. Examples are shown in Figure 6.25. 

(a) (b) 

(c) (d) 

Figure 6.25 Examples of surface void fracture origins. (a) and (b) are in asfired bend bars of injection-molded sintered silicon nitride. (c) and (d) show a 

surface void in an as-fired silicon nitride component with yttria and alumina 

sintering aids. (d) is tilted so that the outer surface is visible. EDX analysis 

showed an excess concentration of aluminum at the site, so the flaw also has 

characteristics of a compositional inhomogeneity. (a, b are courtesy A. Pasto; 

c,d are courtesy T. Lube, W. Harrer, and R. Danzer). 

6-27 
 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

6.7.2 Pits from environmental exposure 

Pits are surface cavities that may form as a result of a reaction between a 

ceramic or glass and the usage environment. Corrosion or oxidation can create 

pits. Examples are shown in Figure 6.26. Richerson24 shows several good 

examples of oxidation and corrosion pits in silicon nitride. Pits have chemistry 

or structural differences that make them distinguishable from pores or surface 

voids. Sometimes the damage can be confined to grain boundaries and the 

flawed regions can be difficult to discern on a fracture surface.

 (a) (b) 

(c) (d) 

Figure 6.26 Examples of pit fracture origins. (a) – (c) are in oxidized hotpressed silicon nitride. (b) is a close-up of (a). (d) is a reaction-bonded 

silicon nitride. 

6-28 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Origins of Fracture 

6.7.3 Handling scratches and polishing scratches 

Handling damage may be scratches, chips, cracks, or other flaws due to 

handling or impact. Examples are shown in Figures 6.27, 6.28, 6.29, and 6.5b 

earlier in this chapter. Scratch damage varies widely depending upon the 

material, the sharpness or bluntness, the speed, depth of cut, and force of the 

scratching device or abrasive. At low force, a simple groove may be created 

that is smooth with mostly plastic deformation and minimal cracking. At 

higher forces and with more abusive tools and conditions, grooves with severe 

cracking may be generated. The cracking will be of several types: a crack that 

penetrates beneath the groove and follows the direction of the scratch, a series 

of shorter jagged cracks that jut out to either side of the scratch, and sometimes 

lateral cracks that are parallel to the surface and cause material to spall off. 

The depth of the cracking can be much greater (ten times or more) than the 

visible groove. Scratches that occur after fabrication and finishing are usually 

easy to identify. They are usually isolated and stand out on a surface, provided 

that the fractographer is careful to illuminate the part from different directions. 

Figure 6.5b and 6.27 show examples. It is easy to overlook a scratch on a 

surface if the lighting is not optimal. If scratches are suspected, then it is wise 

to rotate the part under the illumination or move the illumination around while 

watching through a microscope. A quick momentary reflection may be the 

first sign a fractographer gets of a scratch. Scratches may also be overlooked 

on the SEM since they have very low contrast. If the tensile stresses are 

oriented at an angle to the axis of the scratch, then only be a small portion of 

the scratch may trigger a fracture as shown in Figure 6.27. The origin will be 

seen as an irregular crack on the fracture surface. These cracks may be so 

irregular as to confuse a fractographer who may mischaracterize the origin or 

arbitrarily assume it is machining damage. This is a good example of how, if a 

surface-located flaw controls strength, the fractographer should look at the 

exterior surface for other clues. 

In other cases with biaxially-loaded parts, such as pressurized vessels or plates, 

or laboratory ring-on-ring strength tests, there will always be a tensile stress 

available that is oriented perpendicular to the scratch axis. Figure 6.28 shows 

examples. In these cases the fracture follows the scratch for a distance. These 

can be tricky to diagnose if the fracture follows the entire length of the scratch, 

since the scratch trace on the outer surface may be overlooked or hard to see on 

the exact fracture edge, even if the pieces are put back together. On the other 

hand, fractures often veer off from the scratch axis, since the irregular jagged 

cracks that protrude off to the side of the scratch can redirect the main fracture. 

6-29 
 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

scratch 

Side surface 

(a) (b) 

(c) 

tensile surface 

fracture surface 

Figure 6.27 Examples of scratch origins with the scratch at an angle to the 

tensile stress. The cracks penetrate much deeper than the scratch groove 

depth. Scratches can create damage that may be mistaken for grinding cracks, 

so it is wise to look at the orientation of the scratch relative to the grinding 

striations. Scratches will be at irregular angles and are often isolated. (a) is 

in a SiAlON bend bar with a scratch near the chamfer. (b) is a side view of a 

reaction-bonded silicon nitride bend bar. Scratches often require illumination 

at just the right angle to be seen. (c) shows a scratch on an alumina bend bar 

surface (white arrows) and the continuation of it into the interior that acted as 

the fracture origin in a bend bar (black arrows) (c is courtesy J. Swab). 

6-30 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Origins of Fracture 

It is not uncommon with polished surfaces, to find a scratch-created flaw at the 

fracture origin, but all grooves or exterior surface traces are gone. Additional 

polishing after the scratch flaw was created removed all the surface traces. 

Figure 6.28 shows examples. 

Scratch/dig specifications are used with visual inspections of the outer surfaces 

of glass or transparent ceramic materials. The older MIL-PRF-13830B (1997) 

scratch specification is based on a visual comparison to reference pieces. The 

number is sometimes connected to the width of the scratch surface trace, but 

the literature and web sites are contradictory about the actual values. A #40 

scratch sometimes is listed as a 4 μm wide scratch and in other cases as a 40 

μm wide scratch. The dig specification is more consistent, and is the diameter 

of a spot such as is caused by grinding, impact, or contact damage (see 6.7.4 

and 6.7.5 below). Unfortunately, the outer surface features often have no 

correlation to the extent and size of cracks that may, or may not, lie underneath. 

borosilicate crown glass disk 

(a) (b) 

fracture surface 

polished outer surface 

single crystal sapphire dome 

(c) (d) 

Figure 6.28 Two examples of scratch crack fracture origins with the tensile 

stress normal to the scratch. (a) and (b) show the origin in a polished BK-7 

borosilicate crown glass disk broken in ring-on-ring flexure. The origin is a 

very long, but smooth 7.3 μm deep crack. Figure 5.15 shows the fracture 

mirror. (c) and (d) show matching SEM and optical images of an origin in a 

sapphire dome that broke in a high-temperature wind tunnel test rig. There 

were no traces of the scratches on the outer polished surface in either case. 

6-31 
 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

When scratches are suspected, it is usually a good idea to look for other 

scratches elsewhere on the exterior surface of the part. Figure 6.29 shows an 

example where scratches controlled strength in an engine part. 

Occasionally a fractographer is called upon to examine an intact part and assess 

where a line on the surface is a crack or a scratch. The former may be very 

serious and may prompt a part to be replaced, whereas the later could be a 

simple cosmetic irregularity. How does one decide? In my experience, 

scratches usually have jagged edges when viewed with a 5 to 10 power hand 

loupe and strong lighting. They also will not allow water or alcohol to 

penetrate. Cracks, on the other hand, have very-sharp, narrow edges on the 

surface. If water or alcohol is applied, the surface may dry quickly, but the 

crack will stay wet a while longer. If the material is transparent, one can shine 

a strong flashlight or laser pointer through the part side and look for reflections 

from the interior of the crack. Sometimes cracks may be invisible, however, as 

discussed on page 4-48. 

(a) (b) 

(c) 

Figure 6.29 A fractured silicon nitride pin. Examination of the fracture 

surface (b) alone might lead one to conclude that machining damage was the 

origin flaw, but closer examination of the outer surface (c) shows it actually 

was a scratch. Notice how the origin flaw has the same arc as the scratch. 

6-32 
 

 

 

 

 

 

 

 

 

 

Origins of Fracture 

6.7.4 Sharp object impact or contact 

Figure 6.30 shows examples of origins created by impact or contact of a sharp 

object. Median and radial cracks may penetrate deeply beneath the impact site, 

which is often heavily damaged with small fragments missing. Vickers 

indentations are often used to mimic such flaws. Witness marks on the exterior 

surface may give clues as to the nature of the impactor. The exterior surface 

should be examined for other nearby impacts. Fréchette39 describes sharp 

contact site damage as “pecks.” “Digs” are visually-detected contact damage 

sites in glasses and optical ceramics. Specifications for the allowable dig 

diameter may have some utility, but as the figures in this Guide show, and as 

discussed in the previous section (page 6-31), the surface damage size may or 

may not correlate with the size of the cracking beneath the surface. 

(a) (b) 

(c) (d) 

Figure 6.30 Examles of sharp particle imact origin. in glass. The direction of 

impact is shown by the arrows. All of the impacts have created penetrating 

median – radial cracks. (b) and (d) have pronounced lateral cracks as well. In 

each case the impact site has been pulverized or crushed, unlike blunt impact 

sites.(a-d courtesy J. Varner). 

6-33 
 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

6.7.5 Blunt object impact or contact 

Blunt impacts or blunt contact loadings produce round Hertzian cone cracks as 

illustrated in Figure 6.31. These may be complete or partial rings as shown in 

Figures 3.57 (a C-crack in a ball bearing), and 6.32, 6.33 – 6.35. The cone 

crack initiates as a ring just outside the footprint of the two contacting bodies as 

shown in Figure 6.31a. The size of the contact area depends upon the load, 

geometries, and elastic properties of the two materials. A simple rule of thumb 

is that the contacting object is almost always several times larger than the 

observed ring size on the contacting surface as shown in Figure 6.31. Of 

course, if the blunt object penetrates through the plate and makes a large hole, 

then the impactor may be comparable in size to the hole. When contact 

damage is detected, the exterior surface should be examined for witness marks 

at the contact site, as well as other nearby impacts. 

Blunt contact cracks often leave little trace on an outer surface. A pretest 

inspection for “digs” may not find contact damage cracks. The contact cracks 

might be revealed prior to fracture by very careful examination with adjustment 

of lighting angles. After fracture, they are easier to detect since they leave a 

small telltale arc on the fracture surface plane at the origin site. 

(a) (b) (c) 

Static Static Quasi-static Oblique 20o 

slow sliding 

Partial cones 

behind contact 

Impact Impact Oblique 55o 

250 m/s 150 m/s 

skids 

(d) (e) 

Figure 6.31 Blunt contact or blunt impact cone crack profiles. These are side 

views. Notice the blunt impactor is larger than the size of the contact circle 

and first ring to form at the impact site. 

6-34 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Origins of Fracture 

Ring cracks initially pop in normal to the surface but then turn and propagate 

into the depth with an included angle of 125o to 135o

 for quasi static loading1 in 

glasses as shown in Figure 6.31a. The included cone angle decreases 

markedly with increasing impact velocity as shown in Figure 6.31d. It may be 

possible to estimate the projectile velocity from the included cone angle.25,26,27 

For example, at 250 m/s, 0.8 mm to 1 mm diameter steel balls created 60o to 

80o cone cracks in glass25 and 70o to 90o cracks in silicon carbide.27 During 

quasi static or impact loadings, as load increases the footprint area may expand 

and generate multiple concentric ring and cone cracks. Kirchner and Gruver28 

have described interesting variations of cone cracks from impact, whereby 

cracks can nucleate on one side, then branch, and then even spiral around the 

contact site. 

If a ball that is loaded perpendicular to the surface is slid over the surface at 

quasi static rates, the cone axis tilts such that the rear of the cone dips down 

deeper, and the front portion of the cone disappears (Figure 6.31c).

29 Partial 

cone cracks are the result. This is not surprising since the forward motion of 

the ball creates an additional tensile stress behind the ball, but adds 

compressive stresses ahead of the ball. 

Conditions are different if the contactor is loaded (or strikes at) at nonperpendicular angles. Full cones form, but they tilt so that the backside 

approaches the surface while the front side penetrates deeper into the surface as 

shown in Figures 6.31b,e.30 

Most blunt contact loadings produce a cone crack that penetrates partially or 

completely through a plate or body as shown in Figure 4.17. These may not 

cause component breakage and may merely be a cosmetic fault unless the 

structure is intended to seal the environment. Subsequent loading in tension or 

flexure may cause new crack extension from the cone crack as shown in 

Figures 6.32 (d) and (e). Many structures such as dental crowns may have 

numerous harmless cone cracks in the veneers. They are only revealed when a 

fracture that commenced elsewhere passes through the cone (Figure 6.32f). 

Impacts at moderate or high energy create radial cracks and bending stresses 

that can propagate cracks far beyond the impact site and any cone cracks, 

thereby causing component fracture. Figure 6.34b shows an example. 

A common, serious flaw in ceramic balls and ball bearings is a so called “C 

crack” such as shown previously on page 3.68 in section 3.24 on fluorescent 

Stationary or with slow motion. 

6-35 

1 
 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

fracture surface 

tensile surface 

(a) (b) 

(c) (d) 

(e) (f) 

outer surface 

outer surface 

outer surface core ceramic fracture surface 

veneer 

outer surface 

fracture 

surface 

fracture 

surface 

Figure 6.32 Examples of blunt impact or contact crack damage fracture 

origins. (a) shows the tensile surface of a silicon nitride bend bar with partial 

crack rings that were created by a ball-tipped micrometer. Such rings were 

fracture origins in some other bend bars. (b) is a cone crack fracture origin 

from a blunt contact or impact on a rectangular bend bar. (c) shows a blunt 

contact crack from a loading roller on a cylindrical rod glass flexure specimen. 

Notice that the ring is incomplete. The cone crack can act as a strengthcontrolling flaw if loaded in tension as shown in (d) and (e). (f) shows a cone 

crack in a dental bilayer crown. The cone crack, which initiated from contact 

on the veneer layer, was harmless. Fracture initiated elsewhere and a crack 

simply passed though the cone crack. 

6-36 
 

 

 

Origins of Fracture 

(a) (b) 

(c) (d) 

(e) (f) 

Figure 6.33 Blunt impact origins in electrical insulator porcelains. (a) shows 

the installed insulators. (b) and (c) are blunt impact damage from “BB” gun 

projectiles. (d) is damage from a sling shot projectile. (e) and (f) are impact 

origins from contact with support structures. Note the partially exposed cone in 

(f) (arrow). (Photos courtesy of J. J. Taylor and R. W. Rice) 

6-37 
 

 

 

 Fractography of Ceramics and Glasses 

(a)

 (b) 

Figure 6.34 Blunt impact origins in glass plates. (a) shows a mild – moderate 

impact with cone cracks. The plate did not break into pieces. (b) is a severe 

impact that was sufficient to create cone and radial cracks. The latter caused 

plate fracture. (Specimens courtesy of J. Varner and V. Fréchette, Alfred 

University) 

6-38 
500 μm

 

 

 

 

 

 

 

 

 

Origins of Fracture 

dye penetration. It was a partial cone crack. They are often formed by ceramic 

balls impacting each other! Manufacturers wisely started to package ceramic 

balls more carefully once the problem was recognized. 

Blunt-contact cracks can also be a problem in the glass container and 

pharmaceutical industries. Figure 6.35 shows examples. Case 4 in Chapter 10 

has another example of contact damage cracking in a medicinal vial. Figure 4-

23 on page 4-37 shows another good example in a medicinal vial. 

(a) (b) (c) 

Wallner lines 

vial outer surface 

500 μm 

fracture surface 

syringe outer surface 

fracture surface 

500 μm 

fracture surface 

syringe outer surface 

abrasion damage 

syringe outer surface 

fracture surface 

(d) (e) 

Figure 6.35 Blunt contact damage in pharmaceutical glass. (a) and (b) are a 

vial. The pictorial (a) schematic shows a possible impact direction and is a 

useful aid for explaining a fracture event. (c) - (d) are three glass syringes. 

When arc-shaped contact cracks propagate to fracture, they often have hackle 

lines that emanate from the flaw. Notice how these tilt in opposite directions 

on the flaw sides indicating that fracture radiated outwards. Such hackle at an 

origin site suggests “eyelashes.” Notice the witness mark of abrasion damage 

in (d). (Photos courtesy of F. Maurer) 

6-39 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

6.7.6 Grinding and machining cracks 

Grinding or cutting introduces damage to the surface. The damage can be 

controlled, usually by using progressively finer grinding and polishing steps to 

remove damage introduced by prior steps. Edges are particularly susceptible to 

damage and can easily be cracked or chipped. Bevels, chamfers, and rounding 

are effective treatments for the latter. They are commonly applied to biaxial 

disks and bend bars and are carefully aligned circumferentially or lengthwise 

lest the beveling introduce more cracks. 

Surface grinding procedures use grinding wheels as shown in Figure 6.36. The 

grinding process entails material removal, often by means of microcrack and 

microchip formation. By its very nature, cracking is introduced into the 

surface. Initial grinding is done with coarse-grit wheels and modest depths of 

cut in order to facilitate material removal. Intermediate- and final-stage 

grinding is done with finer-grit wheels and smaller depths of cut to reach final 

part dimensions and surface finish requirements, but also to remove prior 

damage while introducing a minimum amount of new damage. 

Fractographers realize that there often is no correlation between the strength of 

a part and the final surface finish. The reason is simple. Grinding induced 

cracks penetrate far deeper (10 to 20 times) below the surface than the surface 

undulations and striations. Sometimes the crack depths are proportional to the 

depth of the surface grooves, and the surface finish may correlate to strength. 

On the other hand, strength may be controlled by deep grinding cracks from 

earlier steps that are not removed in the final finishing phase which may 

produce a very fine outer surface finish. The following case illustrates the 

point. The author ordered a batch of conventional alumina bend bars to be cut 

and ground out of a billet by a new machine shop. The shop was instructed to 

surface grind longitudinally and not use their preferred Blanchard grinding 

process that creates random swirl marks and attendant cracking. The shop 

ignored the instructions and Blanchard ground the bend bars and then lapped 

them to a beautiful finish. The strength of the bars was quite low, however. 

They fractured from severe, randomly-oriented grinding cracks introduced by 

the Blanchard process. The final lapping did not go deep enough to remove the 

prior damage. 

Figure 6.36, an adaptation of a figure by Rice and Mecholsky,

31 illustrates 

some key aspects of machining flaws. Parallel cracks oriented in the direction 

of grinding are longer and are the same depth or deeper than the orthogonal 

cracks. Thus, the parallel cracks are more deleterious to strength. 

6-40 
 

 

Origins of Fracture 

Figure 6.36 Surface grinding (insert) creates grinding grooves (striations) 

and subsurface cracking. The cracking includes long deep “coplanar parallel 

cracks” or “offset parallel cracks” aligned with the striations in the direction 

of abrasive motion. Orthogonal cracks form at right angles to the striation and 

may be bowed in the direction of abrasive motion. Lateral cracks or spalls that 

are roughly parallel to the surface may also form (not shown). The axis of 

grinding is often chosen to align the worst cracking damage (parallel cracks) 

with a specimen axis as shown on the bottom right. 

6-41 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

The differences between the orthogonal and parallel cracks give rise to strength 

dependence with orientation (if grinding cracks are the controlling flaws). 

Only one striation groove is shown in the schematic and it should be 

remembered that the ground surface actually contains a complex overlapping 

network of such cracks and grooves. 

With some experience it is not too difficult to find parallel grinding crack flaws 

in dense materials, such as glass, single crystals, and either very coarse-grained 

or extremely fine-grained ceramics. Detection of grinding cracks in porous, 

heterogeneous, or mid grain-sized materials is very difficult, however, since the 

machining crack blends in to the background microstructure. This is especially 

true for the short orthogonal cracks. They are hard to find in nearly all 

materials. Figure 6.37 shows some examples of orthogonal cracks, which 

sometimes are described in the literature as “longitudinal-grinding flaws.” 

(That expression means: “the flaws that are created by longitudinal grinding” 

with respect to the specimen’s long axis.) Additional examples are shown in 

some of the papers by Rice and Mecholsky listed in the Bibliography, 

Appendix A. 

Parallel machining cracks are much easier to detect and often can be found with 

routine optical microscopy, but only if low-angle grazing illumination is used. 

Figure 6.38 shows examples. Figure 6.39 shows a schematic of these parallel 

flaws and their telltale signs. Figure 6.40 shows how “V” machining cracks 

form in ground round parts. Several of these grinding crack flaws exhibit 

another form of hackle on fracture surfaces: 

Grinding crack hackle: Hackle formed on the fracture surface originating 

from irregularities and steps in grinding cracks. 

Grinding crack hackle is sometime also known as “shark’s teeth.” The size of 

grinding cracks, and hence their effect on strength, depends upon material 

properties such as hardness, elastic modulus, and fracture resistance or fracture 

toughness. It also depends on the grinding wheel and machine conditions such 

as abrasive type, grit size and concentration, depth of cut, wheel speed, 

grinding forces and machine horsepower. The crack size also depends upon 

the material’s grain size. Deep grinding cracks can form in very coarse-grained 

ceramics since the single crystal fracture toughness is less than the 

polycrystalline fracture toughness. It is beyond the scope of this guide to delve 

into this topic other than to say that abrasive grit size is the dominant parameter 

that controls crack size for typical surface grinding procedures. Figure 6.40 

shows the effect of grit size upon parallel crack depth in a commercial silicon 

nitride.32 

6-42 
 

 

 

Origins of Fracture 

(a) 

(b) 

Figure 6.37 Orthogonal grinding cracks as seen in some very fine-grained 

ceramics. (a) shows a crack in a hot-pressed silicon nitride bend bar. (b) 

shows a crack in a hot-pressed silicon carbide. The bars have been tilted back 

in the SEM to reveal parts of the ground surface “T” which was in tension. 

“fs” means fracture surface. The specimens were ground longitudinally with 

respect to the specimen’s long axis. Slightly larger than normal striations are 

evident on the ground surface in each case. The orthogonal cracks are hard to 

find even with the SEM, since they tend to blend in with the microstructure. 

6-43 
 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a)

 (b) (c) 

(d) (e) 

Figure 6.38 Examples of parallel grinding cracks. (a) and (b) show optical 

images of a “zipper crack” in a sintered silicon nitride bend bar. (d) and (e) 

show SEM images of a crack in a silicon nitride rod. 

6-44 
 

 

 

Origins of Fracture 

 (f)

 (g) 

Figure 6.38 (cont’d) Examples of parallel grinding cracks. (f) shows back-toback halves of a silicon nitride bend bar with coarse grinding (150 grit) cracks. 

A zipper crack is at the origin (large white arrow). The small black arrows 

show that there also is a machining crack “skin zone” along the entire surface. 

(g) is an alumina bend bar with low-angle grazing illumination from the side 

revealing a zipper crack. 

6-45 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) elongated “coplanar parallel crack”

(or coplanar linked semi-elliptical cracks).

A deep striation may or may not necessarily be present.

The fracture mirror may be elongated along the outer

specimen surface.
	

(b) elongated “coplanar parallel crack”

linked with a natural flaw.
	

A step in the fracture origin emanates from the material flaw.
	

(c) “zipper crack” 

This is a series of short semi elliptical cracks, which have linked.

A series of short tails, or “machining crack hackle,” emanate 

from the links or overlaps of the flaws and extend up into the fracture 

mirror. These tails may be tilted to the left or right and help confirm that

fracture originated in the central region of the set. The short tails are 

telltale features of slightly misaligned or overlapping transverse 

machining cracks (or a scratch) and are often easier to see with an 

optical microscope with low angle lighting than with a scanning electron 

microscope. The fracture mirror may be elongated along the specimen 

outer surface or it may have one or two prominent side lobes.

This origin type is common in transversely-ground rectangular

flexure specimens or scratched biaxial disk specimens.
	

(d) coarse “zipper crack” 

This is made up of a series of irregular, less coplanar semi- elliptical 

cracks. Larger tails than in (c ) are created. In severe cases, the tail
	

may extend all the way to the mirror boundary. The fracture mirror may
	 be elongated. This origin is common in transversely ground or

scratched specimens and the markings are sometimes termed “shark’s
	

teeth.”
	

(e) “V machining crack” 

The crack intersects the fracture surface at an angle. Only a portion of

the machining crack or crack series is exposed. A pronounced step 

occurs in the fracture mirror. One or two (shown) tails extend well up 

into the fracture mirror. The machining direction is not quite perpen- dicular to the specimen length and uniaxial stress axis due to grinding 

wheel cross feed. This origin is common in cylindrical specimens
	 prepared by centerless or cylindrical transverse grinding wherein the 

wheel and work piece displace axially relative to each other.
	

(f) “coarse grinding parallel crack”

The origin is a deep machining crack that extends along the entire 

surface. The origin is often bumpy since the origin is comprised of offset
	 parallel cracks. Thin bands of uniform depth extend along the specimen

surface on either side of the fracture mirror. The bands have the same 

depth as the grinding cracks. Short tails, or “machining crack hackle”

which may be in the thin bands are tilted away from the origin. This
	

origin type is common in coarse ground surfaces.
	

Figure 6.39 Schematics of telltale features of parallel machining cracks. 

These origins are common in biaxial disks, pressure-loaded plates, and 

transversely-ground flexural or tension strength specimens. Low-angle 

grazing illumination facilitates their detection with optical microscopy. 

Some have elongated mirors along the surface, an important telltale sign. 

6-46 
 

 

 

Origins of Fracture 

(a) (b) 

Ground or lapped 

surface 

(c) (d) 

Figure 6.40 Round parts may have “V” machining cracks formed from small 

parallel crack segments a, b, c on different planes as shown in (a) and (b) for a 

silicon nitride rod. The steps between the individual crack segments form two 

or more telltale hackle lines, which extend up into a mirror. Examination of 

the ground surface shows striations (grinding grooves) that have a small angle 

relative to the rod axis. V crack flaws are easy to see even at very low 

magnification with low-angle grazing illumination. 

6-47 
 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Figure 6.41 The depth of grinding cracks as a function of grinding wheel grit 

in a silicon nitride. The notes show the strength-controlling flaws. Ref. 32 

The average grit size may not necessarily be the controlling parameter. 

Instances of renegade abrasive grits in a grinding wheel that caused atypical 

damage have been reported.32 Figure 6.5a earlier in this chapter and Figure 

6.42 below show an example of damage from a renegade grit in a silicon 

nitride. Calculations based on the wheel and table speeds and the specimen and 

wheel diameters confirm that a single grit can cause such long striations. 

Striation here means a grinding groove. 

Although it is widely believed that materials with rising R-curves form smaller 

grinding cracks than materials with flat R-curves, there is evidence that just the 

opposite is true.33 Materials with rising R-curves are more difficult to grind, 

but for a given wheel speed and depth of cut, greater grinding forces are 

generated to overcome the removal resistance. The greater forces cause cracks 

to penetrate deeper than in materials with low fracture toughness or flat Rcurves that cut or grind more freely. 

6-48 
 

 

 

 

 

Origins of Fracture 

4 mm 

(a) (b) 

(c) (d) 

Figure 6.42 Evidence of “renegade- grit” damage in a 150 grit transverselyground 4 mm wide silicon nitride bend bar. (a) shows the tensile surface and 

the plane of fracture marked by the arrow. (b) shows a pattern of four 

repeating striations that existed periodically (every 2 mm) along the bar length. 

Every bend bar in the set broke from one of the #4 striations such as shown in 

(c). (d) is a close-up in this slightly-translucent material, illustrating that 

considerable subsurface grinding damage occurred beneath striation #4. 

Figure 6.38e shows a typical fracture surface for this set and shows the 

damage associated with striation #4. 

6-49 
 

 

 Fractography of Ceramics and Glasses 

Machining cracks often link up with each other or with other materials flaws as 

shown in Figure 6.43. We tend to focus only on the ones that we see on a 

fracture surface, and it should be borne in mind that the surface actually 

contains a complex network of cracks as shown in Figure 6.44. 

Figure 6.43 Machining cracks can link with each other or with material flaws. 

Figure 6.44 Ground surfaces contain a dense network of machining cracks. 

This simplified schematic does not show cracks running parallel to the stress 

axis or lateral cracks parallel to the surface. 

6-50 
 

 

 

Origins of Fracture 

ground 

SEM view 

side view of fracture surface 

ground 

Figure 6.45 A different pattern of grinding damage termed “herringbone 

crack pattern” discovered by Kanematsu, Ref. 34. Figure 3.58 shows 

additional examples. (images courtesy W. Kanematsu) 

Kanematsu34 identified a “herringbone pattern” of grinding cracks in silicon 

nitride that has a completely different morphology those shown previously in 

this chapter. Figure 6.45 shows the cracking is made up of long, closely-spaced 

(but not touching) parallel cracks that curve off in the direction of abrasive 

motion. Indeed Figure 3.58 in chapter 3, which showed his innovative dye 

penetration technique applied to machining cracks illustrated, has several of 

these. These grinding cracks were in a fine-grained dense silicon nitride 

ground with 100 grit, 75 concentration resin bonded diamond wheels. 

6-51 
 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) (b) 

(c) 

Figure 6.46 Fracture surface of a feldspathic porcelain bend bar. (a) shows 

the entire surface of the 3 mm x 4 mm bar with a gold coating. The origin is on 

the right side wall (arrow). (b) and (c) show SEM images of the damage along 

the right side wall. Notice the slightly greater roughness of the fracture surface 

along the right side as compared to the fracture mirror centered on it. 

Figure 6.46 shows an interesting example of parallel grinding damage on the 

side wall of a dental porcelain bend bar. Five of twenty-six test bars broke 

from such side wall damage.7 This is unusual since bend bars typically have 

the four main surfaces ground the same way. 

Careful examination of the exterior surfaces of the porcelain bars revealed that 

the bottom tensile surface had been polished adequately to minimize residual 

grinding damage. The tensile surface probably had been initially ground in a 

longitudinal direction. The side walls on the other hand had residual grinding 

damage from transverse grinding. Although the side walls were also polished, 

it was not done deep enough and the damage was not completely removed. 

Most porcelains have a rough fracture surface due to the microstructure, and 

such distinct fracture mirrors and grinding damage may be obscured. This 

particular porcelain, Vita Mark 2, is an exception due to its fine-grain size. 

6-52 
 

 

 

 

 

 

 

 

 

 

 

Origins of Fracture 

6.7.7 Chips 

Blunt or sharp impacts or contact loadings near an edge can chip off a piece as 

shown in Chapter 4. The chip can in turn become a strength limiting flaw on 

subsequent loadings. Figure 6.47 shows examples. 

(a) (b) 

(c) (d) 

Figure 6.47 Examples of chips that acted as origins. (a) is an edge chip that 

acted as the flaw in a bend bar. This view shows the reassembled halves from 

the side. The top arrow shows the impact point that first created the chip. 

When broken in flexure, the fracture initiation site was found to be at the small 

black arrow. Notice the continuity of the chip shape on each side of the final 

fracture. This confirms that the chip was in the bar to begin with, and was not 

a secondary fracture. (b) is an edge chip in a sintered silicon carbide bend bar 

chamfer. (c) and (d) show chips on bend bar chamfers in hot-pressed silicon 

nitride. Close examination of the chamfers revealed that the grinding was not 

done longitudinally, but at an angle causing the chips. Notice the striation 

marked by the white arrow on the chamfer in (d). 

6-53 
 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

6.8 Other Flaws 

6.8.1 Processing and firing cracks in ceramics 

Cracks that form in the green state or during firing are quite variable and 

troublesome. Firing stresses can be very irregular, may be transient, and may 

depend upon thermal gradients in the kiln or furnace. Firing or shrinkage crack 

examples are shown in Figures 6.48 - 6.52. 

Cracks that meander, are segmented, and/or are granular on the fracture 

surface may have formed in the green state or during firing. Since the body is 

not fully dense, the fracture surfaces are very granular and do not have the 

characteristic fracture surface markings. Close examination of interior crack 

surfaces may reveal thermal rounding of the particles and grains. 

(a) (b) 

tensile surface 

Fracture surface half A 

Fracture surface half B 

tensile surface 

tensile surface 

(c) (d) 

Figure 6.48 Examples of processing cracks. (a) shows the tensile surface and 

(b) the fracture surface of two different silicon nitride bend bars which had an 

unstable boundary phase associated with yttrium silicates. (c) shows a bend 

bar that had a moderate strength despite a huge processing crack running 

perpendicular to the tensile axis. The origin was at the arrow. (d) shows 

another example of a firing crack in a bend bar. The origin is marked by the 

arrow. 

6-54 
 

 

 

 

 

 

Origins of Fracture 

(a) (b) 

(c)

 (d) (e) 

CAD/CAM ground occlusal surface 

separated surface 

Figure 6.49 Processing cracks in dental ceramics. (a) and (b) show one tooth 

of a three-unit alumina dental bridge. (a) shows a frontal (buccal) view of the 

end unit which fractured at the narrow connection to the next unit (small white 

arrow). A sintering crack is evident on the face of the crown (large white 

arrow). It did not cause fracture in this instance because it was in a low stress 

area. (b) shows a close-up of the processing cracks. (c) – (e) show a broken 

zirconia crown. A large piece has separated off a cusp on the crown top 

(occlusal) surface (c) and (d). The apparent “fracture surface” is in fact a 

huge processing crack. Portions of the crown separated during firing and what 

is seen are exposed spray-dry agglomerates. The manufacturer claimed the 

fracture was due to patient overload or faulty installation by the dentist. This 

was wrong, since the CAD/CAM blank had the processing crack in it. (c - e 

are courtesy Dr. S. Scherrer and the ETH, the University of Zurich.) 

6-55 
 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) 

 (b) 

(c) 

(d) (e) 

Figure 6.50 A prototype Ford silicon nitride automotive turbocharger rotor 

(a). The material is a reaction-bonded form of silicon nitride and was formed 

by injection molding. The lower view in (a) shows the interior after sectioning 

with a diamond cutoff wheel. The close-up views in (b) and (c) show how 

differential shrinkage during firing caused cracks to form in the thickest 

portion of the part. Notice the lighter-colored areas of low density that envelop 

the cracks. (d) and (e) are of a zirconia watch bracelet. (d) is a close-up from 

inside the processing crack that shows rounded zirconia grains and porosity. 

(d, e are courtesy J. Kübler) 

6-56 
 

 

 

 

 

 

 

 

 

 

 

Origins of Fracture 

(a) (b) 

Figure 6.51 Delaminations revealed on the fracture surfaces of two tape-cast 

PZT bend bars. (a) has severe delamination between tape cast layers. (b) was 

processed much better but still had a very small horizontal delamination 

(arrows). The curved lines are arrest lines on the fracture surface. (courtesy 

C. Wu.) 

Delaminations are a particular type of processing crack that are generally 

planar cracks within a material that arise from separation during green body 

forming, especially from mold or press release operations. They can also occur 

between bonded layers of dissimilar material or joints. Delaminations often 

occur in layered or tape cast structures such as the two lead zirconium titanate 

(PZT) examples shown in Figure 6.51 from reference 35. Diagnosis and 

remedial action for processing cracks usually requires close consultation with 

the processor to review the details of the fabrication and any steps that could 

lead to differential shrinkage, outgassing, or phase changes. Processing cracks 

and their effect on overall breakage and fragmentation are also covered in 

section 4.5. 

On the other hand, cracks that are sharp and planar (not rough and granular) 

may have popped in during or after cool down from the firing temperatures. 

Such cracks may have fracture surface markings such as hackle or arrest lines. 

The microstructure will look fully dense with no roughness or rounding. 

Figure 6.52 shows such an example. Sometimes these are termed “thermal 

cracks.” 

Varner and Fréchette36 identified dunt cracks in whitewares. These are cracks 

that from during cool down when quartz particles undergo a sharp change in 

volume (due to the β - α phase change) when matrix glass is viscoelastic. The 

dunt crack appears as an atypically shiny transgranular fracture. 

6-57 
 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

firing crack 

final fracture 

Starting point 

(c) 

(a) 

Figure 6.52. A sharp firing crack in a dental prosthesis. This is a large single 

piece zirconia five-unit “telescoping denture,” a single-piece insert for a 

patient with substantial tooth and jaw bone loss. Two implants held it in place 

at the left end and the fourth tooth in this view. The cantilevered fifth molar 

tooth broke off after only three months. (a) shows the four surviving units and 

(b) shows the end tooth that broke off. (c) shows the same piece with a gold 

coating to make optical examination easier. The origin of the final fracture 

was the large firing crack that propagated almost two-thirds through the 

massive zirconia cross section. This sharp, smooth and probably very-tight 

crack was unnoticed during fabrication and installation in the patient. It 

started from irregularities in the veneer and core area marked by the yellow 

arrows. Final fracture occurred on a different plane and in stages as attested 

to by multiple arrest lines. One normally does not expect fracture in such a 

massive zirconia piece, but the cantilever loading on the end piece which had 

an enormous initial crack caused it to break. (courtesy M. Karl, U. Lohbauer, 

J. R. Kelly) References 37 and 38. 

(b) 

6-58 
 

 

 

 

 

 

 

 

 

 

 

Origins of Fracture 

6.8.2 Artificial or controlled flaws, glasses and ceramics 

Sometimes it is advantageous to deliberately create strength-controlling flaws. 

This is often done in fracture mechanics studies to create a controlled starting 

crack and then load it to fracture in order to evaluate the material’s fracture 

resistance. Cracks made for fracture mechanics tests are typically called 

“precracks.” Artificial flaws are sometimes made by fractographers to help get 

acquainted with fracture markings on a new or difficult material. Usually a 

microhardness testing machine with a Knoop or Vickers indenter is used to 

make the controlled flaw in a bend bar or disk specimen. The indentation not 

only makes a crack, but significant residual stresses as well. The residual stress 

damage zone extends three to five times the indentation depth beneath the 

surface. The best remedy to deal with the residual stresses is to polish or handgrind the residual stress damage zone off, leaving behind a stress-free 

semielliptical surface crack. 

The Knoop indenter produces a much more controlled “cleaner” crack than the 

Vickers indenter. The Knoop indentation procedure for creating controlled 

flaws is described in more detail in section 7.12. When used to measure 

fracture toughness with bars in bending, the method is known as the surface 

crack in flexure (SCF) method. Figures 3.26, 3.56, 5.23, 5.55, 5.57, and 6.53 

show examples of Knoop surface flaws that were fracture origins in bend bars. 

The Vickers indenter creates a complex “messy” three-dimensional network of 

median, radial, and lateral cracks. There is much shear faulting and 

densification as well. Most fracture mechanics experts prefer simple, straight 

cracks for testing purposes. 

(a) (b) 

Figure 6.53 Examples of nice, clean Knoop indentation flaws. (a) shows a 

precrack in a sintered alpha silicon carbide bend bar. (b) shows a BK-7 

borosilicate crown glass bend bar. (c)-(h) are on the next page 

6-59 
 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

(c) (d) 

(e) (f) 

(g) (h) 

Figure 6.53 continued. (c) and (d) are optical and SEM images of cracks in 

a sintered reaction-bonded silicon nitride. (e)and (f) are an e.max CAD dental 

lithium dislicate glass ceramic with optical green dye(e) and SEM (f) images 

of the same crack. (g) is a dental porcelain designed to be fused to zirconia, 

and (h) is a Vita Enamic ceramic network, resin-infiltrated composite. 

6-60 
 

 

 

 

 

 

 

 

 

 

 

 

Origins of Fracture 

6.8.3 Other Glass Origins 

Examples of glass origins have been shown previously in this chapter: nickel 

sulfide inclusion in Figure 6.17a, scratches in Figures 6.5b, polishing-grinding 

cracks in Figure 6.28a,b, impact and contact cracks in Figure 6.30, 6.32 and 

6.35, and artificial flaws in Figure 6.53. Other flaw types that are common in 

glasses are shown below. Varner39 and Fréchette40 have described the 

character of flaws in glass. 

Flaws from the melt may come in several forms. These flaws are inclusions, 

pores, or compositional inhomogeneities, but have been given more specific 

names by the glass community.39,40,41,42,43 Devitrification inclusions or 

devitrification stones (Figure 6.54) can form if the glass crystallizes. Stones 

are grains of unreacted raw material or refractory material broken off the tank 

or pot walls. Devitrification stones and stones have completely different 

chemistries and causes, so they should not be confused. Gas bubble origins are 

termed blisters if large, and seeds if small. Striae are streaks of 

inhomogeneous, transparent glass. Their refractive index is different than the 

glass bulk. Very intense striae are termed cords and may appear as ropes or 

strands such as shown in Figure 6.56. 

Figure 6.54 Unusually large (≈ 1 cm) glass devitrification inclusions with a 

blossom-like structure as well as smaller seed bubbles. (Specimen courtesy of 

J. Varner and V. Fréchette) 

6-61 
 

 

 

 

 Fractography of Ceramics and Glasses 

(a) 

(b) 

(c) 

Figure 6.55 Stone origins exposed on glass fracture surfaces. (a) is an 

example in a large-screen television picture tube and broke from transient 

thermal stress during manufacture. (specimen courtesy of J. Varner and V. 

Fréchette). (b) and (c) show a stone in the middle of the wall of a heatstrengthened consumer glassware. 

6-62 
 

 

 

 

 

Origins of Fracture 

(a) (b) 

(c) 

cord 

stone and cord 

(d) (e) (f) 

Figure 6.56 Flaws in a borosilicate glass cookware. (a) – (c) show a chord. 

(crossed polarizers highlighted the internal strains in views (a) and (b). (c) is a 

close-up of the cord with a mix of transmitted and reflected illumination. The 

striae - stringers are marked by arrows. (d) shows a close-up of a stone with 

striae - stringers. (e) is a seed (a bubble). (f) is an odd internal hairpin flaw. 

Preston42 described heat sleeks, which are burnished streaks of glass where a 

metallic polishing or grinding wheel makes direct rubbing contact with a glass 

surface if the loose abrasive and water are inadvertently squeezed out. The 

glass locally reaches a high temperature and develops a number of hair-like 

fissures entering the glass normally to the surface and penetrating only a short 

distance. Surface contact damage flaws were previously presented in section 

6.7.4. Sharp contact damage sites are pecks. 

6-63 
 

 

 

 

 Fractography of Ceramics and Glasses 

(a) (b) 

(c) 

(d) 

Figure 6.57 Chill check cracks on the outer surface of unbroken glass plates. 

(a) and (b) show the same chill check. (c) and (d) show two close-up views of 

another chill check. Note how different illumination can dramatically alter the 

appearance and apparent extent of the flaw. Much more of the cracking is 

evident in (d). (Specimens courtesy V. Fréchette and J. Varner) 

6-64 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Origins of Fracture 

Figure 6.57 shows chill checks, a pernicious flaw that is often very difficult to 

detect. A chill check is a small wavy and tight crack that is introduced into a 

hot glass surface by localized contact with a cooler object. Thermal stresses 

create the crack, but since the contact and thermal stresses are localized, the 

cracks are small. Glasses with high thermal expansion coefficients are most 

susceptible. The cracks often are difficult to detect in an unbroken component 

and require lighting at just at the right angle. Sometimes transmitted 

illumination is effective. If chill checks are suspected, then the piece or the 

illumination source should be manipulated a variety of ways and angles so as to 

try to get a momentary reflection from the crack. Chill checks form quickly 

and close back up very quickly. They may close so tightly (less than the 

wavelength of light) that they are not discernable. Nevertheless, some portions 

of the crack may have captured small amounts of gasses or water molecules 

that create slight opening pockets that are just barely detectable. These are 

often brightly colored, since they refract the light and may only be detectable 

from certain viewing angles. Once a momentary reflection is detected, the 

fractographer can zero in on the flaw site. Often the full extent of the crack is 

not visible because the crack is so tight. Fréchette40 observed that chill checks 

are hard to characterize since mechanical effects can also occur at the contact 

site. If a glass object breaks from a chill check during subsequent use and 

stressing, then it is likely the origin site will be irregular and may not be 

recognized as a chill check without examination of the outer surface. This 

underscores an earlier recommendation: if an origin on a fracture surface is 

located at the surface of the component, then it is prudent to examine the 

exterior surface of the part as well as the fracture surface. Varner44 has 

compared features of chill checks and impact origins. Chill checks are curved 

and not flat, often have no visible outer surface damage, and, once the flaw is 

exposed on a fracture surface, have few markings on the flaw itself other than 

some twist hackle. Impact cracks are semicircular and are straighter, and the 

outer surface has chipping and lateral cracks. Once exposed on a fracture 

surface, the impact flaw exhibits tertiary Wallner lines. 

Polishing flaws from a loose abrasive rolled across a surface while under load 

can leave a variety of damage including indentations, scratches, Hertzian full or 

partial ring cracks and gouges according to Fréchette.40 The crack depths are 

typically four to five times larger than the particle diameter. They often are 

very elongated along the surface. Polishing, lapping or grinding glass can 

create are sleeks which are fine, scratch like, smooth-boundary imperfections 

usually caused by a foreign particle during polishing.45 

Chatter marks such as shown in Figures 6.58 and 6.59 often are clues to contact 

cracks underneath. The periodicity is a manifestation of vibration or a stick6-65 
 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

slip interaction when a contacting object slides over the surface. Fracture often 

starts from the first or last chatter crack in a series. Abrasion tracks, which 

often occur around the bottom rim or sides of a glass vessel, are also vulnerable 

sites for localized contact damage as shown in Figure 6.58. Polishing, lapping 

or grinding glass can create miniature crack patterns of chatter cracks. 

Preston45 and Ghering and Turnbull46 showed numerous examples and 

correlated them with the forces, the surface lubrication conditions, and the 

abrasive or contactor types. Preston showed that a blunt contactor or ball 

dragged across a surface creates small partial cones cracks behind the 

contactor. The cracks are created by tensile stresses behind the ball or 

contactor as it moves and the cracks are bowed in the direction of the contactor 

movement such as shown in Figure 6.36. In sharp contrast, a rolling ball 

creates arc cracks that bow in the opposite direction. They arch towards the 

direction from which the ball came.45 

(a) (b) 

Figure 6.58 Chatter cracks from an object that contacted and slid across the 

surface of a heat-strengthened soda lime glass kitchenware. (a) shows the outer 

surface. The black arrow points to the fracture surface and the origin site. (b) 

shows the specimen tilted to show the origin crack on the fracture surface 

(large black arrow) and the other chatter cracks on the outer surface (white 

arrows). 

6-66 
 

 

 

 

Origins of Fracture 

(a) 

(b) 

Figure 6.59 Very tiny chatter cracks from a small object that contacted and 

slid across the surface of a heat-strengthened soda lime glass kitchenware. (a) 

shows the outer surface with an assortment of common scratches and edge 

chips, some of which are secondary edge chips after the main fracture. The 

arrow points to the exact origin site which is seen to be one of a series of 

curved chatter cracks. (b) shows the fracture surface. The multiple secondary 

edge chips may confuse interpretation, but the Wallner lines and a faint pop in 

arrest line are centered on the chatter crack (arrow). 

6-67 
 

 

 

 Fractography of Ceramics and Glasses 

(a)

 (b) 

Figure 6.60 Surface abrasion can create subsurface cracks. (a) shows a piece 

from a broken heat-strengthened kitchenware tilted back to show the fracture 

surface and the elongated mirror and the outer surface with the abrasion track 

(large black arrow). Incidental secondary handling edge chips on the left 

(smaller black arrow) came from handling after fracture. (b) is the fracture 

surface showing an elongated mirror and the origin (white arrow). The origin 

region is dark due to a slight curvature of the fracture mirror surface. 

6-68 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

Origins of Fracture 

6.9 Fracture Oddities 

Innumerable other irregularities can act as strength-limiting flaws. The 

fractographer should expect the unexpected. Rice47 described cases of ball 

mills, thermocouple beads, insect carcasses and feces, dandruff, and hair as 

flaws. Alcock and Riley48 analyzed dust particles that are found in ceramic 

processing environments. Figure 3.19 showed a hair or piece of lint, a 

surprisingly common fracture origin type in laboratory-prepared cast resin 

composite specimens. Figure 6.61 shows some other unusual examples. 

The author once had a surprising experience with very bright retroreflective 

flaws when using a stereoptical microscope with dental highly-filled resinmatrix bend bars. Slight adjustments to the lighting angle eliminated the intense 

reflections. Some of the fracture origins were glassy inclusions or 

agglomerates of glassy filler particles that strongly reflected the incoming light. 

This was just as disconcerting and surprising as when one shines a flashlight at 

the retro-reflective eyes of a cat. Retroreflective glass beads are used with 

roadway marker paints to improve their visibility. Regrettably, the author did 

not take photos at the time of examination. 

(a) (b) 

Figure 6.61 Unusual fractures. (a) shows a vein of silicon that was the 

strength limiting flaw in a siliconized SiC bend bar. This probably was caused 

by silicon infiltration of a seam or crack in the green body. (b) shows the 

fracture surface of a silicon nitride that had an unstable yttria silicate 

boundary phase. Phase changes distorted the bar shape from its original 

rectangular shape. Fracture occurred from the crack marked by arrows. 

T and C denote the tension and compression sides of the bend bar. 

6-69 
 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

5 μm 

(a) (b) 

ceramic 

clay 

ceramic 

clay 

(c) (d) 

(e) (f) 

Figure 6.62 Common contaminants. (a) shows a blue clay smeared on a 

glass microscope slide as photographed on an optical microscope. The tiny 

particles and grease smear are evident. (b) shows an SEM image of the clay. 

(c) shows a fracture surface of an fine-grained alumina bend bar with clay 

swiped across the lower left corner (arrows). The boxed area is shown 

enlarged in (d). (e) shows human skin (courtesy A. Pasto). (f) shows an 

indeterminate contamination on a sintered silicon carbide bend bar. 

6-70 
 

 

 

 

 

 

 

 

 

 

 

Origins of Fracture 

6.10 Contaminants 

Every care should be taken to avoid contaminating the fracture surface. This is 

especially true during preparation of the specimens for SEM examination. 

During an optical examination, cleaning is not too difficult, but it is a different 

story if the specimen is being examined in the SEM. Once a specimen has been 

gold coated and is in the SEM, it can be disappointing to realize the flaw is 

masked by contamination or that a photo of the origin has distracting debris in 

the vicinity. Extraction, recleaning, and preparation are time consuming and 

frustrating. Sometimes contaminants can be recognized if they appear to sit up 

on top of the fracture surface and are obviously not part of the material. This 

criterion is by no means perfect, however, as many contaminants get into the 

fracture surface crevices and blend in diabolically well. As noted previously in 

Chapter 3, modeling clays are especially pernicious since they are fine, blend in 

well on the fracture surface, adhere extremely well, and are next to impossible 

to clean off. An EDX analysis of clay contamination on a fracture surface 

revealed aluminum, silicon, and titanium. Figures 6.63 – 6.64 illustrates some 

common contaminants and how they look in the SEM. 

(a) (b) 

ceramic veneer 

origin 

veneer ceramic 

veneer 

ceramic 

Plaque over outer surface 

origin 

clean outer surface 

origin 

(c) (d) 

Figure 6.63 Dental plaque is a common contaminant on ceramic restorations. 

(a) shows the broken middle (“pontic”) piece of a lithium disilicate bridge that 

broke after 7 weeks in vivo. (b) shows the overall fracture surface. (c) and (d) 

show the origin site before and after cleaning to removal plaque. Plaque can 

be removed with bleach as described on page 3-6. See Figure 6.74 for 

additional images of this bridge. 

6-71 
 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

 (a) (b) 

(c) (d) 

(e) (f) 

Figure 6.64 Contaminants in 3Y-TZP bend bars. (a) and (b) show a clean 

fracture with a well-defined fracture mirror and the common flaw in this 

material: a pore or porous region origin. (c) and (d) show clay smeared over 

the fracture surface and fracture mirror in another specimen. Notice the 

parallel smear traces (arrows) from lower left to upper right in (d). The clay is 

not evident in the low magnification image of the mirror, but it covers the 

origin in (d). It is possible the wrong end of the bend bar initially was put into 

mounting clay! (e) and (f) show a mirror with an unknown contaminant 

covering the origin (arrows). (all courtesy of J. Quinn) 

6-72 
 

 

 

 

 

 

 

Origins of Fracture 

6.11 Combined or Hybrid Flaws 

Sometimes flaws may link up or there simply may be side-by-side flaws at a 

fracture origin such as shown in Figure 6.65 and 6.66. When such mixed cases 

arise, some judgment is required as to which origin is primary. ASTM C 

132249 suggests several options for reporting such hybrid flaws. For example, 

if one origin type is primary, then a reporting code may list it first with the 

secondary flaw separated by an ampersand (&). For example, Figures 6.66a, b 

show P & LG denoting the origin is primarily a pore, but with one or more 

associated large grains. 

In other cases two small flaws can link to form a strength limiting flaw. A plus 

sign may be used to indicate that these origin types linked together to limit the 

strength of the ceramic. For example, P + MD indicates a pore combined with 

machining damage. Origins can be difficult to characterize if they have mixed 

attributes. For example, porous regions often have large pores associated with 

them (Figure 6.66e). References 50, 51, and 52 also discuss flaw linking. 

Figure 6.65 Flaws may link up. 

6-73 
 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) 

(b) (c) 

(d) (e) 

Figure 6.66 Combined (hybrid) flaw fracture origins (a) shows a flaw in a 

sintered 99.9 % alumina that can be described as primarily a pore, but with 

some large grains (P & LG). (b) is a pore/large grain (P & LG) combination 

in a sintered alpha silicon carbide bend bar. Note the good quality of the 

chamfer. Edge damage was not a factor in this case. (c) is a combination of 

an agglomerate and a large grain (A & LG) in sintered silicon carbide. (d) is 

a surface grinding crack that linked with a pore in a sintered silicon nitride (P 

+ MD). (e) shows a flaw in a sintered alumina could be described as either a 

porous region or a pore. 

6-74 
 

 

 

 

 

Origins of Fracture 

6.12 Baseline Microstructure Origins 

As fabricators improve materials by careful process control and eliminate large 

abnormal flaws, ceramics will break from origins that are part of the ordinary 

microstructure. Figure 6.67 illustrates the concept and Figure 6.68 shows an 

example. Once the inclusions are eliminated, grains from the large end of the 

grain size distribution become the strength-limiting flaws. The materials 

processor does not believe these are flaws, but the fracture mechanics expert 

does. Rice47 described such origins as “mainstream microstructural features.” 

A shorter descriptor might be “baseline microstructure” as in: “the origin flaw 

was a large grain from the “baseline microstructure.” 

Figure 6.67 Careful processing may eliminate severe flaws such as inclusions 

or gross pores. In the absence of such flaws, the material may fail from 

“baseline microstructure” origins. 

(a) (b) 

Figure 6.68 Baseline microstructure flaw in a sintered aluminum oxynitride 

four-point flexure bar. (a) shows a single grain that was the fracture origin. 

The origin in this case may actually be the grain boundary around the grain. 

(b) shows a polished section view of the microstructure. 

6-75 
 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

6.13 Geometric Sharp Points or Cusps 

Sharp points are extremely vulnerable to crack pop-in at very low stresses 

during fabrication or subsequent handling. Once the part is put under stress, 

these pop-in cracks can propagate to fracture. There may not necessarily be a 

material flaw as a starter. Sometimes an arrest line will be detected at the end 

of the pop-in crack, but often the final plane of fracture is identical to the initial 

pop-in crack plane, and the latter will not stand out. Figures 6.69a-d show 

exceptionally thin cusps between etch pits in theta strength specimens.

53 

(a) (b) 

(c) (d) 

Figure 6.69 Small cracks may pop-in from sharp points or cusps. The insert 

of (a) shows a miniature (300 μm) silicon theta strength-test specimen. A force 

applied to the top of the ring causes the middle web to stretch and break. 

(a) shows the fracture mirror around the surface pits fracture origin. The pits 

are caused by the deep reactive ion-etched (DRIE) processing. Sharp cusps 

between the pits are an easy path for tiny micro-cracks to pop-in and link. 

(c) and (d) show another theta specimen. The insert in (c) shows the fragment 

and the origin location. The hackle lines in (d) indicate fracture started at a 

single sharp cusp between two pits. The cusp is less than 50 nm in size and 

may be the smallest fracture origin ever recorded. 

6-76 
 

 

 

 

 

 

 

 

 

Origins of Fracture 

Porcelain veneer Alumina core 

origin 

Too 

thin ! 

(a) (b) 

Figure 6.70 A “feather tip” margin in an alumina dental crown that broke in 

a patient’s mouth (“clinical 22”). The margin is only 30 μm wide and was 

extremely vulnerable! Hoop stresses caused fracture from some of the tiny 

flaws at the margin tip. (crown courtesy of M. Øilo) 

Figures 6.69 shows a sharp “feather-tip margin” origin in a dental crown. The 

margin of this crown is much too thin. This type of margin is used by some 

dentists since less reduction of the tooth is needed and the crown blends into 

the gum line very well. Such practice is unwise for all-ceramic crowns. Figure 

10.25g shows another crown with a margin pop-in crack at a feather tip. 

Not all sharp point/cusp origins are in small pieces. The author once found 

such an origin in a ~ 1 m large SiC honeycomb structure for a space mirror. 

6.14 Partially-Exposed Flaws 

Flaws may not necessarily be oriented so that they are perpendicular to the 

principle tensile stress. Only a portion of a flaw may be exposed on a fracture 

surface as shown in Figure 6.71. The cone cracks of Figure 6.32 are examples. 

tensile surface 

fracture surface 

Figure 6.71 The flaw may not be fully-exposed on the fracture surface. 

6-77 
 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

6.15 Glaze Flaws 

Glazes and veneers are glassy and have low fracture toughness and resistance 

to fracture. A crack can form in the vulnerable glaze and then propagate into 

the remainder of the component. Figures 6.72 – 6.74 show some examples for 

three clinical dental fractures. Excessive mismatch in the thermal expansion 

coefficients of the glaze and the body can cause cracking, such as shown in Fig. 

6.72. Glaze or veneer cracking formed by tension is called “crazing.” Sometimes excessive porosity can occur at a boundary, and can be the strengthlimiting flaw as shown in Figure 6.73. 

(a) (b) 

Figure 6.72 Glazes can be preferred sites of fracture initiation. This was an 

in vivo broken three-unit zirconia bridge (courtesy T. Kosmac) with a crazed 

veneer probably from a mismatch of the veneer type to the ceramic. Fracture 

started on the plane on the left insert marked by the yellow arrow. 

(a) (b) 

buccal 

palatal 

zirconia 

core 

Veneer 

margin 

origin 

200 μm 

Figure 6.73 Glaze interfaces can be preferred sites of fracture initiation. This 

is an in vivo broken six-unit zirconia bridge with a poor design. The origin is 

the pore cluster at the interface between the veneer and core apex. (images 

courtesy U. Lohbauer and S. Scherrer) 

6-78 
 

 

 

 

 

 

 

 

 

Origins of Fracture 

Excessive compression in a glaze can cause cracking called “shivering.” 

Buckling of the outer layer fragments, often at an edge or rim, can cause a 

piece to detach. 

Krohn et al.54 showed pore and pigment aggregate fracture origins in enamel 

layers that were applied to float glass. Kobayahsi et al.55 showed excellent 

images of fracture origins that were a series of bubbles at the boundaries 

between an alumina porcelain and a glaze. 

some 

plaque here 

(a) (b) 

plaque was 

removed 

(c) (d) 

Figure 6.74 A gold-coated lithium disilicate glass ceramic bridge pontic unit 

(courtesy S. Scherrer). It broke in vivo on the palatal side of a connector. The 

origin was a sharp contact damage crack in the vulnerable veneer. Notice the 

arrest lines marked by arrows in (c), the helpful wake hackle lines from the 

bubbles in the veneer seen in (c) and (d), and even the concentric tertiary 

Wallner lines typical of impact in (d). See also Figure 6.63. 

6-79 
 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

6.16 Microfault Pockets 

Even if a fracture origin site can be found, there may not be an obvious flaw at 

the origin. Subtle porous regions or grinding damage crack at the surface can 

possibly be present but not detectable. This dilemma is more likely to occur 

with coarse-grained ceramics, porous ceramics, and heterogeneous ceramics, 

wherein a flaw could “hide” against the background microstructure. 

There is another possibility why a discrete flaw may not be detectable at an 

origin site: There may not be one. Emphasis in this chapter up to this point has 

been on flaws that are discrete irregularities larger than the baseline 

microstructure. Alternatives are illustrated in Figure 6.75. Fracture origins 

(a) (b) 

(c) 

Figure 6.75 Schematics of “microfault pockets.” The fracture origin may be 

an array of tiny microstructural faults that are linked by microcracking. 

Tension stress is applied vertically. (a) is similar to a schematic by Sines and 

Okada.56,57 (b) shows an origin site that has a greater local concentration of 

the irregularities. (c) shows random pockets of greater than average 

concentration. Fracture may entail linkage of these pockets by microcracking. 

6-80 
 

 

 

 

 

 

Origins of Fracture 

(a) 

(b) 

(c) 

Figure 6.76 Origin in a magnesia partially-stabilized zirconia split-rod bend 

specimen. Coarse hackle lines lead back to the origin. (a) shows the entire 

fracture surface which is very rough and undulating due to the coarse 

microstructure and irregularities in the microstructure. (b) and (c) show closeups of the origin which is a pocket of micropores and grain boundary faults. 

Tiny wake and twist microhackle radiate away from the origin. The next 

figure shows a polished section of the same material. 

6-81 
 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Figure 6.77 Polished section of the magnesia partially-stabilized zirconia 

shown in the previous figure. Microflaw pockets are marked by the arrows. 

may be regions of slightly or moderately greater concentration of ordinary 

irregularities such as micropores, grain boundary cracks or other tiny faults. 

The regions may be termed “microfault pockets.” By themselves, the tiny 

irregularities are insufficient to case fracture, but they can link by 

microcracking during loading. Sines and Okada56,57 suggested that this could 

account for time-dependent fracture, but there is no reason it cannot apply to 

fast fracture strength as well. It should be borne in mind that a random 

distribution does not mean a uniform distribution. A random distribution of 

such tiny irregularities will have local regions of greater-than-average 

concentration. 

Fracture origins of the types shown in Figure 6.75 are extremely difficult to 

diagnose on fracture surfaces, since the features at the origin look exactly like 

the normal microstructure. Figure 6.76 is an example. Polished specimen 

microstructural analysis helped with the interpretation as shown in Figure 6.77 

The origin may take on the character of porous regions as shown in Figure 

6.12. The literature is sparse on such flaws, no doubt because engineers and 

scientists are loath to show fractographs of origins that fail to show a distinct 

flaw! 

6.17 Polished Microstructural Sections 

Figure 6.78 and previous Figures 6.4, 6.67, and 6.77 show examples of how 

polished microstructural sections can provide complimentary information about 

flaw types and their distributions. Polished section preparation is a staple for 

metals failure analysis, usually to ensure that the microstructure is as expected 

and the correct phases are present. 

6-82 
 

 

 

 

Origins of Fracture 

(a) (b) 

(c) 

(d) 

Figure 6.78 Polished microstructural sections can aid origin interpretation. 

(a) Sintered alpha silicon carbide often has large grain fracture origins. The 

acicular or platelet shape of these and their concentration are revealed by the 

micrograph. (b) reaction-bonded silicon nitride sometimes fails from 

unreacted silicon (arrow). (c) and (d) show irregular pressing flaws on 

fracture surfaces and polished sections of a sintered alumina, respectively. (c 

and d courtesy T. Bernthaler) 

6-83 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

6.18 Flaw Size 

It may suffice to identify the nature of the flaw, but a better characterization 

includes at least a rough estimate of its size. A 100 μm inclusion may not have 

the same source or behavior as a 20 μm inclusion. The size need not be 

measured precisely if the characterization is intended to describe the general 

nature of the origin. On the other hand, there are some instances when a good 

estimate is needed, such as when it is intended to do a fracture mechanics 

analysis of the flaw as described in the next chapter. This may be to confirm 

that the fractographer has found the correct flaw. So, for example, if a 

fractographer has tentatively identified a 2 μm feature as an origin, but fracture 

mechanics suggests the flaw should be 50 μm in size, then the fractographer 

ought to reevaluate his assessment. 

Measure an approximate flaw diameter, or if the flaw is elongated, approximate 

it with a semi-elliptical shape and measure the minor and major axis lengths as 

shown in Figure 6.2. The origin’s true size may not be revealed on the fracture 

surface. Exact fracture mechanics analyses of most origins are not possible due 

to their complex shape. An important exception to this is machining damage, 

wherein the origin size measurement may be very useful for the estimation of 

fracture toughness. When measuring the flaw size of surface-located flaws, 

measure and record the origin depth and the width. Chapter 7 and Figures 7.18 

and 7.19 furnish more guidance. 

6.19 What to Report 

For each component, or laboratory-tested specimen, record the origin type, the 

location in the particular specimen, and the flaw size. If a fracture mirror is 

present, it may also be helpful to measure its size, especially if the fracture is in 

a component. Mirror size measurements in laboratory strength test specimens 

also may be done so that mirror constants can be estimated and published for 

new materials. Figure 6.79 shows images that follow the recommendations in 

ASTM C 1322.49 

Organizing a wealth of fractographic images can at times be difficult. Ideally 

the images should be presented in a coherent fashion with some context. 

Section 5.10 and Figure 5.68 showed how montages can be useful in organizing 

a lot of local detail into a global map for a component failure analysis. Figure 

6.80 shows how a montage can be constructed for laboratory strength test data. 

6-84 
Origins of Fracture 

Figure 6.79 Multiple photos should be taken showing whole fracture surface, 

the mirror region and the origin. The right side shows an example for an 

inclusion in a silicon nitride rod flexural strength specimen. 

6-85 
 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Figure 6.80 Schematic of a fractographic montage. 

Component failure analyses often have some overall photographs of the entire 

structure and close-ups of various critical features. Much time and confusion 

can be saved if the successive images are presented in an orderly, systematic 

way. It is always wise to start with an overall image of the structure and then 

add images at progressively higher magnification at key locations. 

When one jumps from one magnification to a higher magnification, later it may 

be difficult to correlate the two images. In the author’s experience, 

magnification jumps of about 5 are the most one should use. 

At each step the relationship of the images to each other should be clear. 

Figure 5-68 in the previous chapter and Figure 6-81 on the next page show 

good examples for two dental crown failures. 

6-86 
 

 

 

 

 

Origins of Fracture 

(a) (b) 

(c) 

(d) 

2 mm 

dcp 

origin 

(e) 

Figure 6.81 Failure of a first upper-left molar crown. (a) shows a view onto 

the occlusal surface (the crown top) and the chipped porcelain veneer on the 

side of the crown (arrow). The porcelain veneer was inadequately supported 

by the alumina core as shown in the x-ray image (b). (c) – (e) show 

progressive close-ups of the fracture surface on an epoxy resin replica. 

This is the same crown as shown in Figure 3.34 on replica techniques. 

(All courtesy S. Scherrer, Univ. Geneva) 

6-87 
 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Chapter 6 References 

1.	 W. Weibull, “A Statistical Distribution Function of Wide Applicability,” J. Appl. 

Mech., 18 (1951) 293 - 297. 

2.	 G. D. Quinn and R. Morrell, “Design Data for Engineering Ceramics: A Review of 

the Flexure Test,” J. Am. Ceram. Soc., 74 [9] (1991) 2037 - 2066. 

3.	 R. W. Davidge, Mechanical Behavior of Ceramics, Cambridge Univ. Press, 

London, 1979. 

4.	 D. Munz and T. Fett, Ceramics, Mechanical Properties, Failure Behavior, 

Materials Selection, Springer, Berlin, 1999. 

5.	 W. E. C. Creyke, I. E. J. Sainsbury, and R. Morrell, Design with Non-ductile 

Materials, Applied Science Publ., London, 1982. 

6.	 R. W. Rice, “Pores as Fracture Origins in Ceramics,” J. Mat. Sci., 19 (1984) 895 -

914. 

7. G. D. Quinn, K. Hoffman, and J. B. Quinn, “Strength and Fracture Origins of a 

Feldspathic Porcelain,” Dent. Mat., 28 [5] (2012) 502–511. 

8.	 D. Lewis, III, “Fracture Strength and Mirror Size in a Commercial GlassCeramic,” J. Am. Ceram. Soc., 64 [2] (1981) 82-86. 

9 .	 T. Bernthaler, “Defect Imaging and Analysis for Predicting the Reliability of 

Ceramics,” PhD thesis, School of Materials Science and Engineering, University 

of New South Wales, Sydney, Australia and Aalen University, Aalen, Germany, 

April 2011. 

10.	 F. F. Lange, “Processing-Related Fracture Origins: I, Observations in Sintered and 

Isostatically Hot-Pressed Al2O3/ZrO2 Composites,” J. Am. Ceram. Soc., 66 [6] 

(1983) 396-398. 

11. F. F. Lange and M. Metcalf, “Processing-Related Fracture Origins: II, 

Agglomerate Motion and Cracklike Internal Surfaces Caused by Differential 

Sintering,” ibid, 398-406. 

12. F. F. Lange, B. I. Davis, and I. A. Aksay, “Processing-Related Fracture Origins: 

III, Differential Sintering of ZrO2 Agglomerates in Al2O3/ZrO2 Composite,” ibid, 

407-408. 

13.	 M. Hangl, “Processing Related Defects in Green and Sintered Ceramics,” pp. 133 

– 143 in Fractography of Ceramics and Glasses, IV, eds. J. R. Varner and G. D. 

Quinn, Ceramic Transactions, Vol. 122, American Ceramic Society, Westerville, 

OH, 2001. 

14.	 H. R. Baumgartner and D. W. Richerson, “Inclusion Effects on the Strength of Hot 

Pressed Si3N4,” pp. 367 – 386 in Fracture Mechanics of Ceramics, Vol. 1, eds., 

R.C. Bradt, D. P. H. Hasselman, and F. F. Lange, Plenum, NY, 1974. 

15.	 A. G. Evans, “Structural Reliability, A Processing Dependent Phenomenon,” J. 

Am. Ceram. Soc., 65 [3] (1982) 127 - 137. 

16. R. W. Rice, Mechanical Properties of Ceramics and Composites, Grain and 

Particle Effects, Marcel Dekker, NY, 2000. 

17.	 M. G. Gee, and R. Morrell, “Fracture Mechanics and Microstructures,” pp. 1 – 22 

in Fracture Mechanics of Ceramics, Vol. 8, R. Bradt, A. Evans, D. Hasselman, and 

F. Lange, eds., Plenum Press, NY, 1986. 

6-88 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Origins of Fracture 

18. E. Carlstrom and R. Carlsson, “A Fractographic Study of Electrical Porcelains 

Containing Alumina,” J De Physique, 47 [C-1] (1986) 563 – 567. 

19.	 D. Southan, “The Porcelain Jacket Crown,” pp. 207 -230 in Dental Ceramics, 

Proc. First International Symposium on Ceramics, ed. J. McLean, Quintessence 

Publ., London ,1983. 

20.	 J. S. Banda and P. F. Messer, “Fracture-Initiating Flaws in Whitewares,” pp. 363 – 

375 in Fractography of Glasses and Ceramics, Advances in Ceramics, Vol. 22. 

eds, J. R. Varner and V. D. Fréchette, 1988, American Ceramic Society, 

Westerville, OH, 1988. 

21.	 R. W. Davidge and G. Tappin, “The Effects of Temperature and Environment on 

the Strength of Two Polycrystalline Aluminas,” Proc. Br. Ceram. Soc., 15 (1970) 

47 – 60. 

22.	 J. A. Salem, “Transparent Armor Ceramics as Spacecraft Windows,” J. Am. 

Ceram. Soc., 96 [1] (2013) 281-289. 

23.	 J. J. Swab, R. Pavlacka, G. Gilde, S. Kilczewski, J. Wright, and D. Harris,” 

Determining the Strength of Coarse-Grained AlON and Spinel,” J. Am. Ceram. 

Soc., 97 [2] (2014) 592-600. 

24. D. W. Richerson, Modern Ceramic Engineering, Marcel Dekker Inc., NY, 1982; 

and 3rd edition, CRC Press, Taylor and Francis, Boca Raton, FL, 2006. 

25. C. G. Knight, M. V. Swain, and M. M. Chaudhri, “Impact of Small Steel Spheres 

on Glass Surfaces,” J. Mat. Sci., 12 (1977) 1573 – 1586. 

26.	 M. M. Chaudhri, “High-Speed Photographic Investigations of the Dynamic 

Localized loading of Some Oxide Glasses,” pp. 87 – 113 in Strength of Inorganic 

Glass, ed. C. Kurkjian, Plenum, New York, 1985. 

27.	 Y. Akimune, “Hertzian Cone Crack in SiC Caused by Spherical Particle Impact,” 

J. Mat. Sci. Lttrs., 9 (1990) 659 – 662. 

28.	 H. Kirchner and R. Gruver, “Branching of Hertz Cracks,” pp. 959-964 in Fracture 

1977, Vol. 3, Proceedings of the 4th Int. Conf. on Fracture, Waterloo, Canada, ed., 

D. Taplin, Univ. Waterloo Press (1977). 

29	 B. R. Lawn, S. M. Wiederhorn, and D. E. Roberts, “Effect of Sliding Friction 

Forces on the Strength of Brittle Materials,” J. Mat. Sci., 19 (1984) 2561 - 2569. 

30. M. M. Chaudhri and C. Liangyi, “The Orientation of the Hertzian Cone Crack in 

Soda-Lime Glass Formed by Oblique Dynamic and Quasi-static Loading with a 

Hard Sphere,” J. Mat. Sci., 24 (1989) 3441 – 3448. 

31.	 R. W. Rice and J. J. Mecholsky, Jr., “The Nature of Strength Controlling 

Machining Flaws in Ceramics,” pp. 351 – 378 in The Science of Ceramic 

Machining and Surface Finishing II, eds., B. J. Hockey and R. W. Rice, Special 

Publication 562, NBS, Gaithersburg, MD, 1979. 

32.	 G. D. Quinn, L. K, Ives, and S. Jahanmir, “On the Nature of Machining Cracks in 

Ground Ceramics: Part I: SRBSN Strengths and Fractographic Analysis,” Mach. 

Sci. Technol., 9 (2005) 169 - 210. 

33.	 G. D. Quinn, L. K, Ives, and S. Jahanmir, “On the Nature of Machining Cracks in 

Ground Ceramics: Part II: Comparison to Other Silicon Nitrides and Damage 

Maps,” Mach. Sci. Technol., 9 (2005) 211 - 237. 

6-89 
 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

34.	 W. Kanematsu, “Visualization of Subsurface Damage in Silicon Nitride from 

Grinding by a Plasma Etching and Dye Penetration Method,” J. Amer. Ceram. 

Soc., 89 [8] (2006) 2564-2570. 

35.	 B. A. Bender, C. Kim, and C. Cm Wu, “Solid Preform Fabrication of a 

Piezoelectric Ceramic Torsional Actuator Motor,” Ceram. Eng. Sic., Proc., 25 [3] 

(2004) 89 – 94. 

36.	 J. R. Varner and V. D. Fréchette, “Fractography of Whitewares,” pp. 305 – 315 in 

Science of Whitewares, Eds. V. Henkes, G. Y. Onoda, and W. M. Carty, American 

ceramic Society, Westerville, OH, 1996 

37. G. D. Quinn, K. Hoffman, S. Scherrer, U. Lohbauer, G. Amberger, M. Karl, J. R. 

Kelly, “Fractographic Analysis of Broken Ceramic Dental Restorations, pp. 161-

174 in, Fractography of Glasses and Ceramics, VI, Ceramic Transactions, 

Volume 230, eds., J. R. Varner, M. Wightman, Wiley, New York, 2012. 

38.	 G. D. Quinn, “Fractographic Analysis of Broken Ceramic Dental Restorations,” 

Ceram. Eng. Sci. Proc., 35 [1] (2015) 39-51. 

39. J. R. Varner, “The Practical Strength of Glass,” pp. 389 – 406 in Strength of 

Inorganic Glass, ed., C. Kurkjian, Plenum, NY, 1986. 

40. V. D. Fréchette, Failure Analysis of Brittle Materials, Advances in Ceramics, Vol. 

28, American Ceramic Society, Westerville, OH, 1990. 

41. H. Pfaender, Schott Guide to Glass, Chapman and Hall, London, 1996. 

42.	 F. Preston, “A Study of the Rupture of Glass,” J. Soc. Glass Techn., 10 (1926) 

234- 269. 

43. ASTM C 162, Standard Terminology of Glass and Glass Products, Annual Book of 

Standards, Vol. 15.02, ASTM Int., West Conshohocken, PA, 1999. 

44. J. Varner, private communication. 

45. F. W. Preston, “The Structure of Abraded Glass Surfaces,” Trans. Opt. Soc., 23 [3] 

(1921-1922) 141 – 164. 

46. L. G. Ghering and J. C. Turnbull, “Scratching of Glass by Metals,” Bull. Am. 

Ceram. Soc., 19 (1940) 290 – 294. 

47.	 R. W. Rice, “Failure Initiation in Ceramics: Challenges of NDE and Processing,” 

pp. 1057 - 1064 in Ceramic Developments, eds., C. Sorrell, and B. Ben-Nissan, 

Materials Science Forum, Vol. 34–36, Trans. Tech. Publ. Ltd., Switzerland, 1988. 

48.	 J. R. Alcock and F.L. Riley, “A n Investigation of Dust Particles Found in a 

Ceramic Processing Environment,” J. Eur. Ceram. Soc., 6 (1990) 339-350. 

49. ASTM C 1322-96, “Standard Practice for Fractography and Characterization of 

Fracture Origins in Advanced Ceramics,” Annual Brook of Standards, Vol. 15.01, 

ASTM Int., West Conshohocken, PA (1996) 

50.	 A. G. Evans and G. Tappin, “Effects of Microstructure on the Stress to Propagate 

Inherent Flaws,” Proc. Brit. Cer. Soc., 20 (1972) 275-297. 

51.	 R. W. Davidge, Mechanical Behavior of Ceramics, Cambridge University Press, 

Cambridge, UK, 1974, pp.98-100. 

52.	 R. W. Rice, “Machining of Ceramics,” pp. 287-343 in Ceramics for High 

Performance Applications, eds., J.J. Burke, A. E. Gorum, and R. N. Katz, Brook 

Hill Publ., Chestnut Hill, MA, 1974. 

6-90 
 

 

 

 

 

 

Origins of Fracture 

53.	 G. D. Quinn, “Fractographic Analysis of Miniature Theta Specimens,” Ceram. 

Eng. Sci. Proc., 29 [3] (2008) 189 – 199. 

54.	 M. H. Krohn, J. R. Hellmann, D. L. Shellemann, and C. Pantano, “Strength and 

Fatigue of Float Glass Before and After Enameling,” The Glass Researcher, 11 [2] 

(2002) 24 – 27. 

55.	 Y. Kobayashi, M. Mukai, T. Mizuno, O. Ohira, and H. Isoyama, “Effect of 

Cristobalite Formation and Glaze on Bending Strength of α-Alumina Reinforced 

Porcelain,” J. Ceram. Soc. Japan, 113 [6] (2005) 413 – 418. 

56. T. Okada, G. Sines, and D. Green, “Crack Origins and Microcracking in Delayed 

Fracture of Alumina,” Comm. Am. Ceram. Soc., May 1982, C 64 - 65. 

57.	 G. Sines and R. Okada, “Flaws Responsible for Slow Cracking in the Delayed 

Fracture of Alumina,” J. Am. Ceram. Soc., 66 [3] (1983) 228 - 232. {correction, 

ibid, 66 [7] (1983) 538.} 

6-91 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Quantitative Analysis 

7. Quantitative Analysis 

7.1 Introduction 

Quantitative analysis can aid the fractographic analysis and extend its value. It 

may be used to confirm that the right flaw has been identified as a fracture 

origin. It can make good estimates of the stresses in the part at failure. Four 

methods to estimate stress are shown in this chapter: fragmentation analysis, 

branching distance analysis, fracture mirror size analysis, and origin size 

analysis. Some analyses are empirical and others are based on the powerful 

tool of fracture mechanics. One goal of this chapter is to convey a working 

knowledge of fracture mechanics so that it can be used to solve practical 

problems. Table 1 includes some useful conversion factors. 

length

1 in * = 

1 micrometer (micron) = 

.001 inch * = 

force 

1 kgf (kilogram force)* = 

stress or pressure

1 ksi (1,000 psi) * 

1 Pa 

1 MPa 

1 MPa 

1 GPa 

stress intensity

1 ksi √inch * 

= 

= 

= 

= 

= 

25.4 mm 

1 x 10-6 m 

25.4 micrometers 

9.80 N (Newtons ) 

6.895 MPa (Megapascal) 

1 N/m2 

1 MN/m2 = 1 x 106 Pa 

145.04 psi = 0.14504 ksi * 

1000 MPa = 1 x 109 Pa 

example: 100 ksi = 690 MPa 

= 1.10 MPa√m 

 1ksi in  example: 5 MPa m x   = 4.55 ksi in   1.10 MPa m 

energy

1 erg = .0000001 (1 x 10-7) J (Joule) 

1 ft-lbf *(foot pound force) = 1.355 J 

energy/area

erg/cm2 = .001 J/m2 

Length on Photos

1 mm = 1 micrometer on a 1000X photo 

* Not an official SI unit, but included here for completeness. 

Table 7.1 Conversion Factors 

7-1 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

σ < 10 MPa (1,500 psi) 

σ > 10 MPa (1,500 psi) 

Figure 7.1 Schematic of thermal fractures of windows. Low stress fractures 

(σ < 10 MPa, 1,500 psi) do not cause branching. (After Orr2 and Fréchette1

) 

7.2 Stresses from the Extent of Fragmentation 

7.2.1 Annealed parts 

A general assessment of stress levels can be made from the fracture pattern. 

Low stressed parts break into only two pieces. Fréchette1 and Orr2 and have 

suggested that 10 MPa (1,500 psi) is a threshold stress level for branching and 

fragmentation in glass as illustrated in Figure 7.1. Additional fragmentation 

occurs with increasing stress in the part, or increasing stored elastic energy. 

Very highly-stressed parts fragment to such a degree that laypersons are apt to 

describe the fracture as an “explosion.” Retrieval of the two fragments with 

the fracture origin may be problematic. 

Fréchette and his students3

 have shown fragmentation correlations with internal 

pressure of bottles as shown in Figure 7.2. They did a similar analysis for 

impact velocity for impacted bottles.4 While their work may have been 

specific to certain bottle types and shapes, similar testing could be done with 

alternative shapes and sizes to aid in the failure analysis of new service failures. 

Figure 7.3 shows how the fragmentation of ring-on-ring (32 mm and 64 mm 

diameters) tested borosilicate crown glass disks (76 mm diameter x 5.4 mm 

thick) can be correlated to fracture stress. This work was done as part of a 

reliability analysis for a pressurized aircraft window.5 Figures 4.12 and 4.13 

7-2 
 

 

 

 

 

 

 

 

 

Quantitative Analysis 

showed examples of such test specimens and their breakage patterns. Fragmentation in Figure 7.3 was quantified by the number of radial cracks that reached 

the outer rim. This was an arbitrary criterion. Morrell showed a very similar 

graph for the number of fragments as a function of stress in ring-on-ring tests of 

high purity alumina.6 Shand presented a similar graph counting radial cracks 

for 6 cm square window glass plates broken by impact in the middle.7 He 

pointed out that by such counting, one could obtain a stress at failure estimate 

even if the impacted site was lost or destroyed. More recently, Wang8

 did 

stress analyses and counts of the number of pieces in glass plate coupons of 

different sizes when pressure loaded to fracture. Graphs such as these are very 

specific to specimen or component size, shape, and material. 

Stress (MPa) 

0 50 100 150 200 250 300 350 

60 

50 

1.9 L (64 oz) 

40 

1.4 L (48 oz) 

30 

.95 L (32 oz) 

20 

.36 L (12 oz) 

10 

0 1000 2000 3000 4000 5000 

Stress (psi) 

Figure 7.2 Fragmentation of internally pressurized bottles. The number of 

branches (fully developed that extend more than 13 mm beyond the first fork) 

are shown versus the hoop stress in the glass wall for four different glass 

containers. Data points and a solid line are shown only for the 1.4 L (48 oz) 

bottles. Dotted lines show the trends for the three other sizes. Some of the 

scatter is due to variations in the glass wall thickness. (After Fréchette and 

Michalske, Ref. 3) 

7-3 
 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Figure 7.3 was actually used to analyze an unexpected fracture. A similar glass 

disk was impacted by a hailstone that was intended to weaken the disk by 

creating surface damage that could be assessed in a subsequent strength test. 

The hailstone unexpectedly fractured the disk. The fragments were recovered 

and reassembled and the origin determined to be an ordinary polishing scratch 

on the surface opposite the hail impacted side. The hailstone diameter (25 mm) 

was similar to the inner ring diameter in the strength test set. From the number 

of radial cracks, the stress at failure was estimated. The estimate matched the 

dynamic stress estimated from an impact stress analysis, thereby confirming 

that the impact indeed had imparted sufficient energy to flex the disk to fracture 

from the ordinary surface flaw. N, Number of Radial Cracks 

110 

100 

90 

80 

70 

60 

50 

40 

30 

20 

10 

0 

N = 0.74 (Stress in ksi)1.43 

N = 0.047 (Stress in MPa)1.43 

0 5 10 15 20 25 

Inner Ring Stress (ksi) 

0 20 40 60 80 100 120 140 160 180 200 

Inner Ring Stress (MPa) 

Figure 7.3 The number of radial cracks that reach the outer rim as a function 

of fracture stress for 76 mm diameter by 5.4 mm thick borosilicate crown glass 

disk broken in ring-on-ring flexure. Data for this graph included only fractures 

that initiated inside the inner loading ring. 

7-4 

30 
 

 

 

 

 

 

 

Quantitative Analysis 

More often, the fractographer simply makes a qualitative stress assessment with 

a cursory examination. This is often sufficient for a diagnosis. For example, if 

a fracture occurs at an unexpectedly low stress, the fractographer might search 

for an unexpectedly large flaw in the component or specimen. A handling or 

preparation flaw may have weakened it to an unexpected level. 

7.2.2 Fragmentation of tempered parts 

McMaster et al.9 state that the most common method of inferring stress levels 

in tempered glass is by breaking it. They and Shinkai10 cite work by 

Akeyoshi,

11 wherein 10 cm square tempered plates of various thickness were 

broken by sharp impact. The number of fragments in a 5 cm square region 

correlated to the internal tension temper stress as shown in Figure 7.4. 

Figure 7.4 The interior tension stress in tempered soda lime glass plates may 

be estimated from the fragment density when the plate is broken. Each line 

corresponds to the plate thickness shown. (After Akeyoshi et al., Ref. 11) 

7-5 
 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Figure 7.5 The interior tension stress in tempered soda lime sheet and plate 

glass may be estimated from the average fragment mass. Center tension stress 

is plotted versus mass normalized by the plate thickness, t. Data for five 

different thickness plates and sheets from 3.2 mm to 9.5 mm (1/8 in to 3/8 in) 

fall on the same curve. (After Barsom, ref. 13) 

For a particular thickness plate, the data suggests there is a minimum temper 

stress to cause dicing. Blank12 presented some interesting fragmentation work 

on tempered soda lime silica plates, 300 mm by 300 mm in size and with 

thicknesses varying from 3.2 to 7.2 mm. The average fragment size (area) in 

mm2 dramatically decreased with increasing surface compression 

strengthening. 

Of course, other means of assessing temper stresses are available, including 

polariscope inspections and use of the stress optical coefficient, but in this 

Guide we are concerned with fractographic techniques. 

Barsom13 showed that interior tension stress σct was related to the fragment 

mass: 

1  m  4 

σ   = constant (7.1) ct  t 

where m is the fragment weight (mass), t is the plate thickness. The constant 

has units of stress • (mass/thickness)1/4. 

7-6 
 

 

 

 

 

 

 

 

 

 

 

 

 

Quantitative Analysis 

It should be borne in mind that these relationships and trends are altered if there 

are significant externally applied stresses superimposed on the temper stresses. 

The fractographer should be on the lookout for irregular breakage patterns and 

fragment sizes that are signs of nonuniform temper stresses. Fréchette1 showed 

two examples. In one case a glass cylinder was not tempered uniformly during 

the fabrication process (his case 10.18). A second case (his 10.22) featured a 

high-powered lamp that cracked spontaneously as a result of heating in service. 

The fragments were atypically large in the fracture origin area, but elsewhere 

the normal dicing fragmentation pattern was observed. Fréchette concluded 

that the intense in-service heating caused the middle of the lamp to lose its 

temper. Blank12 discussed the use of tempered glass in front automobile 

windshields in Germany and showed a fascinating drawing of a fractured 

windshield that had dramatically uneven fragment sizes and crack patterns due 

to nonuniform heat transfer during tempering. It is commonly assumed that the 

residual stresses are symmetric through the thickness as shown in Figure 5.19, 

but Menčík14 illustrated the effects of asymmetric tempering of plates on the 

final stress distributions. Different residual compression stresses could exist on 

the two plate surfaces. There may even be a net tensile stress on one surface. 

7.3 Stresses from the Branching Distances 

7.3.1 Annealed plates and bars 

The stress at fracture is related to the branching distance away from the origin 

by: 

σ R (7.2) b = Ab 

where σ is the stress, Rb is the branching radius, and Ab is a constant called the 

fracture branching constant. It has units of stress⋅√length. These units are the 

same as for stress intensity which is described in section 7.5. The relationship 

between Ab and fracture toughness is discussed in 7.6. A similar relationship 

applies to mirror sizes as will be discussed in section 7.4. Rearranging shows 

that the branching distance is inversely dependent upon the square of the stress: 

A 2 

R = 

 b  (7.3) b  σ 

This empirical relationship has been found to apply to uniaxially-stressed parts 

such as a tensile strength specimen or a component with tensile stresses 

primarily in one direction. The relationship also can be applied to bend bars or 

7-7 
 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

components loaded in bending provided that the branching distance is 

measured along the surface where the tensile stresses are a maximum. The 

relationship also may apply to biaxially-stressed parts such as a ring-on-ring 

strength test disk or pressurized windows or bottles, again using matching 

lengths and stresses. 

The term Ab may not necessarily be a material constant. It certainly does 

depend upon the material properties such as elastic modulus or fracture 

toughness, but there is evidence that stress state and component geometry affect 

it too. Ab is always greater than the fracture mirror constant that is described 

later in this chapter. Branching distances in biaxially-stressed parts may be less 

than the branching distances in uniaxially-stressed parts. Shetty et al. 15 showed 

that branching can occur just beyond the mirror in highly stressed, biaxially 

loaded disks. Hull16 suggested that the cross section thickness of a component 

might influence when branching occurs. Despite these concerns, reported Ab 

values are tabulated in Appendix C. Very few values are available since most 

investigators have focused on fracture mirror constants instead. 

The branching angle can at least be a qualitative indicator of the stress state as 

shown in Figure 4.4 in chapter 4. The branch angle is actually variable and 

curves at the exact branching location. It is recommended that the branching 

angle be measured close to but somewhat away from the splitting juncture, at a 

point where the branching angle has settled into a more or less constant value. 

7.3.2 Tempered Plates 

Figure 7.6 shows an interesting approach suggested by Soltesz et al.

17 to 

estimate temper stresses in tempered glass plates. The lengths of the first and 

second branches can be correlated to internal temper stresses. 

7-8 
 

 

 

 

 

 

 

Quantitative Analysis 

(a) 

(b) 

Figure 7.6 Branch distances in tempered glass may be used to estimate 

internal tension stresses. (a) shows an example from a 3.3 mm thick tempered 

automobile window. The initial branch - crack run is 2a = 2.6 mm. The 

lengths of the first four branches (not including the segments that ran straight 

from the initial crack) are also labeled. (b) shows the initial branch distance 

2a and the average first branch distance ai,avg versus the center tension stress 

respectively for 6 mm to 10 mm thick plates. All dimensions are in mm. Limits 

to the trends for different thickness plates are shown by the dashed vertical 

lines. Data to the right of these limits at low stresses has high scatter. (b is 

after Soltesz et al., Ref. 17) 

7-9 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

7.4 Stresses from the Mirror Size 

7.4.1 Mirror analysis and constants 

The fracture mirror size is related to the stress at the origin. Orr’s empirical 

mirror size relationship is: 

R = A σ (7.4) 

where σ is the tensile stress at the origin at the instant of fracture, R is the 

mirror radius, and A is the “mirror constant” with units of stress intensity 

(MPa√m or ksi√in) and is considered by most as a material property. (A 

version of this equation with a residual stress term is discussed later in this 

chapter.) This relationship is similar to, but not quite the same, as that for 

stress intensity or fracture toughness as will be discussed in section 7.6 below. 

Figure 7.7 shows more details of a mirror centered on a surface flaw. This 

rendition also depicts the prospect that the initial flaw might grow somewhat 

from a to ac before going critical. 

This author reviewed the evolution of fracture mirror size analysis in a 2006 

paper.18 Leighton Orr of the Pittsburg Plate Glass (PPG) Company was the 

first to systematically use equation 7.4 with the square root dependence of 

mirror size to solve glass fracture problems as far back as the 1940s.2,19,20 

ac 

a 

Ri 

Rb 

Ro 

Ab Ao Ai KIc 

Figure 7.7 Fracture mirror schematic. Ri is the inner or mirror-mist 

boundary, Ro is the mist-hackle boundary, and Rb is the branching distance. 

The fracture toughness KIc (discussed in section 7.5.1), the mirror constants Ai 

and Ao and the branching constant Ab correspond to each of these boundaries. 

All have units of stress √length. 

7-10 
 

 

 

 

 

 

 

 

 

 

 

 

 

Quantitative Analysis 

Johnson and Holloway in 196621 offered a physical interpretation of the 

relationship based on the energetics of crack propagation. Levengood,

20 

Shand,22 Kerper and Scuderi,23,24 and later Kirchner et al.25 gave credit to Orr 

for equation 7.4. Orr did not publish his findings until he retired in 1972,2 but 

discussed his empirical equation at a Glass Division meeting of the American 

Ceramic Society in Bedford, PA in September 1955.26 Orr told me in 200119 

that equation 2 had been extensively used for glass fracture analysis at PPG in 

the 1940s. In the 1950s and 1960s, many people used a general form of 

equation 7.4 with the mirror size raised to an arbitrary power. Shand’s 1954 

and 1959 papers22,27 presented an early argument for the 1/2 power based on 

stress concentrations at the tip of a sharp crack. These two papers incorporated 

elements of what become known later as fracture mechanics. Kerper and 

Scuderi23,24 at the National Bureau of Standards in Washington performed 

meticulous experiments on hundreds of glass laths and rods and showed 

conclusive evidence that Orr’s equation was applicable over a broad range of 

specimen and mirror sizes. (See Fig. 7.10 later in this chapter) Some have 

credited Terao28 with equation 7.4, but a careful reading of his paper fails to 

back the claim since it shows an exponential form relating mirror size to stress 

instead. So, although many associate equation 7.4 with Johnson and Holloway 

21 in 1966, the relationship had already been in use by Orr for 20 years. 

The practical significance of equation 7.4 is that one can calculate the origin 

stress if a fracture mirror size is measured and a fracture mirror constant is 

available. The calculated stress is the net tensile stress acting on the flaw and 

the region around the flaw. It may include several stress sources including 

mechanical, thermal, and residual stresses. It is possible to discern separate 

mirror-mist and mist-hackle boundaries in glasses and each has a corresponding 

mirror constant A. The most common notation is to refer to the mirror-mist 

boundary as the inner mirror boundary and the corresponding mirror constant is 

designated Ai. The mist-hackle boundary is referred to as the outer mirror and 

the corresponding mirror constant is designated Ao. The mirror-mist boundary 

is usually not perceivable in polycrystalline ceramics and only the mirrorhackle boundary is usually measured. 

Equation 7.4 is valid even for large mirrors if loading is in direct tension. If a 

part is in flexural loading, deviations from the trend can occur as discussed in 

Appendix D. 

At this point it is appropriate to reiterate that equation 7.4 is empirical. In 

chapter 5, it was noted that there were competing models to account for the 

mirror formation. Kirchner and colleagues29,30 presented compelling evidence 

that a fracture mechanics criterion based on a critical stress intensity, KIB, 

7-11 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

(described later in section 7.5) gives the best fit to data and the exact shape of 

mirrors in various stress fields. A crucial observation is that the stress intensity 

model fully accounts for the small inward facing cusp at specimen surfaces 

(Figures 5.2 and 7.7), whereas the other models do not. Kirchner et al. 29,30 

argued that a more fundamental material parameter might be KIB, the stress 

intensity factor at branching, rather than the A’s in equation 7.4 since the latter 

do not take into account the free surface, geometry factors, and non-uniform 

stress gradients over the crack surface. There is validity to their argument, but 

in practice the ease of use of 7.4 has led to its widespread adoption. (The size 

and shape predictions based on A or KIB match very closely for the limiting 

cases of small mirrors in tension specimens. This is also true for most of the 

periphery of small semicircular mirrors centered on surface flaws in strong 

tension or flexure specimens.) 

The accuracy of the stress estimate varies depending upon the material, the 

experience of the fractographer, the microscopy and illumination used, and 

component geometry and stress gradient effects. Stress estimates may be 

accurate to within ≈ 10 % with glasses if they form ideal fracture mirrors. 

Kirchner and Gruver31 also showed that the relationship holds for elevated 

temperature fractures in ceramics. The mirror constant was invariant from 

room temperature to 1400 OC for a hot-pressed alumina with a 1 μm grain size. 

This is a remarkable finding since the alumina was fired at 1475 OC. Kerper 

and Scuderi32 also showed that mirror constants were unaffected by elevated 

temperature for eight glasses. Shinkai33 showed the mirror-mist constant, Ai, 

varied little with temperature for three glasses. Ai varied little with temperature 

for the Pyrex, but for the soda lime silica float glass and the 96% silica (Vicor), 

the Ai values at -150o

C were greater (0.03 – 0.04) than at room temperature. 

The trends were comparable to similar increases of the fracture surface energy, 

γf, with decreasing temperature. Shinkai33 also did experiments in various 

environments. There was negligible environmental effect and the Ai’s 

measured at room temperate in air and dry N2 gas were comparable. The lack 

of an environmental sensitivity is not surprising since mirror boundaries form 

when the crack is travelling at or near terminal velocity. 

Crack branching constants Ab have sometimes been reported, but some caution 

is advised, however, as discussed on page 7-7. There is evidence to suggest 

that Ab depends upon the loading mode (uniaxial or biaxial) and the size of the 

mirror relative to a part thickness if the part breaks in bending. The three 

boundaries correspond to an expanding crack front, and hence, Ai < Ao < Ab. 

Appendix C has a comprehensive compilation of fracture mirror and branching 

constants. The table shows there is some variability in the published mirror 

7-12 
 

 

 

 

 

 

 

 

 

 

 

 

Quantitative Analysis 

constants, even for model materials such as fused silica or soda lime silica. 

Some of this is due to the aforementioned metrology and judgment issues, but 

much of it is due to genuine material-to-material variability. Thus, one should 

not expect all silicon nitrides or aluminas to have the same mirror constants. 

Microstructural differences can change the mirror constants considerably. 

The word “boundary” must be used with some caution. It is now clear from a 

number of studies that there probably is not a discrete or sudden transition point 

on the fracture surface corresponding to a mirror boundary. The mode of 

viewing and magnification affects the judgment of the boundaries location. 

The higher the magnification the smaller the mirror seems to be, since fine 

detail that was washed out or not resolvable at lower power can be discerned at 

higher power. The mirror-mist boundary probably corresponds to surface 

roughness features that are of the order of 0.1 μm to 0.2 μm, which is a fraction 

of the wavelength of light and at the threshold of observable features with an 

optical microscope. Section D.3 in Appendix D has a discussion of several 

studies that have used various surface roughness measurement methods to 

measure mirrors, but these methods have yielded different results. 

So, although there is some variability in how different observers judge and 

measure fracture mirror sizes, most fractographers agree that there is a fracture 

mirror that can be measured. Size estimates can vary depending upon the type 

of microscope used, the mode of illumination, the objective power of the lens, 

and the judgment criteria of the fractographer. 

7.4.2 How to measure mirror sizes 

This section gives a brief outline of how to measure fracture mirror sizes. Very 

detailed guidelines are presented in Appendix D. The procedures in this Guide 

were adopted as an ASTM standard34 in 2007. 

In principal, a mirror size measurement is a straightforward exercise, but in 

practice it is not as easy as one would think, even with glasses. In addition to 

the boundary interpretation issue discussed above, component geometry effects, 

gradients in the stress fields, residual stresses, and origin flaw effects can all 

contribute to mirror distortions. Most investigators use simple, common sense 

approaches such as measuring a few mirror radii from the origin. Some have 

suggested more elaborate schemes. The boundary criteria are: 

The mirror-mist boundary in glasses is the periphery where one can discern 

the onset of mist. This boundary corresponds to Ai, the inner mirror constant. 

7-13 
 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

The mist-hackle boundary in glasses is the periphery where one can discern 

the onset of systematic hackle. This boundary corresponds to Ao, the outer 

mirror constant. 

The mirror-hackle boundary in polycrystalline ceramics is the periphery 

where one can discern the onset of systematic new hackle, and there is an 

obvious roughness change relative to that inside the mirror region. This 

boundary corresponds to Ao, the outer mirror constant Ignore premature 

hackle and/or isolated steps from microstructural irregularities in the mirror or 

from irregularities at the origin. 

Figure 7.8 shows an example in a moderately difficult-to-judge ceramic. The 

yttria-stabilized zirconia is a difficult to interpret since, although it has a very 

fine grain size (< 0.5 μm), it has a gradual progression of roughness from well 

within the mirror to the obvious hackle region far from the origin. Keeping in 

mind the criteria for polycrystalline ceramics above, the first step is to look at a 

region well within the mirror and note the typical, inherent roughness due to the 

microstructure. The next step is to observe the obvious new hackle that is 

greater in magnitude than anything in the relatively smooth mirror region. 

(a) (b) 

2R 

Figure 7.8 Example of a mirror in a ceramic material that is difficult to 

precisely measure. This is one of the specimens of the 3Y-TZP for the data in 

Figure 7.11 below. (a) shows the uncoated fracture surface of a 2.8 mm thick 

bend bar with low angle, grazing illumination (486 MPa). (b) shows an 

interpretation for a mirror-hackle boundary whereby new hackle is detected 

(small white arrows) as compared to the roughness inside the mirror. The 

mirror is slightly elongated into the interior due to the stress gradient. 

Judgments are best made while viewing the surface through a microscope. 

7-14 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Quantitative Analysis 

Figure 7.9 shows the two most common graph types that are used to show 

stress - mirror size data. The merits of each approach are discussed in 

Appendix D. This Guide recommends the use of stress versus 1/√R, Fig. 7.9a. 

The linear graph of stress versus 1/√R in Figure 7.9a has the advantage that the 

fracture mirror constant is the slope of the line, and residual stresses can be 

detected by non-zero intercepts as described later. 

The log stress versus log R graph in Figure 7.9b facilitates plotting a broad 

range of stress and mirror sizes. The mirror constant corresponds to the stress 

necessary to create a mirror of unit size, since (from eq. 4): 

log σa = log A - 1⁄2log (R). If R = 1, then log σa = log A, and σa = A. 

•

σ A

a 

large small 

mirrors mirrors 

•

•

• • •

•

• •

σ = A/√1 

-1⁄2 

log 

σa 

R = 1 

Log R = 0 log R 

small 

mirrors 

large 

mirrors 

(a) (b) 

Figure 7.9 Stress and mirror size graphs. (a) shows the preferred graph of 

applied stress, σa, versus inverse square root of mirror radius. The mirror 

constant is the slope of the line. (b) shows log σa versus log radius as 

commonly used in the older literature. 

Two examples of mirror data are shown in Figures 7.10 and 7.11 on the next 

page. The glass data in Figure 7.10 is extraordinary since the mirror size range 

is almost a factor of 100. The 3Y-TZP zirconia data in Figure 7.11 shows that 

measurements on uncoated and gold coated specimens concurred. This was 

gratifying since the uncoated specimens were difficult to interpret due to 

material translucency and difficulty in judging a boundary. Not surprisingly, 

the data scatter was greater for the uncoated specimens. More examples of 

marked fracture mirrors are shown in Appendix D which also has detail how 

residual stresses, stress gradients, the origin flaws, and microstructure can 

affect mirror sizes and shapes. 

7-15 
 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Slope ≈ -0.5 

Ai = 2.08 ± .02 MPa√m (1,891 psi √in) 

n = 259 rods 

Figure 7.10 A superb example of fracture mirror data for glass rods of many 

sizes and strengths. Two hundred and fifty-nine Pyrex 7740 rods varying in 

size by almost a factor of ten (4.1 mm to 38.1 mm) were broken in three-point 

flexure by Kerper and Scuderi at the National Bureau of Standards in 1964.

32 

Fracture Stress (MPa) 

900
	

800
	

700
	

600
	

500
	

400
	

300
	

200
	

100
	

0
	

0 10 20 30 40 50 60 70 80 90 

1 / √ mirror radius (m-1⁄2) 

' bars Ao = 8.6 MPa√m 

± 0.2 MPa√m 

+ uncoated disks 

Ao = 9.7 MPa√m 

± 0.2 MPa√m 

L gold coated disks 

Ao = 9.6 MPa√m 

± 0.2 MPa√m 

Figure 7.11 Fracture mirror size data for a 3 mol % yttria-stabilized 

tetragonal zirconia (3Y-TZP) from Ref. 35. Figure 7.8 shows one of the 

mirrors. Figure 3.24 shows a mirror both coated and uncoated. Bend bars 

produced a smaller mirror constant, either due to a geometric effect or simply 

the scatter in the data. 

7-16 
 

 

 

 

 

 

 

Quantitative Analysis 

7.5 Fracture Mechanics Analysis of the Flaw Size 

7.5.1 Introduction to fracture mechanics 

Fracture mechanics is a powerful analytical tool to aid the fractographer. It can 

be used to estimate critical flaw sizes and can help determine whether a flaw 

the fractographer has detected is of appropriate size to cause fracture. It can 

also be used to estimate the stress in a fractured part. A short introduction of 

fracture mechanics fundamentals is presented below followed by some 

practical examples. Figure 7.12 shows a hypothetical “brick” loaded in tension. 

A force P may or may not cause breakage depending upon the size and shape of 

the object and the material properties. Load divided by the area carrying the 

load is known as stress. The stress is constant across the cross section and 

Figure 7.12 A force P acts upon the object. Fracture occurs at a critical 

force that depends upon the material properties and the object’s size. (b) stress 

is the force normalized by the cross sectional area. (c) discontinuities such as 

a thorough-hole concentrate stress and magnify it in the vicinity of the 

discontinuity. (d) sharp discontinuities such as cracks intensify the stress field 

in front of the discontinuity. 

7-17 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

throughout the brick in this simple example. Real engineering structures 

usually have stress variations or gradients and the stresses are magnified or 

concentrated in the vicinity of geometric discontinuities. So, for example, the 

through hole shown in the Figure 7.12c concentrates stresses at the side of the 

hole since “flow lines of stress” (described below) are diverted around the hole. 

Stress concentrations are dimensionless magnification factors usually 

designated by lower case letter k. For the case of a through hole, k = 3, 

meaning that in the immediate vicinity of the hole, the stress is three times the 

nominal stress elsewhere in the body. 

Simple problems may be analyzed by the so-called “strength of materials” 

analysis, wherein assumptions about the structure (e.g., cross section plane 

sections remain plane, deflections are small, the structure remains elastic to 

fracture) lend themselves to simplified analysis. The maximum stress in threeor four-point beams in bending may be readily derived from “simple beam 

theory” a subset of the strength of materials analysis. More elaborate analyses 

may be solved by the “theory of elasticity” or by computer analysis using the 

“finite element analysis” (FEA). 

The local stress conditions in the vicinity of sharp discontinuities such as 

cracks cannot be adequately modeled using the strength of materials analyses. 

Stress concentration analyses lead to estimates of infinite stress at crack tips. 

Fracture mechanics is a discipline that deals with the stresses and strains around 

sharp flaws and has as its roots the 1920s work of Griffith.36,37 He showed that 

the strength of a brittle material with an slender elliptical flaw through crack of 

size 2c in a uniformly stressed plate as shown in Figure 7.13 is: 

2Eγ f σ f = (7.5) π c 

where σf is the fracture stress (the nominal stress in the plate), E is the elastic 

modulus, and γf is the fracture energy that is the energy to create unit surface 

area. In this Guide, flaw size is denoted as either “a” or “c” following 

conventional practice in the fracture mechanics literature. Usually the 

dimension is the radius or half width if the flaw is in the middle of a plate as 

shown in Figure 7.13a. If the crack is on the side of the plate, then a or c are 

the full crack length. The critical feature of this relationship is that strength is 

inversely proportional to the square root of flaw size. The larger the flaw, the 

weaker is the structure. Griffith’s formula was derived on the premise that 

crack propagation occurs if the elastic strain energy release in a body is greater 

than the creation of surface energy due to the new surfaces if the crack were to 

extend. His formula sets a necessary condition for a decrease in overall energy 

of the system. 

7-18 
 

 

 

 

 

 

 

 

 

 

 

 

Quantitative Analysis 

σa 

σa 

(a) (b) 

Figure 7.13 Through crack in a plate loaded in tension. (a) Long slender 

elliptical cracks of this type were envisioned by Griffith as flaws. If the crack 

tips are atomically sharp, fracture mechanics developed by Irwin38 in the 1950s 

must be used to model the stress state in the immediate vicinity for the crack tip 

as shown in (b). The dashed line shows σy as a function of x directly in front of 

the crack tip at y = 0. The stress intensity KI quantifies the magnitude of the 

stress field acting to open up the crack tip. KI depends upon the far field stress, 

σa and the geometry of the crack and the structure. Fracture occurs when the 

intensity of the stress field reaches a limiting value, KIc, known as the “critical 

fracture toughness, Mode I.” 

Mode I Mode II Mode III 

Opening In plane shear Out of plane shear 

Perpendicular to crack front Parallel to crack front 

Figure 7.14 Loading modes on a crack 

7-19 
 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) (b) 

(c) 

flat 

Figure 7.15 Lines of stress or “flow lines of stress” in a body. They cannot 

be transmitted through the irregularity such as a bubble or a pore. The flaws 

create a stress concentration on the sides where the flow lines converge. 

Sharper irregularities, such as the flattened wide flaw on in (b) create even 

greater stress concentrations. In the limit, for a flat sharp crack, the flow lines 

of stress kink severely, and fracture mechanics must be used as opposed to a 

simple stress concentration factor. (c) shows four flaws that have similar size. 

The left two are not particularly severe flaws despite their size, but the right 

two are more deleterious, especially if the rim-cracked pore (“Rings of 

Saturn”) flaw has a large rim crack, c. The flat penny-shaped flaw (shown 

schematically) on the right is the severest. 

7-20 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Quantitative Analysis 

Cracks may be loaded in other configurations as well, as shown in Figure 7.14. 

Ceramics and glasses usually break due to opening mode I loading. 

Before we go any further, there is a useful concept that may aid those who are 

not mechanical engineers or experts in fracture mechanics. When bricks such 

as shown in Figure 7.12 or 7.13 are loaded by remote uniform tension, the 

opposing forces or stresses at either end are transmitted through the bulk by 

lines of stress or “flow lines of stress” such as shown in Figure 7.15. These 

flow lines kink or bend around notches, grooves, or flaws. Engineers often use 

this concept to qualitatively assess what would happen with notches or grooves 

in parts such as rotating shafts, or a step reduction in diameter in a shaft, or 

even the neck and gage section of a direct tension strength test specimen. The 

more severe the kinking, the greater the stress concentration. 

Many flaws are atomically sharp and fracture mechanics analyses developed by 

Irwin in the 1950s are applicable.38 This approach analyses the stress field in 

front of a crack tip. The intensity of the field is quantified by the term KI, 

which is called the “stress intensity factor, mode I.” Mode I means the opening 

mode, whereby the crack faces are pulled directly apart as shown in Figure 

7.14. This mode is of primary concern for most ceramic and glass failures. As 

figures 7.12 and 7.13 show, the intensity of the stress field in front of the crack 

tip is proportional to KI. The flaw will propagate when the stress intensity 

reaches a critical condition, KIc. For many loading configurations, once the 

crack begins to propagate, it will run unstably and cause fracture. KIc is called 

the “critical stress intensity factor, Mode I,” or simply fracture toughness. For 

a small flaw in a body loaded in tension, σa, the stress intensity acting on the 

flaw is: 

K (7.6) I = Y σ c 

and fracture occurs when: 

KIc = Y σ f c (7.7) 

Y is the stress intensity shape factor, a dimensionless parameter that takes into 

account the shape of the crack and the loading conditions, c is the crack size, 

and σf is the stress at fracture. In this Guide, Y combines all geometric terms 

and π or √π terms. Examples of Y are shown below. Fracture occurs when any 

combination of Y, σ, or c leads to the critical condition. It should be 

emphasized that equations 7.6 and 7.7 only apply to small flaws loaded by 

remote tensile stresses, also known as “far-field stresses.” Most textbooks on 

fracture mechanics show illustrations such as Figures 7.12 and 7.13, whereby 

the stress field is remotely applied (not near the crack) and σa is the nominal 

stress if no crack were present. The flaws may be embedded in the bulk or may 

be on the specimen surface. Cracks may be loaded in other component shapes 

7-21 
 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

or loading configurations that do not create far-field stresses. For example, the 

local stresses from the plastic deformation zone underneath a Vickers indentation (that act on a median crack) are not far-field stresses, nor are the localized 

contact stresses where a sphere presses on a ceramic or glass surface. Furthermore, equations 7.6 and 7.7 may be inapplicable once cracks become very large 

relative to a component size. For example, a bend bar with a crack that reaches 

half way through the bar thickness has a much more complicated equation 

relating force, crack size, and beam dimensions to KI. The point here is that if 

one finds a fracture origin crack in a broken ceramic or glass object, and is 

tempted to apply equation 7.7, decide first whether the stresses really are 

remotely applied and are genuinely far field. Do not apply equations 7.6 or 7.7 

if the stresses are localized, or if the part dimensions are so small that the flow 

lines of stress around the flaw are not well established as in Figure 7.15. 

Note that the stress and flaw size dependencies are the same as the Griffith 

equation. Indeed, the energy and stress intensity approaches are consistent, and 

for plane strain (page 1.5) loading conditions: 

2Eγf (7.8) 

(1−ν2) 

KIc =

where E is the elastic modulus, γf is the fracture energy, and ν is Poisson’s 

ratio. (Note: sometimes the reader may find this equation without the 1-ν2 

term. That is the appropriate form for plane stress conditions, but nearly all 

practical problems for ceramics and glasses are plane stress conditions. In any 

case, since ν is typically between 0.15 and .25 for ceramics and glasses, ν2 is 

small and the denominator is close to 1.) 

Sometimes the fracture mechanics literature uses a related term, GIc, the critical 

strain energy release rate. For plane strain conditions, it is related to KIc by: 

EGIc KIc = (7.9) (1−ν2

) 

and GIc = 2γf . The ceramics and glass technical literature of the 1960s and 

early 1970s used analyses based on γf and the Griffith equation 7.5, but few 

flaws have a simple through-plate elliptical shape. Fracture mechanics analysis 

based upon stress intensity KI is more versatile and can model more complexshaped cracks. It became more common in the 1970s and is dominant today. 

Fracture toughness, KIc, is an important material property that characterizes a 

material’s intrinsic resistance to fracture. It has units of stress √length and in SI 

units is expressed either as MPa√m or MN/m1.5 . (The English system of units 

7-22 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Quantitative Analysis 

Material KIc MPa√m R-curve? SCG at room temperature? 

Glasses 0.75 – 0.90 No, flat Yes 

Sintered fused silica 0.92 No Yes 

Glass ceramics 2.0 – 2.5 No Yes 

Mullite 2.2 ? ? 

Alumina, single crystal 

2.4 to 4.5 

different 

planes 

No Yes 

Alumina (hot pressed 

or sintered, 99.9% 

pure, fine grained) 

3.0 – 4.0 No Yes 

Alumina (sintered, 

glass bonded) 2.5 – 5. No Yes 

Alumina (coarse 

grained) 3.0 – 5.0 Yes Yes, especially if there is a glassy boundary phase 

Boron carbide (hot 

pressed) 2.5 No, flat No 

Silicon carbide (solid 

state sintered) 2.5 – 3.0 No, flat No 

Silicon carbide (liquid 

phase sintered, 

elongated grains 

3. – 5. Yes Yes 

Silicon nitride 

(equiaxed fine 

grained) 

4.5 No No 

Silicon nitride 

(elongated grains, 

glassy boundary 

phase) 

5. – 8.0 Yes Yes 

Zirconia, cubic 

stabilized 3. No, flat ? 

Zirconia, Mg partially 

stabilized 3. – 18 Yes Yes 

Zirconia, Y tetragonal 

zirconia polycrystal 4.5. – 5.0 No Yes 

Cast iron 37. – 45. No No 

Table 7.2 Approximate fracture toughness, KIc, values. Fractographers 

should obtain a more specific value for the material being examined whenever 

possible. See sections 5.9.1 and 7.10 for slow crack growth (SCG). See 

sections 5.9.2 and 7.11 for R-curves. The larger values of KIc are usually 

associated with materials that have rising R-curve behavior. If R-curve 

behavior is suspected for a material, it is best to use a small-crack fracture 

toughness in a flaw size calculation rather than a plateau value. 

7-23 
 

 

 

 

 Fractography of Ceramics and Glasses 

for fracture toughness is ksi√in and the conversion factor to the SI units is only 

1.1 as listed in Table 7.1) Fracture energies are usually expressed as erg/cm2

 or 

as J/m2

. Table 7.2 shows some approximate KIc values. It is much better, 

however, to use a value for a specific material, since fracture toughness 

depends strongly upon composition and microstructure. Figure 7.16 shows 

some crack-microstructure interactions that can affect crack propagation and 

hence the fracture toughness. 

Glass and 

fine grained 

crack ceramics 

Microcracking 

ceramics 

Crack branching 

Crack deflection 

Crack bridging 

Second phase 

whisker reinforcement 

Transformation 

toughening 

Continuous fiber 

reinforcement 

Figure 7.16 Crack interactions with microstructure. 

7-24 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Quantitative Analysis 

Table 7.2 includes two columns that require explanation. The third column is 

for R-curve behavior whereby fracture resistance increases as a crack extends. 

This behavior is caused by several of the crack microstructure interactions 

shown in Figure 7.16 such as transformation toughening around a crack tip or 

bridging behind the crack tip. The fourth column in Table 7.2 is for “slow 

crack growth” which is a phenomenon whereby cracks grow stably at stress 

intensities less than KIc. Both of these phenomena are dependent upon the 

chemistry and microstructure. It is convenient to categorize materials 

according to Table 7.3. For now, the discussion will focus on the simplest 

combination in the top left box: a material that does not have slow crack 

growth and which has a flat R-curve. Such a material has a single value for 

fracture resistance, KIc. Many ceramics (e.g., solid-state sintered silicon 

carbide and many fine-grained materials) and all glasses if tested under inert 

atmosphere conditions fit this category. Slow crack growth and R-curve effects 

are discussed later in this chapter. 

Brittle (no R-curve) 

no SCG 

KIc 

Brittle (no R-curve) 

SCG 

K-V (velocity) curves 

Rising R-curve 

no SCG 

KR – a curves 

Rising R-curves and SCG 

KR – a curves and K-V (velocity) 

curves 

Table 7.3 The fracture resistance of ceramics and glasses may be categorized by 

whether the material does or does not exhibit slow crack growth (SCG) or rising Rcurve behavior. (After Fuller, private communication.) 

Users should be cautioned that much fracture toughness data is unreliable, 

especially the data obtained from simple saw cut notched bend bars or Vickers 

indentation crack length methods. The former method almost always 

overestimates fracture toughness. Vickers indentation fracture toughness data 

are notoriously faulty. Reliable methods such as single-edged precracked 

beam, chevron-notched beam, or surface crack in flexure, have been refined 

and standardized. International round robins and reference materials such as 

NIST Standard Reference Materials, SRM’s 2100, 2100A, and 2100B, have 

confirmed their reliability.39 Slow crack growth and R-curve behavior are 

important phenomena that add complications to fracture analysis, but for now 

7-25 
 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

we focus on a straightforward, simple application of equation 7.7 to aid 

fractographic analysis. 

The stress intensity KI on a flaw depends upon the flaw size and shape, the 

stress magnitude and distribution, and on the component size and shape. A 

remarkable variety of fracture crack shapes and loading configurations have 

been analyzed and the reference handbooks listed in Appendix A may be 

consulted for a specific problem. Figure 7.17 illustrates two examples. Figure 

7.17a shows a flat penny-shaped circular flaw in uniform tension, which is a 

case that is commonly used to simulate internal flaws in ceramics. Equations 

7.6 and 7.7 are applicable and the shape factor Y for this configuration is 1.128 

as shown in the Figure. 

Figure 7.17b shows a different loading configuration on the same flaw. At first 

glance it looks similar, but it is not. Opposing forces P pull on the center of the 

crack faces. This is not a far-field stress and equation 7.7 is inapplicable. KI 

has a completely different form as shown in the figure. Indeed, the larger the 

crack size, a, the less the stress intensity! This can easily be understood by 

Uniform remote tension Point forces pulling on the middle 

(a) (b) 

Figure 7.17 The stress intensity KI for penny-shaped circular cracks for two 

loading configurations. (The three-dimensional axes make the flaws appear 

elliptical in this view.) (a) shows a far-field tensile stress σ, and (b) shows 

opposed forces P pulling directly on the middle of the crack. KI increases with 

crack size in (a), but decreases with crack size in case (b). 

7-26 
 

 

 

 

 

 

 

 

 

 

 

 

 or h

Quantitative Analysis 

considering that the effect of the force P is distributed along the flaw’s 

perimeter. The larger the crack radius, a, the greater the perimeter and the less 

the force per unit length of crack perimeter. (This latter case, or more 

precisely the case for a semicircular surface flaw, is commonly used to 

represent the effect of concentrated residual indentation stresses from a Vickers 

indentation on a median crack flaw.) 

Surface cracks can often be modeled by semielliptical flaws as shown in Figure 

7.18. 

Uniform remote tension Bending 

σ

2c 

a 

t 2W 

a 

(a) (b) 

bending a 

2c 

t 

a 

2c 

Ydepth = 1.509 

Ysurface = 1.252 

(c) (d) 

Figure 7.18 Surface cracks in plates or bars in far-field tension (a) or 

bending (b). The stress intensity factor equations by Newman and Raju40 are 

applicable to these configurations. The Y factors for case (a) and (b) are the 

same for very shallow cracks, a/t → 0. (c) is a view of the cross section. (d) 

shows the same but labeled with the Y factors at two locations for the flaw 

dimensions given in the example on page 7-31 in Section 7.5.2.1. 

7-27 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

The flaw shape affects the Y factor. Y varies around the flaw periphery. A 

flaw goes critical when one portion of the flaw reaches KIc, so the maximum Y 

is of primary concern. The author has seen a number of instances where 

localized fractographic markings confirmed that fracture indeed had started at 

the portion of the flaw periphery that had the maximum Y. Figure 7.19 shows 

maximum shape factors for a variety of elliptical and semi elliptical flaws 

subjected to a far field uniform tensile stress. In each instance, the flaw size for 

equation 7.6 is the depth for surface flaws, or the half minor axis length for 

internal cracks. 

Note that for comparably sized internal or surface flaws, the surface flaw has a 

greater Y factor. This is because it has less supporting material in the vicinity 

to share the forces. The crack opening displacements on a surface connected 

flaw are also greater. 

The model cracks shown in Figures 7.18 and 7.19 may be used with equation 

7.7 for small flaws that cause fracture in ceramics and glasses. The equation 

has four variables: the far-field tensile stress at fracture, σf ; the flaw size, a or 

c; the shape factor, Y; and the fracture toughness, KIc. This simple equation 

may be used in three different ways: 

(1) Flaw size may be estimated if Y, σf, and KIc are known. This approach is 

very useful in laboratory testing, whereby specimens are broken under 

controlled conditions, and break loads and specimen and fixture 

geometries are used to compute failure stresses. 

(2) The failure stress, σf, can be estimated if the flaw size and shape can be 

measured fractographically, and KIc is known. This approach is useful in 

estimating stresses in fractured components. 

(3) Fracture toughness, KIc, may be computed if the flaw size and shape, and 

the failure stress are known. 

Examples will be shown below, but first additional information about the flaw 

shape and stress gradients must be considered so that appropriate Y values may 

be estimated. The reader should be forewarned that there is no consistency in 

the fracture mechanics and materials science literatures. Crack sizes are 

usually described by a, b, c, 2c, or other symbol, but sometimes these same 

symbols are used differently and sometimes they are even used for the 

component dimensions. Sometimes the Y factor includes π or a √π term, but 

often it does not. (One reference even uses Y as the inverse of how it is used 

everywhere else.) In this Guide, and in ASTM standards C 132241 for 

fractography and C 142142 for fracture toughness, all geometric terms for the 

stress intensity shape factor (including π or √π) are combined into Y. 

7-28 
 

 

 

Quantitative Analysis 

Uniform 

tension 

t 

Figure 7.19 The stress intensity shape factors Y for surface and internal 

round and elliptical for internal (volume) cracks and surface cracks for farfield tension stresses. Minimum Y factors for the surface cracks are shown in 

parenthesis. The Y factors for the surface flaws may also be used for parts 

loaded in bending if the flaw depth a is small relative to the specimen thickness 

(a/t is small). 

7-29 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

7.5.2 The Newman-Raju Y factors for semielliptical surface flaws in 

bending 

7.5.2.1 The Newman-Raju formulas 

Ceramics and glasses often have surface origins and the stress intensity factor 

solutions of Newman and Raju40 for both tension and bending stress fields are 

very helpful. Their versatile solutions may be applied to tension, bend bar, and 

biaxial disk specimens, and indeed any component shape wherein a remote 

stress field causes fracture. With reference to Figure 7.18 for dimensions for 

surface flaws in a plate or beam in bending, the Ydepth and Ysurface for locations 

at deepest part of the crack and at the surface are: 

Ydepth = 

M H 

Q 2 π
 (7.10)

M Ysurface = S H1 π

Q (7.11) 
 

where M, Q, H2, H1, and S are geometric terms: 

1.13 



































  

+ 14. 1.-































2 24 4

 0.89 1. 























































a a a a M = -0.09 + − 0.54 + + - 

 























c a t a c t 0.2+ 0.65+ c  c 











0.5 

(7.12)
 



























 





1.05 







a a H1 = 1. 0.34 +0.11
 (7.13) - c t







0.75 





1.5
 





2 































 



 0.55 -

Q = 1.+1.464 































a a a a a H = 1.
 - 1.22 +0.12
 + +0.47 (7.14) 2 c t c c t







1.65 





a (7.15) Φ=
 for a/c ≤ 1

c 

2





































(7.16) a a S =
 1.1 + 0.35
 t c  

7-30 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Quantitative Analysis 

(Equations 7.10 and 7.11 do not include an additional term fw in the original 

Newman-Raju analysis, since for nearly all cases for ceramics and glasses, the 

flaw width 2c is small relative to the component width, and fw is 1.0.) These 

polynomials appear intimidating at first glance, but are easily programmed into 

hand calculators or spreadsheets. Notice how many of the dimensions appear 

as ratios, e.g., a/c, the flaw ellipticity ratio; or a/t, the ratio of the flaw depth to 

the plate thickness). The (1-(a/c))24 term for M in eq. 7.12 is genuine and not a 

typographical mistake. The Newman–Raju Y factors have been widely used in 

many textbooks and are included in several ASTM standards including C 1421 

for fracture toughness of ceramics42 and E 740 for metals,

43 as well as ISO 

standard 18756.44 

The above equations are for surface cracks in bending stress fields. Newman 

and Raju listed alternative formulas for direct uniform tension loadings.40 It is 

not necessary to list them here. Equations 7.10 – 7.16 may be easily used for 

cracks in uniform tension loaded parts by the simple expedient of inputting an 

arbitrary large value (e.g., 1000) for the part thickness, t. 

So for example, if a bend bar is broken and it has a surface crack with a depth 

(a) of 40. x 10-6 m; width (2c) of 144. x 10-6 m; and specimen thickness (t) of 

2.998 x 10-3 m; then M = 1.080, H1 = 0.995, H2 = 0.983, √Q = 1.247, and S = 

0.820. Then, Ydepth = 1.509, Ysurface = 1.252. These are shown in Figure 7.18d. 

In practice, the more elliptical the flaw, the more likely the maximum Y is at 

the flaw depth (unless the crack penetrates deeply into the specimen interior 

and into the stress gradient.) 

Using the larger value for Y and the flexural strength which was 453 MPa, 

equation 7.7 gives KIc = 1.509 x 453 MPa √40 x 10-6 = 4.32 MPa√m. In this 

case, the parameters are used to estimate the material’s fracture toughness. 

As noted on page 7.28, equation 7.7 for KIc can be used two other ways. If 

instead, the fracture toughness and flexural strength were known, the flaw size 

could have been estimated. Similarly, if the crack size had been measured and 

the fracture toughness was known, then the stress of 453 MPa could have been 

estimated. 

7-31 
 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Figure 7.20 Grinding crack in a longitudinally-ground silicon carbide bend 

bar at 649 MPa. This flaw was located directly in the center of a well-defined 

fracture mirror. The depth a is 23 μm, the width 2c is 95 μm, and the bar 

thickness, t, was 2.2 mm. Ymax was at the deepest point and was 1.58. The 

computed fracture toughness is: 1.58 * 649 MPa * √ 23 x 10-6 = 4.9 MPa√m, 

which is reasonably consistent with published average values of 4.0 to 4.5 

MPa√m for the material. This orthogonal grinding crack is subtle and does not 

stand out clearly against the background microstructure. There may have been 

some doubt about its identification, but the fracture mechanics calculation 

shows it is about the right size. 

Figures 7.20, 7.21, and 7.21 show three examples of how the calculation can be 

used to confirm that the right flaw has been identified. Sometimes the 

calculated and measured flaw sizes agree very closely, but there are a number 

of reasons why they may differ as discussed below in section 7.8. ASTM 

standard C 132240 states that if the calculated and measured flaw sizes do not 

agree within a factor of two or three, then the origin should be reexamined to 

verify that the correct feature has been identified. 

The surface crack in the BK-7 borosilicate crown glass disk shown in Figure 

6.28a,b and 7.22 was 7.3 μm deep and the stress at fracture was 146 MPa. For 

a long surface crack, Y is 1.99. The calculated fracture toughness from these 

values is 0.78 MPa√m, which is very close to published values for this 

particular glass. 

7-32 
 

 

 

 

 

Quantitative Analysis 

(a) 

(b) 

10 μm 

Figure 7.21 Fracture origin in a single crystal silicon wafer. (a) shows the 

fracture mirror which has a very faceted shape typical of some single crystals. 

The origin is a grinding crack shown by the white arrows in (b). The specimen 

was tilted back to show both the fracture and outer ground surfaces. A 3000 X 

scanning electron microscope photo showed that the crack was 2 micrometers 

deep, and when combined with a 282 MPa fracture stress, and a Y factor of 

1.99, gives KIc = 0.79 MPa√m, in reasonable good agreement with published 

values for fracture on the {110} plane. 

7-33 
 

 

 

 

 Fractography of Ceramics and Glasses 

(a)

 (b) (c) 

Figure 7.22 Fracture origin crack for the borosilicate crown glass biaxial 

disk specimen that was broken in inert conditions as shown in Figure 5.17. (a) 

is an SEM photo of the distorted mirror. The arrows show the long polishing 

scratch. . (b) and (c) show two portions of the polishing scratch which was 7.3 

mm deep (arrows). Using Y = 1.99 and the breaking stress of 146 MPa, gives 

KIc = 0.78 MPa√m. The optical image (see Figure 5.17) showed a slight bulge 

in the flaw at the center of the mirror suggesting the flaw may have grown 

slightly before going critical. This growth is not evident in the SEM images. 

7-34 
 

 

 

Quantitative Analysis 

7.5.2.1 Graphical curves for the Newman-Raju formulas 

Figure 7.23 may be useful for those who do not wish to program a calculator or 

spreadsheet. Y values may be picked off the graph for cracks in bending 

problems for given a/t and a/c ratios. Notice it demonstrates how Y decreases 

as a flaw penetrates into the stress gradient (larger a/t). (Y values for small 

surface flaws in uniform tension fields may be obtained by reading the graph 

value for an a/t ratio of zero.) 

0.00 0.10 0.20 0.30 0.40 0.50 0.60 0.70 0.80 0.90 1.00 

2.0 2.0 

1.9 

Bending compression side 

t, plate thickness 1.9 

tension side a 

1.8 2c2c 1.8 

1.7 Ymax 1.7 c/a = 20 is at the surface 

is at the 

Ymax

1.6 1.6 deepest point 

c/a = 10 1.5 1.5 

c/a = 6

1.4 c/a = 4 1.4 

c/a = 2 

1.3 1.3 c/a = 1.4 

Ymax

c/a = 1 1.2 is at the surface 1.2 

Ymax 1.1 1.1 

is at the surface 

c/a = 1 semicircle 

1.0 1.0 

0.00 0.10 0.20 0.30 0.40 0.50 0.60 0.70 0.80 0.90 1.00 

Fractional Depth (a/t) . 

Figure 7.23 The maximum stress intensity shape factor Y for semielliptical 

surface cracks in a bending stress field in a plate or bar of thickness t. In the 

limit as the flaw is very small relative to the plate thickness, a/t → 0, the 

solutions are the same as for the uniform tension case. The kink in the loci 

corresponds to where the maximum Y changes from the surface to the deepest 

point in the interior or vice versa. 

Ymax 

7-35 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

7.5.2.3 Newman-Raju formulas for deep (a/c > 1) semielliptical 

surface cracks 

Occasionally the fractographer may find a flaw that is deeper than a semicircle, 

i.e., a/c > 1. Fett devised modified Newman-Raju formulas that permits one to 

compute the Yd and Ys factors in such cases for either tension or flexural 

loadings. The conscientious reader may consult a paper by Fett45 for the 

modified formulas. Rather than listing all the revised formulas or showing 

charts or tables, some approximations are provided here. 

In the cases with semi-ellipses that are deeper than a semicircle, the maximum 

Y is at the surface, Ys. For cracks that are shallow relative to the plate 

thickness (a/t ≤ 0.2) in tension or bending, the surface Ys factor decreases 

linearly from that for a semicircle by only 10% as a/c goes from 1 (the 

semicircle) to 2 (a very elongated crack into the thickness). 

The stress intensity shape factor Yd for the deepest part of the flaw also 

decreases approximately linearly, but by greater amounts. Yd decreases by 

33 % in both tension and bending as a/c goes from 1 to 2. If the flaw is larger 

relative to the plate thickness (e.g., a/t = 0.2), Yd decreases by 33% and 45% in 

tension and bending, respectively, as the a/c ratio goes from 1 to 2. In other 

words, as a very deep semi-ellipse penetrates into the stress gradient in bending 

(a/t = 0.2, or 40% of the way to the neutral axis), Yd diminishes rapidly. 

7.5.2.4 Accuracy of the Newman-Raju formulas and alternatives 

A few words about the accuracy of the Newman-Raju stress intensity factors Y 

are in order here. For most fractographic analyses, an accuracy of a few 

percent is perfectly adequate. Developed by finite element analysis almost 40 

years ago, the Newman-Raju factors are widely used in textbooks, fracture 

mechanics and fractography standards around the world. Fett has reviewed46,47 

them carefully and concluded that for most common configurations, they are 

accurate to within 3%. Fett noted that a solution of Isida et al.,48 obtained by 

means of a modified body-force method exhibits superior accuracy with a 

maximum deviation of less than 1% for Yd. 

Even more accurate and versatile solutions have been developed recently by 

Stobl et al.49. For most common semi-elliptical cracks their Yd solutions 

concur with Newman and Raju’s within a few percent, but the newer solutions 

cover a broader range of Poisson’s ratios and cover surface cracks that have 

peripheries that do not reach the specimen surface at exactly 90o

. Strobl et al.’s 

cracks are sections of a semiellipse and in some cases they are better models. 

7-36 
 

 

 

 

 

 

 

 

 

Quantitative Analysis 

7.5.2.5 Bansal’s approximation for semielliptical surface flaws 

Bansal reviewed the stress intensity shape factors for semi-elliptical surface 

flaws in uniform tension50 and arrived at a simple approximation: 

2.0 KKII = 

1.68

σσ4

√AA 7.17 

where A is the area of the flaw. Since the area of a semi-ellipse is πac/2: 

KKII ≈ 1.33 σσ�√aaaa = 1.33σσ4

√aaaa 7.18 

Sometimes the expression is written as: 

KKII = 1.33 σσ√bb 7.19 

where the effective flaw size b = √ac. 

Figure 7.24 Comparison of Bansal’s approximation to the Newman-Raju 

solution for Y of semielliptical flaws in uniform tension. The approximation is 

only suitable for a/c > 0.2 . 

7-37 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Figure 7.24 shows how this approximation compares to the Newman-Raju 

outcomes, also for the uniform tension case. Bansal’s approximation is quite 

good for the condition of 0.2 < a/c <1.0. It begins to deviate for longer, 

shallower surface cracks, e.g., those with a/c ratios smaller than 0.2 (c/a ratios 

of 5 or greater). In the limit of a long surface crack in a plate in tension, Y 

becomes 1.99 and the Bansal approximation is in serious error. In other words, 

do not use Bansal’s approximation for long shallow surface cracks. Furthermore, his solution is only appropriate for uniform tension, or very small flaws 

in bending such that the crack depth a is small relative to the plate thickness, t 

or h. If the flaw depth is greater, e.g. a/t > 0.05, then the Newman-Raju 

solutions should be used. (This means that the deepest part of the flaw 

penetrates into the bending stress gradient by no more than 10%.) 

7.5.3 Irregularly-shaped surface cracks 

Fracture mechanics reference books have tabulations of stress intensity 

formulas and shape factors for a myriad of crack shapes and loading 

configurations. There are even listings for multiple cracks in proximity to each 

other that show how the stress intensity fields interact. Real flaws in real 

materials often do not have convenient shapes, however. Several 

approximations are very helpful. Figure 7.25 shows a case of an irregular 

planar surface crack loaded by far-field tension. The stress intensity may be 

approximated, with a maximum of a 10% error (Poisson’s ratio of 0.3), by a 

solution in Murakami’s book:

51 

Ki =1.15 σ 4 A (7.20) 

where A is the area of the flaw. For example, if the flaw is semicircular with 

radius a, this simplifies to 

KI = 1.29 σ a (7.21) 

that is nearly the same as equation 7.19. Numerical analysis showed that the 

fourth root of area equation 7.20 is an adequate approximation for stress 

intensity factor solution for such diverse shapes as partial circles, triangles, 

rectangles and semi ellipses. Equation 7.20 is also for the case of uniform 

tension. It may be used for small cracks in bending specimens, provided that 

the crack does not penetrate very deeply into the specimen, a/t < 0.05. The area 

approximation should not be used for very long shallow flaws c/a > 5. 

7-38 
 

 

 

 

 

 

 

 

 

Quantitative Analysis 

Figure 7.25 An irregular planar shaped surface crack in a tension stress field. 

(After Murakami, ref. 50.) 

7.5.4 Three-dimensional, blunt, and inclined flaws 

Real flaws are not always flat and planar. Figure 7.26a shows a sphere that is 

not unlike many pores in ceramics and glasses. Sometimes round pores are 

modeled by a penny-shaped crack, Figure 7.26c. This is a gross oversimplification. Indeed, if the pore is round and crack free, it does not act to 

intensify the stress, but simply concentrates stress around the flaw sides. The 

stress concentration factor for an interior sphere in a part in uniform tension is 

only k = 2. Smooth round pores require much greater force to cause fracture 

than the penny-shaped flaw of the same diameter. Round pores may be quite 

docile and not act as strength-limiting flaws. Figure 7.26b shows another 

variation. This rim-cracked pore (“Rings of Saturn”) flaw is a combined stress 

concentration and fracture mechanics problem. For small rim cracks, the pore 

acts as a concentrator of stresses, and the rim crack size, c, may be regarded as 

the crack size in a fracture mechanics calculation. For a very large rim crack, 

the pore’s effect is diminished, and the penny shaped crack is a very good 

approximation. In practice, it is very difficult to detect a rim crack around a 

pore, much less measure it. It is not unreasonable to expect that rim cracks of 

one or two grain diameters extend from the pore and in fact many pores in 

ceramics do have local microstructural irregularities or microporosity in the 

vicinity. Only very rarely have fractographers attempted to precisely model 

7-39 
 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

such flaws and it is customary to use the penny-shaped cracks as an approx.-

imation, but with full cognizance that the actual flaw is much blunter. 

flat 

(a) (b) (c) 

Figure 7.26 Flaws may be three-dimensional. A flat penny shaped crack (c) 

is a two-dimensional approximation. 

The reader may consult fracture mechanics handbooks (e.g., Bibliography 

Appendix A) for more information about rim-cracked pores (also called 

“annular cracks around a spherical void”) or flaws that are inclined to the stress 

direction. A slight inclination of 5 degrees does not affect the mode I loading 

very much, but beyond 10 degrees flaws experience not only mode I loading 

but varying degrees of modes II and III depending upon the flaw shape and 

inclination angle. Modes II and III are have shearing actions on a crack as 

shown in Figure 7.14. 

7.6 Relationship of KIc and A 

KIc and the fracture mirror and branching constants are related to σ√a or σ√R 

and hence have the same set of units: MPa√m (or ksi·√in). As Figure 7.7 

illustrates, KIc pertains to conditions at the flaw. The A values correspond to 

events away from the origin as the crack propagates and reaches terminal 

velocity. Kirchner and Conway30 showed that mirror shapes can be predicted 

on the basis of stress intensity at the crack front as it expands outwards from 

the origin, so it is not unreasonable to expect that KIc and the A’s are related. 

If they are related, then in principle one can estimate KIc from A or vice versa. 

Mecholsky et al. 52 tabulated A’s and KIc’s for over two dozen ceramics and 

7-40 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

Quantitative Analysis 

glasses and showed there was an approximate correlation. On average the ratio 

of Ao to KIc was 20/7 = 2.86, although there was considerable scatter. Bansal 

and Duckworth53,54 carefully measured Ao to KIc ratios of 2.2 to 2.4 for a glass 

ceramic and a sintered alumina, and 2.3 for float glass and a hot-pressed 

alumina. Subsequent tabulations55 on more contemporary materials, many of 

which have R-curve behavior, have given ratios as low as 1.5 to 2.0. 

Kirchner et al.56 showed that the branching constant Ab varied by over a factor 

of 2 with the degree of transformation toughening in a calcia partiallystabilized zirconia. The greater the toughening, the greater the suppression of 

branching to greater distances; i.e., the Ab/KIc ratio increased with increased 

toughness. 

In a dissenting view, Bradt and colleagues57,58 have suggested that A is not 

really a material constant and noted the discrepancies in A/KIc ratios. In 

summary, variability in the quality and accuracy of the KIc and A data, plus 

genuine microstructural effects, probably mean there is no unique A/KIc ratio. 

7.7 Mirror to Flaw Size Ratios 

A simple but very useful observation is that mirror and flaw sizes are related. 

Krohn and Hasselman59 showed that the mirror size to flaw size ratio (R/c) is 

about 10 for glasses. Mecholsky et al. 52,60 obtained ratios of 10 to 11.6 using 

the mirror-mist (inner mirror) boundary in glasses. They estimated the ratio to 

be about 13 for the mist-hackle (outer) boundary to flaw size ratio for glasses. 

Ratios for ceramics are more problematic and depend upon the microstructure 

and the ability to judge a specific boundary. Mecholsky et al.51 estimated the 

ratio of the hackle boundary (outer mirror) to flaw size was 13 for polycrystalline ceramics and single crystals, but varied from as low as 8 to as large as 

40! 

Rice61 pointed out that the bluntness of the flaw may alter the mirror to flaw 

size ratio. For example, pores often act as blunt flaws and it requires extra 

energy or stress to overcome the bluntness. The excess energy causes a mirror 

to form at shorter distances from the origin than if the flaw were sharp. Thus 

the R/c ratio is smaller. 

Slow crack growth can also alter the apparent ratios.60 If the mirror to flaw 

size ratio (R/c) is much greater than the values listed above, then the flaw 

should be checked for evidence that it grew stably prior to fracture. 

7-41 
 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Residual stresses in the vicinity of the flaw or the mirror also can change the 

ratios. Unexpected mirror to flaw size ratios may be evidence of residual 

stresses acting on the flaw or the whole mirror. 

The mirror to flaw size ratio (R/c) may also be affected by the presence of 

mixed Mode I/II loading. Gopalakrishnan and Mecholsky62, and Rice63 have 

shown that increasing Mode II loading on a crack decreases R/c. 

The R/c ratio is also affected if a material has a rising R curve. The larger the 

fracture toughness, the lower the R/c ratio.61,64 In rising R-curve materials, the 

ratio may even depend upon how large the initial crack was and how far it 

might have stably grown before it went unstable. (Stable crack growth and 

crack instability is discussed more in sections 7.10 and 7.11 later in this 

chapter.) In R-curve materials, the larger the initial crack, the smaller the R/c 

ratio. Gopalakrishnan and Mecholsky’s 2013 paper62 even goes so far as to 

analyze the combined effects of R-curve and mixed Mode I/II behavior on R/c 

ratios. (Note, their discussions are on the ratio, c/R, unlike most other work 

which focuses on the R/c ratio.) 

Ccalc < Cmeas Ccalc > Cmeas Ccalc > Cmeas or Ccalc < Cmeas 

Crack blunting Stable crack extension -SCG High temperature SCG 

Use of 2-D model Stable crack extension – R 

curve 

Multiple flaws nest or 

interact 

Stress gradients Flaw causes a local fracture 

toughness degradation 

Residual stresses 

Crack healing Flaw is within a single grain Flaw is truncated on the 

fracture surface 

Flaw links with other flaws or 

the surface 

Flaw shape irregularity 

Flaw elastic properties are 

different than the matrix 

Faulty fracture toughness 

data 

Table 7.4 Factors that can cause a measured flaw size cmeas to differ from the 

calculated ccalc size. (Quinn and Swab, Ref. 65 and 66.) 

7-42 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Quantitative Analysis 

7.8 Comparing Measured to Calculated Flaw Sizes 

Differences in calculated and measured flaw sizes of 20 % to 50 % should not 

be considered too serious, but variations greater than a factor of 2 or 3 should 

prompt the fractographer to study the origin more carefully. Either the wrong 

feature has been identified as an origin or the origin may be more complicated 

than expected. Either something about the calculation is wrong or the 

fractographically-measured size is wrong. Table 7.4 and Figure 7.27 show 

some of the factors than can cause the discrepancy.

65,66 These can be 

appreciated if Equation 7.7 is rewritten: 

 K 

2 

Ic (7.22) ccalc = 

 



Yσ  f 

where ccalc is the calculated critical crack size at fracture, σf is the fracture 

stress, and KIc is fracture toughness. 

The cases where ccalc < cmeas are considered first. A blunt flaw of actual size c 

requires a greater stress to cause fracture than a comparably-sized sharp flaw. 

Some flaws are blunt to begin with. Other initially-sharp flaws may have been 

blunted or even partially healed by oxidation or other environmental effects. 

Glass cracks can partially heal in vacuum or inert atmospheres, especially if 

they form during dynamic events and are unloaded quickly. Inserting the 

greater stress to break healed or blunted cracks into equation 7.22 leads to a 

smaller ccalc than the actual flaw size c. Assuming that flaws are 2-D penny 

shaped or are flat ellipses when they really are 3-D flaws, usually leads to Y 

estimates that are too large, and this also leads to smaller ccalc than the actual 

flaw size c. Unexpected stress gradients can cause σf to be overestimated. 

That in turn will again cause ccalc to be too small. The stress at the origin site 

should be used in equation 7.22. 

Next are the cases where ccalc > cmeas. Slow crack growth and rising R-curve 

behavior can cause stable crack extension. Using the actual breaking stress in 

equations 7.7 or 7.22 gives a correctly calculated flaw size, but if the 

fractographer detects and measures only the initial flaw size, then ccalc > cmeas. 

In this case it is the fractographic size measurement that is off. Alternatively, if 

stable crack extension due to R-curve behavior does occur, but the 

fractographer uses a long crack (plateau) value of fracture toughness, then the 

fracture toughness is overestimated and again ccalc > cmeas . R-curve behavior is 

discussed in more detail in section 7.11. Inclusions or chemical inhomogeneities can be particularly deleterious if they degrade the fracture toughness 

of the surrounding matrix. If the fractographer uses the baseline KIc value in 

7-43 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

equation 7.22, this also leads to ccalc > cmeas. In this case, the calculated size is 

wrong and the observed crack size is correct. The fracture toughness of a 

single crystal, or a single grain, is almost always less than that of a polycrystal. 

If a flaw resides within a single grain, but a polycrystalline KIc value is used in 

equation 7.22, then ccalc will be too large. Flaws may link up with other flaws or 

a surface weakening the specimen. If the linkage is undetected by the 

fractographer, then his cmeas will be too small compared to the calculated and 

actual critical flaw size. 

Some factors can cause the measured size versus calculated size discrepancy to 

go either way. Flaws are often assumed to exist in isolation, but as Figure 

7.27b shows they can interact causing either an increase or a decrease in the Y 

factors that act on them. It all depends upon the geometries and arrangements. 

This is an instance where the “flow lines of stress” concept can aid 

interpretation. If flaws are lined up in a row parallel to the stress, the flaws will 

shield each other. If they are staggered or lined up sideways, they can intensify 

the KI fields and the Y factors. Flaws may be truncated and only a portion 

detected on the fracture surface. Undetected residual stresses can throw off the 

calculated flaw size as well. Faulty fracture toughness data can also throw off 

the correlations. 

The point here is that fracture mechanics is a valuable tool to aid fractographic 

analysis. It can help confirm that the fractographer has found the correct origin 

flaw. A finding that the calculated and measured sizes are appreciably different 

should prompt further consideration and analysis. What factors caused the 

discrepancy? Is the discrepancy an isolated outcome or part of a pattern? 

(a) Flaw bluntness or 2-D approximations 

Figure 7.27 Factors that cause calculated and measured crack sizes to differ 

(continued on next page) 

7-44 
 

 

 

 

Quantitative Analysis 

(b) Flaw interactions 

(c) Flaw linkage (d) Flaw slow crack growth 

(e) Flaw truncation on the fracture surface 

Figure 7.27 continued Factors that cause calculated and measured 

crack sizes to differ. 

7-45 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

7.9 Crack Velocities from Wallner Line Analysis 

Wallner lines can be used to estimate crack velocities, whether they are created 

incidentally during the fracture event or deliberately by ultrasonic fractography. 

Wallner lines may not be evident in very slow fractures in glass or in ceramics 

due to the latter’s rough fracture surfaces. 

Ultrasonic fractography uses a transducer to generate sonic waves that interact 

with the propagating crack as discussed in sections 5.4.4 (on tertiary Wallner 

lines) and 3.21 (on examination techniques). Usually the transducer is placed 

position perpendicular to the crack propagation plane. The sonic waves create 

tertiary Wallner lines and show the crack front shape. Velocities are easily 

estimated from the pulse frequency and the spacing between the Wallner lines 

on the fracture surface. Richter and Kerkhof67,68 and Michalske et al.69,70 used 

this technique over a broad range of crack velocities from terminal velocity 

down to 10-7 m/sec. Indeed Richter and Kerkhof67 used this technique to show 

that surface cracks that grow by slow crack growth evolve in shape in 

accordance with the Newman-Raju stress intensity shape factors as shown in 

Figure 5.44. They also showed Figure 5.45a which proved a semielliptical 

crack could go unstable from one side as shown in Figure 5.56a,b. 

Crack terminal velocities are about 50 % to 60 % of the transverse elastic wave 

speed and range from 1500 m/s for soda lime silica to 2500 m/s for fused silica 

as listed in Table 5.1. The transverse wave velocity is: 

E (7.23) vt = 2ρ(1+ν )

where E is the elastic modulus, ρ is the density, and ν is Poisson’s ratio. 

Richter66,67 studied cracks running at terminal velocity at stress intensities 

above KIc. 

Even in the absence of externally applied sonic pulses, it is possible to interpret 

naturally created Wallner lines to estimate crack velocities as shown in Figures 

7.28 – 7.30. The simplest case is for a straight crack front that intersects a 

discontinuity such as a bubble or inclusion that creates an elastic pulse. This 

generates a fracture surface feature called a “gull wing.” The radiating 

transverse elastic wave always travels faster than the crack front. Figure 7.28 

shows how the radiating circular-shaped wave front catches up and intersects 

with other portions of the crack front. The locus of intersections generates 

Wallner lines that have a telltale “V” pattern centered on the elastic 

discontinuity. There often is a slight curl at the pulse origin. The net effect is 

7-46 
 

 

 

 

 

 

 

 

 

 

 

 

 

Quantitative Analysis 

to produce the gull wing shape. The included angle (2β) between the wing 

segments may be used to calculate crack velocities. Crack velocity is quite 

simply: 

(7.24) vc = vt cos β

where vc and vt are the velocities of the crack front and the transverse wave. 

Slowly-moving cracks do not advance very far in a time interval ∆t, and the 

elastic pulse quickly overtakes the entire front. Thus, the “V” shape is very flat 

and, if it assumed that the included angle 2β is 180o

 – 2o = 178o, then the 

slowest velocity that can be measured in a glass with vt = 3460 m/s is vc,min = 60 

m/s. Conversely, if the crack is advancing at the terminal velocity of 60 % of 

the transverse wave velocity, then the minimum included angle 2β is 106o

. 

Tsirk71 identified hyperbolic “V” shaped Wallner lines in obsidians, created by 

an elastic pulse generated by secondary fracturing behind the main crack front. 

(b) 

3 2 1 

3 

2 

1 

0 

vt vt 

vc 

2β

Crack 

front 

Wave front 

dcp 

(c) 

vc ∆t vt ∆t β

(a) 

vc ∆t 

cos β = vt ∆t 

vc = vt cos β

Figure 7.28 Crack velocity may be estimated from Wallner line analysis. Gull 

wing lines occur when a straight crack front passes a point discontinuity such 

as a bubble or inclusion. (a) shows gull wings in knapped obsidian (specimen 

courtesy A. Tsirk). The direction of crack propagation is from bottom to top. 

The included angle (2β) is used to estimate crack velocity. Slow-moving cracks 

have flat (horizontal) wings. Fast-moving cracks create wings with a “V” 

shape. (b) and (c) show the geometries and simple trigonometry used to 

estimate crack velocity. 

7-47 
 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Wallner lines such as those inside a mirror shown in Figure 7.29 may be 

analyzed as shown in Figure 7.30. 14,27,72,73 

(a) Radiating 

crack 

Radiating 

wave 

Wallner 

line 

 (b) (c) 

Figure 7.29 Crack velocities can be estimated from Wallner line analysis. (a) 

shows a crack front that radiates outward from the origin at velocity vc (thin 

circular lines) and reaches irregularity A on the left. An elastic pulse is 

generated at A and it radiates a transverse (shear) wave outward from A 

(dashed lines) with a faster speed vt than the crack front. The numbers label 

the crack and elastic pulse fronts at the successive times starting at 0, the 

instant the crack reaches A. The intersections for the specific times are marked 

by dots, but of course this is a continuous process and the solid line shows the 

ensuing Wallner line. The crack front is assumed to be moving at constant 

velocity vc = 0.5 vt in this case. If the crack were accelerating, the Wallner 

line would be less hooked in the beginning and would extend more towards the 

upper right. (b) and (c) show Wallner lines inside mirrors in glass rods. 

7-48 
 

 

 

 

 

 

Quantitative Analysis 



tvt ∆ cos β =

 c ∆tv sin γ =

origin 

(a) (b) 

Figure 7.30 Analysis of Wallner lines to determine the crack velocity if a crack 

radiates outwards from an origin. (a) shows a case where the location of the 

Wallner line elastic wave pulse (site A) and the direction of the crack extension 

are known. (b) shows the case where the precise crack growth direction is not 

known, but the origins (A1 and A2) of the Wallner lines are known. 

In the case where the exact direction of crack propagation (radiating at speed vc 

from the origin) and the initiation site of the Wallner line (A) are known: 

vc sin γ (7.25) = vt cos β

Crack velocity can be determined even if the crack propagation direction is not 

known. Two intersecting Wallner lines may be analyzed as shown in Figure 

7.30b: 

vc sin φ = (7.26) 

v cos 2 β +cos 2 t 1 β 2 + 2cos β1 2 cos β cos φ

and if the Wallner line intersection is symmetric (β1 = β2 = β), then: 

φ  sin   vc  2  (7.27) = vt cos (β )

7-49 
 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Elastic waves interactions from slowly-moving cracks run across the crack 

front fairly quickly and thus the Wallner line corresponds closely with the crack 

front shape. The minimum velocity that can be estimated by Wallner line 

analysis was estimated by Mencík as vc/vt = 0.175, using equation 7.27 with φ

=2o

, corresponding to nearly flat crossing Wallner lines. For glass with vt = 

3460 m/s, then vc,min = 60 m/s. So in general, quantitative Wallner line analysis 

for crack velocities is most suitable for fast moving cracks. 

The spacing between the legs of Wallner lambda lines, discussed in section 

5.4.5 and shown in Figure 7.31, can easily be used to estimate a crack velocity. 

In the time that the crack advances approximately a distance a/2, the elastic 

wave has traveled across the thickness once. For a straight crack front:

a Vc = Vs (7.28) 2 2 2 (a 

2) + t 

The formula can also be used to approximate crack velocities for slightly 

curved crack fronts as well. 

Propagating 

crack front 

a 

t 

Figure 7.31 The spacing between the lambda line legs can be used to estimate 

crack velocity. 

7-50 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Quantitative Analysis 

7.10 Slow Crack Growth 

Cracks may grow stably when loaded at stresses and stress intensities less than that 

necessary for the flaw to become unstable. As discussed in sections 5.9.1 and 5.9.3, 

glasses, oxide ceramics, and ceramics with glassy boundary phase are susceptible. 

The mechanism of slow crack growth at ambient temperature is a stress corrosion 

phenomenon whereby stressed cracks that are open to the environment are attacked 

by water (or another polar molecule.) The water may be in the form of liquid or 

gaseous molecules. The water molecules attack the strained silicate or oxide bonds 

at the crack tip causing them to rupture which in turn leads to stable crack extension. 

Other fluids that promote slow crack growth are listed in section 5.9.1. At elevated 

temperature, oxidation attack, grain boundary softening, or other mechanisms can 

lead to slow crack growth (section 5.9.3). 

Crack velocity depends strongly upon the stress intensity, KI. The larger KI, the 

faster the crack grows. Slight changes in it have dramatic effects upon crack 

velocity. As a crack size increases, it increases KI in accordance with equation 7.6. 

If stress and the crack shape are invariant, the crack will gradually accelerate. Data 

are often graphed on log - log axes as shown in Figure 7.32a. Straight lines imply a 

power law dependence of velocity upon stress intensity: 

N V = A K (7.29) I 

Region I crack velocity behavior often controls lifetime or rate effects on 

strength. (Region II occurs when insufficient amounts of the reactive element 

are available. Region III occurs close to KIc and does not require a reactive 

species. With a few exceptions, during most of the time a crack is growing it 

will be in Region I.) The slope of the line, N, is known as the slow crack 

growth exponent and is a critical parameter for reliability and rate effect 

analyses. Very high values of N (e.g., > 100) indicate considerable resistance 

to slow crack growth and hence, little rate sensitivity. Low values (e.g., 5 to 

30) indicate high susceptibility. Velocity data as shown in Figure 7.32a are 

usually collected with laboratory specimens under controlled testing conditions. 

Figure 7.32b shows how slow crack growth affects the lifetime of specimens 

loaded at constant stress levels below the fast fracture strength. This is referred 

to as “static fatigue” or “stress rupture” testing. The decrease in log strength 

varies with log time to the –1/N power. 

Slow crack growth can even affect outcomes in ordinary strength tests as 

shown in Figure 7.32c. The slower the test is run, the weaker are the specimens 

since the crack can grow a surprising amount during the time of loading. 

7-51 
 

 

 

 

 Fractography of Ceramics and Glasses 

(a) (b) 

(c) 

Figure 7.32 Slow crack growth manifests itself in different ways. (a) shows a 

crack velocity versus stress intensity (K-V) graph showing three regions. 

Usually such data is plotted on log-log axes. Fracture mechanics type tests are 

used to collect such data. (b) shows a graph of log stress – log time to failure, 

also known as a “static fatigue” or “stress rupture” plot wherein a constant 

stress is applied until a piece breaks. (c) shows a graph of strength as a 

function of stressing rate also with log - log axes, also known as a “dynamic 

fatigue” plot. The slow crack growth exponent N plays a prominent role in 

each instance. 

7-52 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Quantitative Analysis 

Figure 5.56 showed an example of a flaw in glass that grew as much as 20 

times its original size due to slow crack growth in water. Figure 5.55 showed a 

flaw in alumina that grew about twice its original size in a simple 30 second 

strength test. 

A comprehensive theoretical justification for the power law relationship 

between KI and V is lacking at the present time. There are alternatives, but the 

power law relationship lends itself to simple integrations and transformations 

and is usually adequate. It should be used with caution for any extrapolations 

beyond the times and velocities for which the data was collected, however. 

Crack-fluid interactions can cause abrupt changes in crack growth that can 

create markings such as scarps on fracture surfaces as shown in section 5.6. 

For example, an accelerating crack may outrun the ability of a fluid to keep up 

with the crack tip and this may cause cavitation and a scarp. Conversely, a 

decelerating crack may slow down enough that water can catch up to the crack 

causing a scarp. Sometimes a jump in the crack from region I behavior to 

region II behavior creates a scarp. 

The Newman-Raju stress intensity shape factors also explain one commonly 

observed trend: surface cracks that grow stably evolve shape and become 

semielliptical. Figure 7.33 illustrates the trend for a beam or plate in bending 

and Figure 5.44b an example. The stress intensity shape factor Y varies around 

the crack front. For a semicircular flaw, Y is greatest at the point where the 

crack intersects the outer surface, point B. At the internal point A, there is 

more material on either side of the crack front that can share the load, and 

hence the stress intensity shape factor is about 13 % less. Thus, the crack 

initially grows faster along the outer surface, but then, as the crack shape 

evolves into a semiellipse, the stress intensity evens itself out around the 

periphery. Newman and Raju40 showed examples of this for fatigue crack 

growth in metals, both for bending and direct tension stress fields. The 

semielliptical shapes differ in the two cases, depending upon the size of the 

crack relative to the specimen thickness. The identical trends have also been 

observed in ceramics and glasses.67,74,75,76,77,78 Small surface cracks in large, 

uniform, direct-tension loaded specimens or plates will assume a constant semi 

elliptical shape with an aspect ratio a/c of about 0.83. Cracks in bending stress 

fields will also gravitate to this shape if they are small relative to the thickness, 

but if they penetrate into the stress gradient, they will become more elongated 

sideways as they grow and the semiellipse ratio a/c will decrease as shown in 

Figures 7.33b and c. 

7-53 
 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a)

2cc 

2co 

ao 

ac A 

B• 

• 

 (b) 

t 

(c) 

(d) 

Figure 7.33 Crack shape often changes with growth. (a) shows an initially 

semicircular crack of size ao that changes to a semi elliptical shape as it grows 

in a bending stress field in a plate of thickness t. (b) shows concentric arrest 

lines from cyclic loading (same as Figure 5.47c) (c) shows tertiary Wallner 

lines from a rapidly moving crack (same as Figure 5.44b); and (d) shows the 

evolution of crack shapes in glass bend specimens converges to a shallow 

semiellipse, irrespective of the starting shape (hollow circles). (b courtesy D. 

Green. c is courtesy H. Richter. d is after Fett et al., Ref. 74) 

7-54 
 

 

 

 

 

 

 

 

 

 

 

 

Quantitative Analysis 

7.11 R-Curve Behavior 

R-curve behavior, which arises from crack-microstructure interactions, is a 

phenomenon that may improve fracture resistance in ceramics. It can give rise 

to stable crack extension prior to fracture. It is not operative in glasses. 

Unfortunately, R-curve behavior does not leave any intrinsic telltale markings 

on the fracture that the fractographer can interpret. 

Figure 7.34 illustrates the general concepts. Flat R-curve materials such as 

glasses and many fine-grained ceramics behave as shown in Figure 7.34a. 

Fracture toughness is independent of the crack size and has a specific value: 

KIc. When the combination of stress, flaw shape, and size reach the critical 

condition, unstable crack extension usually occurs. If there are substantial 

tensile stresses throughout the body, it will break. If, on the other hand, the 

stresses decrease either temporally as in thermal shock cases or spatially as in 

localized contact loading cases, the crack initially may be unstable, but then 

slows down and may even stop. 

Rising R-curve behavior occurs when the microstructure impedes crack 

extension. Several crack-microstructure interactions shown in Figure 7.16 can 

cause this. For example, yttria-stabilized tetragonal zirconia polycrystal (YTZP) is carefully processed so that after firing, the microstructure is composed 

of metastable tetragonal grains. These can transform to the monoclinic phase 

when a crack tip and its stress field approaches. This martensitic phase 

transformation consumes some energy that might otherwise drive the crack 

forward. In addition, the transformed grains expand by ≈ 4% thereby putting a 

compressive constraint on the crack tip. The process zone may only be a few 

grains wide in Y-TZP. R-curves can arise from other crack-microstructure 

interactions. For example, with coarse-grained ceramics, large grains behind 

the crack tip may not necessarily be immediately cleaved. They may act as 

bridges across the crack faces that reduce the stress intensity field on the crack 

tip. They can persist for long crack propagation distances. 

Phenomena such as these can lead to behavior shown in Figure 7.34b. Starting 

with an initial crack size, co, if the stress intensity Ka from an applied stress 

reaches the level of Ko, the crack begins to propagate. As it grows, Ka increases 

in proportion to √c if the stress is held constant as per eq. 7.6. The material’s 

fracture resistance KR may increase at a greater rate thereby retarding or 

arresting further crack propagation. Further propagation requires additional 

applied stress to increase Ka. The R-curve typically reaches a plateau such that 

further toughening is not possible. Fracture may occur before the plateau is 

reached if: 

7-55 
 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) 

(b) 

(c) 

- Ko 

K 

Figure 7.34 R-curves in ceramics. (a) shows the case of a material with a 

constant, or flat R-curve. (b) shows a general schematic of a rising R-curve. 

Crack extension leads to a greater applied stress intensity Ka, but the 

material’s fracture resistance curve KR may also increase. (c) shows Ka versus 

√c, in which case the Ka plots as a straight line with slope proportional to 

applied stress. Fracture occurs at ccrit where the rate of increase of applied Ka 

is greater than the slope of the KR curve. Initial flaws larger than ccrit may not 

exhibit any stable crack extension and may immediately go critical. 

7-56 
 

 

 

 

 

 

 

 

 

 

 

Quantitative Analysis 

(7.30) Ka = KR 

and 

dKa dKR >

dc dc (7.31) 

In other words fracture occurs when the applied Ka reaches the level of the Rcurve and the rate of increase in Ka with further extension is greater than the 

rate of toughening. To better show the relative increases, it is common to plot 

fracture resistance with K2 versus c, or K versus √c as shown in Figure 7.34c. 

In this fashion, the increase in Ka with crack size appears as a straight line with 

slopes that increase with σ. This Ka line can be more easily compared to the KR 

trend. Instability occurs at c = ccrit, a value well before the plateau is reached, 

after only a small amount of crack growth has occurred. 

There is no universal R-curve for a material.

79 The R-curve depends upon the 

starting size of the crack relative to the microstructure and its prior propagation 

history. It also depends upon the size, shape, and mode of loading of the 

specimen. R-curves for large cracks are not the same as those for small cracks. 

In other words, data from large crack fracture mechanics specimens are usually 

not applicable to natural small flaws. Large fracture mechanics specimens 

usually are effective in measuring plateau toughness values. A doublecantilever beam specimen has a dramatically different crack opening 

displacement than a notched beam or a surface flaw in a bend bar. The number 

of bridges behind the crack tip bridges depends upon the crack opening 

displacement and the greater the opening, the fewer the intact bridges. For 

example, a 20 μm grinding-induced crack may have already interacted with the 

microstructure and formed bridges or transformed some material during the 

crack formation. Figure 7.35 shows two possible examples. In contrast, a 20 

μm round pore may not have activated any prior toughening. 

It is extremely difficult to measure R-curves for small natural-sized flaws, and 

especially the initial Ko fracture resistance. The limited data that has been 

collected suggests the initial Ko may be very small, approaching single crystal 

or even grain boundary fracture toughness values. Some materials may have 

such small toughening zones that, for all practical purposes, they can be treated 

as having a set fracture toughness value. Many yttria-stabilized tetragonal 

zirconias (Y-TZPs) may be like this since the transformation zone is only a few 

grains large. On the other hand, magnesia partially-stabilized zirconia (Mg 

7-57 
 

 

 Fractography of Ceramics and Glasses 

(a) 

(b) 

Figure 7.35 Grinding cracks in transversely-ground rods that show evidence 

of stable crack extension (arrows) in a silicon nitride with rising R-curve 

behavior. (a) 657 MPa and (b) 641 MPa. (Ref. 80) 

7-58 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Quantitative Analysis 

PSZ) may experience considerable crack extension prior to fracture.81 

Many studies have used Vickers indentation artificial flaws, but they often do 

not behave like genuine flaws as discussed in the next section. The R-curves 

that are generated are not relevant to natural flaws or are simply wrong. For 

example, one study showed that without proper care, the usual analysis with 

Vickers indentation flaws produced the implausible result that a soda lime 

silica glass has a rising R curve.82 (The problem was that the standard 

indentation strength model fails to account for changes in crack shape, 

especially as cracks penetrate into strong flexural stress gradients.) The most 

credible R-curve data to date (that is applicable to strength limiting flaws) has 

been that collected by careful microscopic examination of crack growth from 

natural flaws in controlled loading experiments.77,78,83,84 

The instability point depends upon the starting crack size and the slope of the 

R-curve at small crack sizes as shown in Figure 7.34c. Indeed, the overall 

plateau value may be inconsequential in many cases. A particular crack may 

grow only a small amount before going unstable79,80,85 or it may immediately be 

unstable. Fett and Munz’s mathematical analysis is quite revealing.85 They 

showed that R-curve behavior led to negligible strength enhancements in an 

alumina that had bridging grains, quite simply because the natural flaws did not 

have a chance to grow very far before going unstable. Their analytical and 

experimental results may easily be understood by the simple realization that in 

strong aluminas the flaws are only 5 to 20 times larger than the average grain 

size. The stable crack extension may involve only a few grains before the flaw 

goes critical and it is impossible to generate many behind-the-crack tip bridges. 

Perceptible strengthening would only occur in very weak specimens (σ < 100 

MPa) with large initial flaws. Fett and Munz also showed that the R-curve 

behavior in their aluminas had negligible effect on the Weibull distribution85 

contradicting predictions of dramatic improvements in Weibull modulus based 

on indentation mechanics analysis.86 

Mecholsky et al.87 and Marshall80 suggest that a simple way to measure an 

effective R-curve is to measure strength and critical flaw size in strength test 

specimens and compute apparent Kc at fracture. Rising R-curves will be 

evident on a plot of Kc versus crack size. They showed evidence that the mirror 

to flaw size ratio changed significantly as a function of crack size. 

7-59 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

The fractographer should keep the above factors in mind. As noted above, Rcurve behavior does not leave any intrinsic telltale markings on the fracture 

surface. R curve behavior can throw off flaw size calculations if they are based 

on a single point value for fracture toughness in equations 7.7 or 7.22. In 

principle, the fractographer may be able to find a region of stable extension 

around a flaw, but in practice this may be difficult especially if the toughening 

phenomena and the microstructure produce a rough fracture surface. One 

indication of potential R-curve behavior may be a systematic ccalc ≠ cmeas trend. 

As noted previously, if the fractographer detects and measures only the initial 

flaw size, then ccalc > cmeas. In this case it is the fractographic size measurement 

that is off. Alternatively, if stable crack extension due to R-curve behavior 

does occur, but the fractographer uses a long crack (plateau) value of fracture 

toughness, then the fracture toughness is overestimated and the calculated crack 

size is wrong, and again ccalc > cmeas. 

A rising R-curve does not necessarily translate to superior mechanical 

properties. Careful microstructural control is usually required. If reinforcing 

agents are not well dispersed and are clumped together, they can act as 

strength-limiting flaws. Alternatively, local regions depleted of the reinforcing 

elements may also have ordinary flaws. Many toughened ceramics have not 

shown commensurate increases in strength or Weibull moduli (which is a 

measure on consistency of strengths as discussed in section 7.15), contrary to 

expectations. The reason is simple. The very microstructural changes that 

enhance crack-microstructure interactions also create non-uniform 

microstructural regions, flaws, and microflaw concentration regions. That is 

why Y-TZP with a modest fracture toughness of 5 MPa√m to 6 MPa√m is 

stronger than Mg-PSZ. The Y-TZP has a very uniform sub-micron grain size 

and small ≈ 20 μm diameter pore flaws (Figure 6.6f). The Mg-PSZ has a 

fracture toughness of up to 10 to 20 MPa√m, but a coarse-grain microstructure 

loaded with grain boundary faults and concentrations of micropores (Figures 

6.75, 6.76, 6.77). The flaw size may be hundreds of μm in size. 

7-60 
 

 

 

 

 

 

-

Quantitative Analysis 

7.12 Indentation Mechanics 

A substantial literature exists on the mechanics of flaws made with Vickers or 

Knoop indentations. Figures 7.36 illustrate some of these “controlled flaws.” 

The indentations create not only a residual impression that is measured in a 

hardness test, but cracks and a concentrated damage zone directly underneath 

the indentation. 

(a) 

Boron carbide 

(b) 

Median 

Crack 

Lateral 

Crack 

BK 7 glass 

Figure 7.36 Indentation generated damage sites. (a) and (b) show common 

crack patterns and examples around Vickers and Knoop indentations. The 

author prefers the Knoop indenter since it makes a simple crack system. 

7-61 
 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) (b) 

(c) 

Figure 7.37 The surface crack in flexure method uses a Knoop indenter to 

create a semi elliptical surface crack in a bend bar (a). (b) shows a section 

view of the indentation and the median crack below it. The residual stress 

damage zone is removed by polishing or hand grinding leaving a controlled 

surface flaw in the test piece. (c) shows a Knoop crack in a silicon nitride 

(arrows). 

7-62 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Quantitative Analysis 

Controlled flaws for fracture mechanics experiments should be made with the 

Knoop indenter.88,89 The flaw geometry is simpler and more easily controlled. 

Loads of 19.6 N to 294 N are used to create a semi-elliptical surface cracks in 

ceramic bend bars (Figure 7.37). Loads of 19.6 N to 29.4 N are used for 

glasses. The indentation residual stress damage zone and any lateral cracks are 

removed by polishing and then the test piece broken in flexure. Fractographic 

techniques are used to find and measure the flaw size. Other precrack 

examples are shown in Figures 3.26, 5.55, 5.57, and 6.53. 

Calculations are simple. The Newman-Raju stress intensity shape factors 

described in section 7.5.2 are used with equation 7.6 to compute stress 

intensities. The name of the method was changed from “controlled surface 

flaw” to the “surface crack in flexure (SCF)” method by the author in 1994 in 

order to avoid confusion with some Vickers indentation methods and to also 

use nomenclature then in use in the fracture mechanics community.90 The SCF 

procedure was the primary method used to prepare the world’s first standard 

39,91 reference material for the property fracture toughness, KIc. It was formally 

standardized by the American Society for Testing and Materials41,92 the 

European Committee for Standards,93 and the International Organization for 

Standards.44 The method does not work on soft, porous, or very tough 

materials, since median cracks will not form. Details on the SCF method are in 

these standards and in References 89 and 90. 

Vickers indentation procedures are widely used since they usually do not 

require fractographic examination. This supposed advantage leads to severe 

drawbacks with the method, however. In 200694 and 200795 the author and 

Prof. R. Bradt wrote about the limitations of Vickers indentations for fracture 

mechanics studies. Vickers generated flaws are much more complex than 

Knoop median cracks, and since the Vickers flaws are tested without removal 

of the residual stresses and damage zones, all sorts of complications arise. 

Rather than a single median crack, Vickers indentations have multiple median, 

radial, and lateral cracks. The cracks also have a complex stress residual stress 

damage associated with them.96,97 Rather than envision these Vickers 

indentation sites as model flaws with discrete median cracks with a modest 

residual stress pulling on the faces, it is actually more appropriate to regard 

them as Vickers damage zones. A semicircular median crack is typically 

assumed to pop in underneath the indentation, although the crack type and 

formation sequence is quite complicated and varies dramatically with 

material.98,99,100 The residual stress damage zone is approximated by a point 

force acting to open the crack. The stress intensity factor KIr acting on the flaw 

7-63 
 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

in this case has the relationship shown in Figure 7.17b, namely: 

P K = χ (7.32) Ir 1.5 c 

where c is the crack size, P is the indentation load, and χ is a constant that 

incorporates geometry and residual stress field factors. 

If a test piece with a Vickers flaw is broken, KIr combines with the stress 

intensity created by a far field tensile stress to produce a total KI: 

P (7.33) KI = Yσ a c + χ 1.5 c 

The first term on the right dominates for large cracks. The second term is 

significant for small crack sizes, but rapidly diminishes as the crack extends 

away from the residual stress damage zone. The net effect of the two terms is 

to cause a crack to extend stably somewhat prior to catastrophic propagation. 

The extent of stable propagation has been estimated to be 2.5 times the original 

flaw size when the testing is done in inert environment and all the assumptions 

in the analysis are upheld.97 

It is clear from above that the local residual stresses dramatically alter the local 

stress intensity field, and that a simple application of equations 7.6 (or eq. 7.33 

without the second term on the right) based on only the far-field tensile 

stresses, will be faulty. Typically ccalc will be larger than cmeas, if the latter is the 

original flaw size and stable crack extension is not detected. 

By the time the crack propagates to the extent that it is forming the mirror 

boundary markings, the effect of the localized residual stress term is 

insignificant.101 Hence, the mirror size is unaffected by the indention residual 

stresses and the stress-mirror size relationship equation 7.4 is unaffected. The 

mirror to initial flaw size ratio is affected, however, and can be larger than 

values for annealed flaws.101 Indentation residual stresses do not affect 

branching distances or the fragmentation patterns. 

Evidence of stable crack extension from indentation localized residual stresses 

may be difficult to detect on fracture surfaces. Figure 5.57 shows an example 

in a fine-grained silicon nitride. Stable extension is most often detected in 

laboratory conditions with careful microscopy on a polished outer surface, or 

through the material if the material is transparent. 

7-64 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Quantitative Analysis 

The observant reader will note that Vickers indentation flaws do not resemble 

the majority of origins shown in Chapter 6. Notwithstanding claims to the 

contrary, it is not at all clear whether indentation damage sites simulate genuine 

flaws in ceramics and glasses. They may model contact damage sites and/or 

machining cracks, but Figure 6.30 in the previous chapter shows that actual 

contact damage sites are often more irregular, with significant amounts of 

material removed by lateral cracks and spalls. The contact-generated residual 

stresses are often partially or completely relieved. Service impact conditions 

are rarely as controlled as those in an indentation process, where an ideal 

diamond slowly applies load perpendicular to the surface. 

The localized indentation residual stress field of a laboratory indentation has a 

very stabilizing influence on a flaw, allowing it to have extensive stable 

extension that may not be experienced by genuine flaws. The previous section 

on R-curve effects mentioned some examples where predictions of enhanced 

Weibull modulus based on indentation flaw analysis were misleading or wrong. 

In summary, there has been an over reliance on Vickers indentation flaws. 

From some of the literature one might think that are they are more important 

than genuine flaws. Fractographers know otherwise. 

7.13 Fractal Analysis 

Fractal analysis is a tool that may be used to characterize irregular surfaces, 

such as fracture surfaces. It is not used to find and characterize fracture origins, 

but to characterize the roughness or unevenness of a surface. Fractal geometry 

is a non-Euclidean geometry that exhibits self-similarity and scale invariance. 

Self-similarity means that a geometric shape in one location is similar to a 

geometric shape elsewhere. Scale invariance means that the geometric shape is 

similar irrespective of its size. Fractal analysis is unlike classical surface 

roughness measurements that characterize the heights between peaks and 

valleys. Fractal analysis measures the extent of the “wiggliness,” “tortuosity,” 

or irregularity of the surface.102,103 

Classical Euclidean geometry has a simple relationship between area and 

length. So for example, a circle of diameter D has the well-known 

relationships in accordance with Euclidean geometry: 

7-65 
 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

π D2 (7.34) Area = 4 

and 

Perimeter = π D (7.35) 

so that: 

Perimeter ∝ Area (7.36) 

This is the case for a perimeter line that is unwavering and follows the circle 

rim exactly. If on the other hand, the perimeter line had perturbations and 

wiggles, the perimeter length is greater. If the perimeter line is magnified and 

studied in greater detail, it is possible that each wiggle or perturbation itself 

may be seen to have yet smaller wiggles creating additional length and so on. 

That is the nature of fractal dimensions. The greater the magnification used, 

and the finer the measuring stick used to measure the lengths, the longer is the 

overall perimeter length. This leads to the somewhat unsatisfying outcome that 

the perimeter length actually varies with magnification and the size of the 

measuring scale. The perimeter of a fractal shape around an area can be 

characterized by the following relationship: 

1−D L = L0 s (7.37) 

where L is the length of the line, Lo is a constant, s is the measuring scale or 

ruler size, and D is the fractal dimension which in this instance is between 1 

and 2. If the line has no wiggles (and is a Euclidean line) then D = 1, and 

L = Lo irrespective of the size of ruler used to measure it. On the other hand if 

the line is fractal, and D is between 1 and 2, then L varies with the size of the 

ruler. The exponent of s is negative and the smaller s is, the larger is L. The 

same type relationship applies if areas were to be used in equation 7.34, in 

which case D would be have a value between 2 and 3. The value of D-1 is 

often written as D*

, the fractional part of the fractal dimension. 

Two ways to measure the fractal dimension of a fracture surface are the 

fracture profile technique and the slit island analysis.103 The fracture profile 

technique simply makes a vertical cross section through a fracture surface. The 

profile with all its ups and downs is analyzed. In the slit island technique, a 

horizontal cross section is made through a fracture surface to reveal islands 

corresponding to the peaks of hills and ridges and their wiggly outlines. This 

process is repeated and measurements remade much like in a classical 

metallographic serial section analysis. Both procedures can be tedious, but 

computer analysis and image analysis software can simplify this task. Confocal 

7-66 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Quantitative Analysis 

optical microscopy would seem to be an ideal tool, since it may be able to 

examine the fracture surface directly and section it automatically without the 

need for repetitive polishing.104,105 A typical fractal analysis measures a 

perimeter length or roughness number with finer and finer measuring intervals. 

A graph of log length versus log measuring interval size is constructed and the 

slope used to estimate the fractal dimension. 

Fractal analysis has usually been used to correlate fracture toughness with the 

fractal dimensions. The processes that enhance fracture toughness often 

increase the roughness of a fracture surface. Passoja, Mecholsky and 

colleagues have pioneered correlations of this type for ceramics.102,103,106,107 

For example, they have shown that: 

KIc = Ko + A(D* )

1

2 (7.38) 

where Ko is a baseline toughness of the material for a smooth (Euclidean) 

fracture surface, D* is the fractional part of the fractal dimension and A is a 

constant that can be related to the elastic modulus and the atomic dimensions of 

the structure of the material. The implication of this relationship is that only a 

small portion of a fracture surface is needed to determine the fracture 

toughness. In principal, this could be used in forensic analyses to determine if 

the material was poorly processed and had a lower than expected fracture 

toughness. The fractal analysis gives no information about the fracture origin, 

however. 

Mecholsky and Freiman107 showed that variations in the mirror to flaw size 

ratio correlated with the fractal dimension of the fracture surface well outside 

the mirror: 

R 1

∝ * (7.39) c D 

where R is either the mirror-mist, mist-hackle, or branching distance, and c is 

the flaw size. In other words, the very processes that contribute to formation 

of roughness at the mirror boundary are related to those that create roughness 

elsewhere in areas remote from the origin. 

At the beginning of this section it was noted that fractal analysis is a tool to 

characterize fracture surfaces. Most fractal analyses have focused on using the 

fractal dimension as a materials science tool to study the microstructure or 

material properties. Additional work showing practical applications is needed. 

7-67 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

7.14 Estimation of Residual Stresses 

7.14.1 Introduction 

Residual stresses may exist on many levels as discussed in section 4.19.5. 

They may be local to a flaw, as in a Vickers indentation or an inclusion with a 

mismatch in properties with the matrix. They may be distributed through the 

thickness, as in surface compression and internal tension in chemically- or 

thermally-tempered glasses. They may be global from differential shrinkage 

during firing. Their effects on flaws may be very difficult to analyze unless 

some assumptions are made about their character and distribution. 

For example, residual stresses from grinding vary with direction, whether 

parallel to or perpendicular to the grinding direction. They are strongly 

compressive right at the surface, but rapidly decrease and become tensile a 

short distance (5 μm to 20 μm) below the surface. A grinding crack tip may be 

in tension or in compression depending upon how deep it is. The net effect of 

the stress gradient on the flaw as a whole is a complex fracture mechanics 

problem, but analytical attempts have been made to determine the effect on 

flaws with mixed success. The reader is referred to the work of Holstein et 

al.108 for a good example. It is often difficult to determine whether the net 

effect is tensile or compressive. The weak point in these analyses is the lack of 

adequate quantitative information about the residual stress spatial profile. 

7.14.2 Estimates of residual stresses from fragmentation 

Empirical estimation of residual stresses in tempered glass plates by fragmentation patterns has been covered in sections 7.2.2 and 7.3.2. 

7.14.3 Estimates of residual stresses from fracture mechanics 

analysis of flaws 

The usual approach in dealing with residual stress problems is to rely on the 

principle of superposition. That is, the effect of stresses on a location or flaw in 

a body is additive, whether the stresses are mechanical, transient thermal, or 

residual. To a first approximation, residual stresses, σr, simply add to or 

subtract from the applied stresses acting on the flaw from known external 

sources, σa, so that the net stress acting on a site or a flaw is: 

σ net = σ a r + σ (7.40) 

Tensile residual stresses add to externally applied tensile stresses and 

7-68 
 

 

 

 

 

 

 

 

 

 

 

Quantitative Analysis 

reduce the stresses or forces necessary to cause fracture from a particular flaw. 

Compressive residual stresses acting on a flaw must be overcome and require 

greater applied stresses. (The usual convention is that compressive stresses are 

negative.) For the case of a flaw loaded with a far-field stress σa with a 

residual stress σr that is constant in the vicinity of the flaw: 

K = Y σ a (7.41) Ic net 

Rearranging: 

K (7.42) σ = Ic 

r − σ a

Y a 

If a convenient flaw is found that has a shape that is conducive to a fracture 

mechanics analyses and KIc, σa, and the flaw size, a, are known, then equation 

7.42 can be used to estimate σr. 

7.14.4 Estimates of residual stresses from fracture mirror size 

analysis 

Residual stresses can cause systematic deviations in the stress - mirror size 

trends presented previously. Trend deviations are shown in Figure 7.38. The 

easiest interpretation is in Figure 7.38a, whereby a nonzero intercept indicates 

the magnitude and sign of the residual stress. A compressive residual stress 

shifts the line upward since a greater applied stress, σa, is necessary to cause 

fracture. If a single mirror is measured and the externally applied stress, σa, is 

(a) (b) 

slope < - 1⁄2 

σr 

-1⁄2 

log 

σ

a 

log R R = 1 

log R = 0 

σ = A / √1 

σr 

R 

1 

σaA 

σr 

Figure 7.38 Systematic deviations in the fracture mirror size trends for 

applied stress, σa, as a function of mirror radius are indicative of residual 

stresses. (a) shows that residual stresses create nonzero intercepts. (b) shows 

that residual stresses alter the slope. 

7-69 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

known, then: (σ a + σ r ) R = A (7.43) 

and 

σ = A − σ (7.44) r

R a 

A good example of the application of this approach of measuring mirror-mist 

and mist—hackle boundaries on glass rods with cladding is in a paper by 

Mecholsky and Drexhage.

109 Similarly, the branching distances were used by 

Conway and Mecholsky 110 to estimate residual stresses in tempered glass 

plates. Additional details by this method are in Appendix D. 

7.14.5 Estimates of residual stresses from indentation crack sizes 

Surface residual stresses may be estimated by measuring the crack lengths 

emanating from the corners of Vickers indentations in a variation of the 

procedure described in section 7.12. If there are surface compression stresses, 

then the indentation crack lengths at a particular indentation load will be shorter 

than those in for an annealed piece of material.111,112 Assuming that the 

indentation crack slows down and stops at KI = KIc, and that slow crack growth 

doesn’t affect its length, then: 

P K = χ + Yσ r a (7.45) Ic 1.5

c 

where χ is the dimensionless contact constant which incorporates details of the 

indenter specimen contact conditions. Testing is usually done over a range of 

loads and crack lengths. This methodology is susceptible to all of the 

shortcomings addressed in section 7.12. If a test piece with residual stresses is 

indented and then broken, say in a plate or bend bar form, then the stress 

intensities acting on the crack sum: 

K = Y (σ +σ ) c + χ

P 

1.5 (7.46) Ic a r c 

where σa is the applied stress from external loading, σr is the surface residual 

stresses at the location of the flaw, and the last term on the right is the stress 

intensity from localized indentation residual stresses. In practice, considering 

all the assumptions in indentation analysis, and the probable stress gradients, it 

may be difficult to reliably solve this equation to obtain σr. 

7.14.6 Estimates of residual stresses by other means 

Qualitative assessments can be made from mirror shapes as discussed in section 

5.2.3 and also by changes in the mirror to flaw size ratio in section 7.7. 

7-70 
 

 

 

 

 

 

 

 

Quantitative Analysis 

7.15 Weibull Analysis 

Ceramic and glass strengths depend upon the size and shape of the specimens 

and the mode of loading. Specimen strengths vary due to variations in the size, 

severity, location, and density of flaws. Strength variability is usually 

analyzed in accordance with the Weibull distribution,

113 which is based on the 

premise that the weakest link in a body controls strength. A strength test 

produces two key bits of information: a strength datum and a fracture origin 

flaw. The flaws are just as important as the strength. “Flaw management” is 

an important strategy for process control and reliability improvements. 

Figure 7.39 Fractographically-labeled Weibull graph. Four-point flexural 

strength data for a SiC whisker-reinforced alumina. “G’ is an abbreviation for 

a large grain origin type, “I” is for inclusion, and “?” is unknown. The slope 

of the fitted line is the important Weibull modulus (m) parameter. It is a 

measure of strength variability. 

7-71 
 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

The Weibull two-parameter distribution is: 

m  σ  a −

 



 σθ  Pf = 1. − exp (7.47) 

where Pf is the probability of fracture, σa is the applied stress, and m and σθ are 

constants called the Weibull modulus and the characteristic strength of the 

specimen, respectively. Strength data is typically ranked from weakest to 

strongest and then plotted on a special set of axes intended to linearize the data 

as shown in Figures 7.39 and 7.40. 

With Weibull analyses, the maximum nominal stress in the body is used. 

Stress is not adjusted for location. If the data fits the Weibull distribution, then 

it should fit on a line with the slope equal to the Weibull modulus, m, and the 

characteristic strength σθ corresponding to the strength of the specific test 

specimen configuration at the 63.2 % level. Figure 7.39 and 7.40 show two 

actual data sets with 100% fractographic analysis. The goodness of fit to the 

Weibull function can be assessed by eyeball by comparing the fitted line to the 

data. Both data graphs are fractographically-labeled and have not only the 

individual strength values, but also the flaw identities. 

Figure 7.39 for the alumina with silicon carbide whisker reinforcements shows 

that the strength-controlling flaws were almost always large alumina grains. 

The whiskers that were added to enhance KIc had chemistries or impurities that 

caused grain growth in the matrix. The good news was that one flaw type 

primarily controlled strength and the Weibull distribution fit quite well. 

In contrast, Figure 7.40 shows a SiAlON material that had many different 

concurrent flaw types and a very low Weibull modulus.114 Brittle materials 

design with such a material would be problematic, since each flaw type has its 

own distribution. Strength scaling with size would be complicated. The tester 

might have wondered why the curve had so many wiggles. The fractographic 

analysis makes it clear why the single Weibull line does not fit very well. 

Weibull originally derived his distribution on semi-empirical grounds. No 

assumptions about flaws or their nature were made. He simply assumed that 

small elements in a body had a local strength. He close a plausible distribution 

for the local variations in strength. He then integrated the risk of rupture for all 

elements to compute the failure probability for the whole specimen assuming 

that fracture would occur at the weakest link or element. His perceptive 

analysis has been theoretically verified by subsequent analyses by a number of 

authors, but most especially by Jayatilaka and Trustrum.115,116 They showed 

7-72 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Quantitative Analysis 

Probability of Failure (%)

99.9 

99 

95 

90 

80 

63.2 

50 

20 

10 

5 

2 

1 

0.5 

300 400 500 600 700 800 900 1000 

MLE Analysis 

ASTM C 1239 

σθ = 602 

mu = 7.9 

ASTM C1322-96 

Stress Identity 

699 ? 

688 Machining 

687 Inclusion 

675 ? 

672 Porous Region 

664 Handling 

641 Pore 

636 Inclusion 

618 Pore 

614 Agglomerate 

602 Porous Seam 

595 Handling 

570 Inclusion 

569 Pore 

554 ? 

545 Inclusion 

545 Pore 

540 Agglomerate 

527 Agglom./Pore 

519 Porous Region 

518 Porous Seam 

517 Pore (linear) 

513 Agglom./Pore 

506 Aggolm./Pore 

505 Handling 

483 Handling 

482 Porous Seam 

466 Pore 

457 Agglom./Pore 

450 Pore 

569 MPa Mean 

76 MPa Std. Dev. 

Sialon 

Stress (MPa) . 

Figure 7.40 Fractographically-labeled Weibull graph. Four-point flexural 

strength data for a SiAlON ceramic. 

that flaw size distributions could have the large flaw size tails matched by an 

inverse power law distribution as shown in Figure 7.41. Using this inverse 

power law distribution and applying classical fracture mechanics (equation 

7.7), and weakest-link theory, they were able to mathematically derive the 

Weibull distribution. The Weibull modulus is directly related the flaw size 

distribution exponent. The Weibull distribution is now based on a solid 

theoretical footing. Danzer has considered more general distributions, but 

concluded that if the frequency of flaws follows the inverse power law, then the 

material may be deemed a “Weibull material.117 In the 1990s and 2000s, a 

number of studies confirmed the large flaw size distribution tails can indeed be 

matched by the inverse power law. The reader is referred to papers by Prof. 

Uematsu in the Origins in Ceramics section of Appendix A for many examples. 

Quinn and Morrell’s paper 118 has a useful review of a number of successful 

applications using Weibull analysis and it lists all of the assumptions necessary 

for Weibull strength scaling with size. 

7-73 
 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) 

α 

m = 2n’ -2 (b) 

Figure 7.41 

distribution shown in (a). H’(c) is the probability for occurrence of a flaw of 

size c. n’ and α are constants. The largest flaws control strength and that 

portion of the distribution can be modeled by an inverse power law with 

exponent n’. (b) shows the Weibull distribution for strength. The Weibull 

modulus m is equal to 2n’ – 2. 

Average bulk properties such as density, surface finish, or grain size often do 

not correlate with strength. Strength often depends upon infrequent or aberrant 

microstructural features, namely flaws, and not on the average microstructure. 

In some cases, the aberrant features do scale with the average microstructure, 

so there may be a correlation between strength and average bulk microstructure 

or surface finish. One important aspect of the Weibull model is that the larger 

the structure or test piece is, the more apt it is to contain a severe flaw. Hence, 

strength usually varies inversely with component size: the larger the component, the weaker it is likely to be. Strengths can be correlated or scaled by 

The Weibull distribution can be derived from the flaw size 

7-74 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

Quantitative Analysis 

comparing effective volumes or effective surfaces. It is beyond the scope of 

this Guide to delve into this topic, but suffice to say that such strength scaling 

analyses depend upon crucial assumptions. For example, a Weibull effective 

volume scaling approach is justified if fractography confirms that strength 

limiting flaws are volume-distributed. One often overlooked assumption in the 

Weibull analysis is that specimens must contain a minimum number of flaws. 

There must be a number of “chain links” present in the specimen for it to have 

strength controlled by a “weakest link.” A specimen with a single flaw has 

only one strength controlling link. It is inappropriate to apply Weibull 

statistics to specimens with a single indentation flaw. It is also dangerous for 

engineers to ignore or discard a few specimens from the low strength tail of a 

Weibull distribution. These “atypical” or “non representative outcomes” may 

reflect a genuine flaw type that will cause trouble if the product is scaled up for 

mass production or if the components are designed to have very low 

probabilities of fracture (e.g., Pf << 1%). 

Strength scaling analyses become very complex if more than one flaw type 

(e.g., pores and inclusions; or pores, inclusions, and machining damage) 

controls strength such as shown in Figure 7.40. Various flaw types may 

suddenly appear or disappear as strength controlling flaws as component sizes 

are scaled up or down. See the Quinn and Morrell review115 for more 

information on these fascinating topics. Sung and Nicholson published results 

on a good study119 where fractography was effectively used to identify five 

different flaw types in a 4 mass % yttria-stabilized, tetragonal polycrystal 

material: fiber burn-out pores, agglomerates, iron inclusions, alumina 

inclusions, and pores. By careful process control and “flaw management,” they 

were able to eliminate two flaw types, and control the remaining three. Small 

pores, small alumina inclusions, and small agglomerates became the primary 

strength-controlling flaws. 

The author is often asked whether the Weibull distribution is the best for 

modeling strength variations, and my answer is always yes, but with some 

caveats. The Weibull distribution is the only one that the author knows of that 

permits size scaling of strength. Indeed, the ultimate proof that the Weibull 

distribution is valid is not whether one set of data can be plotted linearly on a 

Weibull graph, but whether size scaling is successful. In the dental materials 

community, confusion about the validity of the Weibull model prompted Dr. J. 

Quinn and the author to write a review article120 on this topic. Everything in 

the review is applicable to other ceramics and glasses as well. 

An amazing new study on 5,100 alumina flexural strength specimens by Gorjan 

and Ambrožič 121 that confirmed the Weibull distribution was perfectly 

7-75 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

suitable. It was better than the log normal, the normal, or the gamma 

distributions. The author congratulated them on their study and offered to do 

some fractography. Imagine the challenge! It was learned that they broke 

even more and had reached 11,700 tests. The Weibull distribution still held 

true. Regrettably, they did not save any specimens for the author to look at. 

What a pity! 

Chapter 7 References 

1. V. D. Fréchette, Failure Analysis of Brittle Materials, Advances in Ceramics, 

Vol. 28, American Ceramic Society, Westerville, OH, 1990. 

2. L. Orr, “Practical Analysis of Fractures in Glass Windows,” Materials Research 

and Standards, 12 [1] (1972) 21 - 23. 

3. V. D. Fréchette and T. A. Michalske, “Fragmentation in Bursting Glass 

Containers,” Bull. Am. Ceram. Soc., 57 [4] (1978) 427 – 429. 

4. V. D. Fréchette and S. L. Yates, “Fragmentation of Glass Bottles by Impact,” J. 

Am. Ceram. Soc., 72 [6] (1989) 1060. 

5.	 E. R. Fuller, Jr., S. W. Freiman, G. D. Quinn, J. B. Quinn, and W. C. Carter, 

“Fracture Mechanics Approach to the Design of Glass Aircraft Windows: A Case 

Study,” pp. 419 – 430 in Proc. SPIE Symp. on Windows and Dome Technologies 

and Materials, IV, 24-29 July 1994, San Diego, CA, SPIE Volume 2286. 

6. R. Morrell, “Fractography of Brittle Materials,” Measurement Good Practice 

Guide, #15, National Physical Laboratory, Teddington, Middlesex, UK, 1999. 

7. E. B. Shand, “New Information on Glass Fracture,” pp. 117 – 128 in Proc. Amer. 

Scientific Glass Blowers Soc. Symp., June 1969. 

8.	 J. Wang, “Fracture and Stress Pattern Correlations in Glass Plates,” Masters 

Thesis, Texas Tech Univ, August 1986. Avail. on-line from Texas Tech Univ. 

9. R, A. McMaster, D. M. Shetterly, and A. G. Bueno, “Annealed and Tempered 

Glass,” pp. 453 - 459 in Engineered Materials Handbook, Vol. 4, Ceramics and 

Glasses, S. Schneider, ed., ASM, Metals Park, OH, 1991. 

10. N. Shinkai, “The Fracture and Fractography of Flat Glass,” pp. 253-297 in 

Fractography of Glass, eds. R. C. Bradt and R. E. Tressler, Plenum, NY, 1994. 

11.	 K. Akeyoshi, E. Kanai, K. Yamamoto, and S. Shima, Asahi Garasu Kenkyu 

Hokoku, Report of the Research Lab, Asahi Glass Co., 17 [1] (1967) 23 - 26. 

12.	 K. Blank, “Some Practical Aspects of Thermally Strengthened Glass,” pp. 485 – 

499 in Strength of Inorganic Glass, ed. C. Kurkjian, Plenum, NY, 1985. 

13. J. M. Barsom, “Fracture of Tempered Glass,” J. Am. Ceram. Soc., 51 [2] (1968) 

75 - 78. 

14. J. Menčík, “Strength and Fracture of Glass and Ceramics,” Glass Science and 

Technology 12, Elsevier, NY, 1992. 

15. D. Shetty, A. Rosenfield, and W. Duckworth, “Crack Branching in Ceramic Disks 

Subjected to Biaxial Flexure,” Comm. Am. Ceram. Soc., Jan. 1983, C10 - C12. 

16. D. Hull, “Influence of Stress Intensity and Crack Speed on the Fracture Surface 

Topography: Mirror to Mist to Macroscopic Bifurcation,” J. Mat. Sci., 31 (1996) 

4483 – 4492. 

7-76 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Quantitative Analysis 

17. U. Soltesz, H. Richter, and E. Sommer, “Influence of Internal Stresses on the 

Development of Cracks in Glasses,” pp. 2303 – 2310 in Advances in Fracture 

Research, eds. D. Francois et al., Pergamon Press, NY, 1980. 

18.	 G. D. Quinn, “Guidelines for Measuring Fracture Mirrors,” pp. 163 – 190 in: 

Fractography of Glasses and Ceramics, Vol. 5, eds. J. R. Varner, G. Quinn, M. 

Wightman, American Ceramic Society, Westerville, OH, 2007. 

19. Private Com., Letter from L. Orr to G. Quinn on December 15, 2001. 

20. W. C. Levengood, “Effect of Origin Flaw Characterization on Glass Strength,” J. 

Appl. Phys., 29 [5] (1958) 820 – 826. Footnote on page 821. 

21. J. W. Johnson and D. G. Holloway, “On the Shape and Size of the Fracture Zones 

on Glass Fracture Surfaces,” Phil. Mag., 14 (1966) 731 - 743. 

22.	 E. B. Shand, “Breaking Stress of Glass Determined from Dimensions of Fracture 

Mirrors,” J. Am. Ceram. Soc., 42 [10] (1959) 474 – 477. 

23.	 M. J. Kerper and T. G. Scuderi, “Relation of Strength of Thermally Tempered 

Glass to Fracture Mirror Size,” Bull. Am. Ceram. Soc., 44 [12] (1965) 953 – 955. 

24. M. J. Kerper and T. G. Scuderi, “Relation of Fracture Stress to the Fracture 

Pattern for Glass Rods of Various Diameters,” ibid, 45 [12] (1966) 1065 - 1066. 

25. H. P. Kirchner, R. M. Gruver, and W. A. Sotter, “Fracture Stress-Mirror Size 

Relations for Polycrystalline Ceramics, Phil. Mag., 33 [5] 775 – 780 (1976). 

26.	 Glass Industry, 36 [10] (1955) 523. 

27. E. B. Shand, “Experimental Study of Fracture of Glass: II, Experimental Data,” J. 

Am. Ceram. Soc., 37 [12] (1954) 559 – 572. 

28. N. Terao, “Sur une Relation entre la Resistance a la Rupture et le Foyer 

d’Eclatement du Verre,” J. Phys. Soc. Japan., 8 [4] (1953) 545 – 549. 

29. H. P. Kirchner and J. W. Kirchner, “Fracture Mechanics of Fracture Mirrors,” J. 

Am. Ceram. Soc., 62 [3-4] (1979) 198 – 202. 

30. H. P. Kirchner and J. C. Conway, Jr., “Criteria for Crack Branching in Cylindrical 

Rods: I, Tension; and II, Flexure,” ibid, 70 [6] (1987) 413 – 418 and 419 - 425. 

31. H. P. Kirchner and R. M. Gruver, “Fracture Mirrors in Alumina Ceramics,” Phil 

Mag., 27 (1973) 1433 – 1446. 

32. M. J. Kerper and T. G. Scuderi, “Modulus of Rupture of Glass in Relation to 

Fracture Pattern,” Bull. Am. Ceram. Soc., 43 [9] (1964) 622 – 625. 

33.	 N. Shinkai, “Fracture Mirror Constants of Silicate Glasses,” Japan J. Appl. Phys., 

14 [1] (1975) 147-148. 

34.	 ASTM C 1678-07, Standard Practice for Fractographic Analysis of Fracture Mirror 

Sizes in Ceramics and Glasses, Annual Book of Standards, Vol. 15.01, ASTM, West 

Conshohocken, PA, 2007. 

35. G. D. Quinn, J. Eichler, U. Eisele, and J. Rödel, “Fracture Mirrors in a Nanoscale 

3Y-TZP,” J. Am. Ceram. Soc., 87 [3] (2004) 513-516. 

36. A. A. Griffith, “The Phenomena of Rupture and Flow in Solids,” Philos. Trans. R. 

Soc. Lond., A221 (1920) 163 – 198. 

37. A. A. Griffith, “The Theory of Rupture,” in Proc. 1st Int. Congress on Applied 

Mechanics, eds. C. B. Biezeno and J. M. Burgers, (1924), 55. 

38. G. R. Irwin, “Analysis of Stresses and Strains Near the End of a Crack Traversing 

a Plate,” J. Appl. Phys., 24 (1957) 361 – 364. 

7-77 
 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

39. SRM’s 2100, 2100A, and 2100B, “Fracture Toughness of Ceramics,” National 

Institute for Standards and Technology, Gaithersburg, MD, 1999. 

40. J. C. Newman and I. S. Raju, “An Empirical Stress Intensity Factor Equation for 

the Surface Crack,” Eng. Fract. Mech., 15 [1-2] (1981) 185-192. 

41. ASTM C 1322-96, “Standard Practice for Fractography and Characterization of 

Fracture Origins in Advanced Ceramics,” Annual Book of Standards, Vol. 15.01, 

ASTM, West Conshohocken, PA, 1996. 

42. ASTM C1421-99, “Standard Test Method for the Determination of Fracture 

Toughness of Advanced Ceramics,” ibid, 2000. 

43. ASTM E 740-88, “Standard Practice for Fracture Testing with Surface-Crack 

Tension Specimens,” Annual Book of Standards, Vol. 3.01, ASTM, West 

Conshohocken, PA, 1988. 

44. ISO 18756, Fine Ceramics (Advanced Ceramics, Advanced Technical Ceramics) – 

Determination of Fracture Toughness of Monolithic Ceramics at Room Temperature 

by the Surface Crack in Flexure (SCF) Method, International Organization for 

Standards, Geneva, SW, 2003. 

45.	 T. Fett, “An Extension of the Newman-Raju Formula,” Int. J. Fract., 33 (1987) 

R47-R50. 

46.	 T. Fett, “Stress Intensity Factors for Semi-elliptical Surface Cracks in a Plate 

Under Tension Based on the Isida’s solution,” ibid, 48 (1991) 139 – 151. 

47. T. Fett and D. Munz, Stress Intensity Factors and Weight Functions, 

Computational Mechanics Publ., Southampton, UK and Boston, MA, USA, 1997. 

48.	 M. Isida, H. Noguchi, and T. Yoshida, “Tension and Bending of Finite Thickness 

Plates with a Semi-elliptical Surface Crack,” Int. J. Fract., 26 (1984) 157-188. 

49.	 S. Strobl, P. Supancic, T. Lube, and R. Danzer,” Surface Crack in Tension or 

Bending – A Reassessment of the Newman and Raju Formula in Respect to 

Fracture Toughness Measurements in Brittle Materials,” J. Eur. Cer. Soc., 32 

(2012) 1491 – 1501. 

50. G. K. Bansal, “Effect of Flaw Shape on Strength of Ceramics,” J. Am. Ceram. 

Soc., 59 [1-2] (1976) 87-88. 

51. Y. Murakami, Stress Intensity Factors Handbook, Vols. 1 and 2, ed., Pergamon 

Press, NY, 1986. (Case 9.49) 

52. J. J. Mecholsky, Jr., S. W. Freiman, and R. W. Rice, “Fracture Surface Analysis 

of Ceramics,” J. Mat. Sci., 11 (1976) 1310–1319. 

53. G. K. Bansal and W. H. Duckworth, “Fracture Stress as Related to Origin and 

Fracture Mirror Sizes,” J. Am. Ceram. Soc., 60 [7–8] (1977) 304 – 310. 

54. G. K. Bansal, “On Fracture Mirror Formation in Glass and Polycrystalline 

Ceramics,” Phil Mag., 35 [4] (1977) 935- 944. 

55. S. R. Choi and J. P. Gyekenyesi, “Crack Branching and Fracture Mirror Data of 

Glasses and Advanced Ceramics,” NASA Technical Report TM 1998-206536, 

NASA Glenn Research Center, Cleveland, OH,1998. 

56.	 H. P. Kirchner, R. M. Gruver, M. V. Swain, and R. C. Garvie, “Crack Branching 

in Transformation-Toughened Zirconia,” J. Am. Ceram. Soc., 64 [9] (1981) 529-

533. 

7-78 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Quantitative Analysis 

57. A. I. A. Abdel-Latif, R. C. Bradt, and R. E. Tressler, “Dynamics of Fracture 

Mirror Boundary Formation in Glass,” Int. J. Fract., 13 [3] (1977) 349-359. 

58. M. Ramulu, R. C. Bradt, A. S. Kobayashi, and K. H. Yang, “A Dynamic Fracture 

Mechanics Interpretation of Multiple Mist Regions on Soda-Lime-Silicate Glass 

Fracture Features,” pp. 215 – 227 in Fractography of Glasses and Ceramics, 

Advances in Ceramics, Vol. 22, J. Varner, and V. Fréchette, eds., American 

Ceramic Society, Westerville, OH, 1988. 

59. D. A. Krohn and D. P. H. Hasselman, “Relation of Flaw Size to Mirror in the 

Fracture of Glass,” J. Am. Ceram. Soc., 54 [8] (1971) 411. 

60. J. J. Mecholsky, Jr., A. Gonzalez, and S. W. Freiman, “Fractographic Analysis of 

Delayed Failure in Soda-Lime Glass,” ibid, 62 [11-12] (1979) 577 – 580. 

61. R. W. Rice, “Ceramic Fracture Features, Observations, Mechanism and Uses,” 

pp. 5 – 103 in Fractography of Ceramic and Metal Failures, ASTM STP 827, 

ASTM, 1984. 

62. K. Gopalakrishnan and J. J. Mecholsky, Jr., “Quantitative Fractography of MixedMode Fracture in an R-curve Material,” J. Mat. Sci., 48 (2013) 7081 – 7087. 

63.	 R. W. Rice, “Ceramic Fracture Features, Observations, Mechanism and Uses,” 

pp. 5-103 in Fractography of Ceramic and Metal Failures, eds., J.J. Mecholsky, 

Jr, S. W. Freiman, and S. R. Powell, ASTM STP 827 ASTM, Int., West 

Conshohocken, PA, 1984. 

64.	 J. J. Mecholsky, Jr., T. J. Hill, and Z. Chen, “Application of Quantitative 

Fractography to the Characterization of R-Curve Behavior,” pp.152-168 in 

Fracture Resistance Testing of Monolithic and Composite Brittle Materials, eds., 

J. A. Salem, G. D. Quinn, M. G. Jenkins, ASTM STP 1409, ASTM, Int., West 

Conshohocken, PA (2002). 

65. G. D. Quinn and J. J. Swab, “Fractography and Estimates of Fracture Origin Size 

from Fracture Mechanics,” Cer. Eng. Sci. Proc., 17 [3] (1996) 51 - 58. 

66. G. D. Quinn and J. J. Swab, “Comparisons of Calculated and Measured Flaw 

Sizes,” pp. 175 – 192 in Fractography of Glasses and Ceramics IV, Ceramic 

Transactions, Vol. 122 eds. J. Varner and G. Quinn, American Ceramic Society, 

Westerville, OH, 2001. 

67. H. G. Richter and F. Kerkhof, “Stress Wave Fractography,” pp. 75 – 109 in 

Fractography of Glass, eds. R. C. Bradt and R. E. Tressler, Plenum, NY, 1994. 

68. H. Richter, “Crack Propagation in Glass Under Liquids in an Intermediate Range 

of Crack Velocities,” pp. 219 – 229 in Strength of Inorganic Glass, ed. C. R. 

Kurkjian, Plenum Press, NY, 1985. 

69. T. A. Michalske, M. Singh, and V. D. Fréchette, “Experimental Observation of 

Crack Velocity and Crack Front Shape Effects in Double-Torsion Fracture 

Mechanics Tests,” pp. 3 - 12 in Fracture Mechanics for Ceramics, Rocks, and 

Concrete, STP 745, ASTM Int., West Conshohocken, PA, 1982. 

70. T. A. Michalske and J. M. Collins, “Fractographic Determination of Crack-Tip 

Stress Intensity,” pp. 229 – 239 in Fractography of Glasses and Ceramics, 

Advances in Ceramics, Vol. 22, eds. J. R. Varner and V. D. Fréchette, American 

Ceramic Society, Westerville, OH 1988. 

7-79 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

71. A. Tsirk, “Formation and Utility of a Class of Anomalous Wallner Lines on 

Obsidian,” ibid, pp. 57 – 69. 

72. E. F. Poncelet, “The Markings on Fracture Surfaces,” J. Soc. Glass Technol., 42 

(1958) 279T – 288T. 

73. A. Smekal, “Procedure for Measurement of Fracture Surface Propagation 

Velocities on Fracture Surfaces,” Glastech. Ber., 23 [3] (1950) 57 – 67. 

74. T. Fett, D. Munz, and K. Keller, “Determination of Subcritical Crack Growth on 

Glass in Water from Lifetime Measurements on Knoop-Cracked Specimens,” J. 

Mat. Sci., 23 (1988) 798 - 803. 

75. P. J. Dwivedi and D. J. Green, “Determination of Subcritical Crack Growth 

Parameters from In situ Observation of Indentation Cracks,” J. Am. Ceram. Soc., 

78 [8] (1995) 2122 - 2128. 

76. V. M. Sglavo and R. Dal Maschio, “Controlled Indentation-Induced Cracks for 

the Determination of Fracture Toughness in Alumina,” pp. 233 – 244, in Fracture 

Mechanics of Ceramics, Vol. 11, eds. R. C. Bradt et al., Plenum, NY, 1996. 

77. M. Stech and J. Rödel, “Method for Measuring Short -Crack R-Curves Without 

Calibration Parameters: Case Studies on Alumina and Alumina-/Aluminum 

Composites,” J. Am. Ceram. Soc., 79 [2] (1996) 291 - 297. 

78. W. Chen, D. Lupascu, J. Rödel, and C. S. Lynch, “Short Crack R-Curves in 

Ferroelectric and Electrorestrictive PLZT,” ibid, 84 [3] (2001) 393 – 397. 

79. A. H. Heuer, “Transformation Toughening in ZrO2-Containing Ceramics,” ibid, 

70 [10] (1987) 689 – 698. 

80. G. D. Quinn, L. K, Ives, and S. Jahanmir, “On the Fractographic Analysis of 

Machining Cracks in Ground Ceramics: A Case Study on Silicon Nitride,” 

Special Publication SP 996, NIST, Gaithersburg, MD, May, 2003. 

81. D. B. Marshall, “Strength Characteristics of Transformation-Toughened 

Zirconia,” J. Am. Ceram. Soc., 69 [3] (1986) 173 – 180. 

82. S. M. Smith and R. O. Scattergood, “Controlled Crack Shapes for Indentation 

Fracture of Soda Lime Glass,” ibid, 75 [9] (1992) 2593 – 2596. 

83. R. W. Steinbrech and O. Schmenkel, “Crack-Resistance Curves of Surface Cracks 

in Alumina,” ibid, 71 [5] (1988) C271- C273. 

84. T. Fett, D. Munz, J. Siedel, M. Stech, and J. Rödel, “Correlation Between Long 

and Short Crack R-Curves in Alumina Using the Crack Opening Displacement 

and Fracture Mechanical Weight Function Approach,” ibid, 79 [5] (1996) 1189 -

1196. 

85. T. Fett and D. Munz, “Evaluation of R-curve Effects in Ceramics,” J. Mat. Sci., 

28 (1993) 742 – 752. 

86. R. F. Cook and D. R. Clarke, “Fracture Stability, R-Curves and Strength 

Variability,” Acta Metall., 36 [3] (1988) 555 – 562. 

87. J. J. Mecholsky, Jr., T. J. Hill and Z. Chen, “Application of Quantitative 

Fractography to the Characterization of R-Curve Behavior,” pp. 152 – 167 in 

Fracture Resistance Testing of Monolithic and Composite Brittle Materials, 

ASTM STP 1409, eds., J. Salem, et al., ASTM, West Conshohocken, PA, 2002. 

88. J. J. Petrovic, L. A. Jacobson, P. K. Talty, and A. Vasudevan, “Controlled Surface 

Flaws in Hot-Pressed Si3N4,” J. Am. Ceram. Soc., 58 [3-4] (1975) 113 - 116. 

7-80 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Quantitative Analysis 

89. G. D. Quinn, R. J. Gettings, and J. J. Kübler, “Fractography and the Surface Crack in 

Flexure (SCF) Method for Evaluating Fracture Toughness of Ceramics,” pp. 107 -

144 in Fractography of Glasses and Ceramics III, Ceramic Transactions, Vol. 64, 

eds., J. R. Varner, V. D. Fréchette, and G. D. Quinn, ACS, Westerville, OH, 1996. 

90. G. Quinn, R. Gettings, and J. Kübler, “Fracture Toughness by the Surface Crack in 

Flexure (SCF) Method: Results of the VAMAS Round Robin,” Cer. Eng. Sci. Proc., 

15 [5] (1994) 846-855. 

91. G. D. Quinn, K. Xu, J. A, Salem, and J. J. Swab, “SRM 2100: the World’s First 

Fracture Toughness Reference Material,” pp. 499 – 530 in Fracture Mechanics of 

Glasses and Ceramics, Vol. 14, eds., R. C. Bradt, D. Munz, M. Sakai, and K. W. 

White, Klewer/Plenum, NY, 2005. 

92. I. Bar-On, G. D. Quinn, J. Salem, and M. J. Jenkins, “Fracture Toughness Standard 

Test Method C 1421-99 for Advanced Ceramics,” pp. 315-335 in Fatigue and 

Fracture Mechanics, Vol. 32, ASTM STP 1406, ed. R. Chona, ASTM Int., West 

Conshohocken, PA, 2001. 

93. EN ISO 18756:2005, “Advanced Ceramics, Advanced Technical Ceramics); 

Determination of Fracture Toughness of Monolithic Ceramics at Room Temperature 

by the Surface Crack in Flexure (SCF) Method,” European Committee for 

Standards, Brussels. 

94.	 G. D. Quinn, “Fracture Toughness of Ceramics by the Vickers Indentation Crack 

Length Method: A Critical Review,” Cer. Eng. Sci. Proc., 27 [3] (2006) 45 – 62. 

95. G. D. Quinn and R. C. Bradt, “On the Vickers Indentation Fracture Toughness 

Test,” J. Am. Ceram. Soc., 90 [3] (2007) 673 – 680 

96. B. R. Lawn and D. B. Marshall, “Indentation Fractography,” pp. 1 –36 in 

Fractography of Glass, eds., R. Bradt and R. Tressler, Plenum, NY, 1994. 

97. P. Chantikul, G. R. Anstis, B. R. Lawn and D. B. Marshall, “A Critical Evaluation 

of Indentation Techniques for Measuring Fracture Toughness: II Strength 

Method,” J. Am. Ceram. Soc., 64 [9] (1981) 539 – 543. 

98. R. F. Cook and G. M. Pharr, “Direct Observation and Analysis of Indentation 

Cracking in Glasses and Ceramics,” ibid, 73 [4] (1990) 787 – 817. 

99.	 T. Wilantewicz, “Crack Initiation Behavior of Optical Glasses from Vickers 

Indentations,” Ph. D. Thesis, Alfred University, November, 2005. 

100.	 T. Wilantewicz and J. R. Varner, “Vickers Indentation Fracture in Optical Glass 

Compositions,” pp. 131 -151 in Fractography of Glasses and Ceramics, Vol. 5, 

eds. J. R. Varner, G. D. Quinn, M. Wightman, American Ceramic Society, 

Westerville, OH, 2007. 

101. D. B. Marshall, B. R. Lawn, and J. J. Mecholsky, “Effect of Residual Stresses on 

Mirror/Flaw Size Relations,” J. Am. Ceram. Soc, 63 [5-6] (1980) 358 – 360. 

102. D. E. Passoja, “Fundamental Relationships Between Energy and Geometry in 

Fracture,” pp. 101 – 126 in Fractography of Glasses and Ceramics, eds., J. R. 

Varner and V. D. Fréchette, Advances in Ceramics, Vol. 22, American Ceramic 

Society, Westerville, OH, 1988. 

103. J. J. Mecholsky, T. J. Mackin, and D. E. Passoja, “Self-Similar Crack Propagation 

in Brittle Materials,” ibid, pp. 127 – 134. 

7-81 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

104.	 D. A. Lange, H. M. Jennings, and S. P. Shah, “Relationship Between Fracture 

Surface Roughness and Fracture Behavior of Cement Paste and Mortar,” J. Am. 

Ceram. Soc., 76 [3] (1993) 589 – 597. 

105.	 Y. Xin, K. J. Hsia, and D. A. Lange, “Quantitative Characterization of the 

Fracture Surface of Si Single Crystals by Confocal Microscopy,” ibid, 78 [12] 

(1995) 3201 – 3208. 

106. J. J. Mecholsky, D. E. Passoja, and K. S. Feinberg-Ringel, “Quantitative Analysis 

of Brittle Fracture Surfaces Using Fractal Geometry,” ibid, 72 [1] (1989) 60 – 65. 

107.	 J. J. Mecholsky, Jr. and S. W. Freiman, “Relationship Between Fractal Geometry 

and Fractography,” ibid, 74 [12] (1991) 3136 – 3138. 

108.	 T. Hollstein, W. Pfeiffer, M. Rombach, and B. Thielicke, “Analysis of Machining 

Damage in Engineering Ceramics by Fracture Mechanics, Fractography, and Xray Diffraction,” pp. 145-169 in Fractography of Glasses and Ceramics, III, 

Ceramic Transactions, Vol. 64, eds., J. Varner, et al., ACS, Westerville, OH, 1996). 

109. J. J. Mecholsky and M. G. Drexhage, “Comparison of Optical and Fractographic 

Measurements of Residual Stress in Compressively Clad Glass Rods,” J. Am. 

Ceram. Soc., 63 [5-6] (1980) 347-349. 

110.	 J. C. Conway, Jr. and J. J. Mecholsky, Jr., “Use of Crack Branching Data for 

Measuring Near-Surface Residual Stress in Tempered Glass,” ibid, 72 [9] (1989) 

1584-1587. 

111.	 D. B. Marshall and B. R. Lawn, “An Indentation Technique for Measuring 

Stresses in Tempered Glass Surfaces,” J. Am. Ceram. Soc., 60 [1-2] (1977) 86-87. 

112.	 M. V. Swain, J. T. Hagan, and J. E. Field, “Determination of the Surface Residual 

Stresses in Tempered Glasses by Indentation Fracture Mechanics,” J. Mat Sci. 

Lttrs., 12 (1977) 1914-1917. 

113.	 W. Weibull, “A Statistical Distribution Function of Wide Applicability,” J. Appl. 

Mech., September, (1951) 293 – 297. 

114.	 G. D. Quinn, “Static Fatigue and Creep Resistance of a Commercial Sialon,” J. Mat. 

Sci., 22 (1987) 374-380. 

115.	 K. Trustrum and A. De S. Jayatilaka, “Application of Weibull Analysis for Brittle 

Materials,” J. Mat. Sci., 18 (1983) 2765-70. 

116.	 A. De S. Jayatilaka and K. Trustrum, “Statistical Approach to Brittle Fracture,” 

ibid, 12 (1977) 1426-30. 

117. R. Danzer, “A General Strength Distribution Function for Brittle Materials,” J. 

Eur. Ceram. Soc., 10 (1992) 461 – 472. 

118.	 G. D. Quinn and R. Morrell, “Design Data for Engineering Ceramics: A Review of 

the Flexure Test,” J. Am. Ceram. Soc., 74 [9] (1991) 2037 - 2066. 

119.	 J. Sung and P. S. Nicholson, “Strength Improvement of Yttria-Partially-Stabilized 

Zirconia by Flaw Elimination,” J. Am. Ceram. Soc., 71 [9] (1988) 788 - 795. 

120.	 J. B. Quinn and G. D. Quinn, “A Practical and Systematic Review of Weibull 

Statistics for Reporting Strengths of Dental Materials,” Dent. Mat., 26 [2] (2010) 

135-147. 

121.	 L. Gorjan and M. Ambrožič, “Bend Strength of Alumina Ceramics: A 

Comparison of Weibull Statistics with Other Statistics Based on a Very Large 

Experimental Data Set,” J. Eur. Ceram., Soc., 32 (2012) 1221 -1227. 

7-82 

 

 

 

 

 

 

 

 

Single Crystals 

8. Single Crystals 

8.1 General 

Fractographic procedures for single crystals are similar to those for glasses and 

polycrystalline ceramics, but there are some nuances and differences. Fracture 

surfaces may have regions with conchoidal fracture, Wallner lines, and twist 

hackle. Other regions may have dramatic faceting. Fracture resistance often 

varies significantly between the different crystallographic planes, causing 

jagged fracture patterns as cracks radically change directions onto preferred 

cleavage planes. Fracture surfaces may be very difficult to interpret. Some 

examples of the unusual markings are shown in Figure 8.1. 

(a) (b 

(c) (d) 

Figure 8.1 Examples of single crystal fractures. (a) and (b) show fracture 

mirrors on the {100} type planes in MgO, but with two different tensile surface 

orientations. (c) and (d) show a cubic zirconia bend bar with curved linear 

lance like markings emanating from the origin area and the grinding crack 

origin. (courtesy R. Rice) 

8-1 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

The overall fracture plane may not necessarily be perpendicular to the direction 

of principal normal stress, since a preferred cleavage plane may be at an 

irregular angle to that stress. Sometimes the overall fracture extension may be 

approximately perpendicular to the maximum principal stresses, but local 

fracture proceeds in a zigzag fashion on other planes. This crack redirection is 

different than conventional crack branching in glasses or polycrystalline 

ceramics. Branching is much less common in single crystals. If the maximum 

principal stress is perpendicular to a preferred cleavage plane, the fracture 

surface may be very flat and featureless with little or no hackle, and no 

branching at all. 

Fracture mirrors may be present, but may have odd shapes and markings. 

There may be no mist or hackle on preferred cleavage planes. Mirror sizes 

may be extremely difficult or impractical to measure. Cleavage hackle steps 

may be confused with Wallner lines. Twinning may cause fracture. 

8.2 Preferred Cleavage Planes 

Bradt et al.1,2 have reviewed the various criteria for cleavage and concluded 

that fracture toughness is the best indicator of the preferred cleavage plane. 

The fracture toughness can vary dramatically on different planes. The energy 

to propagate cracks along preferred cleavage planes is less than that for lessfavored planes. Nevertheless, there is contradictory information in the 

literature on preferred cleavage planes in some crystals, even for such well 

studied materials as silicon, alumina (sapphire), and magnesium aluminate 

spinel. Some of contradictory information surely is due to experimental error 

or specific aspects of the test method. The direction of crack advance on a 

plane is as important as the plane itself. The speed of crack advance can even 

determine which fracture plane is preferred.3

 Although the double cantilever 

beam method and the constant applied moment variation as shown in Figure 

8.2 are often used, results are not strictly valid if the crack zigzags down the 

guiding groove or is tilted to the intended fracture plane. 

Knoop indentations in bend bars may be used to generate controlled surface 

flaws in the surface crack in flexure (SCF) method as shown in Figure 8.2. The 

specimen axis and cross section and the indentation and precrack plane may be 

aligned with a particular cleavage plane. The local direction of crack extension 

may vary along the crack periphery, however. Such micro flaws, which mimic 

naturally occurring flaws, are valuable in demonstrating just how complex 

fracture resistance can be in single crystals. A good example is from the work 

of Chen and Leopold4 and Xin et al.5

 on silicon as summarized in Table 8.1. 

Xin et al. 5 showed that the residual stresses from the indentation could change 

8-2 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Single Crystals 

the propagation plane. Bradt et al.6,7,8 have used the SCF method on alumina 

and spinel. Figure 8.3 shows a Knoop precrack in sapphire. One unresolved 

issue with the SCF method for single crystals is whether or not indentation 

residual stresses vary with plane and orientation. Akimune and Bradt9

 showed 

that there is considerable hardness anisotropy even on one plane (100) in 

spinel, but negligible variability on the (111). 

Surface Crack in Flexure 

Double Torsion 

Double Cantilever 

Beam 

Applied Moment Double 

Cantilever Beam 

Figure 8.2 Test methods for single crystal fracture toughness determination. 

Reference Plane 

Type 

KIc 

(MPa√m) Fracture characteristics 

Chen and 

Leipold (4) {111} 0.82 Flat and smooth fractures. This is the preferred cleavage plane. 

Chen and 

Leipold (4) {110} 0.90 

Large but asymmetric cathedral mirrors 

centered on the flaws. 

Propagation switched to {111} planes 

and zigzagged at the mirror boundaries. 

Chen and 

Leipold (4) {100} 0.95 A very rough fracture surface. Cracks switched to {111} very quickly. 

Xin, Hsia, and 

Lange (5) {110} 0.95 Residual indentation stresses intact: 

Cracks propagated on {110}. 

Xin, Hsia, and 

Lange(5) {110} - Residual indentation stresses removed 

by annealing. Crack switched to {111}. 

Table 8.1 Fracture toughness of single crystal silicon from Knoop SCF tests. 

8-3 
 

 

 

 

 Fractography of Ceramics and Glasses 

(a) 

(b) 

Figure 8.3 Relatively smooth fracture surface (a) of a surface crack in flexure 

(SCF) sapphire test specimen with a 49 N Knoop indentation as a starter flaw 

(large arrow). The bar tensile surface was the a-plane. Fracture occurred on 

the m-plane. (b) is a close-up which shows grinding cracks along the surface 

(small white arrows), but the indentation crack (large white arrow) was 

dominant. The long curved line leading from the origin to the upper right is a 

“cleavage step hackle” line. (courtesy R. Krause). 

8-4 
 

 

 

 

 

 

 

 

 

 

 

 

 

Single Crystals 

8.3 Fractographic Techniques 

Reflections and translucency in some crystals make photographing single 

crystals difficult. Different lighting and specimen tilting are needed to bring 

out various details. 

Low-angle grazing illumination on cleaved surfaces may be more problematic 

since fracture surfaces are mirror like and reflect light away from the 

microscope objective lens. Silicon is a good example of this and the 

fractographer is confronted with either getting too little or too much light 

reflected into the eyepieces. When using a stereomicroscope with directional 

illumination, it is often effective to tilt the specimen at an angle to reflect more 

light up into the eyepieces. The good depth of field of the stereomicroscope 

can accommodate the specimen tilt at low to moderate magnifications. 

Stopping down an aperture can increase the depth of field. 

A reflected light compound microscope may be very effective with single 

crystals if the fracture surface is not too rough. Dark field illumination may aid 

finding fracture mirrors, since the symmetry of cleavage steps may be 

accentuated as shown in Figure 8.4. 

Confocal microscopes may not be very effective due to the specular reflections. 

Transmitted illumination is often useful for transparent single crystals. 

Polarizers are essential aids to help discern twinning, which may cause fracture 

in some crystals. 

Gold coating of a transparent single crystal may aid in interpretation and may 

make the fracture surface easier to view, but if the surface is very flat, the gold 

coating may make it behave like a mirror, which will pose problems for 

inspection and photography. 

X-ray topography is a unique tool for single crystals. Imperfections in the 

crystal and its surface produce image contrast. Figure 8.5 shows an example 

that reveals handling damage on the polished surface of a sapphire hemispherical dome.10 The full extent of the handling damage was not visible with 

the optical or scanning electron microscopy. 

8-5 
 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) 

(b) 

(c) 

Figure 8.4 Dark-field illumination may aid optical examination. Different 

views of the fracture origin in a silicon wafer tested in a biaxial ring-on-ring 

mode (282 MPa). Fracture initiated from the surface grinding crack (also 

shown later in Figure 8.17), propagated on a {110} type plane, but started to 

zigzag. (a) is a dark field optical image with a strange swirling appearance. It 

nonetheless has a symmetry that is a tip off to the existence of an origin and 

mirror. (b) and (c) are SEM images of the origin with arrows calling attention 

to the origin flaw. 

8-6 
 

 

 

 

 

Single Crystals 

1 

Polished surface 

Fracture surface 

(a) (b) 

(c) 

1 

Figure 8.5 Sapphire hemispherical dome that fractured during a thermal 

stress proof test. (a) shows the fractured dome and the origin (large white 

arrow). (b) shows an optical fractograph with the origin piece tilted back to 

show both the fracture surface and the polished outer surface. The cracking is 

marked by the black arrows. Another apparently isolated intersecting scratch 

is marked with the white arrow and the numeral “1.” (c) is a composite of two 

x-ray topographs of the polished outer surface, showing the matching fracture 

halves where they meet at the origin. Surface and subsurface damage are 

exposed and the scratch labeled “1” in (b) is revealed to be part of a complex 

handling damage (HD) network. (c is courtesy of D. Black) 

8-7 
 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) (b) 

Crack propagation Crack propagation 

Twist hackle from 

final breakthrough 

Wallner line Cleavage step hackle 

Glass Single crystal 

Figure 8.6 Schematics showing (a) a Wallner line and twist hackle in a glass 

plate loaded in bending and (b) “cleavage step hackle” in a single crystal plate 

loaded in bending. In both cases the crack was leading on the bottom surface. 

8.4 Fracture Surface Markings 

Single crystal fracture planes often have small “cleavage step hackle” that may 

be confused with Wallner lines. Figure 8.6 shows the differences. 

Cleavage step hackle: A form of hackle in single crystals caused by the crack 

running on two very close parallel planes. It is in the form of one or more 

single lines, parallel to the local direction of crack propagation. 

Cleavage step hackle is often generated by the origin or other surface 

irregularities and runs parallel to the direction of crack propagation. The 

hackle may suddenly change from curved lines to parallel vertical segments on 

the compression side of a specimen. The step may be very persistent and 

extend over great distances without merging, unlike twist hackle lines. 

Cleavage step hackle is not unlike “step hackle” in glasses as described in 

section 5.3.6, but here with single crystals, the step is caused by the preference 

of a crack to stay on two nearby parallel planes without merging. Figures 8.1c 

and 8.3 show good examples of cleavage step hackle in sapphire. 

Fracture surfaces and fracture mirrors may have extraordinary shapes and 

facets, but usually have some degree of recognizable symmetry that will help 

the fractographer identify an origin site. Mirror shapes may vary dramatically 

within a crystal depending upon the fracture plane. To illustrate these 

fascinating variations, fracture markings for single crystal cubic spinel are 

shown in Figures 8.7 through 8.14. These figures follow the recommendations 

of this Guide and ASTM C 1322.11 They show a picture of the whole fracture 

8-8 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Single Crystals 

surface, the fracture mirror, and a close up of the origin. The fractographic 

findings are summarized in Table 8.2. These images are from Rice and 

colleagues3,12,13,14 who studied fracture in this material using double cantilever 

beam and flexural strength specimens. They found a remarkable variety of 

fracture surfaces with features such as “cathedral mirrors,” “gull wing mirrors,” 

“skewed cathedral mirrors,” and “whisker lances.” Cathedral and gull wing 

mirrors have also been observed in other cubic crystals such as silicon (Figure 

8.4) and fully-stabilized zirconia. Gull wing mirrors look like cathedral mirrors 

that are tilted 90o on the fracture surface. Bear in mind that there may be 

threshold stresses or stress intensities for the formation of markings on certain 

planes. Thus, the flat fractures shown in figures 8.8 and 8.9 may indicate 

fracture on a preferred cleavage plane, or that the stresses were too low to 

generate mirror markings. The preferred cleavage plane in spinel is the {100} 

which has a fracture toughness of 0.9 to 1.2 MPa√m depending upon the 

method of measurement.3,6,7,12,13,14 The fracture toughness of the {110} is not 

much greater at 1.0 to 1.5 MPa√m. Consequently, fracture usually runs on 

{100} or {110} type planes. Fracture is not preferred on the {111} and the 

fracture toughness is much greater (1.9 MPa√m or more). 3,6,7,12,13,14 

Intended 

fracture 

plane type 

Bend bar 

tensile 

surface 

Bend bar 

tensile axis Fracture type Figure 

{100} 

(Case A) 

{100} <100> Cathedral mirror 

Major crack redirection beyond the 

mirror sides 

8.7 

{100} 

(Case D) 

{110} <100> No mirror 

Flat and featureless fracture surface 

8.8 

{110} 

(Case B) 

{100} <110> No mirror 

Flat and featureless fracture surface 

8.9 

{110} 

(Cases 

F1 + F2) 

(Case C) 

{110} <110> Skewed fracture mirror and 

Major crack redirection beyond the 

mirror sides, or 

Gull wing mirror and 

Jagged fracture surface beyond top of 

mirror 

8.10 

8.11 

8.12 

{110} 

(Cases 

G1 + G2) 

{111} <110> Skewed cathedral mirror 

Major crack redirection beyond the 

mirror sides, or flat 

8.13 

8.14 

{111} 

(Case E) 

{110} <111> Gull wing mirror 

Jagged crack propagation beyond top 

of mirror 

8.15 

Table 8.2 Fracture markings in MgAl2O4 spinel bend bars. The brackets {100} 

denote the family of (100) planes and the angles <100> denote the family of 

[100] directions. (After Wu, McKinney, and Rice, Ref. 3) 

8-9 
 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) Bend bar orientation 

(b) 

(c) 

Figure 8.7 Spinel bend bar with a {100} tensile surface and <100> tensile 

axis designed to cause fracture on a {100} plane. (185 MPa) (a) shows the 

whole fracture surface and origin (arrow). (b) shows the “cathedral mirror” 

and the arc-ribs that gradually get progressively severer until a branch or 

major crack redirection occurs (white arrow). (c) shows the grinding crack 

fracture origin. Fracture starts on the {100} but shifts or branches to the {110} 

just beyond the sides of the mirror. (Case A, Table 8.2) (Courtesy R. Rice) 

8-10 
 

 

 

 

Single Crystals 

(a) 

(b) 

(c) 

Figure 8.8 Spinel bend bar with a {110} tensile surface and <100> tensile 

axis designed to cause fracture on the {100} cleavage plane. (170 MPa) (a) 

shows the whole fracture surface which is flat except for the cleavage steps on 

the side and the cantilever curl. Other similar specimens had no cleavage 

steps. (b) shows the origin site and (c) shows the grinding crack fracture 

origin. (Case D, Table 8.2) (Courtesy R. Rice) 

8-11 
 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) 

(b) 

Figure 8.9 Spinel bend bar with a {100} tensile surface and <110> tensile 

axis designed to cause fracture on a {110} type plane. (169 MPa) (a) shows 

the whole fracture surface which is flat and featureless except for the cantilever 

curl. The arrow shows the origin. (b) shows the grinding crack fracture 

origin. Fracture stays on a {110} type plane. (Case B, table 8.2) (Courtesy R. 

Rice) 

8-12
 
 

 

 

 

 

 

Single Crystals 

(a) 

(b) 

(c) 

Figure 8.10 Spinel bend bar with a {110} tensile surface and <110> tensile 

axis designed to cause fracture on a {110} plane. (283 MPa) (a) shows the 

whole fracture surface, (b) shows the mirror with “whisker lances” radiating 

outward, and (c) shows the grinding crack fracture origin. (Case F1, table 8.2) 

(Courtesy R. Rice) 

8-13 
 

 

 Fractography of Ceramics and Glasses 

(a) 

(b) 

(c) 

Figure 8.11 A second example of a spinel bend bar with a {110} tensile 

surface and <110> tensile axis. (172 MPa) The fracture surface (a) is very 

rough as in the previous figure. The mirror (b) is somewhat more symmetric 

than in the previous figure. This mirror was shown in an earlier publication 

(Figure 29 of Ref. 13), but with a typographical mistake identifying it as 

{110}<100>. (Case F2, table 8.2). (Courtesy R. Rice) 

8-14 
 

 

 

 

 

 

 

 

Single Crystals 

(a) 

(b) 

(c) 

Figure 8.12 Spinel bend bar with a {110} tensile surface and <110> tensile 

axis designed to cause fracture on a {110} plane. (361 MPa) (a) shows the 

whole fracture surface, (b) shows the “gull wing mirror,” and (c) shows the 

grinding crack fracture origin. (Case C, table 8.2) (Courtesy R. Rice) 

8-15 
 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) 

(b) 

(c) 

Figure 8.13 Spinel bend bar with a {111} tensile surface and <110> tensile 

axis designed to cause fracture on a {110} plane. (unknown stress) (a) shows 

the whole fracture surface and the major crack redirection on the mirror sides. 

(b) shows the “skewed cathedral mirror,” and (c) shows the grinding crack 

fracture origin from the matching fracture piece. (Case G1, table 8.2) 

(Courtesy R. Rice) 

8-16 
 

 

 

 

Single Crystals 

(a) 

(b) 

(c) 

Figure 8.14 Another example of a spinel bend bar with a {111} tensile surface 

and <110> tensile axis designed to cause fracture on a {110} plane. (261 

MPa) (a) shows the whole fracture surface, (b) shows a “skewed mirror” 

with “whisker lances” and (c) shows the grinding crack fracture origin (black 

arrows) which is at a slight angle to the tensile stress axis and the fracture 

surface). (Case G2, table 8.2) (Courtesy R. Rice) 

8-17 
 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) 

(b) 

(c) 

Figure 8.15 Spinel bend bar with a {110} tensile surface and <111> tensile 

axis designed to cause fracture on a {111} plane. (325 MPa) (a) shows the 

whole fracture surface and the major crack redirection on the top half, (b) 

shows the “gull wing mirror,” and (c) shows the grinding crack fracture 

origin. (Case E, table 8.2) (Courtesy R. Rice) 

8-18 
 

 

Single Crystals 

Another distinctive shape seen in some crystals are long hackle lines that 

extend vertically from the sides of an origin suggesting a “batman mirror” 

shape as shown in Figure 8.16. 

(a) 

(b) 

Figure 8.16 “Batman mirror” in a sapphire bar. (a) shows matching halves 

of the bend bar with a tensile surface parallel to the a-plane and bend bar sides 

oriented so as to cause fracture on the r-plane. The tensile direction was the r 

axis. The origin was grinding damage. (Specimen courtesy R. Krause.) 

8-19 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Mirror shapes seem to correlate with elastic moduli15 or fracture surface 

energy16 variations in different directions in a fracture plane. The very limited 

data on fracture mirror constants that have been published are listed in 

Appendix C. The asymmetry and diversity of fracture mirror shapes and 

difficulties in judging mirror sizes hamper any attempts to accurately correlate 

fracture stresses to mirror sizes by the methods described in chapters 7 and 

appendix D. Recent work by Gopalakrishnan and Mecholsky 17 provides a new 

explanation why mist does not appear around fracture mirrors in single crystals. 

The existence of any mode II loading on a crack tends to suppress mist in 

general and accentuates twist hackle lance lines at a mirror boundary. They 

argue that mixed mode I and II loading is naturally introduced in single crystals 

due to crystal anisotropy. 

Scarps can occur in single crystals and are indicative of the presence of liquid 

during fracture. Figure 5.48b shows an example of scarps in sapphire. 

Fracture surface 

Tensile 

surface 

10 μm 

(a) (b) 

Figure 8.17 Fracture origin of the silicon wafer shown in Figure 8.4. (282 

MPa) (a) shows the origin is a grinding crack that extends only 2 μm below 

the tensile surface and is associated with a grinding striation (groove) that 

seems to be deeper and has caused greater cracking than other nearby 

striations. The grinding direction is slightly different than the final plane of 

fracture, which caused the fracture surface to have a jog at the origin as shown 

in the optical view of the tensile surface (b). The fractographer should be alert 

to subtle features like the jog, since they are important clues. 

8-20 
 

 

 

 

 

 

 

 

 

Single Crystals 

8.5 Origins 

Single crystals are very sensitive to surface contact damage of any type. The 

low fracture toughness and high stiffness means that tiny strength-limiting 

cracks can easily pop into the crystal. There are no grain boundaries or other 

microstructural features to impede or redirect surface initiated cracks. 

Grinding or polishing flaws are common fracture origins. Flexure and tensile 

specimens are very sensitive to edge flaws and special care must be taken with 

the edge beveling or rounding procedures. Figure 8.17 shows the grinding 

crack origin of the silicon wafer shown earlier in Figure 8.4. 

(c) 

Fracture surface 

Polished surface 

(a) (b) 

Figure 8.18 Polishing or grinding crack origin in a sapphire dome. (a) shows 

the entire fractured structure. The origin is marked by the black arrow. (b) 

shows an SEM photo of the piece with the origin marked by the arrow. 

Fracture followed the axis of a polishing crack for a short distance, but then 

quickly shifted to a different plane. (c) is a SEM close-up of the origin. The 

piece has been tilted back to show both the fracture surface and the polished 

surface. There is no trace of damage on the polished surface. 

8-21 
 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Single crystals are also very susceptible to polishing flaws. Subsequent 

polishing or finishing may eliminate any surface trace or scratch associated 

with the flaw, rendering optical post-finishing inspection ineffective. Figure 

8.18 shows an example in a sapphire dome. Figure 8.5 is another example of a 

sapphire dome with subtle surface damage. 

Surface impact or contact loading creates strength-limiting flaws. The contact 

can create cracks that may be extended by subsequent loading. Alternately, 

there may be stresses already present in the part that can propagate the crack. 

Figure 8.19 shows an example where a small high velocity particle “dinged’ the 

surface of a sapphire dome. 

Figure 8.19 Impact crack fracture origin in a hemispherical sapphire dome 

that was impacted by a small particle while the dome was under thermal stress. 

The impact site on the surface is missing a small divot (arrow) and the 

approximately 150 μm deep crack shows concentric semicircular tertiary 

Wallner lines that are telltale signs of impact. See Figure 4-15 on page 4-25 

for views of the entire dome. 

8-22 
 

 

 

 

 

 

 

 

Single Crystals 

Other surface flaws may originate as a result of processing conditions. Figures 

8.20 and 8.21 show fractures in microelectromechanical (MEMS) sized single 

crystal silicon theta specimens.18,19 Deep reactive ion etching (DRIE), a 

lithographic process used to prepare integrated circuits and small structures, 

was used to fabricate the specimens from a standard {100} type silicon wafer. 

The DRIE process created the surface damage that was strength limiting in this 

case. Figure 6.69 on page 6-76 shows other theta specimen origins. 

• {100} 

← <110> →

(a) (b) 

Original (001) wafer surface 

(c) (d) 

Figure 8.20 Tiny silicon theta test specimens. The ring shown in (a) is 

compressed on its rim on the top and bottom in a nanoindenter causing the 7.5 

μm tall by 100 μm thick web section in the middle to stretch in uniaxial tension. 

(b) shows a fractured specimen on a strip. (c) shows that the web fractures 

usually occurred on {111} planes which are the cleavage planes. A few 

initiated on {110} planes but then changed to {111} planes. Origin flaws are 

shown in the next figure. 

{110} type etched surface 

parallel to the strip bottom surface 

{111} Fracture planes {111} Fracture plane 

{110} plane 

8-23 
 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) (b) 

(c) (d) 

(e) (f) 

Figure 8.21 Fracture origins in single crystal silicon theta specimens. (b) is a 

close-up of (a). The “cathedral mirrors” are very clear. The origins are etch 

pits from the DRIE processing. The cusps between them are very sharp and 

probably initiated cracks. Figure 6.69 shows other sharp cusp origins. 

8-24 
 

 

 

 

 

 

 

 

Single Crystals 

Figure 8.22 Internal origin in a flame-polished non-stoichiometric (alumina 

rich) spinel single crystal. The bend bar strength was 238 MPa. (courtesy R. 

Rice) 

Volume-distributed flaws such as pores, inclusions or compositional 

inhomogeneities can also cause fracture. They usually will stand out quite 

clearly on the fracture surface. Newcomb and Tressler 20 and Rice21 showed 

pore examples in sapphire fibers. An example of an internal origin in a 

nonstoichiometric spinel is shown in Figure 8.22. 

Twinning can initiate fractures, especially if twins from different planes 

intersect and nucleate cracks. Figure 8.23 shows twins that controlled the 

strength in a sapphire bend bar tested at high temperature. This failure 

mechanism is especially a problem in sapphire since twins are easily nucleated 

by even small compression stresses at fairly low temperatures (300 o

C to 500 o

C). They cause drastic reductions of compression and flexure strengths in 

some orientations.22,23 They sometimes form with an audible pop.24 It has 

even been reported that twins can form in sapphire from sudden temperature 

changes and hydrostatic compression.25 

8-25 
 

 

 

 Fractography of Ceramics and Glasses 

(a) (b) 

(c) 

Figure 8.23 Twinning in sapphire bend bars in an elevated temperature bend 

bar specimen. The bar was cut with an a-plane tensile surface and an intended 

m-plane fracture surface. The specimen instead broke on the r-plane from 

cracks initiated by twins probably from the compression side of the specimen. 

(a) shows the tensile face under ordinary lighting. (b) shows the same 

specimen, but with common reflected lighting adjusted to reflect off the twins. 

(c) shows the bar between crossed polarizers that make the twins much easier 

to detect. (courtesy R. Krause) 

8-26 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Single Crystals 

Figures 8.24 and 8.25 show two fractured c-axis sapphire tubes with intersecting twins that created fracture origin cracks. The tubes were 38 mm in 

diameter and 400 mm long and fractured during two runs of a prototype plasma 

asher. This is a device used to dry-remove photoresist from large 300 mm 

diameter silicon wafers during manufacture of integrated circuits. The tubes 

contained oxygen and nitrogen gasses which were energized into a plasma state 

by a 2.5 kW microwave chamber which was cycled on and off in 60 s intervals. 

The sapphire tubes reached temperatures of 600 o

C to 700 oC. One tube broke 

during a gas purge step three seconds after the shutdown of the microwave 

energy. The second tube broke two minutes after a shutdown. The tubes 

should have had lifetimes in excess of 400,000 cycles The breakage during 

cool down suggested thermal stresses had caused fracture. The cracks curved 

and wiggled at the tube ends, indicating the cracks slowed down and arrested. 

Some of the cracks crossed over each other, indicating that primary cracks did 

not run all the way through the wall thickness. During optical microscopic 

examination of one of the fracture surfaces, using common reflected light, a 

chance reflection off one surface revealed thin parallel reflective lines that ran 

counter to the fracture surface faceting as shown in Figure 8.24b. This was a 

sign of twinning and the next step was to view the tubes between crossed 

polarizers (as discussed in sections 3.8 and 3.25) with the results shown in 

Figure 8.25. Multiple twins in the hot region triggered crack formation at twin 

intersections. The twins were oriented at 32o to the sapphire tube axis 

corresponding to classic r-plane twinning in sapphire, which can be caused by 

rather low compression stresses at temperatures as low as 300 o to 500 o

C.22 

Fracture of these tubes occurred in two stages. First, axial constraint from a 

new type of seal at the tube ends prevented the free expansion of the tubes 

during the heating portion of the operating cycle. The earlier model tubes that 

did not have breakages used a different seal design that allowed expansion. 

The constraint in the new design created compression stresses. These created 

the intersecting twins that generated cracks. The cracks grew in stages and 

propagated by cool down thermal stresses until the tube ruptured. Although 

initially there was some concern that the sapphire tubes were faulty, they were 

fine. The problem was solved when the seal design was changed to allow for 

expansion of the sapphire tubes. 

8-27 
 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) 

(b) 

Cleavage 

steps 

twins 

Tube 1 Tube 

Figure 8.24 Two broken 38 mm diameter sapphire tubes from an integrated 

circuit plasma asher are shown. (a) shows both tubes. Although the locations 

of twins are marked with a felt tip pen, the twins are not visible in this lighting. 

(b) shows a stereoptical microscope photo close-up of the fracture surface of a 

fragment from the middle of one tube in normal white light illumination at just 

the right angle. Five thin twin bands are visible against the very facetted 

fracture surface. 

8-28 
 

 

Single Crystals 

(a) 

(b) 

Tube 1 

Tube 2 

Figure 8.25 Twin bands became very apparent when the sapphire tubes were 

illuminated by light from crossed polarizers. Cracks initiated at twin band 

intersections (white arrows) and then were propagated by cool down thermal 

stresses. (Photos courtesy R. Pingree) 

8-29 
 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Case study 13 in Chapter 10 is a complex sapphire solar refractor part. It also 

shows twins, but they were incidental and not the cause of fracture. Very tiny 

“micro twins” were also detected in that part. These presumably were the 

nuclei of twin bands that would have grown much larger with time. The 

fractographer should be on the lookout for twins in materials such as sapphire, 

and they may be present, but not necessarily the cause of fracture. 

Earlier in this chapter, it was shown in the spinel examples Figs. 8.8 and 8.9 

that, if a preferred cleavage pane is perpendicular to the maximum principal 

tensile stress, the fracture surface will be flat and smooth and at right angles to 

the tensile stress. If the preferred cleavage plane is not perpendicular to the 

maximum principal tensile stress, then the fracture plane may be at an inclined 

angle to the tensile stress, such as in Figure 8.23 for sapphire, and the zig-zag 

planes of Fig. 8.10, 8.11 and 8.13 for spinel. These illustrate exceptions to the 

law of normal crack propagation discussed in section 4.3 

In some extreme cases, the loading geometries and stresses may be such that 

fracture is forced to occur on an unfavorable plane, such as the basal plane in a 

sapphire tube shown in Figure 8.26. The fracture surface is extremely jagged 

with many zig-zags. Other examples of basal plane fracture in sapphire are 

shown in references 8, 20, and 26. 

Good estimates of the fracture toughness KIc may be obtained from fracture 

mechanics analyses of flaws. For example, the long uniform 2 μm deep surface 

crack in the silicon wafer shown in Figure 8.17a combined with a breaking 

stress of 282 MPa and a shape factor Y = 1.99 produced an estimate of KIc = 

1.99 (282) √2 x 10-6 = 0.79 MPa√m in good agreement with values in table 8.1. 

Rice showed several good examples13,21 and Newcomb and Tressler20 were able 

to measure fracture toughness on the basal plane of sapphire by this approach, 

when other methods usually fail due to the difficulty of controlling fracture on 

this plane. Many groups have effectively used Knoop indentation surface crack 

in flexure tests to study fracture behavior and measure fracture toughness of 

single crystals since the small semi-elliptical flaws mimic the behavior of other 

flaws. 

8-30 
 

 

 

 

Single Crystals 

(a)

 (b) 

(c) 

Figure 8.26 Example of basal plane fracture in a c-axis sapphire tube that 

fractured from severe thermal shock (a). The origin is at the tube mid length 

(large arrow) and is perpendicular to the tube axis. The thermal stresses acted 

in the axial direction. (b) and (c) show the mirror and origin on the outer tube 

surface. 

8-31 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

8.6 Other Sources 

Rice has written several outstanding, richly illustrated papers on single crystal 

fracture and fractography. Reference 13 is a book length article with over 

thirty figures showing single crystal fracture surfaces in MgO, MgAl2O4, Si, 

SiC, Al2O3, ZrO2, CaF2, and W. Reference 14 has six pages of illustrations and 

a discussion of fracture markings in various single crystal ceramics. 

Chapter 8 References 

1. R. A. Schultz, M. C. Jensen, and R. C. Bradt, “Single Crystal Cleavage of Brittle 

Materials,” Int. J. Fract., 65 (1994) 291- 312. 

2. R. C. Bradt, “Cleavage of Ceramic and Mineral Single Crystals,” pp. 355 – 365 in 

Cleavage Fracture, Proceedings of a Symposium at the 1997 TMS Fall meeting, 

Indianapolis, IN, Sept. 15- 17, 1997, ed. K. S. Chan, The Minerals, Metals, and 

Materials Society, Warrendale, PA, 1997. 

3. C. C. M. Wu, K. R. McKinney, and R. W. Rice, “Zig-zag Crack Propagation in 

MgAl2O4 Crystals,” J. Mat. Sci. Lttrs., 14 (1995) 474 – 477. 

4. C. P. Chen and M. H. Leipold, “Fracture Toughness of Silicon,” Bull. Am. Cer. 

Soc., 59 [4] (1980) 469-472. 

5. Y-B. Xin, K. J. Hsia, and D. A. Lange, “Quantitative Characterization of the 

Fracture Surfaces of Silicon Single Crystals by Confocal Microscopy,” J. Am. Cer. 

Soc., 78 [12] (1995) 3210 – 3208. 

6. R. L. Stewart, M. Iwasa, and R. C. Bradt, “Room Temperature KIc Values for 

Single-Crystal and Polycrystalline MgAl2O4,” Comm. Am. Cer. Soc., Feb. 1981, 

C22-23. 

7. R. L. Stewart and R. C. Bradt, “Fracture of Single Crystal MgAl2O4,” J. Mat. Sci., 

15 (1980) 67-72. 

8. M. Iwasa and R. C. Bradt, “Fracture Toughness of Single-Crystal Alumina,” pp. 

767 – 779 in Advances in Ceramics, Vol. 10, ed. W. Kingery, American Ceramic 

Society, Westerville, OH, 1984. 

9. Y. Akimune and R. C. Bradt, “Knoop Microhardness Anisotropy of Single-Crystal 

Stoichiometric MgAl2O4 Spinel,” J. Amer. Cer. Soc., 70 [4] (1987) C84 – C86. 

10. D. Black and G. D. Quinn, “X-ray Topography for Fractography of Single-Crystal 

Components,” J. Fail. Anal. Prev., 6 [3] (2006) 79 - 86. 

11. ASTM C 1322-96, “Standard Practice for Fractography and Characterization of 

Fracture Origins in Advanced Ceramics,” Annual Brook of Standards, Vol. 15.01, 

ASTM int., West Conshohocken, PA (1996). 

12. R. W. Rice, “Fractographic Determination of KIc and Effects of Microstructural 

Stresses in Ceramics,” pp. 509 – 545 in Fractography of Glasses and Ceramics II, 

eds. V. Fréchette and J. Varner, 1991. 

13. R. W. Rice, “Ceramic Fracture Features, Observations, Mechanisms and Uses,” 

pp. 5 –102 in Fractography of Ceramic and Metal Failures, ASTM STP 827, eds. 

J. Mecholsky and Powell, 1984, American Society for Testing and Materials, West 

Conshohocken, PA, 1984. 

8-32 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Single Crystals 

14. R. W. Rice, “Perspectives in Fractography,” pp. 3-56 in Fractography of Glasses 

and Ceramics, Advances in Ceramics, Vol. 22, eds. J. Varner and V. Fréchette, 

American Ceramic Society, Westerville, OH, (1988). 

15. S. W. Freiman, J. J. Mecholsky, and P. Becher, “Fractography: A Quantitative 

Measure of the Fracture Process,” pp. 55 - 78 in Fractography of Glasses and 

Ceramics, II , eds. V. Fréchette and J. Varner, American Ceramic Society, 

Westerville, OH, 1991. 

16. Y. L. Tsai and J. J. Mecholsky, Jr., “Fracture Mechanics Description of Fracture 

Mirror Formation in Single Crystals,” Int. J. Fract., 57 (1992) 167-182. 

17.	 K. Gopalakrishnan and J. J. Mecholsky, Jr., “Quantitative Fractography of MixedMode Fracture in Soda Lime Silica Glass,” J. Am. Cer. Soc., 95 [11] (2012) 3622-

3627. 

18. G. D. Quinn, E. Fuller, D. Xiang, A. Jillavenkatessa, L. Ma, D. Smith, and J. Beall, 

“A Novel Test Method for Measuring Mechanical Properties at the Small -Scale: 

The Theta Specimen,” Cer. Eng. Sci. Proc., 26 [2] (2005) 117 – 126. 

19.	 G. D. Quinn, “Fractographic Analysis of Miniature Theta Specimens,” Cer. Eng. 

Sci. Proc., 29 [3] (2008) 189 – 199. 

20. S. A. Newcomb and R. E. Tressler, “High-Temperature Fracture Toughness of 

Sapphire,” J. Am. Cer. Soc., 77 [11] (1994) 3030 – 3032. 

21. R. W. Rice, “Corroboration and Extension of Analysis of c-Axis Sapphire 

Filament Fractures from Pores,” J. Mat. Sci. Lttrs., 16 (1997) 202 – 205. 

22. F. Schmid and D. C. Harris, “Effects of Crystal Orientation and Temperature on 

the Strength of Sapphire,” J. Am. Cer. Soc., 81 [4] (1998) 885 – 93. 

23. P. F. Becher, “Fracture Strength Anisotropy of Sapphire,” J. Am. Cer. Soc., 59 [1-

2] (1976) 59 –61. 

24. J. A. Salem, “Slow Crack Growth and Fracture Toughness of Sapphire for the 

International Space Station Fluids and Combustion Facility,” NASA TM 214023, 

NASA Glenn Center, Cleveland, OH, 2006. 

25.	 L. M. Belyaev, Ruby and Sapphire, Translated from Russian, Amerind Publ. Co., 

New Delhi, 1980, pp. 286. 

26. R. W. Rice and P. F. Becher, “Comment on Creep Deformation of 0o Sapphire,” J. 

Am. Cer. Soc., 60 [3-4] (1977) 186 – 188. 

8-33 

 

 

 

 

 

 

 

 

 

 

Ceramic and Glass Composites 

9 Ceramic and Glass Composites 

9.1 Particulate, Whisker, or Self-Reinforced Ceramic Composites 

Ceramics or glasses that are reinforced by second phases, whiskers, or 

particulates may be examined using the same equipment and techniques as used 

for monolithic materials. The reinforcing agents act to deflect or bridge cracks 

and create very rough fractures surfaces that may mask classical fracture 

surface features such as Wallner or hackle lines. 

The composite ideally should have a controlled microstructure such that the 

reinforcing phase is uniformly distributed. This is routinely achieved with 

glass ceramics, but is more problematic with fiber or particulate-reinforced 

ceramics. Sometimes the reinforcing agent forms clumps or aggregates that 

can act as strength-limiting flaws as shown in Figure 9.1. Flaws of this type 

may be broadly described as compositional inhomogeneities using the logic in 

chapter 6, but more specific descriptors are probably better (fiber clump, fiber 

tangle, platelet cluster, etc.). 

Figure 9.1 Nonuniform dispersal of toughening agents often causes strengthlimiting flaws in ceramic composites. This figure shows a whisker clump that 

was the strength-limiting flaw in a siliconized SiC bend bar. The material was 

made by mixing chopped carbon fibers with SiC starting powders and then 

infiltrating with silicon. The carbon fibers were intended to react with the 

silicon to form dispersed silicon carbide fibers that would toughen the matrix. 

9-1 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Second phase aggregates can also inhibit local sintering causing pockets of 

porosity that also act as strength limiting flaws. Self-reinforced materials that 

rely on interconnected elongated-grain microstructures may develop 

excessively large grains that control strength as shown in Figures 6.22c and d. 

Very often reinforcing agents such as SiC whiskers can create local chemistry 

imbalances in the matrix leading to exaggerated grain growth. SiC whiskerreinforced alumina for cutting tools is an example. These large grains then can 

act as flaws. Although few would argue against reinforcing ceramics to 

enhance fracture toughness, this approach sometimes comes with a price. The 

very agents used to reinforce the material can create flaws that limit the 

material’s strength. 

9.2 Fiber-Reinforced Composites 

Fiber-reinforced ceramics and glasses may have much greater resistance to 

fracture if stressed in the same direction as the fiber reinforcements. Usually 

the fibers have greater strength and elastic modulus than the matrix. Optimum 

bonding does not necessarily mean perfect bonding, however, since the latter 

would lead to cleavage through the reinforcing fibers. Optimum bonding is 

such that the fiber-matrix interface is able to transmit some load, but then 

fractures or debonds along the interface. The fibers pull out of the matrix with 

additional loading. While pulling out they exert frictional or mechanical 

interlocking resistance so that more energy is consumed. The microstructure is 

designed to shed load from the matrix to the strong fibers and to divert cracks 

in the matrix to run along the matrix fiber interface. The processing must be 

done carefully so that the interface strength is controlled and that the fiber does 

not deteriorate too much from microstructural changes, surface pitting, or 

interactions with the matrix. 

Fractography of fiber-reinforced ceramics and glasses is different than that for 

monolithic materials. Fracture surfaces are often very rough and there often is 

no point looking for a specific fracture origin. There are exceptions. Figures 

9.2 and 9.3 shows a silicon carbide filament-reinforced silicon nitride that did 

have discrete fracture origins. 

Fracture typically entails damage accumulation and coalescence leading to 

rupture. Often the objective is to observe the general fracture mode and to 

ascertain whether it was due to tensile fiber pull out and fiber rupture, shear 

delamination in the matrix parallel to the fibers, compression buckling of 

fibers, or buckling of a protective coating in compression. The failure mode 

depends upon whether the structure is one-, two, - or three-dimensionally 

9-2 
 

 

 

 

 

 

 

Ceramic and Glass Composites 

Figure 9.2 A CVD SiC (SCS-6) filament-reinforced Si3N4 bend bar tilted back 

at an angle so as to view the fracture and tensile surfaces. A series of 

branching cracks radiate outwards from one origin site (arrow). 

(a) (b) 

Figure 9.3 Fracture surface of CVD SiC filament reinforced Si3N4 bend bar. 

(Ref. 1) (a) shows that there was not very much fiber pullout in this case. 

Wake hackle in the matrix behind the filaments are telltale indicators of the 

local direction of crack propagation through the matrix and lead back to the 

origin, which is a porous seam in the matrix shown by the white arrows in (a) 

and (b). 

9-3 
 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

reinforced and the size and shape of the component and the loading conditions. 

Although bend testing is reliable for creating tensile fractures in monolithic 

ceramics and glasses, composites are more susceptible to alternative failure 

modes when loaded in bending. Figures 9.4 and 9.5 shows tensile stress 

fractures with good pull-out in a fiber-reinforced epoxy and a ceramic, 

respectively. 

(a) 

(b) 

Figure 9.4 Fracture surface of an E-glass fiber reinforced epoxy matrix onedimensional composite. The arrow marks the origin in (b). (courtesy of S. 

Scherrer) 

9-4 
 

 

 

 

 

 

 

 

 

Ceramic and Glass Composites 

Figure 9.5 SEM image of the fracture surface of a SiC (Nicalon) ceramic 

fiber – zirconia matrix composite. The arrows mark fracture origins and 

mirrors in several fibers. (Ref. 2, courtesy of R. Rice) 

Images such as these can be used to qualitatively assess the fiber-matrix 

interface bonding and also if any damage has been done to the fibers during 

processing. Very often the fibers have fracture mirrors and it is instructive to 

ascertain whether the fibers have fractured from surface or internal flaws. 

Furthermore, the fracture mirror sizes in the fibers may be analyzed and used to 

obtain good estimates of the stress at fracture in the fibers. This strength can be 

compared to analytical predictions for the composite behavior to determine 

whether the load redistribution to the fibers was as expected. The strength of 

the fibers in the matrix may be compared to the strength of virgin fibers tested 

alone. Processing interactions with the matrix or with a fiber coating often 

weakens the fibers. 3,4,5,6,7,8 

Some material science or material development investigations focus on crack 

interactions with the structure. The scanning electron microscope is the 

preferred tool. Emphasis is placed on cracking in the fibers, the fiber pullout 

lengths, cracking in and around the fiber-matrix interface, or the microcrack 

9-5 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

density in the matrix. Micro-analytical chemical analyses are often used to 

study the chemistry of the fibers and the interfaces. 

The usual desired mode of fracture is one with a very rough fracture surface 

with ample fiber pullout. Observation that portions of a composite structure 

have fractured in a planar fashion with minimal pull out is often a sign that 

some deterioration in the composite has occurred or that the fiber-matrix bond 

is too strong. For example, Glass et al.9 analyzed a spectacular rupture of a 

large fiberglass–epoxy tank that stored pressurized sulphuric acid. Large 

portions of the tank had the classic rough fracture fractures with extensive fiber 

pullout, but the origin region had flat planar fracture zones. The cause was 

traced to spilled acid on the exterior surface that penetrated matrix microcracks 

and reached and degraded the reinforcing E-glass fibers. Fracture mirror 

analysis showed the fibers broke at only 20 % to 40 % of the baseline fiber 

strengths. Progressive fracture from the exterior occurred over a period of time 

as outer fibers broke, matrix microcracks opened up, and more acid penetrated 

and damaged more fibers until the vessel ruptured. 

Composites that are exposed to elevated temperature sometimes fracture with 

planar fracture regions near the surface surrounding a core of rough fibrous 

fracture in the core.10,11,12 Oxidation, corrosion, or surface reactions can alter 

the fiber–matrix bonding or fiber strength causing composite degradation. 

Composites that are tested transversely, that is with tensile stresses 

perpendicular to the fiber directions, are usually very weak. Fractography in 

such cases often is focused on the crack propagation paths around or through 

the fiber and the fiber-matrix interface. Hull shows several examples in section 

7.2 in his book on fractography.13 

9.3 Dental composites 

Although resin-matrix composite materials might seem to be outside the scope 

of this Guide, they are in fact brittle. Many dentists use dental composites as 

an alternative to ceramics, glass ceramics, or conventional porcelain fused to 

metal crowns. Crowns fractures are a serious problem. Fractographic analysis 

is underutilized by the dental materials community, but that is now changing in 

large part due to the leadership of Drs. S. Scherrer, Janet Quinn, K. Anusavice, 

J. Mecholsky, M. Øilo, and J. Robert Kelly as well as the widespread 

distribution of the first edition of this Guide. 

Dental composites in the form of “direct” restoration monomers with ceramic 

or glass fillers can be applied to cavities in a patient’s tooth in a dentist’s office 

9-6 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Ceramic and Glass Composites 

They are cured in situ with a “blue” ultraviolet light to polymerize the material. 

Alternatively, larger “indirect” restorations such as crowns (for more damaged 

teeth) can be made by first making an impression. An outside laboratory then 

fabricates a custom shaped tooth restoration that is cemented onto the reduced 

and prepared tooth. These indirect restoration materials are typically prepared 

as blanks by a manufacturer and then are custom ground to shape. (In some 

instances, an alternative procedure is taken whereby the dentist digitally scans a 

tooth and uses a CAD/CAM machine to mill crowns in his own office.) The 

blanks have a resin matrix with a very high glass or ceramic filler content, and 

are designed to have good strength, wear and fracture resistance, and good 

aesthetics. They behave as brittle materials and are usually linear to fracture 

when strength tested in a laboratory. A simple rule of thumb is that the more 

the filler, the more the material will fractographically resemble glasses and 

ceramics. Filler contents of as much as 70% by volume are not uncommon. 

There are some similarities, but many important differences in the fractography 

of dental composites versus glass ceramics and ceramics. The lower elastic 

modulus of resin composites and often lower strengths mean they have much 

less stored elastic energy when fracture occurs. Thus, there will be less 

fragmentation and branching. Fracture surfaces will have some of the same 

markings described in this book, but sometimes are described in the polymer 

literature with different nomenclature. The reader is referred to Anne RoulinMoloney’s excellent book14 on the fractography of polymers and composites 

for more information about resin-matrix composites in general. Much of the 

fractography is aimed at observing filler-matrix interfaces and what cracks do 

at the interfaces as opposed to finding origin flaws. 

Fracture mirrors may exist, and are described as smooth areas on the fracture 

surface at the origin. With highly-filled composites, the interpretation is the 

same as described in this book: the mirror is a smooth zone surrounding the 

fracture origin flaw. By the time the rough hackle lines form, the crack has 

reached terminal velocities (e.g., 800 m/sec) that are much slower than in 

glasses and ceramics. Figure 9.6 shows fracture surfaces of one such highlyfilled resin matrix composite from Reference15. It has a mean strength of 145 

MPa, a Weibull modulus of about 10, and a fracture resistance of 0.9 MPa√m. 

to 1.3 MPa√m indicative of a material with a slight R-curve. These materials 

have very small glass filler particles and are translucent. Inclusions and 

agglomerates are easy to detect in uncoated bend bars, either with reflected 

illumination or sometimes more effectively with transmitted illumination. 

Once the specimens have been examined in the natural state, they can be gold 

coated. The fracture surfaces will be much easier to interpret. 

9-7 
 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) (b) 

(c) (d) 

(e) (f) 

Figure 9.6 Filled resin-matrix composite fractography. (a) shows the fracture 

mirror in an uncoated Paradigm MZ 100 bend bar, with transmitted light from 

under the bar, (b) shows the same view, but with reflected light after gold 

coating. (c) - (e) show SEM images of the mirror and the flaw, a calcium-rich 

inclusion. (f) shows the original blanks and a machined bend bar. (a-e 

courtesy J. Quinn) 

9-8 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Ceramic and Glass Composites 

On the other hand, mirrors in unfilled resins14,16 might look similar, but they are 

in fact a region of slow crack extension that occurs prior to the crack going 

critical. The “mirror” is in fact the critical crack. So one must be careful in 

interpreting “fracture mirror constants” in the polymers field. They may not be 

the same as those in the ceramic and glass field. 

Chapter 9 References 

1. J. J. Swab, G. D. Quinn, and D. J. Snoha, “Mechanical Behavior of a SiC-Fiber/Si3N4 

Composite,” U. S. Army MTL Technical Note TN 90-2, Watertown, MA, Sept. 1990. 

2.	 R. W. Rice and D. Lewis, III, “Ceramic Fiber Composite Based Upon Refractory 

Polycrystalline Ceramic Matrices,” pp. 117 – 142 in Reference Book for Composite 

Technologies, Vol. 1, ed S. N. Lee, Technomic Publ. Co., Lancaster, PA, 1989. 

3. M. D. Thouless, O. Sbaizero, L. S. Sigl, and A. G. Evans, “Effect of Interface 

Mechanical Properties on Pullout in a SiC-Fiber-Reinforced Lithium Alumino 

Silicate Glass-Ceramic,” J. Am. Ceram. Soc., 72 [4] (1989) 525 - 532. 

4. W. A. Curtin, “In situ Fiber Strengths in Ceramic-Matrix Composites from 

Fracture Mirrors,” J. Am. Ceram. Soc., 77 [1] (1994) 1075 –1078. 

5. I. J. Davies and T. Ishikawa, “Mirror Constant for Tyranno Silicon-TitaniumCarbon-Oxygen Fibers Measured In Situ in a Three Dimensional Woven Silicon 

Carbide/Silicon Carbide Composite,” J. Am. Ceram. Soc., 85 [3] (2002) 691-693. 

6. A. J. Eckel and R. C. Bradt, “Strength Distribution of Reinforcing Fibers in a 

Nicalon Fiber/Chemically Vapor Infiltrated Silicon Carbide Matrix Composite,” J. 

Am. Ceram. Soc., 72 [3] (1989) 455-458. 

7. J. F. Jamet, D. Lewis, and E. Y. Luh, “Characterization of Mechanical Behavior 

and Fractographic Observations on Compglas SiC/LAS Composites,” Cer. Eng. 

Sci. Proc., 5 [7] (1984) 625 – 642. 

8. A. C. Jaras, B. J. Norman, and S. C. Simmens, “The Measurement of Glass Fibre 

Strength in Composites from Studies of their Fracture Surfaces,” J. Mat. Sci., 18 

(1983) 2459 – 2465. 

9. S. J. Glass, E. K. Beauchamp, M. J. Carr, T. R. Guess, S. L. Munroe, R. J. Moore, 

A. Slavin, and N. R. Sorenson, “Failure Analysis of A Fiberglass-Reinforced 

Plastic Pressure Vessel,” pp. 527 – 541 in Fractography of Glasses and Ceramics, 

III, Ceramic Transactions, Vol. 64, J. Varner, Fréchette, V., and G. Quinn, eds., 

American Ceramic Society, Westerville, OH, 1996. 

10. K. M. Prewo, “Tension and Flexural Strength of Silicon Carbide Fibre-Reinforced 

Glass Ceramics,” J. Mat. Sci., 21 (1986) 3590 – 3600. 

11. T. Mah, M. G. Mendiratta, A. P. Katz, R. Ruh, and K. S. Mazdiyasni, “HighTemperature Mechanical Behavior of Fiber-reinforced Glass-Ceramic Matrix 

Composites,” Comm. Am. Ceram. Soc., 68 [9] (1985) C 248-251. 

12. T. Mah, M. Mendiratta, A. Katz, and K. Mazdiyasni, “Recent Developments in 

Fiber-Reinforced High Temperature Ceramic Composites,” Bull Am. Ceram. Soc., 

66 [2] (1987) 304 – 308, 317. 

13. D. Hull, Fractography, Observing, Measuring and Interpreting Fracture Surface 

Topography, Cambridge University Press, Cambridge, 1999. 

9-9 
 

 

 Fractography of Ceramics and Glasses 

14.	 A. C. Roulin-Moloney, Fractography and Failure Mechanisms of Polymers and 

Composites, Elsevier, London and New York, 1989. 

15.	 J. B. Quinn and G. D. Quinn, “Material Properties and Fractography of an Indirect 

Dental Resin Composite,” Dent. Mat., 26 [6] (2010) 589 – 599. 

16. E. H. Andrews, Fracture in Polymers, Elsevier, NY, 1968. 

9-10 

 

 

 

 

 

 

 

 

Case Studies 

10. Case Studies 

The following case studies illustrate the application of fractographic techniques
 

for failure analysis. Many additional cases are tabulated in Appendix B and
 

can be found in the books listed in Appendix A.
 

Case 1 Ruptured Rotor (Ceramic gas turbine rotor)
 

Case 2 Busted Barrel (Silicon carbide gun liner) 

Case 3 Conflicting Carbide Data (Silicon carbide flaws and slow crack
 

growth) 

Case 4 Vulnerable Vials (Broken medicinal vials) 

Case 5 Damaged Dome (Coarse-grained aluminum oxynitride IR dome) 

Case 6 Suffering Setter Plate (Silicon carbide furnace plate) 

Case 7 Ruptured Radomes (Fused silica missile nosecones) 

Case 8 Maligned machinists (Bend bars made by different machine shops) 

Case 9 Modeler’s Match (Fracture origins in MEMS scale micro tensile 

specimens) 

Case 10 Fractious Fractographers (A VAMAS fractography round robin) 

Case 11 Perilous Prostheses (Four ceramic dental crowns) 

Case 12 Mangled Margins (Twenty alumina dental crowns ) 

Case 13 Ruined Refractors (Single crystal sapphire solar lenses) 

Case 1: Ruptured Rotor (Ceramic Gas Turbine Rotor) 

Figure 10.1 shows a model gas turbine rotor made of hot-pressed silicon nitride 

that was designed to fail. It was one of twelve such model structures made by 

the Ford Motor Company’s scientific research laboratory in the late 1970s to 

mid-1980s.1,2 It was part of an ambitious endeavor to incorporate advanced 

ceramics into automotive gas turbine engines. Considerable effort was 

expended on improving ceramic materials, developing reliability codes, 

generating data bases, fabricating parts, and running them in test rigs. Full 

scale engine testing was expensive and risky, so a small project was set up to 

verify the ceramic design and reliability codes by using a realistic model rotor. 

This simulated rotor was mounted in a hot-spin rig shown in Figure 10.2 and 

rotated at 50,000 rpm while heated by hot gasses to a peak rim temperature of 

1260 o

C (2300 oF). These conditions were representative of the gas turbine 

operating conditions and the speed, mass, and shape of the part were designed 

so that some of the rotors would fracture in a time-dependent manner, from 

slow crack growth of preexisting flaws. The goal was to correlate failure times 

to predictions based on life prediction computer models. 

10-1 
 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Twelve model rotors were fabricated and tested. The rotors were designed to 

have a high probability of survival on loading (> 95 %), but to probably fail 

within 25 h. The design analysis assumed that failure would occur due to slow 

crack growth of preexisting volume-distributed flaws. The rotors were 

intended to break in the thin web where the maximum principal steady state 

stress was 131 MPa (19 ksi) and the peak temperature was 1240 o

C (2260 oF). 

Figure 10.1. A Ford silicon nitride 95 mm diameter model gas turbine rotor 

that survived 25 h intact. The outer rim is thicker than a real rotor in order to 

simulate the mass of the vanes that would be attached in a real rotor. 

ceramic rotor 

ceramic spacer 

Figure 10.2 The Ford hot spin test rig (Ref. 2). 

10-2 
 

 

 

 

 

 

 

 

Case Studies 

Six broke in a time dependent manner as intended. Two rotors broke on initial 

startup. Four were run outs and did not fail in the allotted 25 hour times and 

were treated as censored outcomes. 

The rotors were made of a state-of-the-art hot-pressed silicon nitridea that was 

carefully machined to final dimensions. The particular grade was one of the 

most thoroughly-analyzed structural ceramics of all time and was eventually 

used as the world’s first reference material for the property fracture toughness.3 

Eventually enough data was available that a comprehensive fracture 

mechanism map was constructed,4 but that was after the conclusion of the Ford 

study. The finite element and heat transfer models used the best data available 

and temperature-dependent material properties. A substantial amount of 

laboratory test coupon data were available including flexural and tensile 

strengths, elastic properties, and slow crack growth parameters from three types 

of tests: variable stressing rate strength tests (dynamic fatigue), crack velocity - 

fracture mechanics tests (double torsion), and stress rupture (flexural and 

tensile). 

Figure 10.3 shows the reliability as a function of time. The three types of slow 

crack growth data gave very divergent predictions. The six data points show 

the six rotors that failed at times from 0.2 h to 18.6 h. The predictions made 

Figure 10.3 Reliability versus time showing reliability predictions based on 

three different types of laboratory test data (solid lines) versus the actual 

outcomes for six rotors that failed.(Refs. 1,2) The actual failure times matched 

the predictions based on stress rupture data. The two solid points correspond 

to rotors that were fractographically analyzed by the author. 

a Grade NC 132, Norton Co., now St. Gobain Advanced Ceramics, Worcester, Ma. 

10-3 
 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

with the laboratory stress rupture data gave the best correlation to the actual 

lifetimes. The Ford authors concluded that stress rupture data was the best for 

predicting reliability versus time.1,2 

Fractographic analysis revealed that fracture did not occur quite as anticipated. 

The test rig was constructed in such a manner that most of the fragments were 

retrievable. The longest (18.6 h) and third longest (13.8 h) rotors were sent to 

the author for fractographic evaluation. It was felt that these would have large 

slow crack growth zones that would be easy to find. One virtue of NC 132 is 

that it was a very fine-grained fully-dense material with minimal second phase. 

Rotor spacer 

A A

g 

h 

A g 

h 

Primary 

origin 

g 

h 

Figure 10.4 Reconstructed model rotor and spacer. Fracture initiated either 

at location “A” on the left side of the bore of the rotor or at two locations “g” 

or “h” on the opposite side. Fracture started at “h.” Fracture “g” occurred 

later. The lower schematic shows the overall fracture pattern in the rotor. 

10-4 
 

 

 

 

 

 

 

 

 

 

 

 

Case Studies 

Fracture markings were quite clear and the local direction of crack propagation 

could be assessed in every single piece. One rotor (18.6 h) was reconstructed 

piece-by-piece over the course of one week culminating in the assembly shown 

in Figure 10.4. It was very much a puzzle-solving exercise. Many fragments 

from the rim were missing, but they did not matter. The figure also shows that 

the ceramic spacer part had also fractured and its pieces were commingled with 

the rotor, thereby complicating the matter. Nevertheless, the central and web 

portions of the primary rotor were almost completely reassembled. After 

completion, the rotor was taken apart again and a map of crack propagation 

directions was made. The interpretation was clear. Primary fracture 

commenced at the bore or the curvic coupling teeth and cracks branched and 

ran out to the rim as shown in the schematic in Figure 10.4. Every single 

fragment was examined in a futile search for a web origin with a slow crack 

growth zone, but none were found. All primary fractures started from the bore 

or coupling area. There were three origins from the inner part of the rotor. 

Origin “A” was at the exact edge of the bore, which although chamfered, 

nonetheless had surface grinding cracks (Figure 10.5a-c). It could not be 

determined whether slow crack growth had enlarged the grinding cracks. The 

other two origins “g” and “h” were located on the other side of the bore and 

were from the curvic coupling teeth machined into the rotor (figure 10.5d,e. 

These were part of the attachment scheme. One origin, “g”, was a parallel 

machining crack (aligned parallel to the axis of grinding) in the tooth, and the 

other, “h”, was an impact – contact crack also on a tooth. 

Thus, fracture started on one side of the rotor in the bore area at either “g” or 

“h,” then branched into two main crack(s) that ran out to the rim. This opened 

up the disk on that side. The unbalanced forces triggered a rupture on the 

opposite side at site “A.” Origin “A” had a well-defined small fracture mirror 

allowing a stress estimate of a 630 MPa (92 ksi) which is much greater than the 

model estimated bore stress of ≈172 MPa (25 ksi). This, plus the violent 

branching at that site, confirms that site “A” was a secondary fracture. Primary 

fracture occurred on the other side of the bore at site “g” or site “h,” causing 

the rotor to go completely out of balance. Origin “h” was determined to be the 

first break based on crack intersection patterns. The flaw was a large contact 

damage crack with a large fracture mirror. The local stress was estimated to be 

500 MPa (72 ksi) from the flaw size and approximately 580 MPa (84 ksi) from 

the fracture mirror. Origin “g” was determined to have fractured after site “h”, 

since the crack from “g” stopped when it encountered the prior crack from “h.” 

At these locations the computer models indicated there were modest bore 

stresses (≈138 MPa, 20 ksi) and temperatures (1150 o

C, 2200 oF). These did 

not match the fractographic estimates. 

10-5 
 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

origin A 

(a) (b) 

(c) 

Branch 

Fracture 

Mirror 

origin A 

origin h 

(d) (e) 

Figure 10.5 Fracture surfaces showing a fracture mirror and the origin site 

A, which is a surface crack from grinding. (b) and (c) highlight the fracture 

mirror and grinding crack origin at site “A.” The primary fracture origin was 

contact crack damage at site “h” on the curvic coupling teeth as shown in (d) 

and (e). Origin “h” is marked by the larger arrow on the right. Notice how 

the radiating hackle lines converge on it. 

10-6 
 

 

 

 

 

 

 

 

 

 

 

 

Case Studies 

An important fractographic observation helps explain the fracture. When the 

rotor and spacer were assembled in the hot test rig, it was customary to separate 

them with a thin platinum foil to prevent direct ceramic-to-ceramic contact. 

The fractographic examination showed uneven platinum wear traces and even 

some bare spots. Could it be that the foil deformed or crept with time such that 

ceramic-to-ceramic contact eventually occurred? If so, then the rotor stress 

distribution may have become unbalanced or local stress concentrations or 

contact stresses may have contributed to cause an initial contact crack and 

fracture at site “h.” 

This exercise was a good example of one of the author’s laws of fractography: 

“The first one is the hardest.” While at first glance, reassembly of a burst rotor 

may appear to be a formidable task, it was merely time consuming. The second 

rotor took less time to analyze. It also had fracture initiation sites at the bore or 

at the teeth. Years later, the author became aware of comparable work done 

on model silicon nitride rotors at Daimler-Benz’s research center in Stuttgart.

5 

They also reconstructed burst silicon nitride rotors and used fractography to 

find that grinding cracks in the bores sometimes were fracture origins. 

Some lessons learned from this case study were: 

a.	 Fracture occurred from a different cause than expected and modeled. 

b. Stress rupture data may have been the best for reliability estimation, but 

the correlation of failure times was fortuitous. 

c.	 Volume flaws were not found in the two fractured rotors. Surface
 

machining cracks and contact cracks were found and Weibull area 

scaling should also have been included in the reliability model.
 

Case 2: Busted Barrel (Silicon Carbide Machine Gun Liner) 

The U. S. Army has off and on over the past twenty-five years investigated 

ceramic liners to improve gun barrel life and reduce mass. The low density, 

high compression strength, refractoriness, and erosion and wear resistance of 

ceramics could be advantageous. Figures 10.6 and 10.7 show drawings for a 

50-caliber machine gun breech with a ceramic liner from a project conducted 

for the U. S. Army in the 1980s.6 One design placed the ceramic into 

compression by shrink fitting a steel sleeve around the ceramic. The steel 

sleeve was heated and placed over the cool ceramic tube. As the steel cooled, it 

contracted and put the ceramic liner into axial, radial, and hoop compression. 

The dimensions and temperature differentials were chosen so that the ceramic 

liner was always in compression, even when firing a projectile. Several 

ceramics were tried, but most testing was on a sintered α-SiC. Figure 10.7 

shows the maximum axial stresses, σx, in the three components 

10-7 
 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Breech Muzzle →

Figure 10.6 Schematic of the breech end of the 50 caliber (12.7 mm) gun 

barrel (Ref. 6). 

(a)

 (b) 

Figure 10.7 Axial stress distribution, σx, as a function of position x in the 

Ceramic liner 

σx 

X 

Steel jacket 

Steel sleeve 

670 MPa 

assembly along the bore (a). The stresses correspond to the three parts shown 

in a sectional view in (b) which is a simplified version of the assembly shown in 

the previous figure. The stresses included those from shrink fitting and 

projectile firing. (Figure from Ref. 6) 

10-8 
 

 

 

 

 

 

 

 

 

Case Studies 

of the assembly from combining the residual shrink fit and projectile firing 

stresses. The maximum axial stress σx was 670 MPa in compression in the 

ceramic liner mid-portion. This stress tapered to zero at each end. The axial 

stress maximums were only 112 MPa tensile in the steel sleeve and 154 MPa 

tensile in the steel jacket. The maximum radial stress in the ceramic sleeve 

was 345 MPa compression, and the maximum hoop stress, σθ, was 590 MPa 

compression during projectile firing. The environment is severe, but if the 

ceramic was always under compression, perhaps it would not fracture. 

Some assemblies survived as many as a thousand single-shot firings, 

confounding skeptics who felt that the assembly would not survive one shot. 

One assembly that did develop ceramic cracking after a few hundred shots had 

circumferential fractures as shown in Figure 10.8a. The steel sleeve and jacket 

were machined away to allow extraction of the ceramic as shown in Figure 

10.8b. The fracture planes were perpendicular to the axial direction, suggesting 

that hoop stresses from internal pressure during projectile firing were not the 

cause of failure. The latter would have created radial cracking. Fractographic 

analyses showed that every fragment fractured from one or more contact cracks 

that were periodically spaced on the outer rim of the ceramic where it contacted 

the steel sleeve. 

So if the shrink fitting created compressive stresses, where did the tensile 

stresses come from? The plane of the fractured surfaces and also of the initial 

semi-elliptical contact cracks indicated that the tensile stresses were axial: σx. 

Figure 10.7 shows that the axial residual compressive stresses did taper off 

towards the tube ends. The most likely sources of tensile stresses are dynamic 

stress waves generated during the firing. Even if these are initially 

compressive, they can change phase and become tensile when the stress wave 

reflects off of the end faces. Furthermore, the elastic wave velocity and 

impedance of the silicon carbide and the steel sleeve are not matched, and 

stress waves propagated axially at different rates. Hence, a dynamic tensile 

stress could develop at the steel/ceramic interface due to mismatch of the 

transient elastic strains. 

The contact stress cracks often were periodic around the rim. Although the 

parts were machined to tight tolerances, it is likely that slight variations in the 

mating surfaces led to an uneven fit and stress concentration sites that triggered 

the contact cracks. 

So, in this case, the ceramic was designed to always be in compression, but 

fractography showed otherwise. Some design and modeling improvements 

were suggested. The tolerances and surface specifications for the mating parts 

10-9 
 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

could be changed. The elastic properties of the ceramic and the confining 

sleeve could be matched better. More sophisticated stress models could 

examine the transient stress states in the assembly. 

(a) (b) 

(c) 

(c) 

Figure 10.8 Fractured α-SiC gun barrel. (a) is a schematic sketch. (b) 

shows the fracture surface of one of the ring shaped fragments. Contact cracks 

are spaced periodically around the outer rim and appear as shadows in this 

view (arrows). The bore diameter is 12.7 mm. (c) shows a close-up of one of 

the contact cracks. 

10-10 
 

 

 

 

 

 

 

 

 

Case Studies 

Case 3: Conflicting Carbide Data 

(Silicon carbide flaws and slow crack growth)
	

This case is not about a specific component, but about the risk of treating all 

flaws as if they behave the same. As part of a program on characterization of 

structural ceramics for heat engine ceramics in the 1980s, the author conducted 

extensive stress rupture testing of pressureless-sintered α-SiC.b 7,8,9 Stress 

rupture testing entails applying a constant stress to a specimen, at a level below 

that needed to cause fast fracture, and measuring the time to failure. Slow 

crack growth can cause flaw enlargement and time-dependent fracture. 

Sometimes such tests are termed creep rupture if the loading conditions are 

such that bulk creep deformation leads to accumulated damage and fracture. 

The author’s data7,8,9 at 1300 o

C in air was consistent with data published by 

another team.10 Both teams detected time-dependent fractures due to 

intergranular slow crack growth of preexisting flaws as shown in Figure 10.9. 

Figure 10.9 Fracture origin in a 1300 oC stress rupture SiC specimen. The 

origin is a large intergranular slow crack growth region that may have started 

from the side of or beneath the large pore. 

b	 Hexoloy SA, Carborundum, Co., Niagara Falls, NY. Now St. Gobain Advanced 

Ceramics. 

10-11 
 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

A much different story emerged at 1200 oC. This author recorded timedependent fractures such as shown in Figure 10.10 with origin flaws as shown 

in Figure 10.11. The other team did not observe a single time-dependent 

fractures at all in their experiments. Why did the two groups have such 

contrasting outcomes at 1200 o

C? The answer was that they used Knoop 

indented, artificially-flawed specimens, whereas the author used as-machined 

specimens that allowed the material to fail from whichever of its flaws it 

“chose.” 

Figure 10.10 Stress rupture curve at 1200 oC in air for sintered α-SiC.

7-9 

Every time-dependent fracture was from a pore or porous zone and only if it 

was surface connected. The hollow points with arrows to the left are breaks on 

loading and are consistent with the fast fracture strength which was about 350 

MPa. The slow crack growth exponent N was 40.8. 

(a) (b)
 

Figure 10.11 Fracture origins in 1200 oC α-SiC specimens that fractured at
 

863 h (a) and 1460 h (b) from surface-connected porous regions.
 

10-12 
 

 

 

Case Studies 

At 1300 oC and above, easily-detected intergranular slow crack growth caused 

fracture from volume-distributed flaws such as large grains, pores, or 

agglomerates as shown in Figure 10.9. Some of these volume origins were 

located in the bulk. Both teams detected this behavior. 

On the other hand, every one of the 1200 o

C time-dependent fracture origins 

was a pore or porous region connected to the outer surface as shown in Figure 

10.11. These flaws did not have slow crack growth markings. Evidently these 

flaws were susceptible to very localized stress-corrosion crack extension (from 

oxidation) that sharpened or locally extended tiny microcracks on the pore 

periphery. These flaws did not need much change in flaw size or severity for 

them to go critical, since the failures occurred at stresses (250 to 350 MPa) very 

close to the fast fracture strength (350 MPa). Figure 10.12 summarizes the 

fractographic findings. 

Figure 10.12 Fracture origins in α-SiC. 

10-13 
 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

The other testing team did not detect the 1200 oC stress corrosion mechanism 

since they used Knoop artificial flaws in all of their bend specimens.10 

Artificial flaws often are effective tools for studying fracture, but in this 

instance they produced misleading results. At 1300 o

C and above, the Knoop 

flaws did grow by slow crack growth. They did not grow by slow crack growth 

or by stress corrosion at 1200 o

C. The Knoop flaws were not susceptible to the 

same failure mechanism as the pores and porous zones in the as-machined 

specimens. The Knoop flaws reduced the fast fracture strength of the 

specimens to less than 200 MPa. Their stress rupture specimens were loaded at 

stresses from 100 to 200 MPa, a stress well below that necessary to activate the 

pore flaw stress corrosion mechanism. 

This case demonstrates that it is often best to let a material reveal what type of 

flaws it is apt to fail from, rather than rely on artificial indentation flaws. 

Case 4: Vulnerable Vials (Broken Medicinal Bottles) 

A pharmaceutical company had a problem with neck rim cracking in 17 mm 

diameter medicinal vials. The cracking rarely caused the vials to break, but 

trace leakage and loss of the airtight seal due to through-cracks in the rim 

caused great concern. Loss of seal carries a risk of contamination with 

potential fatal consequences (e.g., see case 10.8 in Fréchette, Ref. 11). 

Optical stereomicroscope examination of vials with either intact or severed 

caps revealed a circumferential damage zone in the glass just under the 

aluminum cap seal (Figure 10.13 a,b). Focusing through the glass revealed that 

some cracks had penetrated deep beneath the surface although the full extent of 

the penetration was difficult to assess. Examination of a vial with a severed 

cap confirmed the contact cracks could reach all the way to the interior. No 

defects in the glass were detected. The concentricity of the neck with the vial 

body was within specifications. The source of the fractures was traced to a 

misalignment of the cap crimping machine star-wheel assembly. An incidental 

finding was the beautiful example of scarps, shown previously as Figure 5.48a, 

on one of the fracture surfaces of a severed cap. The scarps were consistent 

with the outer surface cracks having reached the vial interior and the interior 

medicinal fluid, therein aiding the final fracture of the glass neck. 

10-14 
 

 

- 

Case Studies 

Glass vial 

Aluminum cap 

Close up view of 

the neck 

(a) (b)

 (c) 

Figure 10.13 Rim cracking in glass medicinal vials. (a) and (b) show blunt 

contact cracking damage (arrows). (c) shows a top view of the fracture surface 

of a vial with a severed cap. 

10-15 
 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Case 5: Damaged Dome (AlON IR dome) 

This case illustrates what can be accomplished with optical microscopy alone, 

even on very difficult ceramics. A coarse-grained aluminum oxynitride missile 

sensor dome broke during a thermal stress proof test. It broke into several 

pieces, and the branching patterns directed attention to an area near the base as 

shown in Figure 10.14a. One might expect the origin to be at the dome rim, 

possibly at a bevel, but that was not the case. 

The roughness of the microstructure hid nearly all fracture markings as shown 

in 10.14b. Careful adjustment of low-angle grazing lighting at higher 

magnifications revealed local hackle lines and twist hackle within the 

transgranularly-fractured grains as shown in Figure 10.14c and d. It was 

possible to read these lines back to the inside surface of the dome to a location 

well away from the rim. Close-up examination in Figure 10.15 revealed the 

origin site had zipper machining crack hackle lines within one or two grains, 

indicating the origin was grinding or scratch damage that was not completely 

removed by polishing. 

(a) 

(d) (c) 

(b) 

Figure 10.14. AlON infrared missile dome. (a) shows how the branch on the 

right leads away from the origin area. (b) shows the fracture surface with the 

coarse microstructure. 

10-16 
 

 

 

 

 

Case Studies 

Unfortunately, the dome was not available for SEM confirmation or further 

examination of the inside surface for traces of a scratch at the origin site. The 

polished surfaces in general did have an assortment of polishing flaws as shown 

in the insert. Similar grinding/polishing flaws inside single grains of aluminum 

oxynitride have been reported by Swab et al. 12 A similar approach may be 

used to find origins in coarse-grained polycrystalline alumina (PCA) lamps. 

Polished surface 

Figure 10.15 Wake and twist hackle within cleaved AlON grains lead back to 

the inside surface and to tiny zipper machining crack hackle lines within one or 

two large grains. This image was taken at the maximum magnification (125X) 

of the stereoptical microscope that was used for the examination. The insert 

shows the outer polished surface with polishing scratches and a well-ground 

bevel-chamfer. 

10-17 
 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Case 6: Suffering Setter Plate (Silicon Carbide Furnace Plate) 

A large silicon carbide furnace setter plate fractured during a furnace run. Was 

the material faulty or was there some other cause? Figure 10.16 shows one half 

of the broken plate. It has the wavy fracture that is a telltale sign of a centerheated plate fracture as shown in Section 4.12. Figure 10.17 shows the fracture 

surface revealing that fracture initiated from grinding damage at a hole near the 

rim. No material flaws of consequence were detected. The mirror size was 

about 9.3 mm. Using an Ao value of 11 MPa√m for a dense hot-pressed SiC 

(from Appendix C), the stress at fracture was estimated to be only 115 MPa. 

The material was completely satisfactory, but the machining, grinding, and 

attachments were not. This failure analysis required only a one minute visual 

inspection. 

(a) 

(b) (c) 

Figure 10.16 Fractured SiC furnace setter plate (a). The curvy pattern is 

typical of a center-heated plate fracture as shown in Chapter 4. The arrow 

marks the origin which is a hole through the plate. The machined grooves 

were chipped and had ragged edges as shown in (b). The plate was 36 cm 

round and 1.9 cm thick. (c) shows the origin hole from the bottom surface. 

There is spall and evidence of a chemical reaction with a mounting bolt or 

washer. (plate courtesy of B. Mikijelj) 

10-18 
 

 

 

 

Case Studies 

(a) 

(b) 

Figure 10.17 Fracture surface of the SiC setter plate. A fracture mirror at the 

origin (arrow in (a)) attests to a moderate stress level. Hoop tensile stresses on 

the rim were generated by thermal strains. Thermal expansion of the hot plate 

center was resisted by the cooler rim. As the crack propagated towards the 

plate middle it slowed since it reached hot regions that were originally in 

compression. The fracture surface became relatively featureless and the crack 

meandered. The hole through the plate at the origin (b) had chipping and 

grinding damage that weakened the plate. 

10-19 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Case 7: Ruptured Radomes (Fused Silica Missile Nosecones) 

First generation U. S. Army Patriot air defense missiles in the early 1980’s 

utilized fused-silica radomes (Figure 10.18). Radomes are nosecones that are 

radar transmissive. The radome was prepared by slip casting and then sintering 

silica particles to greater than 85% density. This porous structure had a low but 

acceptable strength and, like an insulating firebrick, was more effective in 

arresting cracks than if the radome were a fully-dense glassy body. Although 

sintered fused silica is weak, it was selected since it had a low coefficient of 

thermal expansion and thus would be less susceptible to thermal stresses from 

aerodynamic heating. Selected radomes were deliberately loaded to fracture or 

proof tested as part of the engineering development program. This was done by 

attaching a sling to the side of the radome and pulling it laterally several times 

in different directions. Unexpected fractures occasionally occurred. Most 

fractures occurred in the ceramic where it was joined to a threaded fiber-epoxy 

thread ring attachment part (Figure 10.19a,b). Fractography was difficult in the 

porous weak ceramic since it had a rough fracture surface, but a number of 

fractures were successfully diagnosed. Figure 10.19 shows an instance where 

atypical grinding damage created strength limiting cracks. Better control of the 

surface grinding eliminated the problem. Figure 10.20 shows a different 

example wherein excessive shrinkage of the epoxy bond caused debonding. 

This created a stress concentration site in the ceramic that initiated fracture. 

The remedy in this case was to adjust the bond-cure thermal cycle. 

(a) (b) 

Figure 10.18 U. S. Army first generation Patriot air defense missile. The 

radome was 1.1 m tall with a 40 cm diameter. (b is courtesy Raytheon) 

10-20 
 

 

 

 

 

Case Studies 

Fused 

silica 

Fiber – epoxy thread ring 

Epoxy 

bond 

Polymer 

spacer 

(a) (b) 

ceramic 

bond 

origin crack 

(c) contamination 

(d) 

Figure 10.19 Sintered fused silica radome that fractured during a proof test. 

(a) is a top view with the origin region marked by the arrow. The ceramic is 

bonded on the inside to a fiber epoxy attachment ring. (b) is a schematic 

showing the bottom of the radome where it was attached to the thread ring. 

Fractures in the ceramic typically occurred near the top of the thread ring. (c) 

shows a grinding crack at the origin area (arrows). (d) shows a grinding 

pattern on the inside ceramic surface. 

10-21 
 

 

 

 

 Fractography of Ceramics and Glasses 

(a) (b) 

(c) 

(d) 

Epoxy bond 

Fused silica Debonded 

region 

Figure 10.20 Sintered fused silica radome fracture. (a) shows the origin site 

on the inside wall. (b) is a close-up showing a fracture mirror centered on a 

feature on the inside wall of the ceramic at the top of the bond line. (c) is a 

composite of two SEM photos of the origin area. The origin is less obvious, but 

hackle lines point back (white arrows) to the circled area. (d) shows there is 

no significant material flaw at the origin, but the ceramic and the epoxy are 

debonded at the origin. 

10-22 
 

 

 

 

 

 

 

 

 

 

Case Studies 

Case 8: Maligned Machinists (Bend Bars from Machine Shops) 

This case was part of an investigation whether different machine shops could 

prepare common bend bars to a set of specifications, with minimal damage and 

at a reasonable cost. The work described here was done in the 1980s in support 

of the first standard test method for flexural strength of high performance 

ceramics in the United States, MIL STD 1942(MR).

13 The standard, which was 

adopted by the U.S. Army in 1983, served as a basis for American Society for 

Testing and Materials standard C 1161 in 199014 and (along with elements 

from analogous European and Japanese standards) was used as the basis for 

Inter-national Organization for Standards ISO 14704 in 2000.15 

Standardization has led to dramatic improvements in the quality of data and 

significant cost savings. Bend bar costs decreased as machine shops prepared 

specimens to a common, reliable procedure. 

As part of the standardization work, a major international round robin on 

flexural strength with 11 labs around the world was conducted in 1984 and 

1985.16,17 Several thousand specimens of 99.9 % pressureless-sintered alumina 

and reaction bonded silicon nitride were tested. The round robin showed that 

consistent flexural strength results could be obtained if the standard procedure 

with defined test fixtures and alignments were used. 

It has long been known that grinding may introduce strength-limiting damage. 

A side topic investigated in the round robin was whether different machine 

shops could meet the grinding specifications and prepare damage-free B type 

bend bars (3 mm x 4 mm x 50 mm). Requests for price quotations were sent to 

eight machine shops to prepare 20 specimens from the sintered alumina 

according to the MIL STD 1942. The specifications were for a multi-step 

grinding process ranging from coarse initial grinding to final finishing. Finer 

grinding wheels with reduced depths of cut were used at each step with the 

objective of removing prior damage and minimizing final residual damage. 

One shop did not bid. Two others bid $101 and $112 per bar and were not 

contracted since these prices were exorbitant. In 1984 – 1985 the typical price 

for bend bars from experienced shops were in the $12 to $25 range. Plates 

from the common batch were sent to each of the remaining five shops which 

bid from $15 to $50 per bar. The bars that they machined were then broken in 

four-point flexure on 20 mm x 40 mm spans. The surfaces and edges of all 

bars were inspected beforehand to ascertain whether the bars met specifications 

or if there were any signs of machining damage. The results are shown in 

Table 10.1. The initial visual examination showed that the five venders met all 

of the specifications for the most part. Deviations were usually on isolated 

10-23 
 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

specimens. Machining damage can be hidden, however, and the strength 

testing was intended to reveal the damage. 

A comparison of the average strength values suggested that venders B, C, and 

D damaged their specimens since the strengths of their specimens were much 

lower than those from shops A and E. It would have been tempting to reject 

venders B, C, and D for further jobs on the basis of the strength outcomes, but 

the fractographic analysis revealed a surprising explanation for the outcomes. 

Optical examination of the fracture surfaces revealed that machining damage 

was not the prime factor in any of the five samples sets. Machining cracks did 

cause fracture in a few specimens, but the primary strength-limiting flaws were 

volume-distributed sintering flaws such as pores, porous zones, porous seams, 

agglomerates, and occasional inclusions. Figures 6.6a, 6.13, and 6.66a,e show 

some of these. The crucial difference was that the flaws varied between 

the plates. 

Shop 

Bar 

Price 

$ 

Were the 

specifications 

met? 

alumina 

plate 

Avg. 

MPa 

Std. 

Dev. 

MPa 

Char. 

Str. 

MPa 

Weibull 

Modulus Fracture Origins 

A 15 Yes P 372 42 391 10.3 

Round pores 

Porous zones 

Porous seams 

2 machining 

B 19 Yes, a, b 2 315 22 325 17.1 

Porous zones 

Porous seams 

Agglomerates 

4 machining 

C 20 Yes, c 1 301 30 314 11.5 

Porous seams 

Pores 

Agglomerates 

D 41 Yes, b 3 335 32 350 12.0 

Porous seams 

Porous zones 

Agglomerates 

E 50 Yes, b 4 373 36 389 11.9 

Pores 

Agglomerates 

2 machining 

a Minor edge chips on some 

b Some skip marks or deep striations
	

c Chamfers a bit uneven
	

Table 10.1 Flexural strengths of sintered 99.9% alumina from test sets 

prepared by different machine shops. The flaws that were strength limiting are 

listed in their order of frequency in each set. 

10-24 
 

 

 

 

 

 

 

 

 

 

Case Studies 

Porosity was the most common flaw type, but it manifested itself differently in 

the plates. In some plates the flaws were primarily round discrete pores. In 

other plates, the flaws were equiaxed regions of microporosity, porous seams, 

or porosity associated with inclusions. The plates were prepared from the same 

powder lot by the identical procedure and to all external appearances were 

identical. Only when the specimens were fractured was the true flaw character 

revealed. 

This case illustrates the hazards of interpreting strength results without 

supportive fractographic analysis. The variability in flaw character between 

billets and its effect on strength also underscores a serious problem for 

structural ceramic designers. If the flaws in a material vary between nominally 

identical ceramics pieces, this does not bode well for the success of reliability 

analyses. 

Case 9: Modeler’s Match
	

(Fracture Origins in MEMS Scale SiC Micro Tension Specimens)
	

Testing methodologies must keep pace with emerging technologies for 

miniature devices and structures for microelectromechanical systems (MEMS) 

and even smaller devices. Sharpe et al.18 investigated the strength of miniature 

silicon carbide specimens shown in Figure 10.21. These had cross section sizes 

of ≈ 200 μm or less and were prepared by Beheim at NASA-Glenn by chemical 

vapor deposition (CVD) followed by deep reactive ion etching (DRIE) to final 

shape.19 Specimens with straight, curved, and notched gage sections were 

tested in direct tension and fractured in a special miniature tensile tester. 

Weibull statistics were applied to scale the strengths and to determine whether 

surface or volume flaw scaling gave better correlation. The Weibull analysis 

using area scaling worked reasonably well for the curved and straight 

specimens, but gave a poor correlation for the notched specimens, which were 

much weaker than expected. 

Fractographic analysis was difficult, but productive.19,20 The CVD SiC had a 

coarse microstructure that created a very rough fracture surface that masked 

common fracture markings. Large SiC grains affected the crack propagation 

across the fracture surface, causing significant crack redirection and severe 

roughness as the propagating crack sought out preferred cleavage planes. The 

most helpful features for pinpointing the origin were twist hackle lines on 

cleaved grains and occasional large hackle lines that showed the local direction 

of crack propagation. These markings led back to the origins. 

10-25 
 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

1 mm 

Notched

 Straight 

Curved 

(a) (b) 

(c) 

100 μm 

Figure 10.21 Miniature 3.1 mm long SiC tensile test specimens. (a) shows 

curved, straight and small notched gage sections. The notch is too small to see 

in this view. (b) shows a fractured curved gage length specimen (c) shows the 

fracture surface of a straight section specimen. The arrow shows the origin is 

a 25 μm x 50 μm large grain at the root of an etch groove. (photos a and b, 

courtesy of W. Sharpe) 

The fracture origins in all three specimen types usually were a combination of a 

deep etch groove that combined with a large, favorably-oriented columnar 

grain. Cracks popped-in on a preferred cleavage plane in such grains. Thus, 

most critical flaws were a hybrid surface-volume type flaw: a sidewall groove 

and a large grain as shown in Figure 10.21c. The fractographic results 

supported the Weibull area scaling for comparing the straight and curved 

10-26 
 

 

 

 

 

 

Case Studies 

specimens, and the best correlation was obtained if only the etched sidewall 

areas were used in the calculation. 

The reason the Weibull strength scaling did not work for the intentionallynotched specimens became obvious as shown in Figure 10.22. The notches 

were quite small (15μm to 25 μm radius). They concentrated stress to a very 

small region around the notches. Hence, the Weibull effective volumes or 

areas were tiny, and comparable to the size of some of the single grains. 

Hence, one cannot assume there is a well-distributed set of flaws scattered 

throughout the stressed volume, as is assumed in conventional Weibull 

analysis. Furthermore, an analysis based on continuum mechanics assuming 

the material was isotropic and homogeneous is questionable. 

notch notch 

(a) (b) 

(c) (d) 

Figure 10.22 A broken notched miniature SiC tensile specimen. (a) shows the 

whole specimen. (b) shows a close-up of the gage section notches. Transillumination through the green single crystal SiC grains created the light 

patches. They are individual grains. (c) and (d) show the fracture surface. 

Fracture started at the bottom of the round notch (arrows), from a etch groove 

linked with a single long and large columnar SiC grain that ran from the top to 

the bottom. 

10-27 
 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Case 10: Fractious Fractographers (A Fractography Round Robin) 

Can different fractographers agree on fracture origin analysis? Can they reach 

similar conclusions? How do they exercise their craft? 

These were questions addressed in an international round robin conducted in 

the mid-1990s.21 Round robins are interlaboratory exercises using a common 

procedure. There are many reasons for conducting a round robin, but usually 

the goals are to evaluate whether different laboratories can apply a prescribed 

procedure and obtain consistent results, or to generate test method uncertainty 

data such as repeatability or reproducibility precision estimates. 

This author and Dr. J. Swab of the U.S. Army Materials Technology 

Laboratory (now with U. S. Army Research Lab, Aberdeen, MD) conducted a 

major international round robin on characterization of fracture origins of 

advanced ceramics in 1994 - 1995. Seventeen laboratories participated. The 

participants had experience that ranged from zero to thirty-five years. The 

project was coordinated under the auspices of the Versailles Advanced 

Materials and Standards (VAMAS) program. This work evaluated a new 

(1992) set of guidelines for finding and characterizing fracture origins in 

ceramics: “Fractography and Characterization of Fracture Origins in Advanced 

Structural Ceramics,” Military Handbook MIL HDBK 79022 which was a 

predecessor to ASTM Standard Practice C 1322,23 the world’s first 

fractography standard for characterizing fracture origins. 

Topic 1 of the VAMAS round robin was a photo interpretation exercise with 

focus on detection and characterization of grinding damage in bend bars. 

Photos of one specimen each of silicon nitride, a zirconia-alumina composite, 

and an alumina were furnished. Six photos of each specimen were provided 

showing for each half, a low magnification overall shot, a picture of the 

fracture mirror area, and a close-up of the origin area. Results were mixed. 

There was consensus that the origins were machining cracks, but there was 

considerable variability in how they were marked and measured. Some 

participants disagreed and felt the origins were not grinding damage. Although 

a short paragraph on each specimen with strength and fracture toughness 

information was furnished along with the photos, very few of the participants 

used this information to estimate a critical crack size. Participants were also 

asked to mark and measure the fracture mirror sizes on the photos. The 

organizers were surprised at the scatter in mirror size estimates. Evidently 

many participants were not familiar with measuring fracture mirrors. 

10-28 
 

 

 

 

 

 

 

 

 

Case Studies 

Topic 2 was an actual examination of six bend bar strength specimens. Each 

participant received one example of the specimens listed in Table 10.2. Most 

were carefully prepared (no small feat) by Dr. Swab. Again, results varied. 

Most participants had no trouble identifying the pore as the fracture origin in 

specimen #3. Many had difficulty with the others. For example, the scratch 

that caused cracking damage in specimen #2 was sometimes identified as 

machining damage or even a pore or large grain. Figure 6.25c shows one of 

these scratched specimens. The scratch should have been obvious. Dr. Swab 

had made it with a diamond indenter and it stood out very clearly on the tension 

surface of the bend bar. Evidently some participants did not bother looking at 

the external surface and only looked at the fracture surface. 

In the other specimens, participants looked at only one half. Sometimes this 

was satisfactory, but in just as many cases observers looked at the less clear 

piece. 

Many participants evidently did not read the brief four to six sentences of 

information furnished with each specimen, including treatments if any and the 

breaking stress and fracture toughness. For example, the surface pit origin 

specimens #4 had been exposed to an oxidizing environment at high 

temperatures. Some labeled the origin a pore and ignored the surface condition 

and reaction zones around the pits. Only a few participants estimated flaw 

sizes from fracture mechanics and compared them to their fractographicallymeasured sizes for verification. 

Specimen 

type Material Origin Flaw 

1 Hot-pressed alumina with silicon 

carbide whiskers 

Large grain 

2 Sintered 99.9% alumina Handling damage (scratch) 

3 Sintered 3 mol% yttria stabilized 

zirconia (Y-TZP) 

Pore 

4 Reaction bonded silicon carbide Surface Pit 

5 Hot isopressed silicon nitride Machining damage 

6 Sintered titanium diboride Porous seam or porous region 

Table 10.2 Bend bar specimen types distributed in Topic 2 of the VAMAS 

round robin. 

10-29 
 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

The entire exercise was a learning experience for the organizers. Some 

participants did very well, but many had difficulty. Much was learned by 

studying the causes of the misinterpretations. Even for such mundane objects 

as bend bars, there was considerable variability in how the participants 

analyzed the photos, the specimens, and the ancillary information provided. 

Some simply looked at the fracture surfaces and ignored everything else. The 

organizers came to realize that fractographic analysis (and failure analysis in 

general) is a process whereby an expert integrates all information including, but 

not limited to, fracture surface examination. This realization led the author and 

Dr. Swab to prepare Figures 1.1 (the fractographer as detective) and 1.2 (the 

fractographic analysis puzzle) of this Guide. The organizers also came to the 

realization that they had a huge advantage compared to the participants. The 

organizers could look at many specimens of a type, and were not limited to a 

solitary example. The participants only had one. Furthermore, there is no 

substitute for actually looking at the actual specimen under a microscope, as 

opposed to looking at two-dimensional photographs as in Topic 1. Several of 

the author’s fractographic Rules of Thumb listed in the next chapter stem from 

this exercise. 

Case 11: Perilous Prostheses (Four Ceramic Dental Crowns) 

All-ceramic dental crowns are increasingly being used as alternatives to gold or 

porcelain fused-to-metal restorations (Figure 10.23). A variety of ceramic 

materials have been used including feldspathic or leucite porcelains, glass 

ceramics, alumina, and zirconia. These ceramics function as the main load 

bearing and structural elements of the crown, and are called the “core” or the 

Figure 10.23 A fractured Dicorc glass ceramic molar crown. (courtesy of S. 

Scherrer) 

c Dentsply Int., York, PA. 

10-30 
 

 

 

 

 

 

 

 

 

 

 

 

 

Case Studies 

“foundation.” The crowns also have a glassy exterior veneer layer added for 

cosmetic purposes. Brittle, highly-filled resin matrix composites with ceramic 

or glass fillers are common crown materials as well, but do not need a veneer 

layer. The fracture resistance and durability of the new ceramic and composite 

crowns are a primary concern. Despite the sanguine claims of the manufacturers, premature fractures do occur, much to the consternation of the patient 

and his or her dentist. 

Progress in the field has been typified by a trial and error approach. The usual 

approach has been to statistically analyze the rates of failures and infer or 

guess causes of failures. There are conflicting explanations as to why the 

crowns do or do not fail. It is not even clear what properties are desired and 

what are the best laboratory-scale tests to evaluate these properties. Until the 

late 1990s, postmortem failure analyses were quite rare. 

One key reason for this was that dental ceramic crown fractures are in fact 

among the most challenging problems to solve. Many dental ceramics have 

coarse-grained and/or porous microstructures. They often fracture at low stress 

and classic fractographic markings are often masked by the roughness of the 

fracture surface. High-density aluminas and 3Y-TZP zirconias are exceptions 

and are easier to interpret. Another complication is that stress states in the 

crowns are transient and uneven. Fracture mirrors are almost never observed in 

a clinical fracture since fracture stresses are low, or they are larger than the 

crown wall sizes. Another difficulty is that key fragments are often missing, 

either due to loss in the mouth or damage during crown extraction. Crowns 

usually accumulate damage from multiple events at multiple sites, creating 

complex, conflicting fracture networks. For example, the author has often 

detected Hertzian cone crack damage sites on crown occlusal and facial 

surfaces, but has only seen a few instances where the cone acted as an actual 

crown fracture origin. The Hertzian cones crack damage sites usually are 

dormant. Contact damage does occur and cause clinical failures, but it almost 

always in due to edge chipping. 

Fractographic analysis of all-ceramic dental prostheses is a field where recent 

fractographic analyses have been very productive. 24,25,26,27,28,29 Additional 

references are listed in the Appendix B, Case Studies. The following four 

crown fracture case studies illustrate the progress that has been made. 

Figure 10.24 shows a Cerestored alumina-magnesia spinel fractured molar 

crown that was documented by J. Quinn et al., in Ref. 24. In this and the 

d	 Originally developed by Coors Biomedical, Lakewood, Co.. Available later from 

Ceramco Inc., East Windsor, NJ. 

10-31 
 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

following three examples, gold coating has made the veneer look dark in the 

optical images and the core material lighter. The origin was located at the 

margin (or bottom) of the crown and the vertical-split nature of the fracture 

suggests that hoop stresses around the bottom of the crown, where the crown 

Half B 

Half A 

Half A 

Half B origin 

Missing triangular piece 

(a) (b) 

(c) 

Figure 10.24 a-c Fractured alumina-spinel molar crown. (a) is a top occlusal 

surface view showing the two halves held together. Half A is fully- and half B 

is partially-gold coated. (b) shows the interior looking up into the crown. The 

missing triangular piece was caused by secondary breakage. The origin was 

on the margin, at the bottom of the crown. (c) shows an optical photo of the 

fracture surface with arrows showing the local crack propagation direction. 

The dark outer regions are the gold-coated glassy veneer. The lighter gray 

portions are the gold-coated core ceramic material. The letters d, e, f identify 

regions shown in images (d) – (f) on the next page. (images courtesy J. Quinn) 

10-32 
 

 

 

 

 - 

 

Case Studies 

veneer core veneer 

(d) (e) 

(f) 

Margin origin area 

glass 

veneer 

ceramic 

core 

Figure 10.24 d-f (continued). (d) wake hackle from tiny bubbles in the veneer 

glass. This location is marked on the left of view (c) and indicates the crack 

moved right to left (arrow). (e) shows wake hackle and gull wings from 

bubbles in the veneer on the right side of the crown. At this location the crack 

was running upwards from the margin. Careful scrutiny of the adjacent core 

ceramic material showed that it also had wake hackle generated by pores in the 

ceramic pointing in the same direction. (f) is an optical image from the origin 

area at the margin. There is no obvious flaw. 

10-33 
 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

was the thinnest, caused the breakage. A specific material flaw could not be 

identified. This is not unusual with dental crown fractures: an exact site is 

identified, but a specific flaw is not obvious. It is very likely that an initial 

radial crack has popped-in at some point, and then final fracture occurred on 

the same plane. Sometimes the final fracture will be on a slightly different 

plane, and an arrest line from the pop-in will be apparent as evident in the next 

example given below. By the time the crack propagated over to the opposite 

side of the crown, the two halves hinged apart in bending, creating the 

compression curl. 

Figure 10.25 shows a second example, a similar case from Ref. 24, but for a 

Procera 99.9% purity alumina.e It also split vertically, evidently in response to 

a hoop stress. The fractographic markings again lead to a margin origin site. A 

specific single material flaw is not evident. In the first edition of this Guide, 

the author stated that the origin site coincided with the end of the cement bond 

to the core material, on the inside of the crown. The author reevaluated the 

crown and now believes the origin was in fact a barely-discernable semicircular 

crack in the core material at the end of the margin as shown in Figure 10.28f+g. 

The origin crack was on almost exactly the same plane as the final fracture. It 

was not discernable in SEM photos. From its size (179 μm radius), and using 

Compression 

curl 

Corner 

hackle 

Half A Half A 

Compression 

curl 

Margin 

origin 

Occlusal surface 

(a) (b) 

Figure 10.25 Procera alumina molar crown. The crown split vertically into 

two halves. The fracture surfaces have been gold coated. (b) is a close-up of 

the compression curl on the side away from the origin site. The veneer glass is 

dark and the core ceramic is a lighter shade. The arrows show the dcp from: 

corner hackle around the inside corner, the compression curl, and from veneer 

wake hackle. (c) – (g) are on the next page. 

e Nobel Biocare, Stockholm, Sweden. 

10-34 
 

 

 

 

 

Case Studies 

KIc of 4.0 MPa√m for dense fine-grained alumina, and a Y of 1.4, the hoop 

stress at fracture was estimated to be 213 MPa. The margin in this example 

was well-prepared and over 1 mm thick. 

(c) (d) 

veneer core 

Half A origin site Half B origin Half B origin 

(e) (f) (g) 

Figure 10-25 continued. (c) and (d) are examples of wake hackle from pores 

in the veneer. (c) is an optical image, and (d) is an SEM image. (e)-(g) show 

optical images of the margin. The origin is a subtle margin crack as marked 

clearly in (f) and (g). This origin crack was so flat and coplanar with the final 

fracture surface, that it did not stand out well in SEM images. (all images 

courtesy of J. Quinn and S. Scherrer) 

10-35 
 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) (b) 

(c) 

Cantilever 

curl 

Origin 

dcp 

veneer core 

(d) 

As received Cleaned 

and 

magnified 

edge chip 

veneer 

(e) (f) 

Figure 10.26 Procera alumina molar crown that broke from a margin chip. 

(a-f) are optical images without any coatings. (a) is a clinical view with the 

origin marked. (b) is one half of the fractured crown. (c) shows a fanlike array 

of corner hackle (between the two arrows which show the propagation 

direction) in both the veneer and core emanating from the top inside corner. (d) 

is from near the margin on the left side and shows tiny wake hackle from 

veneer pores. The crack ran from bottom to top. (e) and (f) show the origin 

was an edge chip from a force applied on the bottom of the margin aimed 

upwards (image a courtesy of S. Scherrer). 

10-36 
 

 

 

 

 

 

 

 

 

 

 

 

 

Case Studies 

The third crown, also a Procera alumina, is shown in Figures 10.26 and 10.27. 

It also has a margin origin site that led to crack propagation by hoop stresses. 

Images 10.26 a-f are all optical images of the crown without coatings. As 

such, some of the details were “washed out” and did not photograph well due to 

the material’s translucency and light scattering. Nevertheless, this crown 

failure was diagnosed solely on the basis of an optical examination with a 

stereo optical microscope. Subsequent SEM examination produced the sharper 

images shown in Figures 10.27. A full fractographic montage for this crown 

was shown previously as Figure 5.68. 

Corner hackle around the inside corners of the core, a compression curl, and 

wake hackle again led back to a margin initiated fracture origin. A welldefined edge chip on the veneer was the origin. The edge chip was discolored 

suggesting that it had been present for quite some time. The direction of the 

edge chip was puzzling as well. It appeared to have been from a force directed 

upwards on the crown. Such a loading is not likely in the mouth since the 

margin is protected by the gum line and jaw. The chip was probably created 

during fabrication or installation. 

veneer core 

(a) (b) 

Figure 10.27 SEM images for the Figure 10.26 Procera crown that broke 

from a margin chip (arrows). The final fracture that split the crown started at 

the bottom at the margin, from part of the margin chip, and ran as shown by 

the black arrows. (images courtesy of S. Scherrer) 

The fourth crown shown in Figures 10.28 -10.31 is an Empress II lithium 

disilicate glass ceramic incisor.f The crown had been in service only 4 months 

when it split into two pieces. No information was furnished on how the 

fracture event occurred or on the installation or fabrication of the crown. The 

f Ivoclar, Schaan, Liechtenstein. 

10-37 
 

 

 

 

 

Margin

 Fractography of Ceramics and Glasses 

fracture plane was vertical and perpendicular to the dental arch. The two 

halves fitted closely together. A few small chips were missing, probably from 

secondary fracture, and were inconsequential to the analysis. The lingual 

surface (the side that faces the tongue) had some surface damage and also an 

unusually thick veneer. There were several shallow depressions (≈ 1 mm2 

square area) as well as numerous conchoidal chip fractures, small pits and 

gouges at mid crown height. The shallow depressions had tool marks 

indicating they were from adjustments to the lingual surface by the dentist once 

the crown had been cemented in place and opposing tooth contact checked. 

This adjustment was only partially successful in relieving contact from the 

opposing tooth as evidenced by the numerous pits, chips and gouges. This 

location was eventually determined to be near the fracture origin region. 

The facial surface (the surface that faces outwards) was in good condition and 

had little damage or evidence of abuse. The unusually thin veneer had a few 

secondary spall chips and one harmless Hertzian cone crack that did not 

penetrate into the core material. The chips were noted on one fragment but not 

on the other, indicating that they were secondary fractures that occurred after 

the crown had broken. 

Occlusal surface 

Margin 

Facial 

surface 

(a) (b) 

Figure 10.28 Fractured Empress II glass ceramic anterior (#7) incisor crown. 

The insert (a) shows a facial view of the cleaved crown. (b) is a view looking 

up into the interior with the two halves together. The origin is located between 

the two red dots on the inside surface. 

10-38 
 

Case Studies 

(a) 

(b) 

Figure 10.29 The fracture surfaces of the two halves. (a) shows an optical 

photo of the uncoated halves. (b) shows gold-coated surfaces. Two different 

coaters were used causing the different coloration. The gold coating of the 

glassy veneer reflects more. The veneer thickness is very uneven and out of 

normal guidelines. It is very thick on the lingual side and very thin on the 

facial side. The arrows mark the origin site located at the thinnest part of the 

core glass ceramic material. 

10-39 
 

 

 

 

 Fractography of Ceramics and Glasses 

(a) 

(b) 

core veneer 

Figure 10.30 Details of the incisor crown fracture. (a) is an SEM image wake 

hackle in the glassy veneer from a region near the top (occlusal) surface of the 

A half. The wake hackle shows that the crack was running in the direction of 

the arrow in this region. All the wake hackle in the entire crown showed 

fracture ran from the inside to the outside. (b) is an optical image of the origin 

which originated from a contact crack on the inside of the crown at the thinnest 

point of the core. 

10-40 
 

Case Studies 

(a) 

(b) 

glassy veneer ceramic 

core 

Figure 10.31 SEM images of the origin. Notice all the wake hackle in the 

veneer aiming away from the interior. The crack ran from the inside to the 

outside. The inside surface is uneven and there are mold or tool marks in the 

vicinity. This region is directly opposite a region of extensive localized contact 

damage on the outer veneer surface. 

10-41 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

There was negligible damage on the occlusal (top) surface. Some residual 

porous powdery material was detected on the inside surface of the crown and 

was the remnants of the cement. Tool grinding marks and/or casting 

impression marks were detected on the inside of the crown. 

Green dye staining for an optical examination and subsequent SEM 

examination brought out a number of telltale brittle material fracture features. 

Fractographic markings in the glass ceramic core material were difficult to 

detect, but subtle hackle markings were similar to features observed on bend 

strength test specimens of the same material. Extensive wake hackle was 

detected from bubbles in portions of the glassy veneer as shown in Figure 

10.30a and 10.31,a,b. Seven or eight crack arrest lines were also detected. 

This was a difficult analysis. Initial examinations had only limited success, but 

with repeated careful scrutiny of the fragments using different examination 

techniques (optical, optical green dye coated, optical gold coated, and SEM) 

enabled the author and J. Quinn to find the origin region. Maps of crack 

propagation direction were constructed, primarily through observations of the 

wake hackle markings in the veneer. Fracture initiated in the core material 

(where it was very thin) from an internal surface Hertzian cone crack. This 

location was in the same vicinity where external lingual surface damage has 

also been detected. Fracture then radiated upwards and downwards in response 

to hoop stresses causing the crown to split into two halves. 

Only after repeated examination was it realized that both fracture halves 

showed matching damage features characteristic of blunt contact damage at the 

exact same location on the inside (cementatious) surface of the crown. The 

critical observation occurred when stereo SEM paired images were viewed of 

this area, and the arc-shaped crack that is typical of blunt contact damage was 

revealed. It was further observed that the inside crown surface was uneven and 

had raised hills, ridges, and depressions. 

In summary, the Empress II incisor crown had uneven veneer and core wall 

thickness. Fracture initiated from contact damage on the inside of the crown, 

probably due to an uneven fit or improper placement. Opposing tooth contact 

in the same vicinity on the crown exterior surface (as evidenced by shallow pit 

and chip damage) created the stresses in the origin area. Surface grinding 

adjustments made to the same area suggested a fitting problem. 

These four crowns broke from hoop stresses and split vertically, perpendicular 

to the dental arch. Specific gross material flaws were not be detected in any of 

the four. 

10-42 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

Case Studies 

Case 12: Case of the Mangled Margins (Twenty dental crowns) 

Recent work that typifies progress in the dental fractography field is presented 

here. The failures shown above in Case 11 were isolated cases. Most failure 

analyses have been on such isolated examples or on batches of laboratorytested (in vitro) crowns, which may or may not be accurate representations of in 

vivo failures. Indeed, many “crunch the crown” tests with loading by steel or 

tungsten carbide balls on the occlusal surfaces are ridiculous and clinically 

irrelevant. 

The cases presented here are the first study to identify a single leading cause of 

fracture, with high specificity, for a large number of cases in one particular 

restoration system. Prof. M. Oilo at the University of Bergen, Norway has 

collected dozens of clinical fractures and is organizing them by type. As 

discussed in Chapter 4, she has also devised a laboratory set-up that causes 

crowns to fracture into the same patterns as clinical fractures.30,31 The author 

and she identified the fracture origins in twenty-two crowns of a set of twentyseven retrieved crowns, all from one material system.32 The crowns were all 

high-alumina crowns prepared by CAD/CAM machining. They fractured in 

vivo from times as short as days to years. Twenty-one of the twenty-two 

crowns (95%) broke from one cause: hoop stresses on the margins that split the 

crowns, usually at the area of shortest wall length in the approximal area, with 

the fracture origins being irregularities in the margin including chips, grinding 

cracks and veneer irregularities. Figures 10.32 - 10.34, show several of the 

fractured crowns. 

One common element in all these crowns was irregular grinding damage on 

feather margins that were as thin as 100 μm (0.1 mm). Compare these margins 

to the nice, thicker ones in Figures 10.24 and 10.25. It is unwise to make thin 

ceramic margins since they are vulnerable to all sorts of machining and 

handling damage and chipping. 

Dental material fractography is making progress. We now can analyze clinical 

fractures by the dozens. Patterns are emerging. Analysis in this alumina crown 

project was still time consuming, but it went quicker the more that we did. 

Over two thousand images were taken and organized and assembled into one 

gigantic PowerPoint file, which served as our record-keeping and organization 

tool. Of note is that the origin area was identified in all crowns by light 

microscopy. We learned that gold coating the specimens dramatically speeded 

up the process. Green dye worked, but gold coating was better. It was then not 

too difficult to get the appropriate piece and area in the SEM and focus on the 

origin itself. 

10-43 
 

 

 

 Fractography of Ceramics and Glasses 

(a) 

(b) 

(c) 

veneer 

core 

Figure 10.32 Alumina central incisor that broke in vivo after 1 year. Fracture 

started at the margin from cracking in both the veneer and core. The margin is 

too thin, and even worse, the core portion is exceptionally thin. The origin is 

chipping damage at the margin. (clinical case 2) 

(a) 

(b) 

(c) veneer core 

Figure 10.33 Alumina incisor that broke in vivo at an unreported time. 

Fracture started at the margin from cracks and chipping in both the veneer and 

core. The crown split in half from the mesial to distal side. The margin is too 

thin and the core portion is exceptionally thin. (clinical case 8). 

10-44 
 

 

 

 

 

 

 

Case Studies 

(a) (b) veneer core 

Figure 10.34 Alumina molar #37 that broke in vivo at 7 months in vivo. It 

split in half from the mesial to distal side. Fracture started at the margin from 

cracks and chipping in the veneer which extends to the very tip and even onto 

the inside surface as evident in (a) Notice the tiny wake hackle lines in the 

veneer in (a) that show the origin was at the margin. (a) also shows chipping 

elsewhere on the margin. (clinical case 9). 

(c) core veneer (a) 

(b) 

Figure 10.35 Alumina molar that broke in vivo at 5.5 year in vivo. It split into 

two pieces, but additional fractures were made during extraction. The initial 

split was from mesial to distal side. The origin is at the margin of piece C. 

The SEM image in (c) shows excess veneer that has come into the inside 

surface and tiny wake hackle lines from the veneer bubbles(marked by the 

arrow showing the dcp, and the exceedingly small core thickness at the margin 

tip. Ceramic margins should never be made this thin. (clinical case 12). 

10-45 
 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Case 13 Case of the Ruined Refractors (Sapphire Lenses) 

This case is of two large elaborate components made of single crystal sapphire. 

Sapphire is a harsh mistress. She has alluring physical and mechanical 

properties such as good transmission of optical and infrared light, strengths of 

1000 MPa or more, and good high temperature properties. She also has 

vulnerabilities, such as high variability in strength, sensitivity to contact and 

machining damage, and a shocking weakness at 300 o

C - 500 o C due to 

compression-induced twinning. 

NASA is laying the groundwork for spacecraft of the future that could harness 

solar energy. One concept is to use a large primary mirror solar collector that 

focuses light to a secondary lens collector. The collector focuses and 

concentrates energy by total internal reflection into a heat exchanger.33 Two of 

these full-scale (30 cm long by 9.5 mm diameter) c-axis sapphire lenses were 

tested in the giant solar thermal vacuum facility that simulates outer space at 

the Glenn Research Center. They fractured during high-temperature trials. 

This fractographic analysis was done by Dr. J. Salem and myself.33 

Figure 10.36 shows one of the fractured lenses, which has changed to a rosebrown color from solarization at the 1300o C peak temperature of the test. The 

fracture surface had a mix of cleavage and conchoidal fracture and some crack 

bifurcations. The fracture origin was on the lens face. The origin shown in 

Figure 10.37 was a large, partial Hertzian cone crack that extended from one of 

two partially polished impact divots or “bruises” on the polished surface. From 

the size of the ring crack, it was deduced that the impact object had a diameter 

greater than 0.5 mm. Thermal stresses subsequently caused the lens to fracture. 

This structure also had twins, but they were incidental and not a cause of 

fracture. Figure 3.59b in the chapter on equipment shows one twin in a 

polariscope. 

Figure 10.38 shows the origin in the second broken refractor. It also broke 

from an origin on the curved incident surface, again from thermal stresses, but 

the origin was a scratch. Unusual “micro twins” were found near the origin 

site, but they were not associated with the fracture. 

Both of these refractors, and some of the earlier IR domes shown in this Guide 

typify the vulnerability of sapphire to residual grinding cracks, polishing and 

handling damage. The thermal gradients created both tensile and compression 

stresses, and it is very likely that with additional high-temperature exposure and 

cycling, more twins would have formed with the possibility they would have 

intersected at some point leading to crack formation and eventually fracture. 

10-46 
 

 

 

Case Studies 

(a) 

(b) 

Figure 10.36 Fractured sapphire solar light secondary concentrator which 

broke into four pieces. The origin was on the rounded surface where incoming 

light is focused into the body. The blue insert in (a) shows an intact lens prior 

to testing. It is a complex but beautiful piece. 

10-47 
 

 

 

 

 

 

 

 

c

 

 

 Fractography of Ceramics and Glasses 

(a) (b) 

Figure 10.37 The origin of the first refractor was a blunt contact damage 

crack shown on the fracture surface (a) and on the lens face in (b). 

(c) (d) 

(b) 

origin 

scratch 

(a) polished 

surface 

polished 

surface 

outer rim 

bevel 

fracture plane 

Fracture 

origin 

micro 

twin 

scratch polished 

surface 

polished 

surface 

Figure 10.38 The second refractor lens also broke from damage on the 

rounded surface, but from a scratch. (a) shows the fracture surface. (b) shows 

the scratch on the outer polished surface. It curves at the origin site. With 

careful lighting in (c) and (d), “micro-twins” near the origin site are revealed. 

10-48 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Case Studies 

Chapter 10 References 

1.	 R. R. Baker, L. R. Swank, and J. C. Caverly, “Ceramic Life Prediction 

Methodology - Hot Spin Disk Life Program,” AMMRC TR 83-44, Army Materials 

and Mechanics Research Center, Watertown, MA, August, 1983. 

2.	 L. R. Swank, R. R. Baker, and E. M. Lenoe, “Ceramic Life Prediction 

Methodology,” pp. 323 - 330 in Proc. of the 21st Automotive Technology 

Development Contractors Coordination Meeting, SAE, Warrendale, PA, 1984. 

3.	 G. D. Quinn, K. Xu, R. J. Gettings, J. A. Salem, and J. J. Swab, “Does Anyone 

Know the Real Fracture Toughness? SRM 2100: The World’s First Ceramic 

Fracture Toughness Reference Material,” pp. 76 - 93 in Fracture Resistance 

Testing of Monolithic and Composite Brittle Materials, ASTM STP 1409, eds., J. 

A. Salem, G. D. Quinn, and M. G. Jenkins, ASTM Int., West Conshohocken, PA, 

2002. 

4.	 G. D. Quinn, “Fracture Mechanism Maps for Advanced Structural Ceramics, Part I, 

Methodology and Hot-Pressed Silicon Nitride,” J. Mat. Sci, 25 (1990) 4361 - 4376. 

5.	 K. D. Mörgenthaler and E. Tiefenbacker, “Bauteile Aus Keramik Für 

Gasturbinen,” (Ceramic Components for Gas Turbines), Final Report of Phase III, 

Daimler-Benz AG, Technical Report, Stuttgart, 1983. 

6	 E. J. Bunning, D. R. Claxton, and R. A. Giles, “Liners for Gun Tubes – A 

Feasibility Study,” Ceram. Eng. and Sci. Proc., 2 [7-8] (1981) 509 -519. 

7.	 G. D. Quinn, “Characterization of Turbine Ceramics After Long Term Environmental 

Exposure,” AMMRC TR 80-15, Army Materials and Mechanics Research Center, 

Watertown, MA, May, l980. 

8.	 G. D. Quinn, “Review of Static Fatigue in Silicon Nitride and Silicon Carbide,” 

Ceram. Eng. and Sci. Proc., 3 [1-2] (1982) 77 - 98. 

9.	 G. D. Quinn and R. N. Katz, “Time Dependent High Temperature Strength of 

Sintered Alpha SiC,” J. Amer. Ceram. Soc., 63 [1-2] (1980) 117 - 119. 

10.	 J. Coppola, M. Srinivasan, K, Faber, and R. Smoak, “High Temperature Properties 

of Sintered Alpha Silicon Carbide,” presented at the International Symposium on 

Factors in Densification and Sintering of Oxide and Non-Oxide Ceramics, 

Hakone, Japan, October, 1978. 

11.	 V. D. Fréchette, Failure Analysis of Brittle Materials, Advances in Ceramics, Vol. 

28, American Ceramic Society, Westerville, OH, 1990. 

12.	 J. J. Swab, R. Pavlacka, G. Gilde, S. Kilczewski, J. Wright and D. Harris,” 

“Determining the Strength of Coarse-Grained ALON and Spinel,” J. Amer. 

Ceram. Soc., 97 [2] (2014) 592-600. 

13.	 MIL STD 1942 (MR), “Flexural Strength of High Performance Ceramics at Ambient 

Temperature,” U. S. Army Materials Technology Laboratory, Watertown, MA, 21 

Nov., 1983. 

14	 ASTM C 1161-90, “Standard Test Method for Flexural Strength of Advanced 

Ceramics at Ambient Temperature,” ASTM Int., West Conshohocken, PA, 1990. 

15.	 ISO 14704, “Fine Ceramics (Advanced Ceramics, Advanced Technical Ceramics) – 

Test Method for Flexural Strength of Monolithic Ceramics at Room Temperature,” 

Int. Org. for Standards, Geneva, SW, 2000. 

10-49 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

16.	 G. D. Quinn, “Fractography and the Army Flexure Standard,” pp. 319 - 334 in 

Fractography of Glasses and Ceramics, eds., J. Varner and V. Fréchette, American 

Ceramic Society, Westerville Ohio, (1988). 

17.	 G. D. Quinn, “Flexure Strength of Advanced Structural Ceramics: A Round Robin,” 

J. Am. Ceram. Soc., 73 [8] (1990) 2374 - 2384. 

18.	 W. N. Sharpe, Jr., O. Jadaan, N. N. Nemeth, G. M. Beheim, and G. D. Quinn, 

“Fracture Strength of Silicon Carbide Microspecimens,” J. Microelectromechanical Systems, 14 [5] (2005) 903 - 913. 

19.	 G. D. Quinn, W. Sharpe, Jr., G. Beheim, N. Nemeth, and O. Jadaan, “Fracture 

Origins in Miniature Silicon Carbide Structures,” pp. 62 - 69 in Fractography of 

Advanced Ceramics, II, eds., J. Dusza, R. Danzer, and R. Morrell, TransTech 

Publ., Zurich, 2005. 

20.	 G. D. Quinn, “Design and Reliability of Ceramics, Do Modelers, Designers and 

Fractographers See the Same World?” Ceram. Eng. and Sci. Proc., 26 [2] (2005) 

239 - 252. 

21.	 J. J. Swab and G. D. Quinn, “The VAMAS Fractography Round Robin: A Piece of 

the Fractography Puzzle,” pp. 55 -70 in Fractography of Glasses and Ceramics III, 

Ceramic Transactions, Vol. 64, eds., J. R. Varner, V. D. Fréchette, and G. D. Quinn, 

ACS, Westerville, OH, 1996. 

22.	 MIL HBK 790, “Fractography and Characterization of Fracture Origins in Advanced 

Structural Ceramics,” U. S. Army Materials Technology Laboratory, Watertown, MA, 

June, 1992. 

23.	 ASTM C 1322-96, “Standard Practice for Fractography and Characterization of 

Fracture Origins in Advanced Ceramics,” Annual Book of Standards, Vol. 15.01, 

ASTM Int., West Conshohocken, PA. 

24.	 J. B. Quinn, G. D. Quinn, J. R. Kelly, and S. S. Scherrer, “Fractographic Analyses 

of Three Ceramic Whole Crown Restoration Failures,” Dent. Mater., 21 (2005) 

920 - 929. 

25.	 S. S. Scherrer, J. B. Quinn, G. D. Quinn, and J. R. Kelly, “Failure Analysis of 

Ceramic Clinical Cases Using Qualitative Fractography,” Int. J. Prosthodont., 19 

[2] (2006) 151 - 158. 

26.	 J. B. Quinn, G. D. Quinn, J. R. Kelly, and S. S. Scherrer, “Useful Tools for Dental 

Failure Investigation,” pp. 36 - 53 in Proc. Conf. on Scientific Insights and Dental 

Ceramics and Photopolymer Networks, ed. D. C. Starrett, Trans. Academy of 

Dental Materials, Vol. 16, 2004. 

27.	 J. R. Kelly, “Fractography of Dental Ceramics,” pp. 241 – 251 in Fractography of 

Glasses and Ceramics, IV, eds., J. R. Varner and G. D. Quinn, American Ceramic 

Society, Westerville, OH, 2001. 

28.	 J. B. Quinn, “Failure Analysis of a Broken Tooth,” J. Fail. Anal. Prev., 4 [1] 

(2004) 41 - 46. 

29.	 J. Y. Thompson, K. J. Anusavice, A. Naman, and H. F. Morris,” Fracture Surface 

Characterization of Clinically Failed All-Ceramic Crowns,” J. Dent. Res., 73 [12] 

(1994) 1824 – 1832. 

30.	 M. Øilo, A. K. Kvam, J. E. Tibballs, and N. Gjerdet, “Clinically Relevant Fracture 

Testing of All-Ceramic Crowns,” Dent. Mater., 29 (2013) 815-823. 

10-50 
 

 

 

 

 

 

 

 

Case Studies 

31.	 M. Øilo, N. R. Gjerdet, “Fractographic Analysis of All-Ceramic Crowns: A Study 

of 27 Clinically-Fractured Crowns,” Dent. Mater., 29 (2013) e78-e84 

32.	 M. Øilo and G. D. Quinn, “Fracture Origins in Twenty-two Dental Alumina 

Crowns,” J. Mech. Beh. Biomed. Mater., 53 (2016) 93 – 103. 

33.	 J. A. Salem and G. D. Quinn, “Fractographic Analysis of Large Single Crystal 

Sapphire Refractive Secondary Concentrators,” J. Eur. Cer. Soc.,34 (2014) 3271-

3281. 

10-51 

 

 

 

 

 

 

 

 

 

 

Conclusions 

11. Conclusions
	

Figure 11.1 

The fractographer as a 

detective. 

This Guide is intended to introduce engineers and 

scientists to the science of fractographic analysis of 

brittle materials as a tool to solve fracture 

problems. This tool may be applied to a broad 

range of practical problems including component 

failure analyses, materials processing refinements, 

routine materials characterization, laboratory scale 

mechanical testing, and reliability and design. 

Fractography is the examination of fracture 

surfaces, the examination of the general crack 

patterns, other part surfaces, other specimens, and a 

review of processing, exposure, and testing 

conditions. Much like the master detective, the 

fractographer starts with an open mind, integrates 

and analyzes the available information and relates 

it to similar episodes from his own experience and 

to what is published in the literature. The 

fractographer applies deductive logic to arrive at a 

conclusion. 

We are in the midst of microscopy revolution, as the marriage of computers 

with digital cameras and glass lenses expands the fractographer’s tool box 

dramatically. Quantitative, virtual, three-dimensional imaging is already 

available. Nonetheless, no matter how powerful these tools become, they will 

never supplant the fractographer as an expert system. No computer or artificial 

intelligence is likely to ever replace the skill of a fractographer who knows how 

to tilt a specimen or light source under a stereoptical microscope and recognize 

a fleeting reflection that unlocks a fractographic puzzle. No computer will 

know how to improvise how to look at something differently or how to stain a 

specimen on a hunch. 

My fractographic experiences over forty-seven years have been distilled into 

the following rules of thumb for fractography: 

1. The first one is the hardest. 

The more examples of a particular fracture problem that are examined the 

better. This is true for both experienced and novice fractographers. Patterns 

may be recognized that are not apparent in a single example. An unfamiliar or 

subtle feature may be overlooked in one example. When the fractographer sees 

a curious or unfamiliar marking in multiple examples, he is less likely to 

dismiss it as a random artifact. One of the most gratifying experiences a 

11-1 
 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

fractographer has is when he or she first recognizes and understands a new 

marking and adds it to his or her experience base. 

2. The more experience the fractographer has the better. 

Fractography is a cumulative learning experience. The more one learns, the 

easier and faster fractography becomes. The fractographer should look at as 

many different test specimens, components, and materials as possible. 

3. Photos almost never show all there is to see. 

Photos are two-dimensional representations of a three-dimensional world and 

perspective is inevitably lost. 

4. Good equipment makes the job easier or even possible, but is no 

substitute for skill and experience. 

One must know where to look and great pictures of the wrong things are 

worthless. Many published close-up SEM images fit this category. 

5. Good fractography under laboratory testing situations begins even 

before the specimen is broken. 

The specimen should be tested in clean conditions to minimize contamination 

and secondary fractures of the fracture surfaces. 

6. Fractography is more than looking at fracture surfaces. 

7. A good fractographer is patient. 

Don’t be afraid to reexamine the pieces. Look at all the pieces, even the ones 

that may not seem to be relevant. Sometimes, in a group of laboratory-tested 

specimens, a key feature is discerned or recognized on the very last specimen. 

Fractographers in a rush often overlook important details. 

8. It’s better to let the material tell you why it failed than the other way 

around. Theoretical modelers cause the most trouble. They already know why 

something should fail since it is in their computer model. Indentation 

mechanics people come next. They have more interest in indentations than in 

the real flaws in a material. 

9. A good fractographer does not overreach, guess, or force an 

interpretation. Don’t let other people put words in your month. 

The good fractographer is not afraid to say: “I don’t know.” It is often a good 

idea to say up front: “I’ll take a look, but I make no promises.” The credibility 

of the fractographer is a precious commodity. 

11-2 
 

 

 

 

 

Conclusions 

10. Fractography may seem subjective to the uninitiated, but it is very 

objective. 

One last metaphor to summarize the fractographic craft is offered. When a 

fractographer first sees broken fragments, he or she sees pieces that are trying 

to tell a tale. With the tools of the trade and with the fractographic skills 

summarized in this Guide, a fractographer can listen to the pieces, interpret 

their story, and unlock the secrets of the fracture. 

Finally, the astute reader will have noted that fracture surfaces can be objects of 

beauty. Figure 11.2 is one such example. This figure, from an unpublished 

study by Dr. J. Quinn of dental porcelains with varying amounts of crystallization and leucite content, illustrates many of the features described in this 

book. 

Figure 11.2 Fracture surface of a dental glass-porcelain, flexural strength 

bar. One can see an origin that is a row of grinding cracks that interacted with 

a near-surface bubble. One also sees the fracture mirror, Wallner lines, gull 

wings, wake hackle, mist, and velocity hackle. (unpublished, found in Dr. J. 

Quinn’s computer after her death.) 

11-3 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Appendix A Bibliography 

APPENDIX A BIBLIOGRAPHY
	

Books on Glass and Ceramic Fractography 

V. D. Fréchette, Failure Analysis of Brittle Materials, Advances in Ceramics, Vol.28, 

American Ceramic Society, Westerville, OH, 1990. 

A must for the serious fractographer. This book covers all aspects of the 

fractography of glasses including fundamental markings on crack surfaces (Wallner 

lines, hackle, and so forth), crack forking, failure origins, and estimates of stress at 

fracture and fractographic techniques. Superbly illustrated with a number of service 

failures and amazing case histories. 

Fractography of Glass, Ed. R. C. Bradt and R. E. Tressler, Plenum, NY, 1994. 

Eight technical articles on glass fractography. 

Fractography, D. Hull, Cambridge University Press,1999. 

A superbly illustrated and laid out textbook on the fractography of metals, ceramics, 

and polymers. Strong on the physics underlying the fracture markings, but weak on 

practical applications. 

Fractography, Vol. 12 ASM Handbook, ASM Int., Materials Park, OH, 1987. 

The seminal work for metals fractography. Expensive. Includes a large atlas of 

metal failures, but only a few examples for ceramics or glasses. Includes excellent 

chapters on the history of fractography, metal models of fracture, photographic 

equipment and techniques, optical and electron microscopy that are directly relevant 

to ceramics and glasses. Caution: Some cleaning and preparation techniques such 

as surface coatings, replicating tapes, replicating tape stripping, and aggressive 

detergent cleaning prescribed for metals are not recommended for ceramic fracture 

surfaces. 

Failure Analysis and Prevention, Vol. 11, ASM Handbook, ASM Int., Materials Park, 

OH, 2002. 

Includes sections on fractography, and fractographic equipment and techniques, but 

deals more with the broader issues of failure analysis: why a component fails to 

perform its intended function. Fracture is but one of many failure modes. Includes 

sections on failure analysis philosophies and approaches. Has one brief chapter on 

ceramics. The fracture mechanics chapter is tailored towards metals fracture and not 

ceramic or glass fracture. The quantitate fractography chapter by Gokahle is 

excellent and has good information about quantitative microcopy and fractals and 

their limitations. 

A-1 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Fractography and Failure Mechanisms of Polymers and Composites, ed. A. RoulinMoloney, Elsevier, 1989. 

This is by far the best book on fractography of polymeric materials and composites. 

There are chapters on microscopy, analytical techniques, the materials including 

unfilled and filled composites, and failure mechanisms. Some of the more brittle 

composites and polymers have similar fracture markings to those on glasses and 

ceramics, and this book is a very complementary tool to this Guide. Out of print, this 

is an excellent book on fractography of polymers and polymeric composites that has 

value to dental restoration researchers. 

Conference Series Books on Fractography of Glasses and Ceramics 

Fractography of Glasses and Ceramics, Advances in Ceramics, Vol. 22, J. Varner, and 

V. Fréchette, eds., American Ceramic Society, Westerville, OH, 1988. 

Proceedings of the first of a quadrennial-quinquennial conference series by the same 

name held at Alfred University in 1986. Sections on: fundamental phenomena, hightemperature fracture, fractography and fracture mechanics, fractography in materials 

development and testing, and component failures. 

Fractography of Glasses and Ceramics, II, Ceramic Transactions, Vol. 17, V. 

Fréchette, and J. Varner, eds., American Ceramic Society, Westerville, OH, 1991. 

Twenty-six papers from the second Alfred University conference, held on campus in 

1990. 

Fractography of Glasses and Ceramics, III, Ceramic Transactions, Vol. 64, J. Varner, 

V. Fréchette, and G. D. Quinn, eds., American Ceramic Society, Westerville, OH, 

1996. 

Twenty-five papers from the third Alfred University conference, held on campus in 

1995. 

Fractography of Glasses and Ceramics, IV, Ceramic Transactions, Vol. 122, J. Varner, 

and G. D. Quinn, eds., American Ceramic Society, Westerville, OH, 2001. 

Thirty-two papers from the fourth Alfred University conference, held on campus in 

2000. Special session on edge chipping. 

Fractography of Glasses and Ceramics, V, Ceramic Transactions, Vol. 199, J. Varner, 

G. D. Quinn, and M. Wightman, eds., American Ceramic Society, Westerville, OH, 

2007. 

Thirty-six papers from the fifth Alfred University conference, held in 2006 but this 

time in Rochester, NY. Papers by Quinn on fracture mirror size measurements; 

Wilantewicz and Varner on indentation damage evolution under Vickers indentations 

in glass and the importance of the ratio of G/K; by Hecht-Mijic and Richter on 

fractography of hip joint balls; Morrell on component fractures including SiC seal 

rings, CaF2 windows, and alumina hip balls; Cleary and Nichols on automotive 

windows; and Engelder about controversies in interpreting mesoscopic cracks in rock 

fractures in the earth. 

A-2 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Appendix A Bibliography 

Fractography of Glasses and Ceramics, VI, Ceramic Transactions, Vol. 230, J. Varner 

and M. Wightman, eds., American Ceramic Society and Wiley, Westerville, OH, 2012. 

Twenty-five papers from the sixth Alfred university conference, held in 2011 this 

time in Jacksonville, FL. Notable papers by Quinn on the history of fractography; 

Glaesemann et al., on fractography in the development of ion-exchanged cover 

glasses (e.g., Gorilla glass); Gross on crack evolution in alkali aluminosilicate glasses 

for cover glass; Bradt on an energy criterion for mirror formation; Hecht-Mijic on 

dental implants; Maurer on a glass pharmaceutical syringe fracture. 

Fractography of Advanced Ceramics, ed., J. Dusza, Trans Tech Publ., Zurich, 2002. 

First (2001) of a new conference series in Stará Lesná, Slovakia. Thirty-six papers, 

many of which are on mechanical testing and have no fractography. Excellent papers 

including a bioceramics fracture paper by H. Richter and a paper by R. Morrell and J. 

Kübler on the background of the then new CEN prEN 843-6 fractography standard. 

Fractography of Advanced Ceramics II, eds. J. Dusza, R. Danzer, and R. Morrell, 

Trans Tech Publ. Zurich, 2005. 

Second (2004) of the Stará Lesná, Slovakia conference series. Sixty one short 

papers. Papers by Quinn et al. on machining damage detection and characterization, 

Wang et al. on ZnO varistor failures, Supanic et al., on fractures in stacked PZT 

piezoceramic actuators, and Tatami et al. on nanofractography of alumina by 

scanning probe microscopy. Many of the papers have negligible fractography, 

however, and deal more with mechanical properties, processing, and indentation. 

Fractography of Advanced Ceramics III, eds. J. Dusza, R. Danzer, R. Morrell, and G. 

D. Quinn, Trans Tech Publ. Zurich, 2009. 

Third (2008) of the Stará Lesná, Slovakia conference series. Sixty-three short 

papers. Papers by Quinn on the history of fractography of brittle materials; Bradt on 

some legal aspects and challenges; Varner on use of replicas. (The fourth conference 

in this series was held in 2013, in Smolenice castle outside Bratislava in Slovakia. 

Dozens of papers were presented, but only 23 were included in a special edition of 

the Journal of the European Ceramic Society in November, 2014.) 

Fractography of Ceramic and Metal Failures, J. Mecholsky, Jr., and S. Powell, Jr., 

eds., ASTM STP 827, ASTM, Philadelphia, PA, 1984. 

Eight papers on ceramics from a 1982 symposium at ASTM headquarters in 

Philadelphia. Includes an outstanding, comprehensive, and well-illustrated review 

paper by Rice, and papers by Pantano and Kelso, and Healy and Mecholsky (cited 

below). 

Fracture Mechanics of Ceramics, Vol. 1, R. Bradt, D. Hasselman, and F. Lange, eds., 

Plenum Press, NY, 1974. 

Proceedings of a conference at Pennsylvania State University in 1973 with twentythree papers on fracture mechanics applied to origin detection and fractography in 

ceramics. The later volumes of this quadrennial series also have some fractography 

papers. 

A-3 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Microscopy Techniques 

C. G. Pantano, and J. F. Kelso, “Chemical Analysis of Fracture Surfaces,” 

Fractography of Ceramic and Metal Failures, ASTM STP 827, ASTM, 1984, pp. 139– 

156. 

The applicability of various instrumental techniques for chemical analysis of fracture 

surfaces is reviewed. The relative merits and spatial and depth resolutions of Auger 

microscopy and energy or wavelength dispersive electron microscopy are given. 

J. T. Healy, and J. J. Mecholsky, Jr., “Scanning Electron Microscopy Techniques and 

Their Application to Failure Analysis of Brittle Materials,” Fractography of Ceramic 

and Metal Failures, ASTM STP 827, ASTM, 1984, pp. 157–181. 

Discusses cleaning, coating, and other procedures for SEM specimens. The merits of 

secondary and backscattered electron imaging are presented. 

C. R. Brooks and B. L. McGill, “The Application of Scanning Electron Microscopy to 

Fractography,” Mater. Char., 33 (1994) 195-243. 

An excellent, well-illustrated review of the application of scanning electron 

microscopy for topographical and chemical analysis of fracture surfaces of ceramics, 

metals, and polymers. Includes a good discussion of stereo SEM fractography. 

Fractography of Ceramics and Glasses - Overview Papers 

J. J. Mecholsky, Jr., and S. W. Freiman, “Determination of Fracture Mechanics 

Parameters Through Fractographic Analysis of Ceramics,” pp. 136 – 150 in Fracture 

Mechanics of Ceramics Applied to Brittle Materials, S. Freiman, ed., ASTM STP 678, 

ASTM, 1979. 

A short but useful overview of the utility of fractography as a quantitative tool to 

determine strength-limiting origins, the stress at failure, and critical fracture 

toughness. 

R. W. Rice, “Fractographic Identification of Strength Controlling Flaws and 

Microstructure,” pp. 323 – 345 in Fracture Mechanics of Ceramics, Vol. 1, R. Bradt, 

D. Hasselman, and F. Lange, eds., Plenum Press, NY, 1974. 

A short but valuable discussion of several key origins (pores, pore groups, and large 

grains) and their relationship to fracture energy. The fracture energy can either be a 

single-crystal or polycrystalline value depending upon the relative sizes of origin and 

microstructure. 

G. D. Quinn, J. J. Swab, and M. J. Slavin, “A Proposed Standard Practice for 

Fractographic Analysis of Monolithic Advanced Ceramics,” MTL TR 90-57, 

November 1990, NTIS Access No. ADA-231989. 

Discusses MIL HDBK 790 the predecessor of ASTM C 1322 in 1996. Discusses 

essential background information and the rationale for consistency in characterization. A standard nomenclature and origin characterization scheme were proposed. 

A-4 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Appendix A Bibliography 

R. W. Rice, “Topography of Ceramics,” pp. 439 – 472 in Surfaces and Interfaces of 

Glass and Ceramics, V. Fréchette, W. LaCourse, and V. Burdick, eds., Plenum Press, 

NY, 1974 . 

A very helpful introduction describes the role of unaided eye, hand lens, optical, 

scanning, and transmission electron microscopy. Fracture surface features such as 

transgranular and intergranular fracture, crack microstructure interactions, crack 

branching, mirrors, and single crystal fractography are discussed. 

R. W. Rice, “Ceramic Fracture Features, Observations, Mechanism and Uses,” 

Fractography of Ceramic and Metal Failures, ASTM STP 827, ASTM, 1984, pp. 5– 

103. 

An extraordinary book-length article with many useful tips and analyses. Many 

single crystal examples. 

T. A. Michalske, “Quantitative Fracture Surface Analysis,” pp. 653 – 662 in Ceramics 

and Glasses, Engineered Materials Handbook, Vol. 4, ed. S. Schneider, ASM, Metals 

Park, OH, 1991. 

J. R. Varner, “Descriptive Fractography,” pp. 635 - 644 in Ceramics and Glasses, 

Engineered Materials Handbook, Vol. 4, ed. S. Schneider, ASM, Metals Park, OH, 

1991. 

J. S. Wasylyk, “Special Terminology Used in Fractography,” pp. 632 – 634 in 

Ceramics and Glasses, Engineered Materials Handbook, Vol. 4, ed. S. Schneider, 

ASM, Metals Park, OH, 1991. 

R. C. Bradt, “The Fractography and Crack Patterns of Broken Glass,” J. Fail. Anal. 

Prev., 11 (2011) 79-96. 

R. J. Parrington, “Fractography of Metals and Plastics,” J. Fail. Anal. Prev., 2 [5] 

(2002) 16-19, 44-46. 

A short article with some interesting history of fractography. Metals and plastic 

fracture modes and patterns are compared. 

Fractography Round Robins 

J. J. Swab and G. D. Quinn, “Fractography of Advanced Structural Ceramics, Results from 

the VAMAS Round Robin Exercise,” U.S. Army Technical Report, ARL-TR-656, U.S. 

Army Research Laboratory, Aberdeen, MD, Dec. 1994. 

An international Versailles Advanced Materials and Standards (VAMAS) Round Robin 

for identification of fracture origins in test coupons and photographs. Lessons learned. 

J. J. Swab and G. D. Quinn, “Fractography of Advanced Structural Ceramics: Results 

From Topic #2 of the VAMAS Round Robin Exercise”" Ceram. Eng. Sci. Proc., 16 [5] 

(1995) 929-938. 

A short summary of part of the 1994 VAMAS round robin on ceramic fractography. 

A-5 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

J. J. Swab and G. D. Quinn, “Results of a Round Robin Exercise on the Fractography of 

Advanced Structural Ceramics,” Cer. Eng. Sci. Proc., 15 [5] (1994) 867-876. 

A short summary of findings from the 1994 VAMAS fractography round robin. 

J. J. Swab and G. D. Quinn, “The VAMAS Fractography Round Robin: A Piece of the 

Fractography Puzzle,” pp. 55 - 70 in Fractography of Glasses and Ceramics III, eds. J. 

Varner and G. D. Quinn, American Ceramic Society, Westerville, OH, Ceramic 

Transactions, Vol. 64, (1996). 

Seventeen laboratories analyzed photos of grinding cracks. They also analyzed 

actual specimens with different flaws. Results varied. Much was learned by the 

organizers from this exercise. How do fractographers sift through and use the clues? 

Origins in Ceramics 

H. Kirchner, R. Gruver, and W. Sotter, “Characteristics of Flaws at Fracture Origins 

and Fracture Stress—Flaw Size Relations in Various Ceramics,” Mater. Sci. Eng., 22 

(1976) 147–156. 

A concise but useful report on strength-limiting origins in alumina, silicon nitride, 

and silicon carbide with a detailed tabulation of different types of origins. Emphasis 

is on porosity, large grains, and machining origins. An important observation (Fig. 

1b) is that origins in the center of fracture mirrors may intersect the fracture surface 

at an angle and a true view of the origin may not be seen. 

H. Baumgartner and D. Richerson, “Inclusion Effects on the Strength of Hot Pressed 

Si3N4,” Fracture Mechanics of Ceramics, Vol. 1, 1974, pp. 367–386. 

Good characterization of machining damage and inclusions in silicon nitride. The 

inclusions are much smaller than expected due to local-degradation of the matrix KIc. 

M. G. Gee and R. Morrell, “Fracture Mechanics and Microstructures,” pp. 1 – 22 in 

Fracture Mechanics of Ceramics, Vol. 8, R. Bradt, A. Evans, D. Hasselman, and F. 

Lange, eds., Plenum Press, NY, 1986. 

Principally a discussion of the application of fracture mechanics theories to strength. 

Microstructural influences significantly complicate matters and may limit utility to 

qualitative issues. The nature of strength-limiting origins and their severity is 

discussed. In some instances, sharp cracks will not form until the stress is applied. 

A. G. Evans, “Structural Reliability, A Processing Dependent Phenomenon,” J. Amer. 

Ceram. Soc, 65 No. [3] (1982) 127–139. 

Evans didn’t spend much time looking through microscopes, but he was the 

preeminent mathematic modeler of mechanical behavior in ceramics and glasses in 

the 1970s - 2000s. This paper and the two below were important contributions 

pertaining to flaws and their effect upon strength and reliability. This paper has an 

emphasis on the micromechanics of fracture with a good discussion of the effect of 

thermal and mechanical property mismatches between an origin and the matrix. 

Graph of stress versus origin size for silicon nitride showing the relative severity of 

different origins (WC, Fe, Si, C inclusions, porosity, and machining damage). 

A-6 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Appendix A Bibliography 

A. G. Evans, “Non-Destructive Failure Prediction in Ceramics,” pp. 595-626 in 

Progress in Nitrogen Ceramics, ed. F. L. Riley, Matirnus Nijhoff Publ., The Hague, 

1983. 

Includes the graph of stress versus origin size for silicon nitride showing the relative 

severity of different origins (WC, Fe, Si, C inclusions, porosity, and machining 

damage). 

A. G. Evans, M. E. Meyer, K. W. Fertig, B. I. Davis, and H. R. Baumgartner, 

“Probabilistic Models for Defect Initiated Fracture in Ceramics,” J. Nondest. Eval., 1 

[2] (1980) 111 – 122. 

Includes some of the actual data used to construct the strength versus flaw size graph 

for different flaws. Many of the flaws were intentionally seeded into hot-pressed 

silicon nitride. Additional data were obtained from work done at the U.S. Army 

Materials and Mechanics Research Center at Watertown, MA and documented in 

U.S. Army Technical Report TR78-11 by Brockelman and Hansen in 1978. 

G. D. Quinn, L. K. Ives, and S. Jahanmir, “On the Fractographic Analysis of 

Machining Cracks in Ground Ceramics: A Case Study on Silicon Nitride,” Special 

Publication SP 996, NIST, Gaithersburg, MD, May, 2003. 

An exhaustive, well-illustrated treatment of the nature of machining flaws in ground 

ceramics and how to find them. Photos and many schematics. 

G. D. Quinn, L. K. Ives, and. S. Jahanmir, “On the Nature of Machining Cracks in 

Ground Ceramics: Part I: SRBSN Strengths and Fractographic Analysis,” Mach. Sci. 

Technol., 9 (2005) 169 - 210. 

G. D. Quinn, L. K. Ives, and S. Jahanmir, “On the Nature of Machining Cracks in 

Ground Ceramics: Part II: Comparison to Other Silicon Nitrides and Damage Maps,” 

Mach. Sci. Technol., 9 (2005) 211 - 237. 

G. D. Quinn, L. K. Ives, and S. Jahanmir, “Machining Damage Cracks: How to Find 

and Characterize Them by Fractography,” Ceram. Eng. Sci. Proc., 24 [4] (2003) 383-

394. 

A short summary of the work in the previous three references. 

G., D. Quinn, L. K. Ives, and S. Jahanmir, “Machining Cracks in Finished Ceramics,” 

pp. 1-14 in Fractography of Advanced Ceramics, II, eds. J. Dusza, R. Danzer, and R. 

Morrell, TransTech Publ., Zurich, 2005. 

A short summary of the work in the previous three references. 

R. W. Rice and J. J. Mecholsky, “The Nature of Strength Controlling Machining Flaws 

in Ceramics,” pp. 351-378 in The Science of Ceramic Machining and Surface Finishing 

II, eds. B. Hockey and R. Rice, National Bureau of Standards, Gaithersburg, MD, 

Special Publication SP562. 1979. 

A-7 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

R. W. Rice, “Failure Initiation in Ceramics: Challenges of NDE and Processing,” and 

“Ceramic Developments,” C. Sorrell, and B. Ben-Nissan, eds., Materials Science 

Forum, Vol. 34–36, Trans. Tech. Publ. Ltd. Switzerland, 1988, pp. 1057–1064. 

A comprehensive, well-illustrated review of failure-initiating origins. Nearly an 

encyclopedia of flaws. Origins include: agglomerates, pores, large grains, inclusions, 

machining damage, handling damage, thermocouple beads, ball mills, dandruff, 

insects, feces, inadequate mixing of constituents, etc. 

R. W. Rice, “Processing Induced Sources of Mechanical Failure in Ceramics,” pp. 303– 

319 in Processing of Crystalline Ceramics, eds., H. Palmour, R. Davis, and T. Hare, 

Plenum Press, NY, 1978. 

A short, well-illustrated review of origins. A good starting point. 

R. W. Rice, J. J. Mecholsky, Jr., and P. F. Becher, “The Effect of Grinding Direction on 

Flaw Character and Strength of Single Crystal and Polycrystalline Ceramics,” J. Mat. 

Sci., 16 (1981) 853–862. 

Machining damage in a variety of ceramics is well illustrated by nine figures. 

J. J. Mecholsky, Jr., S. W. Freiman, and R. W. Rice, “Effects of Grinding on Flaw 

Geometry and Fracture of Glass,” J. Amer. Ceram. Soc., 60 [3–4] (1977) 114–117. 

Two primary sets of cracks result from surface grinding. These are schematically 

shown and complemented by SEM photos and related to fracture mechanics 

parameters. 

R. W. Rice, “Pores as Fracture Origins in Ceramics,” J. Mat. Sci., 19 (1984) 895–914. 

A well-illustrated examination of pores in glassy and polycrystalline materials. Pores 

tend to be “sharper” in the latter than in the former. 

R. W. Rice, S. W. Freiman, and J. J. Mecholsky, Jr., “The Dependence of StrengthControlling Fracture Energy on the Flaw Size to Grain Size Ratio,” J. Am. Ceram. 

Soc., 63 [3-4] (1980) 129-136. 

The flaw-size to grain-size ratio determined whether the relevant fracture toughness 

is that for a polycrystalline material or that for a single crystal. Small flaws within 

very large grains are ruled by single crystal fracture toughness values which are 

much less than for polycrystals. Large grains are susceptible to machining cracks. 

D. Munz, O. Rosenfelder, K. Goebells, and H. Reiter, “Assessment of Flaws in 

Ceramic Materials on the Basis of Non-Destructive Evaluation,” pp. 265–283 in 

Fracture Mechanics of Ceramics, Vol. 1, eds. R. Bradt, D. Hasselman, and F. Lange, 

F., Plenum Press, NY, 1986 . 

Six different flaw types were characterized in reaction bonded and sintered silicon 

nitrides. Some flaws were artificially created to support a fracture mechanics 

analysis. Pores have a different effect upon strength than inclusions. 

A-8 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Appendix A Bibliography 

J. E. Ritter, Jr., and R. W. Davidge, “Strength and Its Variability in Ceramics with 

Particular Reference to Alumina,” J. Am. Ceram. Soc., 67 [6] (1984) 432-437. 

A good study of strength variability in alumina with flaws deliberately introduced by 

adding glucose or corn starch to starting powders, or by grit blasting surfaces. Flaw 

linking was observed. 

F. F. Lange, “Processing-Related Fracture Origins: I, Observations in Sintered and 

Isostatically Hot-Pressed Al2O3/ZrO2 Composites,” J. Am. Ceram. Soc., 66 [6] (1983) 

396-398. 

Four companion articles about flaws generated during processing one system.
 

Differential shrinkage of agglomerates created cracking.
 

F. F. Lange and M. Metcalf, “Processing-Related Fracture Origins: II, Agglomerate 

Motion and Cracklike Internal Surfaces Caused by Differential Sintering,” ibid, 398-

406. 

F. F. Lange, B. I. Davis, and I. A. Aksay, “Processing-Related Fracture Origins: III, 

Differential Sintering of ZrO2 Agglomerates in Al2O3/ZrO2 Composite,” ibid, 407-408. 

F. F. Lange, B. I. Davis, and E. Wright, “Processing-Related Fracture Origins: IV, 

Elimination of Voids Produced by Organic Inclusions,” J. Am. Ceram. Soc., 69 [1] 

(1986) 66-609. 

M. Hangl, “Processing Related Defects in Green and Sintered Ceramics,” pp. 133-156 

in Fractography of Glasses and Ceramics, IV, Ceramic Transactions, Vol. 122, J. 

Varner, and G. D. Quinn, eds., American Ceramic Society, Westerville, OH, 2001. 

An excellent article about spray-dried agglomerates in varistor grade sintered zinc 

oxide. Many photographs show the agglomerates from their initial state through 

green body and final sintered stages. Agglomerates can be hollow or donut shaped. 

The final fracture surface will reveal the spray-dried pattern if one looks carefully. 

S. Nakamura, S. Tanaka, R. Furushima, K. Sato, and K. Uematsu, “Estimation of 

Weibull Modulus from Coarser Defect Distribution in Dry-pressed Alumina 

Ceramics,” J. Ceram. Soc. Japan, 117 [6] (2009) 742-747. 

Starting in the 1990s, Profs. Keizo Uematsu and Satoshi Tanaka at Nagaoka 

University did exceptional work characterizing flaws in ceramics, starting from the 

earliest processing stages and monitoring their evolution during subsequent 

processing. Innovative microscopy techniques were devised to see the flaw 

population in ceramic bodies, as well as the flaws exposed by conventional polished 

section work and fractographic analysis of test specimens. The flaws were closely 

coupled to the final strength distribution. Their work confirmed analytical 

predictions by Jayatilaka and Trustrum in 1977 and 1983. The latter work plus the 

Nagaoka University work established beyond any doubt that the Weibull distribution 

is appropriate for brittle materials. There are dozens of good papers by this group, 

and only a few are listed here. 

A-9 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

H. Abe, M. Naito, T. Hotta, N. Shinohara, and K. Uematsu, “Flaw Size Distribution in 

High-Quality Alumina,” J. Am. Ceram. Soc., 86 [6] (2003) 1019-1021. 

Y. Saito, S. Tanaka, N. Uchida, and K. Uematsu, “Direct Observation of ThreeDimensional Structure in Powder Compact,’ pp. 657-660 in Key Engineering 

Materials, Vols. 206-213, Proceedings of the 7th European Ceramic Society, Brugge, 

Belgium, 2001, Transtech Publ., Zurich, Switzerland, 2002. 

Confocal laser scanning fluorescence microscopy with a special immersion fluid 

was used to see the three dimensional flaw distribution in alumina. 

Y. Saito, S. Tanaka, N. Uchida, and K. Uematsu, “Direct Evidence for Low-Density 

Regions in Compacted Spray-Dried Powders,” J. Am. Ceram. Soc., 84 [10] (2001) 

2454-2456. 

K. Sato, S. Tanaka, N. Uchida, and K. Uematsu, “Characterization of Large Defects in 

Alumina,” Bull. Am. Ceram. Soc., 82 [4] (2003) 39-42. 

Transmission optical microscopy was used to study the flaw population through a 

thin 200 μm piece of alumina. 

K. Sato, S. Tanaka, N. Uchida, and K. Uematsu, “Morphological Change of Large 

Pores in Alumina Ceramics in the Final Stage of Densification,” J. Ceram. Soc. Japan, 

111 [7] (2003) 525-527. 

Transmission optical microscopy was used to study the flaw population through a 

thin 200 μm piece of alumina. 

K. Sato, S. Tanaka, N. Uchida, and K. Uematsu, “Direct Study on the Formation of 

Large Defects in Final Stages of Sintering in Alumina Ceramics,” pp. 225 – 228 in Key 

Engineering materials, Vols. 264-268, Proceedings of the 8th European Ceramic Society 

conference, Trans Tech Publ., Zurich, 2004. 

The formation and development of defects in sintered alumina was studied with 

transmission optical microscopy and the SEM.
 

N. Shinohara, M. Okumiya, T. Hotta, K. Nakahira, M. Naito, and K. Uematsu, 

“Morphological Changes in Process-Related large Pores of Granular Comp-acted and 

Sintered Alumina,” J. Am. Ceram. Soc., 83 [7] (2000) 1633-1640. 

Exceptional pictures and characterization of pores of different shapes formed as a 

result of power agglomerate collapse during processing. Compacts sintered at 

different temperatures had different flaws. 

H. Takahasi, N. Shinohara, K. Uematsu, and T. Junichiro, “Influence of Granule 

Character and Compaction on the Mechanical Properties of Sintered Silicon Nitride,” J. 

Am. Ceram. Soc., 79 [4] (1996) 843-858. 

N. Shinohara, M. Okumiya, T. Hotta, K. Nakahira, M. Naito, and K. Uematsu, 

“Seasonal Variation of Microstructure and Sintered Strength of Dry-Pressed Alumina,” 

J. Am. Ceram. Soc., 82 [12] (1999) 3441 – 3446. 

A-10 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Appendix A Bibliography 

An exceptional study that showed strength varied with the season of the year. The 

variation was traced to subtle flaw variations in the powder compacts that were 

caused by differences in the flowability of the spray-dried powders and their 

deformability in winter and summer due to temperature and humidity differences in 

the processing environment! 

H. Abe, M. Naito, T. Hotta, N. Shinohara, and K. Uematsu, “Flaw Size Distribution in 

High-Quality Alumina,” J. Am. Ceram. Soc., 86 [6] (2003) 1019 -1021. 

Transmission optical microscopy characterized the flaws in three-dimensional pieces 

of three aluminas. The exponents of the flaw size distributions did not correlate well 

to the actual four–point strength distributions in this case, probably since only flaws 

in a small 0.5 mm3 slice were measured by transmission optical examination. In 

contrast, the bend bars had Weibull effective volumes of 1.9 – 2.5 mm3

. The flaws 

were also round pores and not penny-shaped cracks. 

Y. Zhang, M. Inoue, N. Uchida, and K. Uematsu, “Characterization of Processing
 

Pores and Their Relevance to the Strength in Alumina Ceramics,” J. Mater. Res., 14 [8] 

(1999) 3370-3374.
 

T. Bernthaler, “Defect Imaging and Analysis for Predicting the Reliability of
 

Ceramics,” PhD Thesis, School of Materials Science and Engineering, University of
 

New South Wales, Sydney, Australia, and the University of Aalen, Germany, April
 

2011
 

A remarkable study under the guidance of Profs. Mark Hoffman and Gerhard 

Schneider, comparing flaw size distributions measured with automated image 

analysis on polished sections and strength distributions and fractographic analysis in 

alumina, alumina with seeded defects, dense and porous silicon carbide, and an 

alumina-zirconia composite. A variety of flaw types were identified and 

characterized. The size distributions matched an inverse power law function. 

Origins in Glass 

J. R. Varner, “The Practical Strength of Glass,” pp. 389 – 406 in Strength of Inorganic 

Glass, ed. C. Kurkjian, Plenum, NY, 1986. 

V. D. Fréchette, Failure Analysis of Brittle Materials, Advances in Ceramics, Vol.28, 

American Ceramic Society, Westerville, OH, 1990. 

Fracture Mirrors 

J. J. Mecholsky, Jr., S. W. Freiman, and R. W. Rice, “Fracture Surface Analysis of 

Ceramics,” J. Mat. Sci., 11 (1976) 1310–1319. 

A detailed correlation of origin size, fracture mirror sizes and characterization, and 

fracture mechanics parameters for single and polycrystalline ceramics. A table of 

mirror constants is given for a range of ceramics, and it is demonstrated that the outer 

A-11 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

mirror (hackle) to origin size ratio is about 13 to 1. The inner mirror (mist) ratio is 

between 6 to 1 and 10 to 1. 

J. J. Mecholsky, Jr., R. W. Rice, and S. W. Freiman, “Prediction of Fracture Energy 

and Flaw Sizes in Glasses from Measurements of Mirror Size,” J. Amer. Ceram. Soc., 

57 [10] (1974) 440–443. 

Details of fracture mirror features are discussed and related to fracture mechanics 

parameters for glasses. A table of mirror constants for glasses. 

H. P. Kirchner, R. M. Gruver, and W. A. Sotter, “Fracture Stress—Mirror Size 

Relations for Polycrystalline Ceramics,” Phil. Mag., 33 [5] (1976) 775–780.
 

Many mirror constants for a range of ceramics.
 

H. P. Kirchner, and J. C. Conway, Jr., “Fracture Mechanics of Crack Branching in 

Ceramics,” pp. 187–213 in Fractography of Glass and Ceramics, Advances in 

Ceramics, Vol. 22, American Ceramic Society, Westerville, OH, 1988. 

Analysis that fracture mirror features are controlled by stress intensity. 

J. J. Mecholsky, Jr., and S. W. Freiman, “Determination of Fracture Mechanics 

Parameters Through Fractographic Analysis of Ceramics,” pp. 136 – 150 in Fracture 

Mechanics Applied to Brittle Materials, ASTM STP 678, S. Freiman, ed., ASTM, 1979. 

A short discussion of fracture mirrors and mirror constants with a comparative table 

of mirror constants. Comments on useful techniques to measure mirror parameters. 

Fracture Mechanics—Estimates of Flaw Size 

D. W. Richerson, Modern Ceramic Engineering, Marcel Dekker Inc., NY, 1982. 

Chapter 3 is a good primer on strength and fracture toughness measurements and 

their applicability to fractographic analysis. Several numerical examples are given 

for estimating the strength of a specimen on the basis of a fracture mechanics 

calculation using the measured flaw size. 

H. K. Baumgartner and D. W. Richerson, “Inclusion Effects on the Strength of Hot 

Pressed Si3N4,” Fracture Mechanics of Ceramics, Vol. 1, eds. R. Bradt, D. Hasselman, 

and F. Lange, Plenum Press, NY, 1974, pp. 367–386. 

Applies fracture mechanics to one class of flaws with several numerical examples. 

The strength-limiting inclusions are smaller than expected from fracture mechanics, 

suggesting that the fracture toughness is altered in the vicinity of the inclusions. 

J. J. Mecholsky, Jr., S. W. Freiman, and R. W. Rice, “Fracture Surface Analysis of 

Ceramics,” J. Mat. Sci., 11 (1976) 1310–1319. 

Compares measured flaw sizes to fracture mechanics estimates for a range of 

ceramics and glasses. 

A-12 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Appendix A Bibliography 

H. P. Kirchner, R. M. Gruver, and W. A. Sotter, “Characteristics of Flaws at Fracture 

Origins and Fracture Stress-Flaw Size Relations in Various Ceramics,” Mater. Sci. 

Eng., 22 (1976) 147–156. 

Measured flaw sizes were compared to fracture mechanics estimates for several
 

different types of flaws in alumina, silicon nitride, and silicon carbide.
 

A. G. Evans, and G. Tappin, “Effects of Microstructure on the Stress to Propagate 

Inherent Flaws,” Proc. Brit. Ceram. Soc., 20 (1972) 275–297. 

Discusses flaws in alumina ceramics and compares the stress actually needed to 

cause fracture to estimates based on fracture mechanics. Microstructural factors such 

as flaw linking prior to catastrophic fracture are discussed. 

D. Munz, O. Rosenfelder, K. Goebells, and H. Reiter, “Assessment of Flaws in 

Ceramic Materials on the Basis of Non-Destructive Evaluation,” pp. 265–283 in 

Fracture Mechanics of Ceramics, Vol. 1, eds. R. Bradt, D. Hasselman, and F. Lange, 

Plenum Press, NY, 1986. 

A superb, comprehensive fracture mechanics analysis of six different flaw types in 

two silicon nitrides. Fractographic size measurements agreed with fracture 

mechanics estimates for some flaw types, but not others. Over 100 specimens. 

Discusses the different crack models that can be used to simulate real flaws as well 

as the shortcomings of such models. Includes Raju-Newman and elliptical integral 

flaw stress intensity factor solutions. 

G. D. Quinn and J. J. Swab, “Fractography and Estimates of Fracture Origin Size from 

Fracture Mechanics,” Ceram. Eng. Sci. Proc., 17 [3] (1996) 51–58. 

Fracture mechanics should be used routinely in fractographic analyses to verify that 

the correct feature has been identified as the fracture origin. This paper reviews the 

factors that may cause calculated and measured flaw sizes to differ. 

Fracture Mechanics—Stress Intensity Factors 

Y. Murakami, Stress Intensity Factors Handbook, Vols. 1 and 2, Pergamon Press, NY, 

1986. 

A collection of stress intensity factors for various cracks under different loading 

conditions. 

H. Tada, P. C. Paris, and G. R. Irwin, The Stress Analysis of Cracks Handbook, 3rd 

edition, ASM International, Metals Park, OH, 2000. 

An update of their 1973 collection of stress intensity factors for various cracks under 

different loading conditions. 

T. Fett, and D. Munz, Stress Intensity Factors and Weight Functions, Wessex Institute 

of Technology, Southhampton, UK, 1997. 

A collection of stress intensity factors for various cracks under different loading 

conditions. 

A-13 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

D. P. Rooke, and D. J. Cartwright, Compendium of Stress Intensity Factors, Her 

Majesty's Stationary Office, London, 1976. 

A collection of stress intensity factors for various cracks under different loading 

conditions. 

G. C. Sih, Handbook of Stress Intensity Factors, Lehigh University, Bethlehem, PA, 

1973. 

An older but still useful collection of stress intensity factors for various cracks under 

different loading conditions. 

I. Bar-on, “Applied Fracture Mechanics,” Engineered Materials Handbook, Vol. 4, 

Ceramics and Glasses, S. Schneider, ed., ASM, Metals Park, OH, 1991, pp. 645–651. 

A good introduction to the application of fracture mechanics analysis to idealized 

crack configurations. Stress intensity shape factors are given for through slits, 

surface cracks, and pores with rim cracks. 

J. C. Newman, Jr., and I. S. Raju, “An Experimental Stress-Intensity Factor Equation 

for the Surface Crack,” Eng. Fract. Mech., 15 [1–2] (1981)185–192. 

Presents an widely used equation for the calculation of the stress intensity shape 

factors (Y) for surface semicircular or semielliptical cracks. Y is given for the 

location where the origin meets the surface and at the deepest point of the origin. 

Both direct tension and flexural solutions are presented. 

S. Strobl, P. Supancic, T. Lube, and R. Danzer, “Surface Crack in Tension or Bending 

– A Reassessment of the Newman and Raju Formula in Respect to Fracture Toughness 

Measurements in Brittle Materials,” J. Eur. Cer. Soc., 32 (2012) 1491–1501. 

Also presents shape factor (Y) for surface semicircular or semielliptical cracks, but 

for more general cases than the Newman-Raju solutions. The Strobl et al. solutions 

can be considered refined solutions. They allow different values of Poisson’s ratio 

and crack intersection angles at the tensile surface to be input. For most surface 

cracks the new solutions are only a few percent different than Newman and Raju’s. 

The solutions match Newman and Raju’s for ν = 0.3 and for semielliptical surface 

cracks (where the crack surface angle is 90o

). 

F. Baratta, “Stress Intensity Factor Estimates for a Peripherally Cracked Spherical Void 

and a Hemispherical Surface Pit,” J. Am. Ceram. Soc., 61 [11-12] (1978) 490-492. 

F. Baratta.” Refinement of Stress Intensity Factor Estimates for a Peripherally Cracked 

Spherical Void and a Hemispherical Surface Pit,” J. Am. Ceram. Soc., 64 [1] (1981) 

C3-C4. 

D. J. Green, “Stress Intensity Factor Estimates for Annular Cracks at Spherical Voids,” 

J. Am. Ceram. Soc., 63 [5-6] (1980) 342-344. 

T. Fett, “Stress Intensity Factors and Weight Functions for a Void with an Annular 

Crack,” Int. J. Fract., 67 (1994) R41-R47. 

A-14 

 

 

 

 

 

 

 

 

 

 

 

 

Appendix B Case Studies 

APPENDIX B FRACTOGRAPHIC CASE STUDIES
	

Case studies for components and structures are organized by topic. Glass cases come 

first, followed by ceramics and glass ceramics, fibers, single crystals, gemstones, and 

rocks and lithics. 

Glass Windows and Plates of All Sizes 

Automotive and building windows 

N. Shinkai, “The Fracture and Fractography of Flat Glass,” pp. 253 – 297 in 

Fractography of Glass, eds. R. C. Bradt and R. E. Tressler, Plenum Press, NY, 1994. 

A superb practical article on annealed, tempered, and laminated glass for 

automobiles or building windows. 

V. D. Fréchette and M. Donovan, “Some Effects of the Glue Chipping Process on 

Strength,” pp. 407 – 411 in Fractography of Glasses and Ceramics II, Ceramic 

Transactions Vol. 17, eds. V. D. Fréchette and J. R. Varner, American Ceramic 

Society, Westerville, OH, 1991. 

Fractures of hundreds of very large float glass windows in a Boston skyscraper. 

V. D. Fréchette, Failure Analysis of Brittle Materials, Advances in Ceramics, Vol. 28, 

American Ceramic Society, Westerville, OH, 1990, pp. 119 -120. 

Fractures in large float glass windows in a skyscraper. Glue chips. 

T. Sakai, M. Ramulu, A. Ghosh, and R. C. Bradt, “A Fractal Approach to Crack 

Branching (bifurcation) in Glass,” pp. 131 –146 in Fractography of Glasses and 

Ceramics II, eds., V. D. Fréchette and J. R. Varner, Ceramic Transactions, Vol. 17, 

American Ceramic Society, Westerville, OH, 1991. 

Fracture patterns in laminated safety glass for bus passenger shelters. 

S. T. Gulati, R. Akcakaya, and J. Varner, “Fracture Behavior of Tin vs. Air Side of 

Glass,” pp. 317 – 325 in Fractography of Glasses and Ceramics, IV, eds., J. R. Varner 

and G. D. Quinn, American Ceramic Society, Westerville, OH, 2001. 

S. T. Gulati, J. D. Helfinstine, T. A. Roe, “Strength Degradation of Automotive 

Windshield from Manufacturing to On-Road Service,” Glass Technol., 43C (2002) 303 

– 308. 

Testing of singlets and windshield coupons. Edging, scoring, solder tabs. 

T. Cleary and S. Gulati, “Influence of Glass Score and Seam Orientation on Edge 

Strength of Multi-Layered Glass Articles,” pp. 327 – 341 in Fractography of Glasses 

and Ceramics, IV, eds. J. R. Varner and G. D. Quinn, American Ceramic Society, 

Westerville, OH, 2001. 

Fracture patterns in laminated safety glass (sandwich of two glass plates and a 

central plastic sheet). 

B-1 
 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

R. A. Allaire and T. Ono, “Fracture Analysis of the Glass Scoring Process,” pp. 467 – 

471 in Fractography of Glasses and Ceramics, IV, eds. J. R. Varner and G. D. Quinn, 

American Ceramic Society, Westerville, OH, 2001. 

Fracture patterns on scored glass plates. 

T. M. Cleary and R. T. Nichols, “ Adhesive Induced Fracture of Automotive Glass,” 

pp. 369-380 in Fractography of Glasses and Ceramics V, eds. J. R. Varner, G.D. 

Quinn, M. Wightman, Ceramic Transactions, Vol. 199, American Ceramic Society, 

Westerville, OH, 2007. 

Glue chips and adhesive and thermally-induced fractures in both front laminated 

and tempered rear windows. 

R. Huet, J. M. Wolf, P. D. Moncarz, “Delayed Fracture of Tempered Glass Panels Due 

to Nickel Sulfide inclusions,” pp. 431-433 in Handbook of Case Studies in Failure 

Analysis, Vol. 1, ed. K. A. Esakul, ASTM Int. Materials park, OH, 1992. 

C. C. Hsiao, “Spontaneous Fracture of Tempered Glass, “ pp. 985 -992 in Fracture 

1977, Vol. 3, ed. D. Taplin, International Congress on Fracture 4, Waterloo, Canada, 

1977. 

S. A. Batzer, “Failure Mechanisms of Automotive Side Glazing in Rollover 

Collisions,” J. Fail. Anal. Prev., 6 [3] (2006) 6-11, 33,34. 

R. C. Bradt, M. E. Barkey, S. E. Jones, M. E. Stevenson, “Projectile Impact – A Major 

Cause for Fracture of Flat Glass,” J. Fail. Anal. Prev., 3 [1] (2003) 5-11. 

R. C. Bradt, “Macro- and Microfracture Patterns of Thermally Tempered Plate Glass 

Falling from Nickel Sulfide Inclusions,” pp. 417 – 426 in Fractography of Glasses and 

Ceramics, eds. J. D. Varner and V. D. Fréchette, Advances in Ceramics, Vol. 22, 

American Ceramic Society, Westerville, OH, 1988. 

Cell phone cover glass 

G. S. Glaesemann, T. M. Gross, J. F. Bayne, and J. J. Price, “Fractography on IonExchanged Cover Glass,” pp. 85 – 93 in Fractography of Glasses and Ceramics VI, 

eds. J. Varner and M. Wightman, Ceramic Transactions, Vol. 230, American Ceramic 

Society, Wiley, Westerville, OH, 2012. 

T. M. Gross, “Scratch Damage in Ion-Exchanged Alkali Aluminosilicate Glass: Crack 

Evolution and the Dependence of Lateral Cracking Threshold on Contact Geometry,” 

pp. 113-122 in Fractography of Glasses and Ceramics VI, eds. J. Varner and M. 

Wightman, Ceramic Transactions ,Vol. 230, American Ceramic Society, Wiley, 

Westerville, OH, 2012. 

Sight glass windows, water hammer 

V. D. Fréchette, Failure Analysis of Brittle Materials, Advances in Ceramics, Vol. 28, 

American Ceramic Society, Westerville, OH, 1990, p. 112. 

B-2 
 

 

 

 

 

 

 

 

 

 

 

 

 

Appendix B Case Studies 

Tempered glass spontaneous failure 

R. C. Bradt, “Macro- and Microfracture Patterns of Thermally Tempered Plate Glass 

Falling from Nickel Sulfide Inclusions,” pp. 417 – 426 in Fractography of Glasses and 

Ceramics, Advances in Ceramics, Vol. 22, ed. J. D. Varner and V. D. Fréchette, 

American Ceramic Society, Westerville, OH, 1988. 

C. C. Hsiao, “Spontaneous Fracture of Tempered Glass,” pp. 985 – 992 in Fracture 

1977, Vol. 3, International Congress on Fracture 4, Waterloo, Canada, 1977. 

Plate glass mirrors for solar cells 

V. D. Fréchette, Failure Analysis of Brittle Materials, Advances in Ceramics, Vol. 28, 

American Ceramic Society, Westerville, OH, 1990, p. 117. 

Hail stones, thermal stresses, and edge cracks. 

Laser glass slabs 

J. E. Marion, “Fracture Mechanisms of Solid Sate Slab Lasers,” pp. 307 – 318 in 

Fractography of Glasses and Ceramics, eds. J. D. Varner and V. D. Fréchette, 

Advances in Ceramics, Vol. 22, American Ceramic Society, Westerville, OH, 1988. 

J. E. Marion, “Fracture of Solid State Laser Slabs,” J. Appl. Phys., 60 [1] (1986) 69 – 

77. 

Borosilicate Glass Flowmeter 

J. Kübler, “Failure Analysis on a Flowmeter,” pp. 65-71 in Fractography of Advanced 

Ceramics III, eds. J. Dusza, R. Danzer, R. Morrell, and G. D. Quinn, Trans Tech Publ. 

Zurich, 2009. 

Home cocktail table 

A. A. Johnson and R. J. Storey, “A Fatal Accident Involving a Glass-Topped Cocktail 

Table,” J. Fail. Anal. Prev., 14 (2014) 267-271. 

A college-graduating student fell over a glass-topped cocktail table which 

shattered. A long sharp shard pierced his heart and killed him. The untempered 

glass was supported by eight tabs, only four of which had plastic inserts. 

Glass Containers, Lamps, and Pressure Vessels 

Pharmaceutical Glass Vials and Syringes 

F. Maurer, “Application of Fractography in Pharmaceutical Industry,” pp. 289 – 298 in 

Fractography of Glasses and Ceramics, VI, eds. J. R. Varner, M. Wightman, Ceramic 

Transactions, Vol. 230, Wiley, New York, 2012. 

Glass syringe compliant specimen. 

G. D. Quinn, This Guide, Chapter 10. Medicinal Glass Vials 

Neck cracking from the cap sealing operation. 

B-3 
 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Bottles and flasks 

V. D. Fréchette, Failure Analysis of Brittle Materials, Advances in Ceramics, Vol. 28, 

American Ceramic Society, Westerville, OH, 1990, pp. 94 – 101, 112, 118, 123. 

Internal pressure, water hammer, thermal shock, and impact fractures. Various 

bottles including intravenous medical bottles, wine bottles, frozen citrus fruit 

concentrate bottles, malformed bottles. 

F. W. Preston, “Bottle Breakage – Causes and Types of Fractures,” Am. Ceram. Soc. 

Bull., 18 [2] (1939) 35 – 60. 

An early paper, but with an exceptional treatment of a variety of failure causes. 

V. D. Fréchette and T. A. Michalske, “Fragmentation in Bursting Glass Containers,” 

Bull. Amer. Ceram. Soc., 57 [4] (1978) 427 – 429. 

Four different sizes of internally-pressurized, water-filled bottles. 

V. D. Fréchette and S. L. Yates, “Fragmentation of Glass Bottles by Impact,” J. Am. 

Ceram. Soc., 72 [6] (1989) 1060. 

Three different sizes of water-filled bottles were swung at different velocities 

against a stationary steel rod. 

J. B. Kepple, and J. S. Wasylyk, “Fracture of Glass Containers,” pp. 207 –252 in 

Fractography of Glass, eds. R. C. Bradt and R. E. Tressler, Plenum Press, NY, 1994. 

An excellent, multifaceted treatment. Shows fracture causes and patterns. 

R. E. Mould, “The Behavior of Glass bottles Under Impact,” J. Amer. Ceram. Soc., 35 

[9] (1952) 230 – 235. 

A well illustrated paper on controlled breakages of bottles. This article (and 

Dimmick’s that followed it) proved that hinge origins away from the impact site 

were sources of breakage in some cases. 

C. Colin and J. L. Heitz, “Failure Analysis: Glass Bottle Breakage,” Verre, 16 [4] 

(2010) 33-36. 

Sparkling wine bottle broke spontaneously in boxes in a warehouse. Glass-to-glass 

contact damage from improper packing caused cracks that grew to failure. 

D. D. Cannon, C. S. Musso, J. C. Williams, and T. W. Eagar, “Analysis of Brittle 

Fracture of Soda Glass Bottles Under Hydrostatic Pressure,” J. Fail. Anal. Prev., 45 [5] 

(2004) 72-77. 

Soda glass bottles were fractured in the lab by increasing pressure after a scratch 

had been made. Other bottles were impacted in various manners. 

Fluorescent lamps 

D. Johnson, “Arc Induced Fractures from Lampmaking Test Equipment,” pp. 517 – 

525 in Fractography of Glasses and Ceramics, III, Ceramic Transactions, Vol. 64, J. 

R. Varner, V. D., Fréchette, and G. D. Quinn, eds., American Ceramic Society, 

Westerville, OH, 1996. 

B-4 
 

 

 

 

 

 

 

 

Appendix B Case Studies 

Cathode Ray Tube 

A. Ghosh, C. Y. Cha, Vaidyanathan, and R. C. Bradt, “Finite Element Stress Analysis 

and Crack Path Prediction of Imploding CRT,” pp. 1- 24 in Fractography of Glasses 

and Ceramics II, eds. V. D. Fréchette and J. R. Varner, Ceramic Transactions, Vol. 17, 

American Ceramic Society, Westerville, OH, 1991. 

Glass cooking thermometer 

V. D. Fréchette, Failure Analysis of Brittle Materials, Advances in Ceramics, Vol. 28, 

American Ceramic Society, Westerville, OH, 1990, p. 120. 

Tempered glass cylinders 

V. D. Fréchette, Failure Analysis of Brittle Materials, Advances in Ceramics, Vol. 28, 

American Ceramic Society, Westerville, OH, 1990, pp. 120 –121, 123 –124. 

Uneven thermal temper in a structural cylinder and also a high power gymnasium 

lamp that lost its temper due to heating. 

Glass neutron beam guideline 

G. D. Quinn, O. Zilcha, J. M. Rowe, and D. J. Pierce, “Failure Analysis of a 41 Meter 

Long Neutron Beam Line Guide,” 4th International Conference on Fractography of 

Advanced Ceramics,” Smolenice Castle, Oct. 2013, J. Eur. Cer. Soc., 34 (2014) 3263-

3270. 

Glass Bulletproof or Impact Resistant Windows 

M. E. Stevenson, S. E. Jones, and R. C. Bradt, “Fracture Patterns of a Composite Safety 

Glass Panel During High Velocity Projectile Impacts,” pp. 473 – 488 in Fractography 

of Glasses and Ceramics, IV, eds. J. R. Varner and G. D. Quinn, American Ceramic 

Society, Westerville, OH, 2001. 

D. L. Ahearn, J. L. Ladner, S. E. Jones, R. E. Wright and R. C. Bradt, “Fracture 

Patterns of Impact Resistant Glass Panel Laminates with Annealed and Heat 

Strengthened Glass Plates,” pp. 383 – 395 in Fractography of Glasses and Ceramics V, 

eds. J. R. Varner, G.D. Quinn, M. Wightman, Ceramic Transactions, Vol. 199, 

American Ceramic Society, Westerville, OH, 2007. 

H. Yoshimura and R. G. Morrone, “Analysis of Projectile Impact Damage in 

Bulletproof Glasses,” pp. 397 – 406 in Fractography of Glasses and Ceramics V, eds. 

J. R. Varner, G.D. Quinn, M. Wightman, Ceramic Transactions ,Vol. 199, American 

Ceramic Society, Westerville, OH, 2007. 

J. LK. Ladner, D. L. Ahearn, and R. C. Bradt, “Laminate Design Effects on the 

Fracture Patterns of Impact Resistance Glass Panels,” pp. 281-287 in in Fractography 

of Glasses and Ceramics, VI, eds. J. R. Varner, M. Wightman, Ceramic Transactions, 

Vol. 230, Wiley, New York, 2012. 

B-5 
 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Ceramic Windows, IR Domes, Radomes 

Magnesium Fluoride (MgF2) IR window test coupons 

R. W. Rice, “ Failure Analysis of Ceramics,” pp. 369 – 388 in Fractography of Glasses 

and Ceramics, IV, eds. J. R. Varner and G. D. Quinn, American Ceramic Society, 

Westerville, OH, 2001. 

Calcium Fluoride (CaF2) windows in a laser flash diffusivity apparatus 

R. Morrell, “Dealing with Component Failures,” pp. 353 -368 in Fractography of 

Glasses and Ceramics V, eds. J. R. Varner, G.D. Quinn, M. Wightman, Ceramic 

Transactions, Vol. 199, American Ceramic Society, Westerville, OH, 2007. 

48 mm diameter single crystal disks that broke from thermal stresses activating 

edge damage or impact damage. 

Glass ceramic missile radome, thermal shock 

R. W. Rice, “Failure Analysis of Ceramics,” pp. 369 – 388 in Fractography of Glasses 

and Ceramics, IV, eds. J. R. Varner and G. D. Quinn, American Ceramic Society, 

Westerville, OH, 2001. 

Fused silica Missile radome 

This Guide. Chapter 10, case 7. 

Grinding flaws and joint debonding. 

Polycrystalline Alumina (PCA) Lamp Tubes 

F. Lofaj, V. Ivančo, and P.P Varga, “Fracture of Thin walled Translucent 

Polycrystalline (PCA) Tubes,” pp. 223 – 230 in Fractography of Advanced Ceramics 

III, eds. J. Dusza, R. Danzer, R. Morrell, and G. D. Quinn, Trans Tech. Zurich, 2009. 

Microwave tube (alumina) 

R. W. Rice, “Failure Analysis of Ceramics,” pp. 369 – 388 in Fractography of Glasses 

and Ceramics, IV, eds. J. R. Varner and G. D. Quinn, American Ceramic Society, 

Westerville, OH, 2001. 

Ceramic Seals 

Tetragonal zirconia polycrystalline (TZP) hydraulic seal disks 

R. W. Rice, “Failure Analysis of Ceramics,” pp. 369 – 388 in Fractography of Glasses 

and Ceramics, IV, eds. J. R. Varner and G. D. Quinn, American Ceramic Society, 

Westerville, OH, 2001. 

Mechanical seal rings – silicon carbide 

R. Morrell, “Fractography of Brittle Materials,” R. Morrell, Measurement Good 

Practice Guide 15, National Physical Laboratory, Teddington, Middlesex, United 

Kingdom, 1999, pp. 76 – 85. 

Large open rotating rings with shoulders and notches. Impact damage - chipping, 

preexisting cracks, grinding cracks, handling damage. 

B-6 
 

 

 

 

 

 

 

Appendix B Case Studies 

Large (300 mm) silicon carbide seal rings 

R. Morrell, “Dealing with Component Failures,” pp. 353 -368 in Fractography of 

Glasses and Ceramics V, eds. J. R. Varner, G.D. Quinn, M. Wightman, Ceramic 

Transactions, Vol. 199, American Ceramic Society, Westerville, OH, 2007. 

Fractures during over-speed poof testing by spinning the disks. Good sketches of 5 

disk fractures, illustrating how multiple crack patterns can be deciphered to find 

the primary fracture. Origins were edge chamfer damage, contact damage cracks, 

and machining cracks. 

Ceramic Containers and Pressure Vessels 

Large porcelain tubes 

V. D. Fréchette, Failure Analysis of Brittle Materials, Advances in Ceramics, Vol. 28, 

American Ceramic Society, Westerville, OH, 1990, pp. 125 – 126. 

“Spontaneous” fracture during end-face machining. Internal residual stresses. 

Beta alumina hollow cylinders, burst pressure loading 

B. J. McEntire, R. H. Snow, J. L. Huang, L. Viswanathan, and A. V. Virkar, 

“Characterization of Processing Flaws in Beta Alumina,” pp 335 – 349 in 

Fractography of Glasses and Ceramics, eds. J. D. Varner and V. D. Fréchette, 

Advances in Ceramics, Vol. 22, American Ceramic Society, Westerville, OH, 1988. 

Beta alumina electrolyte tubes, service failures 

E. K. Beauchamp, “Beta Alumina failures in Sodium/Sulfur Batteries,” pp. 377 – 387 

in Fractography of Glasses and Ceramics, eds. J. D. Varner and V. D. Fréchette, 

Advances in Ceramics, Vol. 22, American Ceramic Society, Westerville, OH, 1988. 

Ceramic Armor Plates 

S. Winkler, H. Senf, and H. Rothenhäusler, “High Velocity Fracture Investigation in 

Alumina,” pp. 165 – 183 in Fractography of Glasses and Ceramics II, eds. V. D. 

Fréchette and J. R. Varner, Ceramic Transactions, Vol. 17, American Ceramic Society, 

Westerville, OH, 1991. 

High-speed photography of alumina and glass plates hit on end. 

M. E. Stevenson, S. E. Jones, and R. C. Bradt, “Fracture Patterns of a Composite Safety 

Glass Panel During High Velocity Projectile Impacts,” pp. 473 – 488 in Fractography 

of Glasses and Ceramics, IV, eds. J. R. Varner and G. D. Quinn, American Ceramic 

Society, Westerville, OH, 2001. 

J. J. Swab, G. A., Glide, P. J. Patel, A. A Wereszczak, J. W. McCauley, and J. D. 

Risner, “Fracture Analysis of Transparent Armor Ceramics,” pp. 489 – 508 in 

Fractography of Glasses and Ceramics, IV, eds. J. R. Varner and G. D. Quinn, 

American Ceramic Society, Westerville, OH, 2001. 

Biaxial strength tests on sapphire, aluminum oxynitride, and magnesium alumina 

spinel. 

B-7 
 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

R. L. Woodward, R. G. O’Donnell, B. J. Baxter, B. Nicol, and S. D. Pattie, “Energy 

Absorption in the Failure of Ceramic Composite Armors,” Mater. Forum, 13 (1989) 

174 – 181. 

M. J. Slavin, “Fractographic Analysis of Long Rod Penetrator-Armor Ceramics 

Interactions,” U. S. Army Technical Report, TR 89-93, U.S. Army Materials 

Technology Laboratory, Watertown, MA, Oct. 1989. 

C. Tracy, M. Slavin, and D. Viechnicki, “Ceramic Fracture During Ballistic Impact,” 

pp. 295 – 306 in Fractography of Glasses and Ceramics, eds. J. D. Varner and V. D. 

Fréchette, Advances in Ceramics, Vol. 22, American Ceramic Society, Westerville, 

OH, 1988. 

S. Rodriguez, V. B. Munoz, E. V. Esquivel, L. E. Muir, and N. L. Rupert, “Microstructural Characterization of TiB2 Armor Targets,” J. Mat. Sci. Lttrs., 21 (2002) 1161 – 

1666. 

V. D. Fréchette and C. F. Cline, “Fractography of Ballistically Tested Ceramics,” J. 

Am. Ceram. Soc., 49 [11] (1970) pp. 994 – 997. 

V. D. Fréchette, “Fractography and Quality Assurance of Glass and Ceramics,” pp. 

227 236 in Quality Assurance in Ceramic Industries, eds. V. D. Fréchette, L. D. Pye, 

and D. E. Rase, Plenum, NY, 1978. 

Hot-pressed boron carbide blanks fractured during machining. Internal residual 

stresses were created by too rapid a cool down from the firing temperature. 

R. N. Katz and W. A. Brantley, “Fractography of High Boron Ceramics Subjected to 

Ballistic Loading,” pp. 271 – 282 in Materials Science Research, Vol. 5, ed. W. W. 

Kriegel, Plenum, NY, 1971. 

D. G. Christie and H. Kolsky, “The Fractures Produced in Glass and Plastics by the 

Passage of Stress Waves,” J. Soc. Glass Technol., 36 (1952) 65 - 73. 

Ceramic Ball Bearings 

J. J. Swab and M. P. Sweeney, “Fracture Analysis of An All-Ceramic Bearing System,” 

U. S. Army Research Laboratory Technical Report, ARL TR 512, September 1994. 

Y. Wang and M. Hadfield, “Rolling Contact Fatigue Failure Modes of Lubricated 

Silicon Nitride in Relation to Ring Crack Defects,” Wear, 225-229 (1999) 1284 – 1292. 

Y. Wang and M. Hadfield, “The Influence of Ring Crack Location on the Rolling 

Contact Fatigue Failure of Lubricated Silicon Nitride: Experimental Studies,” Wear, 

243 (2000) 157 – 166. 

B-8 
 

 

 

 

 

 

 

 

 

 

 

 

 

Appendix B Case Studies 

T. Lube, S. Witsching, P. Supancic, R. Danzer, O. Schöppl, “The Notched Ball Test – 

Characterization of Surface Defects and Their Influence on Strength,” pp. 225 – 234 in 

Fractography of Glasses and Ceramics, VI, eds. J. R. Varner, M. Wightman, Ceramic 

Transactions, Vol. 230, Wiley, New York, 2012. 

P. Supancic, R. Danzer, W. Harrer, Z. Wang, S. Witschnig, and O. Schöppl, “Strength 

Tests on Silicon Nitride Balls,” pp. 193 – 200 in Fractography of Advanced Ceramics 

III, eds. J. Dusza, R. Danzer, R. Morrell, and G. D. Quinn, Trans Tech Publ. Zurich, 

2009. 

Cellular Ceramics, Catalytic Converters, Filters 

D. J. Green and R. Brezny, “Fractographic Determination of Strut Strength in Cellular 

Ceramics,” pp. 199 – 129 in Fractography of Glasses and Ceramics II, eds. V. D. 

Fréchette and J. R. Varner, Ceramic Transactions, Vol. 17, American Ceramic Society, 

Westerville, OH, 1991. 

J. Kübler, R. Baechtold, G. Blugan, L. Lemster, and S. Fuso, “Failure Analysis on a 

De-Nox Catalyst of a Large Waste Burner,” pp. 78 - 85 in Fractography of Advanced 

Ceramics II, eds. J. Dusza, R. Danzer and R. Morrell, Transtech, Zurich, 2005. 

Special Ceramic Components 

MEMS structures 

S. J. Glass, D. A. LaVan, T. E. Buchheit, and K. Jackson, “Strength Testing and 

Fractography of MEMS Materials,” pp. 227 –240 in Fractography of Glasses and 

Ceramics, IV, eds. J. R. Varner and G. D. Quinn, American Ceramic Society, 

Westerville, OH, 2001. 

G. Quinn, W. Sharpe, G. M Beheim, N. Nemeth and O. Jadaan, “Fracture Origins in 

Miniature Silicon Carbide Structures,” pp. 62 –69 in Fractography of Advanced 

Ceramics II, eds. J. Dusza, R. Danzer and R. Morrell, Transtech, Zurich, 2005. 

Theta Specimens 

G. D. Quinn, “Fractographic Analysis of Very Small Theta Specimens,” pp. 201-208 in 

Fractography of Advanced Ceramics III, eds. J. Dusza, R. Danzer, R. Morrell, and G. 

D. Quinn, Trans Tech Publ., Zurich, 2009. 

Ceramic automotive valves 

R. Danzer, M. Hangl, and R. Paar, “Edge Chipping of Brittle Materials,” pp. 43 – 55 in 

Fractography of Glasses and Ceramics, IV, eds. J. R. Varner and G. D. Quinn, 

American Ceramic Society, Westerville, OH, 2001. 

M. J. Andrews, A. A. Wereszczak, K. Breder, and T. P. Kirkland, “Fractographic 

Analysis Applied to Ceramic Component Life Prediction,” pp. 301 – 315 in 

B-9 
 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Fractography of Glasses and Ceramics, IV, eds. J. R. Varner and G. D. Quinn, 

American Ceramic Society, Westerville, OH, 2001. 

Ceramic gas turbine model rotor 

G. D. Quinn, This Guide, Chapter 10, and 

G. D. Quinn, “Design and Reliability of Ceramics, Do Modelers, Designers and 

Fractographers See the Same World?” Ceram. Eng. Sci. Proc., 26 [8] (2005) 239 - 252. 

α - silicon carbide heat exchanger tube 

K. Breder and J. R. Kaiser, “Failure Analysis of an α-SiC Tube Subjected to Thermal 

Cycling in and Oxygen-Steam Atmosphere,” pp. 301 –315 in Fractography of Glasses 

and Ceramics, III, eds. J. R. Varner, V. D. Fréchette, and G. D. Quinn, Ceramic 

Transactions, Vol. 64, American Ceramic Society, Westerville, OH, 1996. 

Silicon carbide machine gun barrel 

G. D. Quinn, This Guide, Chapter 10, and 

G. D. Quinn, “Design and Reliability of Ceramics, Do Modelers, Designers and 

Fractographers See the Same World?” Ceram. Eng. and Sci. Proc., 26 [8] (2005) 239 -

252. 

Contact cracks from shrink fitting and stress reverberations 

Silicon nitride rolling mill rollers 

M. Lengauer, R. Danzer, D. Rubesa, W. Harrer, and W. Zleppnig, “Failure Analysis of 

Si3N4 Rolls for Wire Hot Rolling by Numerical Simulation of Thermal and Mechanical 

Stresses,” pp. 94 -101 in Fractography of Advanced Ceramics II, eds. J. Dusza, R. 

Danzer and R. Morrell, Transtech, Zurich, 2005. 

R. Danzer, M. Lengauer, D. Rubesa, and W. Harrer, “Silicon Nitride Tools for Hot 

Rolling of High-alloyed Steel and Superalloy Wires,” pp. 43 – 54 in Fractography of 

Advanced Ceramics III, eds. J. Dusza, R. Danzer, R. Morrell, and G. D. Quinn, Trans 

Tech Publ. Zurich, 2009. 

Lead zirconate titanate (PZT) sonar rings 

R. W. Rice, “Failure Analysis of Ceramics,” pp. 369 – 388 in Fractography of Glasses 

and Ceramics, IV, eds. J. R. Varner and G. D. Quinn, American Ceramic Society, 

Westerville, OH, 2001. 

Epoxy bonding stresses 

Silicon Nitride foundry tubes and rolls 

W. Harrer, R. Danzer, K. Berroth, “Thermal Shock Behavior of Si3N4 Specimens – 

Influence of Annealing and Edges,” pp. 245 – 254 in Fractography of Glasses and 

Ceramics, VI, eds. J. R. Varner, M. Wightman, Ceramic Transactions, Vol. 230, 

Wiley, New York, 2012. 

B-10
 
 

 

 

 

 

Appendix B Case Studies 

Beryllium oxide hollow cylinder 

J. R. Varner and V. D. Frechette, “Fractography of Whitewares,” pp. 305 – 315 in 

Science of Whitewares, eds. V. Henkes, G. Onoda, and W. Carty, American Ceramic 

Society, Westerville, OH, 1996. 

Hollow cylinder loaded at high temperature under laboratory conditions. 

Electronic Ceramics 

Porcelain railroad electrical insulator 

J. Woodtli, K. Beroth, and T. Luthi, “Combination of Fractography and Computed 

Tomography for the Determination of the Cause of Fracture,” pp. 257 – 271 in 

Fractography of Glasses and Ceramics, III, eds. J. R. Varner, V. D. Fréchette, and G. 

D. Quinn, Ceramic Transactions, Vol. 64, American Ceramic Society, Westerville, 

OH, 1996. 

Porcelain insulator, high tension power line 

J. R. Varner and V. D. Frechette, “Fractography of Whitewares,” pp. 305 – 315 in 

Science of Whitewares, eds. V. Henkes, G. Onoda, and W. Carty, American Ceramic 

Society, Westerville, OH, 1996. 

Material fault, a clump of feldspar grains 

Silicon wafers 

R. E. Moore, P. G. Hansen, W. Carty, and J. W. Ha, “Fracture of Thermally Shocked 

Silicon Disks,” pp. 351 – 361 in Fractography of Glasses and Ceramics, eds. J. D. 

Varner and V. D. Fréchette, Advances in Ceramics, Vol. 22, American Ceramic 

Society, Westerville, OH, 1988. 

Stacked PZT piezoelectric actuator material 

P. Supancic, Z. Wang, W. Harrer, K. Reichmann, and R. Danzer, “Strength and 

Fractography of Piezoceramic Multilayer Stacks,” pp. 46 –53 in Fractography of 

Advanced Ceramics II, eds. J. Dusza, R. Danzer and R. Morrell, Transtech, Zurich, 

2005. 

PTC (Positive Temperature Coefficient) thermistors – barium titanate 

A. Platzer, P. Supancic, C. Lembacher, U. Theiszi, and R. Danzer, “ Thermography and 

Simulation as a Combined Method for Failure Analysis of PTCs,” pp. 54 - 61 in 

Fractography of Advanced Ceramics II, eds. J. Dusza, R. Danzer and R. Morrell, 

Transtech, Zurich, 2005. 

R. Danzer, A. Platzer, P. Supoancic, Z. Wang, “Fractography of Thermistors,” pp. 231 

– 241 in Fractography of Glasses and Ceramics V, ,eds. J. R. Varner, G.D. Quinn, M. 

Wightman, Ceramic Transactions, Vol. 199, American Ceramic Society, Westerville, 

OH, 2007. 

B-11
 
 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

ZnO varistors 

Z. Wang, P. Supancic, F. Aldrian, A. Schriener and R Danzer, “Strength and 

Fractography of High Power Varistors,” pp. 358 - 365 in Fractography of Advanced 

Ceramics II, eds. J. Dusza, R. Danzer and R. Morrell, Transtech, Zurich, 2005. 

R. W. Rice, “Failure Analyses of Ceramics,” pp. 369 - 388 in Fractography of Glasses 

and Ceramics, IV, eds. J. R. Varner and G. D. Quinn, American Ceramic Society, 

Westerville, OH, 2001. 

Components suffer dielectric breakdown. 

Sold Oxide Fuel Cells 

J. Malzbender, R. W. Steinbrech, E. Wessel, “Brittle Fracture Studies of Solid Oxide 

Fuel Cells,” pp. 81-93 in Fractography of Advanced Ceramics III, eds. J. Dusza, R. 

Danzer, R. Morrell, and G. D. Quinn, Trans Tech Publ. Zurich, 2009. 

L. Blanco, S. Taylor, and K. Wiggins, “Unexpected Thermal Fracture of a Ceramic 

Sensor,” J. Fail. Anal. Prev., 11 (2011) 478-480. 

A ceramic substrate, to which an electronic infrared sensor was bonded, fractured 

when thermally cycled due to excessive force from a polymer pad used to clamp 

the assembly together. 

Bioceramics 

Ceramic hip joint ball heads 

H. G. Richter, “Fractography of Bioceramics,” pp. 157 - 180 in Fractography of 

Advanced Ceramics, ed. J. Dusza, Trans Tech Publ., Zurich, 2002. 

S. Hecht-Mijic and H. G. Richter, “Fractography of Bioceramic Components for Hip 

Joint Replacement,” pp 313 – 328 in Fractography of Glasses and Ceramics V, ,eds. J. 

R. Varner, G.D. Quinn, M. Wightman, Ceramic Transactions, Vol. 199, American 

Ceramic Society, Westerville, OH, 2007. 

R. Morrell, L. Byrne, and M. Murray, “Fractography of Ceramic Femoral Heads,” pp. 

253 – 266 in Fractography of Glasses and Ceramics, IV, eds. J. R. Varner and G. D. 

Quinn, American Ceramic Society, Westerville, OH, 2001. 

A. Walter, “Fracture Phenomena in Orthopedic Alumina,” pp. 403 – 414 in 

Fractography of Glasses and Ceramics, eds. J. R. Varner and V. D. Fréchette, 

Advances in Ceramics, Vol. 22, American Ceramic Society, Westerville, OH, 1988. 

R. Morrell, “Dealing with Component Failures,” pp. 353 -368 in Fractography of 

Glasses and Ceramics V, eds. J. R. Varner, G.D. Quinn, M. Wightman, Ceramic 

Transactions, Vol. 199, American Ceramic Society, Westerville, OH, 2007. 

Includes general observations of femoral head fractures, plus two specific alumina 

cases. One in-vivo ball fracture had an unusual, circular, shiny region of slow 

B-12
 
 

 

 

 

 

 

 

 

 

 

 

 

 

Appendix B Case Studies 

crack growth. The second case was an in vitro fracture that broke from cracks in 

the bore created by contact with a Co-Cr stem. 

R. Morrell, R. Danzer, I. Milošev, R. Trebše, “An Assessment of in vivo Failures of 

Alumina Ceramic Total Hip Joint Replacements,” J. Eur. Cer. Soc., 32 (2012) 3073-

3084. 

Four in vivo cases. Femoral heads may be damaged by the surgeon, or
 

contaminated during or prior to installation.
 

Dental Ceramics and Glass Ceramics 

J. R. Kelly, “Fractography of Dental Ceramics,” pp. 241 –251 in Fractography of 

Glasses and Ceramics, IV, eds. J. R. Varner and G. D. Quinn, American Ceramic 

Society, Westerville, OH, 2001. 

J. B. Quinn, “Failure Analysis of a Broken Tooth,” J. Fail. Anal. Prev., 4 [1] (2004) 41 

– 46. 

A patient’s natural tooth fractured days after open-heart surgery. Was the 

laryngoscopy by the anesthesiologist done correctly? 

S. S. Scherrer, J. B. Quinn, G. D. Quinn, and J. R. Kelly, “Failure Analysis of Ceramic 

Clinical Cases Using Qualitative Fractography,” Int. J. Prosthodont., 19 [2] (2006) 151 

– 158. 

Five dental crown restorations that failed in service were analyzed. The materials 

were: Procera alumina, Cerestore alumina-spinel, In-ceram glass in-fused alumina, 

and porcelain fused to metal. Fractography identified the failure causes. These 

were correlated to clinical observations and crown design. 

J. B. Quinn, G. D. Quinn, J. R. Kelly, and S. S. Scherrer, “Fractographic Analyses of 

Three Ceramic Whole Crown Restorations,” Dent. Mat., 21 (2005) 920 – 929. 

Three dental crown restorations that failed in service were analyzed. The materials 

were Procera alumina, Cerestore alumina-spinel, and Empress II lithium disilicate. 

Origin sites were found in each case. 

S. S. Scherrer, J. B. Quinn, G. D. Quinn, and H. W. A. Wiskott, “Fractographic 

Ceramic Failure Analysis Using the Replica Technique: Two Case Studies,” Dental 

Materials, 23 [11] (2007) 1397-1404. 

S. S. Scherrer, J. B. Quinn, G. D. Quinn, and H. W. A. Wiskott, “Descriptive 

Fractography on All Ceramic Dental Crown Failures,” pp. 339 – 352 in Fractography 

of Glasses and Ceramics, V, eds., J. R. Varner, G. D. Quinn, M. Wightman, Ceramic 

Transactions, Vol. 199, American Ceramic Society, Westerville, OH, 2007. 

S. S. Scherrer, G. D. Quinn, and J. B. Quinn, “Fractographic Failure Analysis of a Procera 

AllCeram Crown Using Stereo and Scanning Electron Microscopy,” Dent. Mater., 24 

(2008) 1107 – 1113. 

B-13
 
 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

J. B. Quinn, S. S. Scherrer, and G. D. Quinn, “The Increasing Role of Fractography in 

the Dental Community,” pp. 253 - 270, Fractography of Glasses and Ceramics, V, eds, 

J. R. Varner, G. D. Quinn, M. Wightman, Ceramic Transactions, Vol. 199, American 

Ceramic Society, Westerville, OH, 2007. 

J. B. Quinn, G. D. Quinn, J. R. Kelly, and S. S. Scherrer, “Useful Tools for Dental 

Failure Investigation,” pp. 36 - 53 in Proc. Conference on Scientific Insights and 

Dental Ceramics and Photopolymer Networks, ed. D. C. Starrett, Transactions of the 

Academy of Dental Materials, Vol. 16, 2004. 

S. S. Scherrer, J. B. Quinn, and G. D. Quinn, “Fractography of Dental Restorations,” in 

Fractography of Advanced Ceramics, III, ed. J. Dusza, R. Danzer, R. Morrell and G. 

Quinn, TransTech Publ., Zurich, 2009. Key Engineering Materials, 413 (2009) 72-80. 

B. Taskonak, J. J. Mecholsky, Jr., and K. J. Anusavice, “Fracture Surface Analysis of 

Clinically Failed Fixed Partial Dentures,” J. Dent. Res., 3 (2006) 277 – 281. 

J. Y. Thompson, K. J. Anusavice, A. Naman, and. H. F. Morris, “Fracture Surface 

Characterization of Clinically Failed All-Ceramics Crowns, J. Dent. Res., 73 [12] 

(1994) 1824 - 1832. 

Ten Dicor crown were fractographically analyzed in accordance with MIL HDBK 

790 and the principles outline by Fréchette. Fracture initiated on the inside surfaces 

from abrasive damage or failure of the cement at the restoration/cement interface. 

Twelve Cerestore crowns were more difficult and appeared to break from the 

porcelain/core interface or inside the core material. Biaxial disk specimens 

furnished valuable property and fractographic information. 

G. D. Quinn, K. Hoffman, S. Scherrer, U. Lohbauer, G. Amberger, M. Karl, J. R. 

Kelly, “Fractographic Analysis of Broken Ceramic Dental Restorations, pp. 161-174 in 

Fractography of Glasses and Ceramics, VI, eds. J. R. Varner, M. Wightman, Ceramic 

Transactions, Vol. 230, Wiley, New York, 2012. 

Three case studies: an alumina 3-unit bridge; a 6-unit zirconia bridge; a 5-unit 

zirconia telescoping-denture. 

G. D. Quinn, “Fractographic Analysis of Broken Ceramic Dental Restorations,” 

Ceramic Engineering and Science Proceedings, 35 [1] (2015) 39-51. 

S. Hecht-Mijic, “Fractography of Ceramic Dental Implants,” pp 175 – 191 in 

Fractography of Glasses and Ceramics, VI, eds. J. R. Varner, M. Wightman, Ceramic 

Transactions, Vol. 230, Wiley, New York, 2012. 

Fractures in ceramic dental implants. 

M. Borba, H. N. Yoshimura, J. A. Griggs, P. F. Cesar, A. Della Bona, “Qualitative and 

Quantitative Fractographic Analysis of All-Ceramic Fixed Partial Dentures,” pp. 205 – 

212 in in Fractography of Glasses and Ceramics, VI, eds. J. R. Varner, M. Wightman, 

Ceramic Transactions, Vol. 230, Wiley, New York, 2012. 

B-14
 
 

 

 

 

 

 

 

 

 

Appendix B Case Studies 

U. Lohbauer, R. Belli, G. Arnetzl, S. Scherrer, and G. D. Quinn, “Fracture of a VeneeredZrO2 Dental Prosthesis from an Inner Thermal Crack,” Case Studies in Engineering 

Fracture Analysis, 2 (2014) 100-106. 

U. Lohbauer, G. Amberger, G. D. Quinn, S. S. Scherrer, “Fractographic Analysis of a 

Dental Zirconia Framework: A Case Study on Design Issues, J. Mech. Behav. Biom. 

Mat., 3 (2010) 623-629. 

M. Øilo, A. K. Kvam, J. E. Tibballs, and N. Gjerdet, “Clinically Relevant Fracture 

Testing of All-Ceramic Crowns,” Dent. Mater., 29 (2013) 815-823. 

M. Øilo, N. R. Gjerdet, “Fractographic Analysis of All-ceramic Crowns: A Study of 27 

Clinically-fractured Crowns,” Dent. Mater., 29 (2013) e78-e84. 

M. Øilo, G. D. Quinn, “Fracture Origins in Twenty-two Dental Alumina Crowns,” J. 

Mech. Beh. Biomed. Mater., 53 (2016) 93 – 103. 

Whitewares and Kitchenware 

J. S. Banda and P. F. Messer, “Fracture-Initiating Flaws in Whitewares,” pp. 363 - 375 

in Fractography of Glasses and Ceramics, eds. J. D. Varner and V. D. Fréchette, 

Advances in Ceramics, Vol. 22, American Ceramic Society, Westerville, OH, 1988. 

P. B. Adams and S. E. DeMartino, “Glass-Ceramic Cookware Failure Analysis,” pp. 

669 – 673 in Ceramics and Glasses, Engineered Materials Handbook, Vol. 4, ed. S. 

Schneider, ASM, Metals Park, OH, 1991. 

V. D. Fréchette, Failure Analysis of Brittle Materials, Advances in Ceramics, Vol. 28, 

American Ceramic Society, Westerville, OH, 1990, pp. 113 - 114. 

Teapots on stoves – thermal stresses. 

V. D. Fréchette, Failure Analysis of Brittle Materials, Advances in Ceramics, Vol. 28, 

American Ceramic Society, Westerville, OH, 1990, pp. 114 - 115. 

Lasagna dish, cone crack – impact origin crack, subsequently propagated by 

thermal stresses. 

J. R. Varner and V. D. Frechette, “Fractography of Whitewares,” pp. 305 – 315 in 

Science of Whitewares, eds. V. Henkes, G. Onoda, and W. Carty, American Ceramic 

Society, Westerville, OH, 1996. 

Porcelain Coffee Storage Jars – contact damage from a metal band cover seal 

Porcelain Bushing, 2.7 m long by 0.9 m diameter – residual stresses from firing 

Porcelain Bushing, small - screw holes acted as origins 

V. D. Fréchette, Failure Analysis of Brittle Materials, Advances in Ceramics, Vol. 28, 

American Ceramic Society, Westerville, OH, 1990, p. 113. 

Porcelain faucet handles 

B-15
 
 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

M. D. Hayes and D. L. Ahearn, “Thermomechanical Analysis of a Ceramic Cooker,” J. 

Fail. Anal. Prev., 13 (2013) 383-388. 

Escape of hot gasses through a crack in a ceramic cooker/grill caused a house fire. 

The cause of the crack was not determined. 

R.C. Bradt and R. L. Martens, “Shattering Glassware,” Bull. Amer. Ceram. Soc., 91 [7] 

(2012) 33-38. 

Consumerwares. Thermal shock fractures of heat-strengthened soda lime silica 

versus borosilicate glass. 

Porcelain Toilet Bowls 

D. L. Ahearn, “Residual Failure Due to Incomplete Sintering of Vitreous China 

Plumbing Fixtures, pp 273 – 279 in Fractography of Glasses and Ceramics, VI, eds. J. 

R. Varner, M. Wightman, Ceramic Transactions, Vol. 230, Wiley, New York, 2012. 

J. R. Varner and V. D. Frechette, “Fractography of Whitewares,” pp. 305 – 315 in 

Science of Whitewares, eds. V. Henkes, G. Onoda, and W. Carty, American Ceramic 

Society, Westerville, OH, 1996. 

Toilet Bowl – water hammer
 

Toilet Tank – mounting-bolt hole failure from mechanical overload
 

Other Ceramic Components 

Porcelain faucet handles 

V. D. Fréchette, Failure Analysis of Brittle Materials, Advances in Ceramics, Vol. 28, 

American Ceramic Society, Westerville, OH, 1990, p. 113. 

Earthenware sewer pipe 

V. D. Fréchette, Failure Analysis of Brittle Materials, Advances in Ceramics, Vol. 28, 

American Ceramic Society, Westerville, OH, 1990, p. 116. 

Uranium dioxide fuel element – bushing 

V. D. Fréchette, Failure Analysis of Brittle Materials, Advances in Ceramics, Vol. 28, 

American Ceramic Society, Westerville, OH, 1990, pp. 120-121. 

Thermal fracture. 

Beryllium oxide hollow cylinder 

J. R. Varner and V. D. Frechette, “Fractography of Whitewares,” pp. 305 – 315 in 

Science of Whitewares, eds. V. Henkes, G. Onoda, and W. Carty, American Ceramic 

Society, Westerville, OH, 1996. 

Hollow cylinder loaded at high temperature under laboratory conditions. 

B-16
 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Appendix B Case Studies 

Ceramic machine base (Electrically insulating base) 

R. Morrell, “Fractography of Brittle Materials,” R. Morrell, Measurement Good 

Practice Guide 15, National Physical Laboratory, Teddington, Middlesex, United 

Kingdom, 1999, pp. 67 – 69. 

Bending failure from bolt holes. 

Ceramic ball valve – alumina 

R. Morrell, “Fractography of Brittle Materials,” R. Morrell, Measurement Good 

Practice Guide 15, National Physical Laboratory, Teddington, Middlesex, United 

Kingdom, 1999, pp. 70- 71. 

150 mm alumina ball with slots and large hole, internal pressure. 

Gas valve plate - alumina 

R. Morrell, “Fractography of Brittle Materials,” R. Morrell, Measurement Good 

Practice Guide 15, National Physical Laboratory, Teddington, Middlesex, United 

Kingdom, 1999, pp. 72 -74. 

Possible grinding or lapping damage near a complex hole. 

Steel continuous casting nozzles – Carbon-bonded alumina, zirconia, magnesia, 

graphite. 

A. Mašelejová and A. Leško, “Fracture Characteristics of Isostatic Pressed Submerged 

Entry Nozzle,” pp. 370 – 373 in Fractography of Advanced Ceramics II, eds. J. Dusza, 

R. Danzer and R. Morrell, Transtech, Zurich, 2005. 

Fibers - Glass - Optical 

H. C. Chandan, R. D. Parker, and D. Kalish, “Fractography of Optical Fibers,” pp. 143 

– 184 in Fractography of Glass, Ed. R. C. Bradt and R. E. Tressler, Plenum Press, NY, 

1994. 

W. R. Wagner, “Failure Analysis of Fiber Optic Connectors,” pp. 389 – 402 in 

Fractography of Glasses and Ceramics, Advances in Ceramics, Vol. 22, ed. J. D. 

Varner and V. D. Fréchette, American Ceramic Society, Westerville, OH, 1988. 

J. J. Mecholsky, “Fracture Surface Analysis of Optical Fibers,” pp. 663 – 668 in 

Ceramics and Glasses, Engineered Materials Handbook, Vol. 4, ed. S. Schneider, 

ASM, Metals Park, OH, 1991. 

P. G. Simpkins and J. T. Krause, “Dynamic Response of Glass Fibers During Tensile 

Fracture,” Proc. R. Soc. London A, 350 (1976) 253-265. 

High-speed photos of strong (> 2.8 GPa) glass fibers tested in tension reveal that 

secondary fractures occur as the result of fiber whipping and flexural and 

longitudinal wave interactions. Fibers disintegrate into a particulate cloud. 

B-17
 
 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Fibers – Nonoptical
	

P. K. Gupta, “Fractography of Fiberglass,” pp. 185 – 206 in Fractography of Glass, Ed. 

R. C. Bradt and R. E. Tressler, Plenum Press, NY, 1994. 

G. V. Srinivasan and V. Venkateswaren, “Fractographic Investigation of Flaws in 

Sintered SiC Fiber,” pp 317 –337 in Fractography of Glasses and Ceramics, III, eds. J. 

R. Varner, V. D. Fréchette, and G. D. Quinn, Ceramic Transactions, Vol. 64, American 

Ceramic Society, Westerville, OH, 1996. 

C-T. Li and N. R. Langley, “Development of a Fractographic Method for the Study of 

High-Temperature Failure of Ceramic Fibers,” pp. 177 – 184 in Fractography of 

Glasses and Ceramics, eds. J. D. Varner and V. D. Fréchette, Advances in Ceramics, 

Vol. 22, American Ceramic Society, Westerville, OH, 1988. 

T. Clark, “Fracture Properties of Thermally Aged Ceramic Fiber Produced by Polymer 

Pyrolysis,” pp. 279 – 305 in Fractography of Glasses and Ceramics, eds. J. D. Varner 

and V. D. Fréchette, Advances in Ceramics, Vol. 22, American Ceramic Society, 

Westerville, OH, 1988. 

Fiber Composites 

Silicon carbide fibers in a glass matrix 

A. S. Fareed, M. J. Koczak, F. Ko, and G. Layden, “Fracture of SiC/LAS Ceramic 

Composites,” pp. 261 – 293 in Fractography of Glasses and Ceramics, eds. J. D. 

Varner and V. D. Fréchette, Advances in Ceramics, Vol. 22, American Ceramic 

Society, Westerville, OH, 1988. 

Glass fibers in a fiberglass-reinforced plastic pressure vessel 

S. J. Glass, E. K. Beauchamp, M. J. Carr, T. R. Guess, S. L. Munroe, R. J. Moore, A. 

Slavin, and N. R. Sorenson, “Failure Analysis of A Fiberglass-Reinforced Plastic 

Pressure Vessel,” pp. 527 – 541 in Fractography of Glasses and Ceramics, III, eds. J. 

R. Varner, V. D. Fréchette, and G. D. Quinn, Ceramic Transactions, Vol. 64, American 

Ceramic Society, Westerville, OH, 1996. 

Violent rupture of a pressurized tank holding sulphuric acid. Occasional and 

incidental spillage penetrated microcracks in the matrix and caused progressive 

degradation in glass fibers. 

Single crystal ceramic components 

Sapphire secondary refractor lens for spacecraft 

J. A. Salem and G. D. Quinn, “Failure Analysis of Sapphire Refractive Secondary 

Concentrators,” NASA Technical Report: TM-2009-215802, NASA Glenn Research 

Center, Cleveland, OH, Dec. 2009. Also in this Guide, chapter 10. 

B-18
 
 

 

 

 

 

 

 

 

 

 

Appendix B Case Studies 

J. A. Salem, G. D. Quinn, “Fractographic Analysis of Large Single Crystal Sapphire 

Refractive Secondary Concentrators,” second keynote presentation and paper, 4th 

International Conference on Fractography of Advanced Ceramics,” Smolenice Castle, 

Oct. 2013, J. Eur. Cer. Soc., 34 (2014) 3271-3281. 

Large single crystal lens designed to concentrate sunlight in space. Fracture 

occurred from thermal stresses that propagated cracks from polishing and impact 

flaws from handling. Although strong and hard, sapphire is vulnerable to surface 

damage. Also this Guide, chapter 10. 

Sapphire Tubes, Plasma Asher 

This Guide, chapter 8. 

Calcium Fluoride (CaF2) windows in a laser flash diffusivity apparatus 

R. Morrell, “Dealing with Component Failures,” pp. 353 -368 in Fractography of 

Glasses and Ceramics V, eds. J. R. Varner, G.D. Quinn, M. Wightman, Ceramic 

Transactions, Vol. 199, American Ceramic Society, Westerville, OH, 2007. 

48 mm diameter single crystal disks that broke from thermal stresses activating 

edge damage or impact damage. 

Coatings and Glazes 

Diamond coatings 

H. A. Hoff, K. A. Snail, A. A. Morrish, and J. E. Butler, “Fractography and Fracture 

Mechanics of Combustion Grown Diamond Thin Films,” pp. 25 – 54 in Fractography 

of Glasses and Ceramics II, eds. V. D. Fréchette and J. R. Varner, Ceramic 

Transactions, Vol. 17, American Ceramic Society, Westerville, OH, 1991. 

Whiteware glazes 

V. D. Fréchette, Failure Analysis of Brittle Materials, Advances in Ceramics, Vol. 28, 

American Ceramic Society, Westerville, OH, 1990, pp. 94 – 101. 

Internal pressure, water hammer, thermal shock, and impact fractures. 

Stoneware reactor catalyst media balls 

R. W. Rice, “Failure Analysis of Ceramics,” pp. 369 – 388 in Fractography of Glasses 

and Ceramics, IV, eds. J. R. Varner and G. D. Quinn, American Ceramic Society, 

Westerville, OH, 2001. 

Gemstones and Jade 

Emerald gemstone (3.66 carat) 

R. W. Rice, “ Failure Analysis of Ceramics,” pp. 369 – 388 in Fractography of Glasses 

and Ceramics, IV, eds. J. R. Varner and G. D. Quinn, American Ceramic Society, 

Westerville, OH, 2001. 

Impact on ceramic kitchen tile countertop caused cracking. 

B-19
 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Jade 

D. J. Rowcliffe and V. Frühauf, “The Fracture of Jade,” J. Mat. Sci., 12 (1977) 35 – 42. 

Gemstones 

J. E. Field, “Brittle Fracture: Its Study and Application,” Contemp. Phys., 12 [1] (1971) 

1 - 31. 

Rocks, Lithic Materials 

J. B. Quinn, J. W. Hatch, and R. C. Bradt, “The Edge Flaking Test as an Assessment of 

the Thermal Alteration of Lithic Materials, Bald Eagle Jasper,” pp. 73 – 85 in 

Fractography of Glasses and Ceramics, IV, eds. J. R. Varner and G. D. Quinn, 

American Ceramic Society, Westerville, OH, 2001. 

A. Tsirk, “An Exploration of Liquid-Induced Fracture Markings,” pp. 87 – 101 in 

Fractography of Glasses and Ceramics, IV, eds. J. R. Varner and G. D. Quinn, 

American Ceramic Society, Westerville, OH, 2001. 

Obsidian tools. 

A. Tsirk, “Fractographic Evidence for Liquid on Obsidian Tools,” J. Archeol. Sci., 27 

(2000) 987 – 991. 

A. Tsirk, “Formation and Utility of a Class of Anomalous Wallner Lines on Obsidian,” 

pp. 57 - 69 in Fractography of Glasses and Ceramics, eds., J. D. Varner and V. D. 

Fréchette, Advances in Ceramics, Vol. 22, Amer. Ceram. Society, Westerville, OH, 

1988. 

A. Tsirk, “Fractures in Knapping,” Archaeopress, Oxford, UK, 2014. 

A comprehensive compilation of the author’s lifelong work in fractures of rocks 

and glasses. This book was published very shortly before the author’s death. Well 

illustrated and with links to modern fracture mechanics. 

Geological, Tectonic Structures 

D. Bahat, Tectono-fractography, Springer-Verlag, Berlin, 1991. 

Bahat’s book is a good starting point, but readers should be warned that some of 

the interpretations are controversial and are not settled. 

B. R. Kulander, “Hackle Plume Geometry and Joint Propagation Dynamics,” pp. 85 – 

94 in Proceedings of the International Symposium on Fundamentals of Rock Joints, 

Björkilden, 15-20 September 1985. 

O. H. Muller, “Fractography Applied to Large Scale (m to km) Cracks in the Earth,” 

pp. 427 – 438 in Fractography of Glasses and Ceramics, eds. J. D. Varner and V. D. 

Fréchette, Advances in Ceramics, Vol. 22, American Ceramic Society, Westerville, 

OH, 1988. 

Similar features observed over nine orders of magnitude of size. 

B-20
 
 

 

 

 

 

 

 

 

Appendix B Case Studies 

A. I. Younes and T. Engelder, “Fringe Cracks: Key Structures for the Interpretation of 

the Progressive Alleghanian Deformation of the Appalachian Plateau,” Geolog. Soc. of 

Amer. Bull., 111 [2] (1999) 219 – 239. 

D. Bahat, T. Bankwitz, and E. Bankwitz, “Preuplift Joints in Granites: Evidence for 

Subcritical and Postcritical Fracture Growth,” Geolog. Soc. of Amer. Bull., 115 [2] 

(2003) 148 – 165. 

D. T. McConaughy and T. Engelder, “Joint Initiation in Bedded Clastic Rocks,” J. 

Struct. Geol., 23 (2001) 203 – 221. 

T. Engelder, “Tectonic Implication Drawn from Differences in the Surface Morphology 

on Two Joint Sets in the Appalachian Valley and Ridge, Virginia,” Geology, 32 (2004) 

413 - 416. 

T. Engeleder, “Propagation Velocity of Joints: A Debate Over Stable vs. Unstable 

Growth of Cracks in the Earth,” pp.457-482 in Fractography of Glasses and Ceramics 

V, eds. G. D. Quinn, M. Wightman, and J. R. Varner, Ceram. Trans., Vol. 199, 

American Ceramic Society, Westerville, OH, 2007. 

B-21
 

 

 

 

 

 

 

 

 

 

 

 

 

 

Appendix C Mirror Constants 

APPENDIX C FRACTURE MIRROR AND BRANCH CONSTANTS 

The following table includes all the published values of mirror and branching 

constants that the author could find up to late 2015 when the second edition of 

this Guide was in preparation. The original reference is listed for each entry. 

The mirror and branching constants were analyzed assuming: 

σ R = A	 (C.1) 

Mirror-mist (or inner mirror) Ai and mist-hackle (or outer mirror) Ao constants 

are listed for glass and glass ceramics. Mist is difficult or impossible to discern 

in most ceramics and the mirror constant Ao that is listed corresponds to the 

mirror-hackle boundary. 

Entries where a modified version of equation C.1 was used are noted by a 

footnote. 

No judgment is made of the veracity of the listed values. Some of the values 

may be inaccurate. Even for ideal materials such as fused silica and soda lime 

silica there are unsettling variations in reported values. The variations 

probably are caused by: 

1.	 Differences in viewing mode, including type of microscope, magnifycations, illumination techniques, and magnification accuracy and precision. 

2. Differences in observer judgment. 

3.	 Differences in the radii measured (e.g., along the surface, into the depth, or 

other) and the accuracy and precision of the length measurements. 

4.	 Differences in test method (e.g., flexure, tension, biaxial plate) and 

accuracy and precision of the strength data. 

5.	 Differences in whether or not corrections were made for stress gradients 

and the stress at the origin location. 

6.	 Unaccounted for residual stresses. 

7.	 Differences in the analysis used including whether data were plotted and 

regressed as log stress versus log mirror radius, or linear stress versus 

inverse square root of radius. 

8.	 Variations in density, microstructure, or fracture toughness between 

nominally identical materials. 

9. Variations in the number of specimens tested. 

One is tempted to conclude that most of the discrepancies are due to viewer 

judgment, but this cannot account for the all of the variability. The mirror 

C-1 
 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

constant is only sensitive to the square root of the measured radii. In other 

words, a 10% systematic error in radii measurements causes only a 5% error in 

the mirror constant. 

All values in the table are listed with the same number of significant figures as 

in the original reference, but only up to a limit of three significant figures. 

Uncertainties (± one standard deviation) are listed when available from the 

original reference. Multiple entries in a cell denote estimates by different 

microscopy techniques or analysis. For polycrystalline ceramics, the mirror 

constants taken from the reference sources are assumed to be for the misthackle boundary unless otherwise stated. 

If the user cannot find the mirror constant for a particular material, he or she 

may consider values for comparable materials within the same class. For 

example, if a particular glass is not listed, then use the value for a glass of 

similar composition. For polycrystalline ceramics, microstructure plays an 

important role in determining a mirror constant. Conscientious readers should 

check the references listed for more details on a particular material. 

As discussed in chapter 7, the fracture mirror constants do not vary 

significantly with temperature or environment. Any variability with 

temperature usually follows trends in fracture surface energy, γf, or fracture 

toughness, KIc. 

It is hoped that the adoption of the Guidelines in Appendix D, and the adoption 

of ASTM C 1678, Standard Practice for Fractographic Analysis of Fracture 

Mirror Sizes in Ceramics and Glasses in 2007, will lead to improved 

consistency and accuracy in future mirror size constant determinations. 

. 

C-2 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Appendix C Mirror Constants 

Material Technique 

Mirror-Mist 

Ai 

(MPa· m ) 

Mist-Hackle or 

Mirror-Hackle 

Ao 

(MPa· m ) 

Branching 

Ab 

(MPa· 

m ) 

Reference 

Glasses 

Flint (Kimble R6 soda 

lime) Flexure (Rods) 2.0 Kirchner, 8 

Flint (Kimble R6 soda 

lime) Flexure (Rods) 1.9 Kirchner, 9 

Flint (Kimble R6 soda 

lime) Flexure (Rods) 2.3 Kirchner, 10 

Soda-Lime Silicate – 

window glass 

Flexure (biaxial ring-on 

ring, large) 2.09 Orr, 5 

Soda-Lime Silica 

(window, sheet, and 

plate) 

Flexure (Bars or laths) 2.05 Shand, 4 

Soda-Lime Silicate – 

float, plate, sheet 

Flexure (biaxial ring-onring and bend bars) 2.09 Barsom, 62 

Soda-Lime Silicate – 

window glass 

Pressurized windows, 

large 1.96 Reed, 39 

Soda-Lime Silicate Flexure (Bars, 3-point) 1.74 Bansal, 14 

Soda-Lime Silicate A 

– plate glass 

Soda-Lime Silicate B 

– plate glass 

Flexure (Bars – large) 

Room Temperature to 

Strain Point 

1.86 ± 0.66 

1.82 ± 0.91 

Kerper, Scuderi 

6 

Soda-Lime Silica 

Float Flexure (Bars) 1.80 ± 0.15 * 

1.81 ± 0.25 

2.42 ± 0.16 * 

2.29 ± 0.24 Ball, 40 

Soda-Lime Silica 

Float (G.E.C. – X8) 

Tension (Rods) 

Flexure (Bars) 

1.89 ± 0.06 2.04 ± 0.06 

2.09 

Johnson, 

Holloway, 1 

Soda-Lime Silica 

Float 

Flexure (Bars) 

Flexure–Delayed failure 

(Bars) 

1.92 

2.0 ± 0.1 

2.21 

2.2 ± 0.1 

Mecholsky, 

18,21 

Mecholsky, 21 

Soda-Lime Silica Flexure 1.8 2.0 2.3 Mecholsky, 22 

Soda Lime Silica Float 

Flexure (Bars, large and 

small, 3-point and 4-

point) 

2.06 ± 0.07 2.29 ± 0.19 Duckworth, 16 

Soda-Lime Silica Tension (Plates) 1.2 – 1.6 Congleton, 

Petch, 2 

Soda-Lime Silica 

Flexure (Bars) 

Flexure (Biaxial ring-onring plates) 1.81 ± 0.28 

3.54 ±

0.64 Choi, 33 

Soda-Lime Silica Tension 1.9 Clark, 34 

Soda-Lime Silica Pressurized Tube 2.0 Aoki, 35 

Soda-Lime Silica 

Flexure (Biaxial ring-onring disks) 

3 environments, 

Vickers indented 

1.82 - 1.94 2.03 – 2.13 2.28 – 

2.42 Marshall, 19 

Soda-Lime Silica Flexure (Biaxial, 

pressure on ring disks) 

2.1 ± 0.1 NZ (10 ± 4) Shetty, 15 

Soda-Lime Silica Flexure (rods) 1.8 Terao, 63 

Soda-Lime Silica Flexure (rods) 

1.85 ± 0.21 

3.20 ⊕

2.82 ⊕

Varner, 68, 70 

Soda-Lime Silica 

Flexure (3-point rods) 

Flexure (4-point rods) 

Tension (Fibers) 

3.15 ± 0.11 

2.74 ± 0.15 

1.72 ± 0.28 

Abdel-Latif, 49 

Soda-Lime Silica Flexure (3-point, laths) (2.9 ***) Levengood, 64 

Soda-Lime Silica Flexure, plates, 2 sizes 1.47 NZ (19.5) 

1.92 NZ (7.0) 

2.88 NZ (18.) 

3.1 NZ (5.7) 

Gaume. 73 

C-3 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Material Technique 

Mirror-Mist 

Ai 

(MPa· m ) 

Mist-Hackle or 

Mirror-Hackle 

Ao 

(MPa· m ) 

Branching 

Ab 

(MPa· 

m ) 

Reference 

Soda-Lime Silica Flexure (4-point, laths) 

1.80 ± 0.07 

(air, 20oC) 

2.08 ± 0.04 

(air, -150OC) 

1.89 ± 0.05 

(water, 17oC) 

1.92 ± 0.04 

(dry N2,24oC) 

Shinkai, 69 

Borosilicate A 

(P 3235) 

Borosilicate B 

(C 7740) 

Flexure (Bars – large; 

Room Temperature to 

Strain Point) 

1.98 ± 0.46 

2.04 ± 0.75 

Kerper, Scuderi, 

6 

Borosilicate 

(C 7740) 

Flexure (Rods, many 

diameters) 2.08 ± 0.02 Kerper, Scuderi, 7 

Borosilicate 

(C 7740) Flexure (Bars) 1.87 ± 0.3 2.10 Mecholsky, 17,18 

Borosilicate 

(C 7740) 

Flexure (Bars) and 

Biaxial disks 1.9 ± 0.3 Mecholsky, 20 

Borosilicate 

(C 7740) Flexure (Rods) 2.2, 2.35 Shand, 3 

Borosilicate 

Pyrex Flexure (4-point, laths) 

1.98 ± 0.09 

(air, 20oC) 

2.08 ± 0.11 

(air, -150OC) 

2.02 ± 0.08 

(water, 17oC) 

1.98 ± 0.12 

(dry N2,24oC) 

Shinkai, 69 

Borosilicate Flexure and Tension 

(Rods) 1.9* Alarcón et al. 56 

Borosilicate crown 

(Schott BK-7) 

Flexure (Biaxial, ring on 

ring disks)* 

1.98 ± 0.02 NZ 2.11 ±0.03 NZ 

2.3 

2.28 ±0.03 NZ Quinn, 25 

Aluminosilicate 

(C 1723) Flexure (Bars) 2.14 2.40 Mecholsky, 18 

Aluminosilicate A 

(P 6695) 

Aluminosilicate B 

(C 1723) 

Flexure (Bars –large) 

Room Temperature to 

Strain Point 

2.31 ± 0.76 

2.34 ± 0.97 

Kerper, Scuderi, 

6 

Alkali-borosilicate Tension (Fibers) 1.33 * Jaras et al., 12 

Alkaline-earth 

boro aluminosilicate 

(C1737) 

Flexure (Biaxial ring-onring) 2.07 ± 0.01 Gulati et al., 57 

Barium silicate 

3BaO-5SiO2 NR (Flexure Bars?) 1.3 1.5 Mecholsky, 22 

E glass 

CaO-Al2O3-B2O3 

alumina borosilicate 

Tension (Fibers) 1.47 * Jaras et al., 12 

Lead silicate 

(G.E.C. L1) Tension (Rods) 1.71 ± 0.06 Johnson, Holloway, 1 

Lead silicate Flexure (Bars) 1.61 1.78 Mecholsky, 18 

Leaded silicate, 

6.7 g/cm3 Flexure (Rods) 1.4 ± 0.1 1.6 ± (0,1) Salem et al., 77 

Lithium silicate 

Li2O-2SiO2 

Flexure Bars (3-point) 2.2 2.6 Mecholsky, 22 

Zinc silicate Flexure Bars 2.1 2.6 3.2 Mecholsky, 22 

Zirconia silicate 

(Cem-FIL AR) Tension (Fibers) 2.37 * Jaras et al., 12 

Yttrium-Alumino-Silica 

Oxyntride 

(2Y-Al-Si-O-N) 

Flexure (Bars) 1.5 ± 0.2 Coon, 48 

Fused Silica 

(C 7940) Flexure (Bars) 2.23 2.42 2.7 Mecholsky, 18, 22 

C-4
 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Appendix C Mirror Constants 

Material Technique 

Mirror-Mist 

Ai 

(MPa· m ) 

Mist-Hackle or 

Mirror-Hackle 

Ao 

(MPa· m ) 

Branching 

Ab 

(MPa· 

m ) 

Reference 

Fused Silica 

(C 7940) Flexure (Bars – large; 

Room Temperature to 

Strain Point) 

1.89 ± 0.51 Kerper, Scuderi, 

6 

Fused silica 

(Vitreosil) Tension (Rods) 2.33 ± 0.06 Johnson, Holloway, 1 

Fused silica Flexure (Rods) 2.20 ± 0.33 Choi, 33 

Fused silica Flexure and Tension 

(Rods) 1.2 NZ Alarcón et al. 56 

Fused silica fibers Tension 2.10 Choi, 33 

Fused silica clad 

fibers Tension 1.96 ± 0.13 Baker, 29 

Fused silica fibers Tension 1.83 Castilone, 30 

Fused silica fibers, 

bars, disks 

Tension (Fibers) 

Flexure (Bars) 

Flexure (Biaxial, piston 

on 3 balls) 

2.2 ± 0.5 

2.3 ± 0.5 

2.4 ± 0.3 

Mecholsky, 20 

Fused silica fibers Tension 2.224 Chandan, 37 

Leached High Silica 

(C 7930) Flexure (Bars) 0.91 1.19 Mecholsky, 18 

96% Silica 

(C 7900) 

Flexure (Bars – large; 

Room Temperature to 

Strain Point) 

1.84 ± 0.65 Kerper, Scuderi, 

6 

96% Silica 

Vycor Flexure (4-point, laths) 

1.77 ± 0.06 

(air, 20oC) 

2.14 ± 0.06 

(air, -150OC) 

1.80 ± 0.05 

(water, 17oC) 

1.83 ± 0.07 

(dry N2, 24oC) 

Shinkai, 69 

Glassy Carbon Flexure (Bars) 1.17 1.67 Mecholsky, 

17,18 

Glassy Carbon Flexure (Bars) 2.1 Bullock, Kaae, 

61 

Chalcogenide 

As-S-Se and 

Ge-AS-Se-Te 

Flexure (Fibers, 2 point 

bending) 

1.58 

1.77 ± 0.14 NZ 

( - 66.4 ± 45.5) 

Quinn et al., 75 

As2Se3 chalcogenide 

glass, untreated 

As2Se3 chalcogenide 

glass, UV irradiated 

Tension (Fibers) 

0.69 

0.35 irradiated 

0.77 

0.38 irradiated 

Hulderman, 45 

As2S3 Flexure (Bars) 0.56 0.65 Mecholsky, 18 

Ge33As12Se55 Flexure (Bars) 0.55 0.65 Mecholsky, 18 

0.3PbSe -

0.7Ge1.5As0.5Se3 

Flexure (Bars) 0.48 0.55 Mecholsky, 18 

ZBLA (halide glass) Flexure (Bars) 0.8 Mecholsky, 22 

Lead Borate glass 

30PbO-70B2O3 (mol) 

40PbO-60B2O3 

50PbO-50B2O3 

60PbO-40B2O3 

70PbO-30B2O3 

Flexure (Bars) 

1.7 ± .05 

1.45 ± .05 

1.15 ± .05 

0.85 ± .01 

0.65 ±.01 

Shinkai, 

Ishikawa, 60 ** 

Glass Ceramics 

Pyroceram 9608 

(Li, Mg, Al silicate) NR 2.8 Adams + DeMartino, 27 

Pyroceram 9607 

(Li, Mg, Zn, Al silicate) NR 2.1 NZ Adams + DeMartino, 27 

Pyroceram 9606 

(Cordierite, Mg, Al 

silicate) 

Flexure 3.6 6.5 Mecholsky, 17 

C-5 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Material Technique 

Mirror-Mist 

Ai 

(MPa· m ) 

Mist-Hackle or 

Mirror-Hackle 

Ao 

(MPa· m ) 

Branching 

Ab 

(MPa· 

m ) 

Reference 

Pyroceram 9606 

(Cordierite, Mg, Al 

silicate) Flexure (Bars) 6.5 Lewis, 44 

Pyroceram 9606 

(Cordierite, Mg, Al 

silicate) 

Flexure (Bars, 3-point 

and 4-point) 5.7 Bansal, 13,14 

Pyroceram 9606 

(Cordierite, Mg, Al 

silicate) 

Flexure (Bars) and 

Flexure (Biaxial, piston 

on 3 balls) 

6.3 Mecholsky, 20 

Pyroceram 9606 

(Cordierite, Mg, Al 

silicate) 

Flexure (Biaxial, 

pressure on ring, and 

ball on ring) 

3.1 ± 0.2 NZ (84 ± 6) Shetty, 15 

Pyroceram 9606 

(Cordierite, Mg, Al 

silicate) 

Flexure (Rods) 4.8 Shand, 3 

Li2O-SiO2 (NPL glass 

ceramic, 2 grades) Flexure (Bars) 3.3, 3.8 4.5, 5.4 Mecholsky, 17 

Dental Ceramics and Filled Composites 

Dicor (dental, tetra 

silica fluoromica) 

glass ceramic 

Flexure (Bars) 0.97 Kelly, 47 

Feldspathic Porcelain 

(alumina filled, 

Vitadur N 338) 

Flexure (Bars) 2.82 Kelly, 47 

Feldspathic Leucite 

Porcelain, Optec 

OPC, 

Flexure (Bars) 2.1 Fischer, 52 

Empress I Leucite 

glass ceramic 

Flexure (Bars) 1.7 Fischer, 52 

Empress II Lithium 

disilicate glass 

ceramic 

Flexure (Bars) 

3.9 Fischer, 52 

Omega - Opaker Flexure (Bars) 1.3 Fischer, 52 

Cerec Mark II 

Porcelain 

Flexure (Bars) 1.6 Fischer, 52 

Glass infused 

Alumina, In-Ceram 

Flexure (Bars) 6.6 Fischer, 52 

Dental resin, 85 wt% 

zirconia-silica filler in 

bisGMA-TEGDMA 

Paradigm MZ100 

Flexure (Bars, 4-point) 2.6 Quinn, 59 

3Y-ZTP, Lava Flexure (Bars, 4-point) 10.7 Quinn, 58 

Alumina 

β-Al2O3 Flexure ~ 6.5 Mecholsky, 17 

Hot-pressed 

(1 μm, Cer. Fin.)) 

Flexure (Rods, 4-point) 

25 C to 1400 C 10.4 Kirchner, 11 

Hot-pressed 

(99+% pure, Cer. Fin.) 

Flexure (Rods, 4-point) 

Flexure-Delayed 

Fracture (Rods) 

10.3 

9.9 

Kirchner, 8,9 

Kirchner, 8 

Hot-pressed 

(99+% pure,Cer. Fin.) Flexure (Rods-3-point) 9.1 Kirchner, 10 

Hot-pressed 

(99+% pure) Flexure 5.2 12 Mecholsky, 17 

Hot pressed Flexure (Rods) 10.4 Mecholsky, 11 

Hot-pressed 

(Avco, 1-2 μm) 

Flexure (4-point, 2 

sizes) 9.8 Bansal, 14 

Sintered 

(Lucalox) 

Tension (Plates) 7.3 Congleton, 

Petch, 2 

C-6
 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Appendix C Mirror Constants 

Material Technique 

Mirror-Mist 

Ai 

(MPa· m ) 

Mist-Hackle or 

Mirror-Hackle 

Ao 

(MPa· m ) 

Branching 

Ab 

(MPa· 

m ) 

Reference 

Sintered (96%) 

(Alsimag 614) Flexure (Rods) 8.5 Kirchner, 9 

Sintered (96%) 

(Alsimag 614) 

Flexure (Rods) 

Flexure-Delayed 

Fracture (Rods) 

8.3 

8.9 Kirchner, 8 

Sintered (96%) 

(Alsimag 614) 

Flexure (Bars, 3-point 

and 4-point, 2 sizes) 9.0 Bansal, 13,14 

Sintered (96%) 

(Alsimag 614) Flexure (Rods) 9.1 Kirchner, 10 

Sintered (96%) 

(Alsimag 614) Flexure (Bars) 13.1 Mecholsky, 17 

Sintered (96%) 

(Alsimag 614) 

Flexure (Bars), 

Flexure (Biaxial ball on 

ring), 

Flexure (Biaxial ring on 

ring) 

7.6 ± 0.5 7.4 ± 0.6? 

7.2 ± 0.7? 

Choi, 33 

Sintered 96% 

(Alsimag 614) 

Flexure (Biaxial, 

pressure on ring disks) 

4.0 ± 0.3 

NZ (164 ±

11) 

Shetty, 15 

Hot isopressed 

Biolox forte Flexure (Bars, 4-point) 7.7 Hecht-Mijic, 76 

Hot isopressed 

Biolox delta 

ZTA 

Flexure (Bars, 4-point) 9.5 Hecht-Mijic, 76 

Silicon Carbide 

Sintered SiC 

(Hexoloy SA) Flexure (Bars) 5.39 Quinn, 23 

Sintered SiC 

(Hexoloy SA) 

Flexure (Biaxial, ring on 

ring plates) 6.3 ± 0.5 Salem, 31 

Sintered SiC 

(Hexoloy SA) 

Flexure (Biaxial, ring-onring plates) 5.5 ± 0.3 Choi, 33 

Sintered SiC 

(Hexoloy SA) C-ring 5.50 ± 0.94 8.20 ± 0.54 10.5 ± 1.0 Conway et al. 67 

Sintered SiC 

toughened 

(Hexoloy SX) 

Tension (Rods) and 

Flexure (Bars) 7.0 7.0? Srinivasan, 42 

Sintered with Al 

(Rioceram SiC) C-ring 5.67 ± 0.60 8.16 ± 0.54 11.6 ± 0.8 Conway et al. 67 

Sintered 

(Carolt S) 

Flexure (Bars, optical, 

SEM) 6.1, 6.8 Woodtli, 36 

Hot-pressed SiC 

(NC-203) 

Flexure (Rods) 

Flexure-Delayed 

Fracture (Rods) 

11.4 

11.9 

Kirchner, 8 

Hot-pressed SiC 

(NC-203) Flexure (Rods) 11.5 Kirchner, 9 

Hot-pressed SiC 

(ACE) Flexure (Rods) 10.8 Kirchner, 10 

Siliconized SiC (KT) Flexure 10.7 Mecholsky, 17 

Reaction bonded SiC 

(Coors SCRB205) C-ring 4.11 ± 0.31 5.22 ± 0.23 6.4 ± 0.4 Conway et al. 67 

Silicon Nitride 

Sintered Reaction 

Bonded 

(Ceralloy 147-31N) 

Flexure (Rods) 

Flexure (Bars) 

8.47 ± 0.07 

7.79 ± 0.02 Quinn, 26 

Sintered (SSN-500 

yttria/alumina) Flexure (Bars) 5.8 Quinn, 23 

Sintered 

(SN 220) 

Flexure (Biaxial, ring on 

ring disks) 

8.1 ± 2.4 Choi, 33 

C-7 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Material Technique 

Mirror-Mist 

Ai 

(MPa· m ) 

Mist-Hackle or 

Mirror-Hackle 

Ao 

(MPa· m ) 

Branching 

Ab 

(MPa· 

m ) 

Reference 

Sintered 

(AS 44) 

Flexure (Biaxial, ring on 

ring disks) 10.9 ± 2.7 Choi, 33 

Hot-pressed 

(Ceralloy 147A) Flexure (Bars) 7.83 Quinn, 23 

Hot-pressed 

(NC-132) Flexure (Rods) 9.2 Kirchner, 9 

Hot pressed 

(NC-132) 

Flexure (Rods) 

Flexure-Delayed 

Fracture (Rods) 

8.9 

9.2 Kirchner, 8 

Hot-pressed 

(NC-132) Flexure (Rods) 14.3 Kirchner, 10 

Hot-pressed 

(NC-132) 

Flexure (Bars, 1100OC) 

Flexure (Biaxial ring on 

ring, 1100OC) 

9.4 ± 1.2 

7.9 ± 2.1? 

Choi, 33 

Hot-pressed 

(HS-130) Flexure 18.1 Mecholsky, 17 

Hot-pressed 

(HS-130) Flexure (Rods) 9.1 Kirchner, 8 

Hot-isopressed 

(NT 154) Flexure (Bars) 5.9 ± 0.14 Choi, 32,33 

Hot-isopressed + 

30vol% SiC whiskers Flexure (Bars) 6.6 ± 0.11 Choi, 32,33 

Hot-isopressed 

(GN-10) 

Flexure (Biaxial ring-onring) 

Tension (Rods) 11.8 ± 1.4 

10.3? 

Choi, 33 

Reaction Bonded 

(NC 350) Flexure (Bars) 3.89 Messier, 24 

Reaction Bonded 

(NC 350) Flexure (Bars) 3.19 Larsen, 28 

Reaction Bonded 

(AME A25B) Flexure (Rods) 4.2 Kirchner, 8 

Zirconia 

Ytttria stabilized 

(3Y-TZP) Flexure (Bars) 9.95 Morrell, 43 

Ytttria stabilized 

(3Y-TZP) 

Flexure (Bars) 

Flexure (Biaxial piston 

on 3 ball) 

8.6 ± 0.23 

9.6 ± 0.23 

Quinn, 50 

Yttria stabilized 

(3Y-TZP, Lava) Flexure (Bars) 10.7 Quinn, 58 

Yttria stabilized 

(3.5Y-TZP) 

Flexure (Biaxial, ring on 

ring disks) 11.5 ± 1.5 Choi, 33 

Calcia partially 

stabilized PSZ 

As fired - quenched 

As fired 

Under aged, 1300oC 

Peak aged 

Peak aged 

Over aged 

Flexure (Bars, 3-point) 9.9 

15.5 

19.2 

22.1 

25.3 

-

Kirchner et al, 80 

Zircar 

(Alfred-Union Carbide, 

0.4 μm, 5-6 mol% Y 

TZP) 

Flexure (Bars) 15.2 Mecholsky, 17 

Zyttrite (AFML, 10 

μm, cubic) Flexure (Bars) 7.4 Mecholsky, 17 

Single crystals 

Ammonium 

diphosphate, 

single crystal Flexure (Bars) 0.5 Mecholsky, 17 

C-8
 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Appendix C Mirror Constants 

Material Technique 

Mirror-Mist 

Ai 

(MPa· m ) 

Mist-Hackle or 

Mirror-Hackle 

Ao 

(MPa· m ) 

Branching 

Ab 

(MPa· 

m ) 

Reference 

MgAl2AlO4 Spinel, 

single crystal Flexure (Bars) 2.6 Mecholsky, 17 

Quartz 

(112�0) plane with 

tension in [0001] 

direction 

Tension (Plates) 1.64 Ball, Payne, 78 

MgO ? 4.3 Ball, Payne, 78 

Sapphire (average of 

several planes) Flexure 6.1 Mecholsky 17 

Sapphire 

(Tyco filaments, 

c-axis parallel to fiber 

axis) 

Tension 

Flexure 

5.5 

10.0 Abdel-Latif, 38 

Sapphire 

(Ruby rods, c axis 

~60o off rod axis) 

Flexure (Rods) 3.3 Abdel-Latif, 38 

Sapphire 

Filament, c-axis 

parallel to fiber axis 

Tension, 800 to 1000oC 2.4 3.2 Rice, 51 

Sapphire, a plane Flexure (Bars, 4-point) 

7.3 ± 0.6 

7.1 ± 0.3 NZ (11) 

Salem, 70 

Sapphire ? 7.3 Ball, Payne 78 

Sapphire 

filament, c-axis 

parallel to fiber axis 

Tension, 800 to 1000oC 2.4 3.2 Rice, 51 

Silicon, single crystal Flexure (Bars) 

2.01 ± .14 

{110} plane 

<110> dir. 

1.61 ± .14 

{110} plane 

<100> dir. 

Tsai and 

Mecholsky, 72 

Silicon, single crystal ? 5.8 Ball, Payne, 78 

Other 

3BaO-SiO2 Flexure (Bars) 3.9 6.0 Mecholsky, 17 

B4C hot-pressed Flexure (bars) 4.8 9.27 Mecholsky, 17 

BaTiO2 (2 grades) Flexure (Bars) 5.0, 5.4 Mecholsky, 17 

Chromium ? 4.5 Ball, Payne, 78 

Mullite Flexure 6.1 Mecholsky, 17 

MgO Flexure 9.6 Mecholsky, 17 

MgO Tension (Plates) 4.3 Congleton, 

Petch, 2 

MgO, single crystal Flexure (Bars) 5 Mecholsky, 17 

MgF2 (Kodak) Flexure (Bars) 1.8 3.1 Mecholsky, 17 

MgF2 (Kodak, 

IRTRAN 1) 

Flexure (Bars) and 

Biaxial Disks 4.4 Mecholsky, 20 

MgAl2AlO4 Spinel Flexure (Bars) 4.0 7.8 Mecholsky, 17 

MgAl2AlO4 Spinel, 

single crystal Flexure (Bars) 2.6 Mecholsky, 17 

PMMA ? 8.5 * 8.5 * Ball, Payne, 78 

Porcelain, aluminous, 

electrical insulator 

grades 

Flexure, (Rods, 3-point) ~ 3.5 Carlström, 74 

PZT Flexure 1.7 3.7 Mecholsky, 17 

Graphite (POCO) Flexure 3.32 Mecholsky, 17 

SrZrO3 Flexure (Bars) 4.4 6.0 Mecholsky, 17 

Steatite (Mg silicate 

insulator, DC -144) 

Flexure (Rods) 4.8 

4.5 

Kirchner, 8,9 

Kirchner, 10 

C-9 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Material Technique 

Mirror-Mist 

Ai 

(MPa· m ) 

Mist-Hackle or 

Mirror-Hackle 

Ao 

(MPa· m ) 

Branching 

Ab 

(MPa· 

m ) 

Reference 

WC-Co 

8 wt % (H8 grade) 

12 wt % (H12) 

10 wt % (G10) 

15 wt % (G15) 

30 wt % (E30) 

10 wt % (J10) 

15 wt % (WCS15) 

20 wt % (WCS20) 

25 wt % (WCS25) 

Flexure (Bars, 4-point) 

24.4 ± 5 

41.0 ± 2 

36.8 ± 2 

51.7 ± 10 

75.0 ± 3 

29.6 ± 2 

58.8 ± 4 

71.9 ± 2 

86.4 ± 1 

Luyckx, 41 

WC-Co 6% Ni 

(Kennametal) 

Flexure 

(Biaxial, disks 

pressure-on-ring (POR) 

POR + ring-on-ring 

8.2 NZ 

⊗ (1154) 

14. NZ ⊗

(329) 

Salem, et al, 79 

WC-Co 

2% Ni, 1% Co 

(Cercom) 

Flexure 

(Bars?- 4-point) 

3.6 NZ ⊗

(493) Salem, et al, 79 

WC (no Co) Flexure 10 ⊗ Swab, 46 

Yttrium Aluminum 

Garnet, polycrystal. 

(2.2 μm) 

Flexure (Bars) 2.15 ± 0.08 Mezeix and 

Green, 66 

Yttrium Aluminum 

Garnet, single crystal 

(111) 

Flexure (Bars) 2.20 ± 0.06 Mezeix and 

Green, 66 

Zircon Porcelain 

(Alsimag 475) Flexure (Rods) 4.0 Kirchner, 8,9 

ZnSe Flexure (Bars) 1.7 Mecholsky, 17 

ZnSe (44 μm) Flexure (Biaxial, ring-onring disks) 0.8, 1.0 # Salem, 65 

Fibers 

Alkali-borosilicate Tension (Fibers) 1.33 NZ Jaras et al., 12 

Chalcogenide 

As-S-Se and 

Ge-AS-Se-Te 

Flexure (Fibers, 2 point 

bending) 

1.58 

1.77 ± 0.14 NZ 

(-66.4 ± 45.5) 

Quinn et al., 75 

Chalcogenide 

As2Se3 Tension (Fibers) 

0.69 

(0.35 if UV 

irradiated) 

0.77 

(0.38 if UV 

irradiated) 

Hulderman et al., 

76 

E glass 

CaO-Al2O3-B2O3 

alumina borosilicate 

Tension (Fibers) 1.47 NZ Jaras et al., 12 

Fused silica fibers Tension 2.10 Choi, 33 

Fused silica clad 

fibers Tension 1.96 ± 0.13 Baker, 29 

Fused silica fibers Tension 1.83 Castilone, 30 

Fused silica fibers Tension (Fibers) 2.2 ± 0.5 Mecholsky, 20 

Fused silica fibers Tension 2.224 Chandan, 37 

Si-Ti-C-O 

(Tyranno LoxM) fiber 

Fibers in 3-D SiC matrix 

composite 2.50 ± 0.09 Davies, 53 

Si-C-O fibers 

(Nicalon) 

Si-C-N-O fibers 

Si-N-C-O fibers 

Tension ~2 Sawyer, 54 

Si-C-O, (Nicalon) Fibers in 2-D CVI SiC 

matrix composite 2.51 Eckel + Bradt, 55 

Sapphire 

filament, c-axis 

parallel to fiber axis 

Tension, 800 to 1000oC 2.4 3.2 Rice, 51 

Zirconia silicate 

(Cem-FIL AR) Tension (Fibers) 2.37 NZ Jaras et al., 12 

C-10 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Appendix C Mirror Constants 

NR		 Not reported 

NZ		 A non-zero intercept was detected on the graph of stress versus inverse square root radius. The intercept 

value in MPa is listed in parenthesis. Mirror or branching constants calculated with non-zero intercepts 

are usually different than those calculated with intercepts through the ordinate. Consult the original 

reference for more information. 

*		 It is not clear whether Ball and Payne, ref. 78 listed this as Ao for PMMA, or Ab. Their other listed glass 

and polycrystalline ceramic Ab values are actually Ao values from Kirchner or Johnson and Holloway. 

** Additional mirror constants for the lead borate glasses are available at liquid nitrogen temperature (-196 

oC) in the original reference. 

*** Mostly likely an overestimate since stresses were not adjusted for failure location in the 3-point rod 

specimens. 

# Alternative analyses (besides a graph of stress versus inverse square root radius) are in the original 

reference. 

?		 These values (from the original references) for Ab cannot be correct since they are less than or equal Ao. 

⊗	 Large internal residual stresses have been reported in WC-Co materials, often lead to large non-zero 

intercepts. Consult the original reference for guidance. 

⊕ Tempered soda lime silica rods 

REFERENCES 

1.	 J. W. Johnson, and D. G. Holloway, “On the Shape and Size of the Fracture Zones 

on Glass Fracture Surfaces,” Phil. Mag., 14 [130] (1966) 731 - 743. 

2. J. Congleton, and N. J. Petch, “Crack Branching,” Phil. Mag., 16 [142] (1967) 749 

- 760. 

3.	 E. B. Shand, “Breaking Stress of Glass Determined from Dimensions of Fracture 

Mirrors,” J. Amer. Ceram. Soc., 42 [10] (1959) 474 - 477. 

4.	 E. B. Shand, “Breaking Stresses of Glass Determined from Fracture Surfaces,” The 

Glass Industry, April 1967, 90-194. 

5.	 L. Orr, “Practical Analysis of Fractures in Glass Windows,” Materials Research 

and Standards, 12 [1] (1972) 21 - 23. 

6.	 M. J. Kerper, and T. G. Scuderi, “Modulus of Rupture of Glass in Relation to 

Fracture Pattern,” Bull. Amer. Ceram. Soc., 43 [9] (1964) 622 - 625. 

7.	 M. J. Kerper, and T. G. Scuderi, “Relation of Fracture Stress to the Fracture 

Pattern for Glass Rods of Various Diameters,” Bull. Am. Ceram. Soc., 45 [12] 

(1966) 1065 - 1066. 

8.	 H. P. Kirchner, R. M. Gruver, and W. A. Sotter, “Fracture Stress-Mirror Size 

Relations for Polycrystalline Ceramics,” Phil. Mag., 33 [5] (1976) 775 - 780. 

9.	 H. P. Kirchner, R. M. Gruver, and W. A. Sotter, “Use of Fracture Mirrors to 

Interpret Impact Fracture in Brittle Materials,” J. Am. Ceram. Soc., 58 [5–6] 

(1975)188 - 191. 

10. H. P. Kirchner, and R. M. Gruver, “Fracture Mirrors in Polycrystalline Ceramics 

and Glass,” pp. 309 – 321 in Fracture Mechanics of Ceramics 1, eds. R. C. Bradt, 

D. P. H. Hasselman, and F. F. Lange, Plenum, NY, 1973. 

11. H. P. Kirchner, and R. M. Gruver, “Fracture Mirrors in Alumina Ceramics,” Phil. 

Mag., 27 [6] (1973)1433 - 1446. 

12. A. C. Jaras, B. J. Norman, and S. C. Simmons, “The Measurement of Glass Fibre 

Strength in Composites from Studies of Their Fracture Surfaces,” J. Mat. Sci., 18 

(1983) 2459 - 2465. 

C-11 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

13. G. K. Bansal, and W. H. Duckworth, “Fracture Stress as Related to Origin and 

Fracture Mirror Sizes,” J. Am. Ceram. Soc., 60 [7–8] (1977) 304 -310. 

14. G. K. Bansal, “On Fracture Mirror Formation in Glass and Polycrystalline 

Ceramics,” Phil. Mag., 35 [4] (1977) 935 - 944. 

15. D. K. Shetty, A. R. Rosenfeld, and W. H. Duckworth, “Crack Branching in 

Ceramic Disks Subjected to Biaxial Flexure,” Comm. of the Amer. Ceram. Soc., 

Jan. 1983, C10 - C13. 

16. W. H. Duckworth, D. K. Shetty, and A. R. Rosenfeld, “Influence of Stress 

Gradients on the Relationship Between Fracture Stress and Mirror Size for Float 

Glass,” Glass Technology, 24 [5] (1983) 263 - 273. 

17. J. J. Mecholsky, Jr., S. W. Freiman, and R. W. Rice, “Fracture Surface Analysis of 

Ceramics,” J. Mat. Sci., 11 (1976) 1310 - 1319. 

18. J. J. Mecholsky, R. W. Rice, and S. W. Freiman, “Prediction of Fracture Energy 

and Flaw Size in Glasses from Measurements of Mirror Size,” J. Amer. Ceram. 

Soc., 57 [10] (1974) 440 - 443. 

19. D. B. Marshall, B. R. Lawn, and J. J. Mecholsky, “Effect of Residual Contact 

Stresses on Mirror/Flaw Size Relations,” J. Amer. Ceram. Soc., 63 [5-6] (1980) 

358 - 360. 

20. J. J. Mecholsky, Jr., and R. W. Rice, “Fractographic Analysis of Biaxial Failure in 

Ceramics,” in Fractography of Ceramics and Glass Failures, ASTM STP 827, 

eds. J. Mecholsky and S. Powell, ASTM, Westerville, OH, 1984 185-193. 

21. J. J. Mecholsky, A. C. Gonzalez, and S. W. Freiman, “Fractographic Analysis of 

Delayed Failure in Soda Lime Glass,” J. Am. Ceram. Soc., 62 [11-12] (1979) 577 -

580. 

22. J. J. Mecholsky, “Quantitative Fractographic Analysis of Fracture Origins in 

Glass,” pp. 37 – 173 in Fractography of Glass, eds. R. C. Bradt and R. E. Tressler, 

Plenum, NY, 1994. 

23. G. D. Quinn, D. R. Messier, and L. J. Schioler, “Characterization of Ceramic 

Vane Materials for 10 kW Turboalternator,” U.S. Army Materials and Mechanics 

Research Center Technical Report TR 83-18, April 1983. 

24. D. R. Messier, L. J. Schioler, and G. D. Quinn, “Fracture Behavior and Strength of 

Reaction-Bonded Si3N4 Turbine Shrouds,” Bull. Am. Ceram. Soc. Bulletin, 60 [8] 

(1981) 812 - 817. 

25. J. B. Quinn, “Extrapolation of Fracture Mirror and Crack-Branch Sizes to Large 

Dimensions in Biaxial Strength Tests of Glass,” J. Am. Ceram. Soc., 82 [8] (1999) 

2126 - 2132. 

26. G. D. Quinn, L. K. Ives, S. Jahanmir, and P. Koshy, “Fractographic Analysis of 

Machining Cracks in Silicon Nitride Rods and Bars,” in Fractography of Glasses 

and Ceramics, IV, Ceramic Transactions, Vol. 122, eds., J. R. Varner and G. D. 

Quinn, American Ceramic Society, Westerville, OH, 2001 343 - 365. 

27. P. B. Adams and S. E. DeMartino, “Glass-Ceramic Cookware Failure Analysis,” 

in Engineered Materials Handbook, Vol. 4, Ceramics and Glasses, ASM, Metals 

Park, OH, 1991 669 – 673. 

28. D. C. Larsen, and G. C. Walther, “Property Screening and evaluation of Ceramic 

Turbine Engine Materials,” IIT Research Institute, Chicago, IL, Interim Technical 

Report No. 5, IITRI D 6114-ITR-30, 1978. 

C-12 
 

 

 

 

 

 

 

 

 

 

 

 

Appendix C Mirror Constants 

29. L. K. Baker, and G. S. Glaesemann, “Break Source Analysis, Alternate Mirror 

Measurement Method,” Proceedings of the International Wire and Cable 

Symposium, Philadelphia, PA, 1998, 933-937. 

30. R. J. Castilone, G. S. Glaesemann, and T. A. Hanson, “Relationship Between 

Mirror Dimensions and Failure Stress for Optical Fibers,” pp. 11 – 20 in Optical 

Fiber and Fiber Component Mechanical Reliability and Testing II, eds. J. 

Matthewson and C. R. Kurkjian, Proc. of SPIE, Vol. 4639, (2002). 

31. J. A. Salem, S. R. Choi, and L. M. Powers, “Toughened Ceramics Life Prediction,” 

in Ceramic Technology Project Semiannual Progress Report for October 1992 

Through March 1993, Oak Ridge National Laboratory Technical Report TM 

12428, September, 1993, 307-316. 

32. S. R. Choi, and J. A. Salem, “Indentation Flaw Formation and Strength Response 

of Silicon Nitride Ceramics at Low Indentation Loads,” J. Mat. Sci. Lettrs, 11 

(1992) 1398–1400. 

33. S. R. Choi, and J. P. Gyekenyesi, “Crack Branching and Fracture Mirror Data of 

Glasses and Advanced Ceramics,” NASA Technical Report TM 1998-206536, 

1998. 

34. A. B. J. Clark, and G. R. Irwin, “Crack Propagation Behaviors,” Experimental 

Mechanics, 6 [6] (1966) 321 - 330. 

35. S. Aoki, and M. Sakata, “Crack Bifurcation Under Hydrostatic Pressure,” 

Engineering Fracture Mechanics, 13 [3] (1980) 491 – 499. 

36. J. Woodtli, “Quantitative Fractographie in Keramik: Auswertung von 

Spiegelzonen in SiC,” Swiss Federal Laboratories for Materials Testing and 

Research, Report 157361, Dec. 1997. 

37. H. C. Chandan, and R. D. Parker, “Fractography of Optical Fibers,” pp. 143- 184 

in Fractography of Glass, eds. R. C. Bradt and R. E. Tressler, Plenum Press, NY, 

1994. 

38. A. I. A. Abdel-Latif, R. E. Tressler, and R. C. Bradt, “Fracture Mirror Formation 

in Single Crystal Alumina,” in Fracture 1977, Vol. 3, ed. D. Taplin, University of 

Waterloo Press, 1977, 933 - 939. 

39. D. A. Reed, and R. C. Bradt, “Fracture Mirror-Failure Stress Relations in 

Weathered and Unweathered Window Glass Panels,” Comm. Am. Ceram. Soc., 

Nov. 1984, C 227 – C 229. 

40. M. J. Ball, D. J. Landini, and R. C. Bradt, “Fracture Mist Region in a Soda-Lime 

Silica Float Glass,” pp. 110-120 in Fractography of Ceramic and Metal Failures, 

ASTM STP 827, J. Mecholsky, Jr., and S. Powell, eds., ASTM, 1984. 

41. S. Bartolucci Luyckx, and A. Sannino, “Crack Branching and Fracture Mirrors in 

Cemented Tungsten Carbide,” J. Mat. Sci., 23 (1988) 1243 - 1247. 

42. G. V. Srinivasan, J. Gibson, and S. K. Lau, “The Origin of Strength Limiting 

Defects in a Toughened SiC (Hexoloy SX-SiC), in Fractography of Glasses and 

Ceramics III, Ceramic Transactions Vol. 64, eds, J. R. Varner, V. D. Fréchette, 

and G. D. Quinn, American Ceramic Society, Westerville, OH, 1996, 181-192. 

43. R. Morrell, L. Byrne, and M. Murray, “Fractography of Ceramic Femoral Heads,” 

in Fractography of Glasses and Ceramics, IV, Ceramic Transactions, Vol. 122, 

eds., J. R. Varner and G. D. Quinn, American Ceramic Society, Westerville, OH, 

2001, 253 - 266. 

C-13 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

44. D. Lewis, III, “Fracture Strength and Mirror Size in a Commercial GlassCeramic,” J. Am. Ceram. Soc., 64 [2] (1961) 82 - 86. 

45. F. Hulderman, J. S. Sanghera, and J. D. Mackenzie, “The Effect of UV Radiation 

on the Mechanical Strength of As2Se3 Glass Fibers,” J. Non-Cryst. Sol., 127 (1991) 

312 - 322. 

46. J. J. Swab, J. C. LaSalvia, and W. A. Gooch, “Characterization of WC-Based 

Armor Ceramics,” unpubl. priv. com., U.S. Army Research Laboratory, Aberdeen 

Proving Ground, MD, 21005. 

47. J. R. Kelly, “Fractography of Dental Ceramics,” pp. 241-251 in Fractography of 

Glasses and Ceramics, IV, eds. J. R. Varner and G. D. Quinn, Ceramic 

Transactions Vol. 122, American Ceramic Society, Westerville, OH 43081, 2001. 

48. D. N. Coon, “Mirror/Mist Radius Constant for Y-Al-Si-O-N Glass,” J. Mat. Sci. 

Lettrs., 21 (2002) 1121-1122. 

49. A. I. A. Abdel-Latif, R. C. Bradt, and R. E. Tressler, “Dynamics of Fracture Mirror 

Boundary Formation in Glass,” Int. J. Fract., 13 [3] (1977) 349 - 359. 

50. G. D. Quinn, J. Eichler, U. Eisele, and J. Rödel, “Fracture Mirrors in a Nanoscale 

3Y-TZP,” J. Am. Ceram. Soc., 87 [3] (2004) 513 - 516. 

51. R. W. Rice, “Corroboration and Extension of Analysis of c-Axis Sapphire 

Filament Fractures from Pores,” J. Mat. Sci. Lttrs., 16 (1997) 202 - 205. 

52. H. Fischer, J. Gröbel, and R. Marx, “Fraktograpische Schadensanalyse an 

Dentalkeramik,” Deutsche Zahnarztl. Zeitschrifft, 55 [10] (2000) 667 - 670. 

53. I. J. Davies, and T. Ishikawa, “Mirror Constant for Tyranno Silicon-TitaniumCarbon-Oxygen Fibers Measured in Situ in a Three-Dimensional Woven Silicon 

Carbide/Silicon Carbide Composite,” J. Am. Ceram. Soc., 85 [3] (2002) 691 -693. 

54. L. C. Sawyer, M. Jamieson, D. Brikowski, M. I. Haider, and R. T. Chen, “Strength, 

Structure, and Fracture Properties of Ceramic Fibers Produced from Polymeric 

Precursors: I: Base-Line Studies,” J. Amer. Ceram. Soc., 70 [11] (1987) 798 - 810. 

55. A. J. Eckel, and R. C. Bradt, “Strength Distribution of Reinforcing Fibers in a 

Nicalon Fiber/Chemically Vapor Infiltrated Silicon Carbide Matrix Composite,” J. 

Am. Ceram. Soc., 72 [3] (1989) 455 - 458. 

56. O. E. Alarcón, R. E. Medrano, and P. P. Gillis, “Fracture of Glass in Tensile and 

Bending Tests,” Metall. Mat. Trans A,, 25A, May (1994) 961 - 967. 

57. S. T. Gulati, J. F. Bayne, W. R. Powell, and J. D. Helfinstine, “Mirror Constant for 

AMLCD Glass,” Bull. Am. Ceram. Soc., 83 [5] (2004) 14 - 15. 

58. J. B. Quinn, D. Cheng, R. Rusin, and D. Suttor, “Fractographic Analysis and 

Material Properties of a Dental Zirconia,” International Association for Dental 

Research (IADR) paper, 0560, Baltimore, MD, March, 2005. 

59. J. B. Quinn and G. D. Quinn, “Materials Properties and Fractography of an Indirect 

Dental Resin Composite,” Dent. Mater., 26 (2010) 589-599. 

60. N. Shinkai, and H. Ishikawa, “Crack Branching of Binary PbO-B2O3 Glasses,” J. 

Noncryst. Solids., 52 [1-3] (1982) 385 - 394. 

61. R. E. Bullock, and J. L. Kaae, “Size Effect on the Strength of Glassy Carbon,” J. 

Mat. Sci., 14 (1979) 920 - 930. 

62. J. M. Barsom, “Fracture of Tempered Glass,” J. Am. Ceram. Soc., 51 [2] (1968) 75 

- 78. 

C-14 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Appendix C Mirror Constants 

63. N. Terao, “Sur ene Relation entre la Résistance à la Rupture et le Foyer 

d’Eclatement du Verre,” J. Phys. Soc. Jap., 8 [4] (1953) 545 - 549. 

64. W. C. Levengood, “Effect of Origin Flaw Characteristics on Glass Strength,” J. 

Appl. Phys., 29 [1] (1958) 820 - 826. 

65. J. A. Salem, “Mechanical Characterization of ZnSe Windows for Use With the 

Flow Enclosure Accommodating Novel Investigations in Combustion of Solids 

(FEANICS) Module,” NASA Technical Report TM-2006-214100, NASA-Glenn 

Center, Cleveland, OH, 2006. 

66. L. Mezeix and D. J. Green, “Comparison of the Mechanical Properties of Single 

Crystal and Polycrystalline Yttrium Aluminum Garnet,” Int. J. Appl. Cer. 

Technol., 3 [2] (2006) 166 - 176. 

67. J. C. Conway, Jr., J. J. Mecholsky, O. M. Jadaan, and D. L. Shelleman, “Test 

Methodology for Ceramic Tubular Components,” Center for Advanced Materials 

Research Report, Pennsylvania State University, 1989. 

68. J. R. Varner and H. J. Oel, “Einfluss von Oberflächenbeschädigungen auf die 

Festigkeit von Glasstäben, ” Grastechn. Ber., 48 [5] (1975) 73 – 78. 

69. N. Shinkai, “Fracture Mirror Constants of Silicate Glasses,” Japan J. Appl. Phys., 

14 [1] (1975) 147-148. 

70. J. A. Salem, Private communication, March, 2008. 

71. J. R. Varner and H. J. Oel, “Surface Defects:Their Origin, Charcaterization and 

effects on Strength, ”J. Non-Cryst. Sol., 19 (1975) 321-333. 

72. Y. L. Tsai and J. J. Mecholsky, Jr, “Fractal Fracture of Single Crystal Silicon,” J. 

Mater. Res., 6 [6] (1991) 1248-1263. 

73. O. Gaume and S. Pelletier, “Analysis of Fracture Featrures for Uni-axial Bent 

Glass Plates,” in 2002 Glass Odyssey, proc. 6th ESG conf., June 26, 2002, 

Montpellier, France. 

74. E. Carlström and R. Carlsson, “A Fractoraphic Study of Electrical Porcelains 

Containing Alumina,” J. De Physique, 47 (C-1) (1986) 563-567. 

75. J. B. Quinn, V. Q. Nguyen, J. S. Sanghera, I. K. Lloyd, P. C. Puzera, R. E. Miklos, 

and I. D. Aggarwal, “Strength and Fractographic Analysis of Chalcogenide AS-SSe and Ge-As-Se-Te Glass Fibers,” J. Non-Cryst. Sol., 325 (2003) 150-157. 

76, S. Hecht-Mijic, “Fracture Mirror Constants of Bioceramics for Hip Joint 

Replacement: Determination and Application, ” pp. 303 -312 in Fractography of 

Glasses and Ceramics V, eds. J. R. Varner, G.D. Quinn, M. Wightman, Ceramic 

Transactions Vol. 199, American Ceramic Society, Westerville, OH, 2007. 

77. J. A. Salem, N. Smith, and A. Ersahin, “Mechanical Properties of a High Lead 

Glass Used in the Mars Organic Molecule Analyzer,” NASA Glenn Research 

Center, Cleveland, OH, Technical Report, TM 2015-218847, August 2015. 

78. A. Ball and B. W. Payne, “The Propagation and Bifurcation of Cracks in Quartz,” 

pp. 971 – 978 in Fracture 1977, Vol. 3, Waterloo, Canada, 1977. 

79. J. A. Salem, S. Slingluff, R. Rogers, and J. J. Swab, “Stress Gradients and Crack 

Branching in Tungsten Carbide,” to be publ. 2015 

80. H. P. Kirchner, R. M. Gruver, M. V. Swain, and R. C. Garvie, “Crack Branching in 

Transformation-Toughened Zirconia,” J. Am. Ceram. Soc., 64 [9] (1981) 529-533. 

C-15 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Appendix D Guidelines for Fracture Mirrors 

APPENDIX D GUIDELINES FOR MEASURING FRACTURE 

MIRRORS 

D.1 Introduction 

This appendix presents a set of guidelines for measurement of fracture mirrors 

in glasses and ceramics. These guidelines were originally presented by the 

author at the Fifth Conference on the Fractography of Glasses and Ceramics in 

Rochester New York in 2006. 1 That paper reviewed the historical 

development (going back as far as 1894) of fracture mirror analysis as a 

quantitative tool for glasses and ceramics. These guidelines also were included 

in the first edition of this NIST Guide in 2007. They were adopted as an 

ASTM standard C 1678 2

 in 2007. 

Fracture mirror measurements are a powerful diagnostic tool for quantitative 

analysis. Mirrors are circular in ideal loading conditions such as uniform 

tension specimens. Otherwise, they can be elongated or distorted due to stress 

gradients or geometrical effects. There is a strong subjective element to 

estimating the location of the mirror boundary. Johnson and Holloway3 said: 

“The position assigned to the boundary between mirror and mist zones 

depends upon the illumination and the magnification at which the fracture 

is examined, even within the range of the optical microscope. .... 

However, under given conditions a reproducible position for the boundary 

can be assigned.” 

Since many of the early mirror measurements were made while viewing 

through the optical microscope, it is safe to say that the first perception of a 

mirror boundary was where the surface roughness was a fraction of the 

wavelength of light (0.39 μm - 0.80 μm). Threshold levels of detectability 

have been estimated to be as small as a few tens of nanometers to as much as 

0.25 μm. Attempts to make objective determinations of the mirror boundary by 

quantitative surface roughness characterization are discussed in section D.3 

below, but a single threshold roughness value cannot be specified. 

The following Guidelines are intended to bring some consistency to measurement procedures. Some of the following recommendations are based on 

common sense. Others are based upon the experiences of the author and other 

experts. The empirical mirror relationship for a residual stress free part is: 

σ R = A (D.1) 

D-1 
 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

where σ is the stress at the origin, R is the mirror radius, and A is the 

appropriate mirror constant as described below. Leighton Orr recognized as 

early as 1945 that stress varied inversely with the square-root dependence of 

the mirror radius, so the author now refers to equation D.1 as “Orr’s equation.” 

The boundary criteria are: 

The mirror-mist boundary in glasses is the periphery where one can discern 

the onset of mist. This boundary corresponds to Ai, the inner mirror constant. 

The mist-hackle boundary in glasses is the periphery where one can discern 

the onset of systematic hackle. This boundary corresponds to Ao, the outer 

mirror constant. 

The mirror-hackle boundary in polycrystalline ceramics is the periphery 

where one can discern the onset of systematic new hackle and there is an 

obvious roughness change relative to that inside the mirror region. This 

boundary corresponds to Ao, the outer mirror constant. Ignore premature 

hackle and/or isolated steps from microstructural irregularities in the mirror or 

irregularities at the origin. 

In coarse-grained or porous materials, it may be impossible to identify a mirror 

boundary. In polycrystalline ceramics, it is highly unlikely that a mirror-mist 

boundary can be detected due to the inherent roughness created by the crackmicrostructure interactions even within the mirror. Another reason a boundary 

may not be evident is that in very weak materials, the mirror may be larger than 

the specimen or component. 

If the mirror is being measured for a component failure analysis, and if the 

mirror constant A is known, follow steps 1 through 9 in the next section and 

compute the stress at the origin in accordance with equation D.1. Note that this 

origin stress may or may not necessarily be the maximum stress in the part. 

On the other hand, if the fracture mirror constant A is being evaluated by 

means of testing laboratory specimens and the origin stresses are known, follow 

steps 1 through 12. 

Examples of how to judge the boundaries are shown in Figures D.1 through 

D.6. Low power images provide an overview. Matching unmarked and 

marked images are shown so that the reader can make a judgment for 

comparison. The boundaries were assessed while looking through a 

microscope and the digital images in Figures D.1 to D.6 were marked. Mirrors 

D-2 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Appendix D Guidelines for Fracture Mirrors 

should be evaluated while looking in a microscope. Although some of the 

mirrors are approximately circular or semicircular, examples of common 

alternative shapes are presented. Kirchner et al.4,5 presented compelling 

evidence that fracture mechanics analyses account for stress gradient and 

geometric effects upon mirror shapes. 

hackle 

mirror 

mist 

(a)

 (b) 

Figure D.1 Fused silica rod broken in flexure (122 MPa). The origin is a 

surface flaw located at the bottom of the specimen (the insert in a) where the 

stress was a maximum. A stress correction is unnecessary. The mirror-mist 

boundary is small relative to the cross section size and is approximated by a 

circle in (b). The inward cusps at the surface are ignored as discussed in 

section D.3. The mist – hackle boundary is slightly elongated towards the top. 

Close examination of the fracture surface in the vicinity of the flaw (not shown) 

showed that fracture started from the deepest part of the flaw. (c) and (d) on 

the next page show SEM images of the same mirror. 

D-3 
 

 

 

 

 

-

-

 Fractography of Ceramics and Glasses 

 (c)

 (d) 

mirror mist 

mist hackle 

Figure D.1 continued (c) is an SEM image of the same mirror and at the 

same magnification as (a) and (b). The mist is indistinct in the SEM image. (d) 

is a composite of two SEM images showing the transition from mirror to mist to 

hackle. The locations of the boundaries as assessed by optical microscopy (as 

shown in (b))are marked by dashed lines on the SEM images. 

D-4
 

 

 

 

 

 

 

Appendix D Guidelines for Fracture Mirrors 

(a) (b) 

(c) 

(d) 

Figure D.2 Silicon carbide tension strength specimen (371 MPa) with a 

mirror centered on a compositional inhomogeneity flaw shown in Figure 

6.20a,b. Note how clear the mirror is in the low power images. The mirror 

boundary (arrows in d) is where systematic new hackle forms and there is an 

obvious difference compared to the roughness inside the mirror region. 

D-5 
 

 

 

 

 Fractography of Ceramics and Glasses 

(a) (b) 

(c) 

(d) 

2R 

Figure D.3 Silicon nitride bend bar with a Knoop surface crack (449 MPa, 

Ceralloy 147-31N, Refs. 6,7). The mirror is incomplete into the stress 

gradient, but the mirror sides can be used to construct boundary arcs in (d). 

Radii are measured in the direction of constant stress along the bottom. 

D-6 
 

 

 

 

 

Appendix D Guidelines for Fracture Mirrors 

(a)

 (b) 

2R 

Figure D.4 Example of a mirror in a fine-grained 3 mol % yttria-stabilized 

tetragonal zirconia (3Y-TZP) polycrystal (Bosch, Ref. 8). The mirror is 

difficult to mark in this material. (a) shows the uncoated fracture surface of a 

2.8 mm thick flexural strength specimen (486 MPa) with low-angle grazing 

illumination. (b) shows an interpretation for a mirror-hackle boundary where 

systematic new hackle is detected (small white arrows) as compared to the 

roughness inside the mirror. The marked circle is elongated somewhat into the 

depth due to the stress gradient. The radius was 345 μm. 

D-7 
 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) (b) 

(c)

 (d) 

Figure D.5 Fracture mirror in a 3Y-TZP zirconia flexural strength specimen 

(3M-ESPE Lava, 798 MPa). (a) and (b) are uncoated and gold coated images 

of the whole fracture surface, respectively. (Images courtesy J. Quinn) 

D-8 
 

 

 

 

 

 

Appendix D Guidelines for Fracture Mirrors 

(a) (b) 

(c) (d) 

Figure D.6 Fracture mirror in another 3Y-TZP bend bar. (3M ESPE Lava, 

1059 MPa maximum stress, 1024 MPa at the origin). (a) and (b) are stereo 

optical microscope images. (c) and (d) are SEM images. Images (b) and (c) 

are at the same magnifications. The flaw is a pore. Fine hackle lines in the 

mirror run right to the origin. They are created by the flaw itself and the 

microstructure. The mirror boundary is where systematic new hackle forms 

and there is an obvious difference compared to the roughness inside the mirror 

region. This is a very difficult example. One interpretation is shown in the 

lower magnification SEM image. (SEM images (c) and (d) courtesy J. Quinn) 

D-9 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

D.2 Measure the Mirror Size 

Use the following steps to measure mirror sizes. These steps are repeated, but 

with detailed notes, clarifications, and illustrations in section D.3. 

1. Use an optical microscope whenever possible. A compound optical 

microscope is best for glasses. A stereoptical microscope is best for 

ceramics. A thin coating may be applied to translucent or transparent 

ceramics. A scanning electron microscope may be used if optical 

microscopy is not feasible. 

2. The fracture surface should be approximately perpendicular to the 

microscope optical path or camera. 

3. Optimize the illumination to accentuate topographical detail. 

Ceramics may be coated with gold or carbon. 

4. Use a magnification such that the fracture mirror area occupies about 

75 % to 90 % of the width of the field of view for glasses, and 

approximately 33 % to 67 % of the width of the field of view for 

ceramics. An additional, lower-power image may be helpful for ceramics. 

5. Measure the mirror size while viewing the fracture surface with an 

optical microscope whenever possible. 

Use either calibrated reticules in the eyepieces or traversing stages with 

micrometer-positioning heads. Alternatively, measurements may be made 

on digital images on a high-resolution computer monitor, preferably while 

the fracture surface can be simultaneously viewed through the microscope 

eyepieces in order to aid judgment. 

Measurements from photos may be used as a last resort if the steps above 

cannot be followed. This may be necessary for very small specimens or 

very strong specimens with tiny mirrors and a scanning electron 

microscope must be used to photograph the mirror. 

Measurements from other devices may be used provided that the criterion 

used for identifying the mirror boundary is carefully documented. 

6. Measure radii in directions of approximately constant stress whenever 

possible. A mirror diameter may be measured and halved to estimate the 

radius if the origin site is indistinct or complex. 

D-10
 
 

 

 

 

 

 

 

 

 

 

 

Appendix D Guidelines for Fracture Mirrors 

7. Exercise caution when mirrors are large relative to the specimen crosssection size. 

8. Show at least one photo with arrows or lines marking the mirror size. 

9. Report how the mirrors were measured. 

Additional steps for the determination of fracture mirror constants: 

10. Use the stress at the origin site. 

Correct the stress for location in specimens with stress gradients. 

11. Evaluate the fracture mirror constants by regressing stress on inverse 

square root of mirror radius. (Preferred method) 

Use linear regression methods to obtain A in accordance with equation D.2 

with a forced zero intercept as shown in Figure D.7 Use units of MPa for 

stress and 1/√m for the graph axes. (It is a major nuisance to have to 

correct the numerical value of the slope if the x axis is in units of mm.) 

The mirror constant A is the slope of the regression line. 

A 

σ =	 (D.2) a R 

where σa is the stress at the origin site, computed from the known applied 

stresses. R is the mirror radius in the direction of constant stress. 

Use some judgment in the regression analysis since fracture mirror data 

frequently has moderate scatter. If the data does not appear to fit a trend 

that has a zero intercept, regress the data with a non-zero intercept and 

equation D.3. Report the intercept if it deviates significantly (> 10 MPa) 

from zero. Investigate possible residual stresses or specimen size or shape 

issues if the intercept deviates significantly from zero. 

σ a = A −σ r	 (D.3) R 

where σa is the stress at the origin from known applied stresses, σr is the 

residual stress at the origin location. A negative σr is a compressive 

residual stress and a positive σr is tensile, i.e., a positive intercept on the 

graph corresponds to a compressive residual stress and a negative intercept 

corresponds to a residual tensile stress. 

D-11
 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

12. Mirrors sizes should be collected over a broad range of sizes and 

fracture stresses if possible. Data from different specimen types and 

sizes may be combined.

 (a)	 (b) 

Figure D.7 Plot of applied stress σa (at the origin) versus 1/√R. (a) shows the 

trend for residual stress-free parts. (b) shows it for parts with residual 

stresses. Compressive residual stresses move the locus up with a positive 

intercept σr, but with the same slope. Tensile residual stresses shift the data 

downwards with a negative intercept (not shown). 

D.3 Clarifications, Additional Notes, and Illustrations 

1. Use an optical microscope whenever possible. A compound optical 

microscope with bright field illumination is best for glasses. A stereo optical 

microscope is best for ceramics. A thin coating may be applied to translucent 

or transparent ceramics. A scanning electron microscope may be used if 

optical microscopy is not feasible. 

Differential interference contrast (DIC, also known as Nomarski) mode 

viewing with a research compound microscope originally was not recommended by this Guide. It is only suitable for glasses. It is not suitable for 

rough ceramic fracture surfaces. There are complications with glass fracture 

surfaces. There is no question that DIC mode viewing can discern very subtle 

mist features in glasses, but the threshold of mist detectability is highly 

dependent upon how the polarizing sliders are positioned. Hence, DIC 

measured radii are quite variable. DIC measured radii can be substantially 

D-12
 
 

 

 

 

 

 

 

 

 

 

 

 

 

Appendix D Guidelines for Fracture Mirrors 

smaller than those obtained with conventional viewing modes. During a review 

process in 2010, paragraph 6.1.3 of standard ASTM C 16782 was revised to 

allow DIC mode for glasses, provided that the polarizing elements are used in a 

consistent manner in the grey mode (with the lambda plate for color control 

removed). 

Dark-field illumination may be used with glasses, but dark-field images may 

lose a little resolution with glasses and radii may be slightly larger as a result. 

It is very effective with highly-reflective mirror surfaces in single crystals. 

Confocal optical microscopes and optical interferometers have been 

occasionally used to examine fracture mirrors, but I am unaware of any 

published systematic study to correlate apparent mirror sizes from these tools to 

those measured with conventional optical microscopes. The quantitative 

surface roughness capabilities of these instruments in principle could be used to 

correlate an average or root mean square roughness to the mist or hackle 

boundaries. 

Scanning electron microscope images of mirrors are not recommended for 

measuring the mirror-mist boundary in glasses since the boundary is usually 

indiscernible at the magnifications needed to see the overall mirror. SEM 

images often appear flat and do not have adequate contrast to see the fine mist 

detail. SEM images may be used to measure mirror hackle boundaries with 

very small mirrors that would be difficult to see with optical microscopy, e.g., 

as in high strength optical fibers. 

Scanning electron microscope images may be used for ceramics if necessary, 

but steps should be taken to enhance contrast and shadowing to produce images 

such as shown in Figures D-6c and d. 

Attempts to correlate the mirror boundaries with a simple surface roughness 

parameter have produced mixed results. The work has been limited to glasses. 

Duckworth et al.9 carefully studied mirror sizes in float glass using optical 

photographs and a conventional surface profilometer. They obtained a good 

correlation with optical boundary estimates when the surface roughness 

reached a level of 0.25 μm for the mirror- mist Ai boundary, and 5 μm for the 

mist-hackle Ao boundary. Kuluwansa et al.

10 used scanning tunneling 

microscopy (STM) to study mirror features and suggested the transition from a 

nano-mist region to the mist-hackle region with its rougher features (that were 

observable in the SEM) may be sharp. The characteristic scale of crack 

branching in the mirror to mist-hackle transition region ranged from 50 nm (the 

size of typical nano-mist features) to about 50 μm (the size of typical hackle 

D-13
 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

features. Hull11,12 and Wünsche et al.13 used atomic force microscopy to show 

that the mist region in glass or brittle epoxy has a roughness of as small as a 

few tens of nanometers, which is much less than the wavelength of visible light 

(390 nm to 800 nm). The scanned regions were quite small, however, and both 

groups noted that if the AFM scans a small region between hackle or river line 

steps, then the measured roughness is much less than if the latter are included. 

Hull pointed out that the greater undulations from Wallner lines need to be 

factored out when evaluating the intrinsic mist roughness. His study showed 

that roughness increased continuously and there were no dramatic jumps in 

roughness at the boundaries, but the rate of change of roughness did change 

significantly at the mirror-mist boundary. Surface roughness measurements 

should be taken perpendicular to the direction of crack propagation since highresolution transmission microscope images14,15,16 of the mist and hackle show 

the surface features are elongated in the direction of crack propagation. 

Hull12 pointed out that different surface roughness characterization devices 

such as atomic force microscopes (AFMs), mechanical profilometers, and laser 

optical profilometers all have different advantages, disadvantages, sensitivities 

and scanning zone sizes. AFM’s can measure tiny regions with very high 

sensitivities, but may miss large hackle steps in a mist or hackle zone. These 

latter features can dramatically alter the average or root mean square roughness. 

A mechanical stylus profilometer or laser profilometer with a 1 μm spot size 

may miss the small undulations and be more sensitive to larger hackle steps on 

the fracture surface. Mist and hackle regions may have different roughness at 

different scales. Hull discuses these various scales of roughness in some detail 

in his book.12 

The author is not aware of any systematic studies comparing fracture mirror 

sizes as a function of the viewing mode or microscope type. 

2. The fracture surface should be approximately perpendicular to the 

microscope optical path or camera. 

This simple and fairly obvious requirement is intended to avoid the 

foreshortening that can occur if the specimen is tilted. A small amount of 

tilting is acceptable in order to get a favorable reflection in a glass piece. 

A compound optical microscope is best for glasses. A stereo optical 

microscope is best for ceramics. 

The requirement poses a small problem if the mirrors are examined with stereo 

binocular microscopes. These have two different tilted optical paths. If 

viewing with both eyes in a stereomicroscope, the specimen should be flat and 

D-14
 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Appendix D Guidelines for Fracture Mirrors 

facing directly upwards. The observer’s brain will interpret the image as 

though he is facing it directly. 

Alternatively, if a camera is mounted on one light path of the stereomicroscope 

and it is used to capture or display the mirror, then the specimen should be 

tilted so that the camera axis is normal to the fracture surface. For example, tilt 

the specimen to the right if the camera is attached to the right optical path. 

3. Optimize the illumination to accentuate topographical detail. 

The mist and hackle features should be accentuated. Glasses may either be 

illuminated from directly down onto a fracture surface or by low-angle grazing 

illumination. Ceramics should not be directly illuminated since the light will 

reduce contrast, especially in translucent or transparent materials. Ceramics 

should be illuminated with low-angle grazing illumination. Stereoptical 

microscopes are strongly preferred for ceramics. Low-angle illumination is 

less convenient with compound light microscopes. The observer should 

experiment with whatever illumination options are available to accentuate 

subtle surface roughness and topography features. Contrast and topographic 

detail should be emphasized with the SEM if it is necessary to use this mode of 

examination because the mirror is too small to be measured optically. 

4. Use a magnification such that the fracture mirror area occupies about 

75 % to 90 % of the width of the field of view for glasses, and approx.-

imately 33 % to 67 % of the width of the field of view for ceramics. An 

additional, lower-power image may be helpful for ceramics. 

Observers usually mark the mirror boundaries closer to the origin at greater 

magnifications than they would at lower magnifications. This is because mist or 

micro hackle markings are easier to see at distances closer to the origin at high 

magnification. This is particularly the case with glasses. Conversely, at very 

low magnifications, much detail is lost and observers typically overestimate the 

mirror size. 

Fracture mirrors are reasonably easy to see in glasses and magnifications 

should be used such that they nearly fill the field of view. 

Mirror interpretation is more problematic with polycrystalline ceramics. 

Excessive magnification often leads to confusion as to where the boundary is 

located. Even though a mirror may be obvious at low or moderate 

magnifications, at higher magnifications it may be impossible to judge a 

boundary. It is more practical to view the mirror region (and the natural 

microstructural roughness therein) relative to the hackle roughness in the 

D-15
 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

regions outside the mirror. “Stepping back” and using the 33 % to 67 % rule 

should help an observer better detect the topography differences. Images 

recorded at these magnifications are also more convincing when shown to other 

fractographers or engineers. Even lower magnification images may also be 

made to aid interpretation such as shown in Figures D-2b, D-3b, and D-6b. 

The images should not be more than 5 times different in magnification; 

otherwise it is difficult to correlate features in one image to another. 

Sometimes the microstructure of polycrystalline ceramics creates even worse 

judgment problems in ceramic matrix composites (particulate, whicker or 

platelet) or self-reinforced ceramics whereby elongated and interlocking grains 

impart greater fracture resistance. These difficulties were experienced in 

studies of the fracture surfaces of a self-reinforced silicon nitride6,7 and also a 

yttria-stabilized tetragonal zirconia polycrystal.

8 Mirrors were plainly evident 

at low magnifications, but accurate assessment of their size was difficult. The 

mirror region itself was somewhat bumpy in the self-reinforced silicon nitride, 

so some judgment as to what was the mirror boundary was necessary. The 

zirconia had intrinsic micro hackle lines well within the mirror. The criterion 

for the mirror boundaries was as follows: The mirror boundary was judged to 

be the point where systematic new radiating hackle commenced and there was 

an obvious roughness change relative to the inside-mirror region. 

The word systematic requires some elaboration. Mirror boundary hackle lines 

are velocity hackle lines created after the radiating crack reaches terminal 

velocity. Premature, isolated hackle can in some instances be generated well 

within a mirror, however. It should be disregarded when judging the mirror 

boundary. Wake hackle from an isolated obstacle inside the mirror (such as a 

large grain or agglomerate) can trigger early “premature” hackle lines such as 

shown in Figure 5.5 in the main text. Steps in scratches or grinding flaws can 

trigger hackle lines that emanate from the origin itself. 

5. Measure the mirror size while viewing the fracture surface with an 

optical microscope whenever possible. 

A compound optical microscope is best for glasses. A stereoptical microscope 

is best for ceramics. Use either calibrated reticules in the eyepieces or 

traversing stages with micrometer-positioning heads. For routine 

measurements for ceramics, the author uses a stereoptical binocular microscope 

with a traversing X-Y stage with digital micrometer heads that read out to 

0.001 mm (1 μm). Alternatively, measurements may be made on digital 

images projected onto a high resolution computer monitor preferably while the 

fracture surface can be simultaneously viewed through the microscope 

eyepieces in order to aid judgment. 

D-16
 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Appendix D Guidelines for Fracture Mirrors 

Measurements from photos may be necessary for very small specimens or very 

strong specimens with tiny mirrors, such as in fibers or microelectromechanical system (MEMS) devices. Scanning electron microscope images 

may be used. Again, the fractographer should take an overall framing 

photograph or image shot in accordance with the 75 % to 90 % rule for glasses 

and 33 % to 67 % rule for ceramics. Higher and lower magnification images 

may be used to help aid in interpretation. 

Mirror size measurements from photographs are usually less accurate. They 

frequently overestimate mirror sizes unless conditions are carefully optimized 

to accentuate contrast and topographic detail. Two-dimensional photographic 

renditions of a three-dimensional fracture surface usually lose much of the 

topographic detail discernable by the eye with a compound optical or 

stereomicroscope. 

Mirror size measurements made on computer monitor screens are also subject 

to inaccuracies, also because they are two-dimensional renditions of a threedimensional fracture surface. Video cameras should not be used to capture 

mirror images since they lack adequate resolution. High-resolution cameras 

and monitors are beginning to match the capabilities and accuracy of an 

observer peering through the optical microscope. 

6. Measure radii in directions of approximately constant stress whenever 

possible. A mirror diameter may be measured and halved to estimate the 

radius if the origin site is indistinct or complex. 

There is no consensus on how many mirror radii measurements should be taken 

and in how many directions. Ideally, measurements should be taken from the 

center of the mirror region, but some judgment may be necessary. A common 

practice is to make a judgment whether a mirror is indeed approximately 

semicircular or circular. If it is, then multiple radii measurements may be made 

in different directions and averaged to obtain the mirror size estimate. 

The center of the mirror may not necessarily be the center of the flaw at the 

origin. Careful inspection of tiny localized fracture surface markings (Wallner 

lines and micro hackle lines) right around the origin may reveal that fracture 

started at one spot on a flaw periphery. For example, fracture from grinding or 

impact surface cracks in glass often starts from the deepest point of the flaw 

and not at the specimen outer surface. Figure D.1 shows an example. Large 

pores often trigger unstable fracture from one side. An example is the pore in 

Figure D.6d, where fracture seems to have started on the left side of this 

D-17
 
 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

internal flaw. If an exact mirror center cannot be determined, it is adequate to 

measure a mirror diameter and halve the measurement. This is commonly done 

for semicircular mirrors centered on irregular surface-located flaws whereby 

the mirror center may be difficult to judge. 

Circular embedded mirrors are easiest to interpret (such as in Figure 5.7). 

Small semicircular mirrors on the surface of a part, such as in a bend bar or a 

flexurally loaded plate, are also not too difficult to interpret. The mirror 

relationship holds up remarkably well in glass optical fibers tested in tension 

for mirror radii almost as large as the fiber diameter.17 The mirror radius should 

simply be measured from the origin to the mirror-mist or mist-hackle boundary 

on the opposite side of the fiber, Rd as shown in Figure D.8. 

Rd 

Rd 

Rd 

Figure D.8 Mirrors surrounding surface origins in rods or fibers loaded in 

direct tension. (after Ref. 17) Measure both the mirror-mist radius (shown) 

and mist hackle radii into the depth as shown. 

2Ri 2Ri 2Ri 

(a) (b) (c) 

Figure D.9 Elongated mirrors in bending stress fields. If the mirror is small 

relative to the part size (i.e., a strong part), then the mirror may be 

semicircular as shown in (a). Weaker parts have larger mirrors that flare into 

the interior and are incomplete as shown in (b) and (c). Measure the mirror 

size (Ri or 2Ri for the mirror-mist in the illustrations here) in the direction of 

constant stress. (Note: special guidance on how to deal with small inward 

bending cusps at the surface is shown in Figure D.12 below.) 

D-18
 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Appendix D Guidelines for Fracture Mirrors 

glass ceramic 

(a) (b) 

Figure D.10 Grinding cracks and scratches can cause mirror elongations 

along the surface, even in bend bars with stress gradients. (a) shows a 

schematic of such a mirror with the mist-hackle boundary marked in glass, and 

(b) shows a comparable image in a polycrystalline ceramic. It has some 

intrinsic microstructural roughness inside the mirror and the mirror-hackle 

boundary is marked. Use an average radius: Ravg = {(R1 + R2 + Rd}/3. 

Quarter circular mirrors centered on an edge or chamfer should be measured 

from the origin. 

On the other hand, mirror shapes are commonly affected by stress gradients in a 

plate or a beam. Mirror radii are elongated in the direction of decreasing stress. 

Examples are shown in Figure D.9b and c. In such cases, measure the mirror 

radius along the tensile surface where the stress is constant. Do not measure 

the mirror radii into the gradient. Even with this precaution, there is 

considerable evidence that the data begins to depart from the stress-mirror size 

curves and the relationship in equation D.1 when the sizes approach the plate, 

beam, or rod thickness. For mirrors radii larger than the plate thickness, data 

points fall above the trend on a log stress – log mirror size plot which means 

that mirrors are larger than they would otherwise be if the part were loaded in 

uniform tension. 

A trend for mirrors to elongate the opposite way, along the external surface of a 

specimen, was detected by the author in recent work on the fractographic 

analysis of grinding flaws in structural ceramics.6,7 Long surface cracks often 

caused mirrors to have perceptible deviations from a semicircular shape as 

shown in Figure D.10. In such cases, measure an average radius: 

Ravg = (R1 + R2 + Rd) / 3. (D-4) 

D-19
 
 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

R1 

R1 + R2 

Rd 

R1 

R2 

R1 + R2 

R2 

Rh Rh 

(a) (b) 

(c) (d) 

Figure D.11 Fracture mirrors in rods tested in flexure. The load was applied 

in a vertical direction and the maximum tensile stresses are at the bottom 

center. Fractures started at flaws part way up the sides of the rods causing the 

mirrors to have unequal radii. Schematics (a and b) are similar to two glass 

rods (c and d). The rod shown in (c) was sufficiently strong that a nearly 

semicircular mirror formed, but with unequal radii due to the stress gradient. 

Use R = Rh if the origin and mirror center is distinct. Otherwise use Ravg = (R1 

+ R2 + Rd) / 3 if the mirror is nearly semicircular. Use Ravg = (R1 + R2) / 2 if 

the mirror is elongated into the interior and Rd is large. (d) shows a weaker 

glass rod. Use R = Rh if the origin and mirror center are distinct, otherwise 

use Ravg = (R1 + R2) / 2. 

D-20
 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Appendix D Guidelines for Fracture Mirrors 

In some cases, it may be difficult to measure mirrors in directions of constant 

stress. The two sides of a mirror may have unequal lengths since the stresses 

are different on either side of the mirror. Figure D.11 shows examples from the 

author’s research of round rods broken in flexure. Many origins were not at the 

rod bottom where the stresses were a maximum, but part way up the side of the 

specimen. The specimen orientation was easily determined from observation 

of the compression curl. The maximum tensile stress on the bottom of the 

specimen was on the rod directly opposite the cantilever curl. It was a simple 

task (see section 9 below) to compute the actual stress at the origin location. 

The mirror radii had obviously different lengths due to the stress gradient. Use 

a radius in the direction of constant stress, Rh, as shown in Figure D.11, if the 

mirror is centered on a well-defined origin site. If there is any doubt, then an 

average radius may be computed. Use Ravg = (R1 + R2 + Rd) / 3 if the mirror is 

nearly semicircular. Use Ravg = (R1 + R2) / 2 if the mirror is elongated into the 

interior and Rd is large or is incomplete. 

For origins located in the interior of a rod broken in flexure, only use the radii 

in the direction of constant stress. As a check, and only if the mirror is a 

complete ellipse, one may measure two orthogonal mirror diameters and 

compute an average diameter and halve this to obtain an average radius. 

There is one important detail about mirror sizes that warrants discussion. 

Mirrors located on a specimen external surface have small cusps at the 

intersection with the outer surface as shown in Figure D.12a. Cusps are often 

detected in glass mirrors, but they are rarely if ever discerned in polycrystalline 

ceramic mirrors. The small cusp is a consequence of fracture mechanics. A 

small element of material near the tip of a crack at the specimen exterior 

surface experiences greater stress intensity than a similar element buried in the 

interior whereby neighboring elements can “share the load.” Kirchner et al.4,5 

discussed the shapes of fracture mirrors that intersect outer surfaces and they 

showed that the local enhancement of the stress intensity KI accounts for the 

cusps. The slightly-greater stress intensity at the surface triggers the mirror 

markings a bit sooner than for interior elements. The usual convention, and the 

one adopted in this Guide, is to truncate the cusp. Extend the semicircular (or 

other mirror shape) arcs as shown in Figure D.12c. Another reason to be 

wary of measurements right along the surface is that surface roughness, 

machining damage, or other surface irregularities may trigger mist or hackle 

formation a bit earlier than in the interior. Others have noted that 

measurements taken right on an exterior surface are slightly different than those 

D-21
 
 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

(a) 

(b) 

(c) 

cusp 

Figure D.12 Fracture mirror in a fused silica rod (a), illustrating the cusps in 

the mirror near the outer surface (b). Mirror measurements should not 

include the inward bend of the mirror and may be made as shown in (c). 

D-22
 
 

 

 

 

 

 

 

 

 

 

 

 

 

Appendix D Guidelines for Fracture Mirrors 

taken into the interior. Even Shand18 recommended in 1959 that size 

measurements be taken 0.1 mm (.004 in) beneath the exterior surface. 

Mecholsky and Freiman19 in 1979 recommended ignoring the cusp at the 

surface on the basis of fracture mechanics considerations: “In measuring the 

mirror-mist and mist-hackle boundaries, these should be projected to the tensile 

surface to compete a circular arc, since there is curvature at the surface due to 

free surface effects.” 

A dilemma occurs when mirrors are large in plates or beams broken in flexure 

as shown in Figure D.9c. Cusps cannot be detected. In this case the only 

plausible way to measure a mirror radius is directly along the surface. In such 

cases, the general warning of step 7 applies. 

It is worth reiterating that equations D.1 and D.2 are empirical. Kirchner and 

colleagues4,5 showed a more rigorous fracture mechanics analysis (based on a 

critical stress intensity for branching) completely explained the mirror shapes 

and distortions including the cusp mirrors in various stress fields. Despite the 

merits of their rigorous fracture mechanics analysis, there is significant 

practical advantage to using the simpler mirror size - stress relationships D.1 

and D.2 even if some rigorousness is sacrificed. In addition, for small mirror 

sizes relative to the specimen dimensions, the simpler approach is almost as 

rigorous. 

(a) (b) (c) 

Figure D.13 Surface residual stresses also may alter a mirror shape. σa is 

the applied stress to cause fracture and σf is the fracture stress in an annealed 

plate. (a) shows a surface mirror in an annealed plate. (b) shows the mirror 

shape in a plate with surface compression stresses that decrease into the 

interior and become tensile. (c) shows a mirror in a plate with surface tensile 

stresses that diminish into the interior and become compressive. 

D-23
 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Residual stresses alter mirror shapes. If the mirror is very small relative to the 

stress gradient, the mirror shape may remain circular or semicircular if along 

the surface. On the other hand, if the mirror is larger or the stress gradient is 

steep, then the gradient alters the mirror shape as shown in Figure D.13. Figure 

D.13a shows an annealed plate that requires an applied stress of σa = σf to 

cause fracture. Figure D.13b shows the case where the same plate has 

residual surface compression stress σr = σc from ion exchange or thermal 

tempering, so that an applied stress to cause fracture is σa = σf + σc. In other 

words, the applied stress must be increased to overcome the residual surface 

stress. Nevertheless, the net stress at the surface at the moment of fracture is 

σ = σa - σc = (σf + σc) - σc = σf, the same stress as in the annealed plate. Hence 

the mirror radii along the surface are unchanged compared to the annealed 

plate. In contrast, in the direction into the interior, tensile stresses combine 

with the applied tensile stress to cause the mirror markings to form sooner, at a 

shorter radius into the interior than in the annealed plate. In this example, the 

mirror shape is flattened to a semiellipse. Mirror radii should be measured only 

along the surface (or just beneath the surface if there is a cusp) in these cases. 

Figure D.13c shows that surface residual tensile stresses have the opposite 

effect: mirror radii are elongated into the interior. Mirror radii again should 

only be measured along the surface, since again the net stress to cause fracture 

is σ = σf. 

There are two possible paths for data analysis if there are residual stresses: 

(a) The mirror is measured on a component. The applied stress and the 

residual stresses are unknown. In this instance the net stress σ at the 

origin can be evaluated from Ravg and equation D.1. 

(b) The mirrors are collected in laboratory conditions with multiple 

specimens such that the apparent origin stresses σa, from applied 

external stresses are known. In this instance, one or more matched pairs 

of σa and R are obtained. Graphical analysis as shown in Figures D.1 or 

D.2 reveals the existence and magnitude of the residual stresses. 

7. Exercise caution when mirrors are large relative to the specimen crosssection size. 

At some point, one can expect departures from of the stress - mirror size 

relationship. The point where the departure occurs depends upon the loading 

geometry and the stress state. 

D-24
 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Appendix D Guidelines for Fracture Mirrors 

Pronounced deviations occur once the mirror size approaches or is greater than 

the component thickness in plate or beam bending fractures. Experimentally 

measured radii are much greater than predicted by equation D.1. Shand 

recommended that the maximum mirror size should be no more than 15% of 

the rod diameter (or 30% of the distance to the centroid) for 

flexure tests.18,20 Kirchner and Conway5 warned about limitations in the 

fracture mechanics models for mirror radii exceeding 20 % of a rod diameter 

tested in flexure. On the other hand, Castilone et al.17 had success with mirrors 

that were almost as large as the fiber diameters for fibers tested in tension. 

There is merit to measuring and recording mirrors even if they are large relative 

to a cross section size. The data may have value for use with analyzing genuine 

component fractures. 

Mecholsky and Freiman19 warned that systematic deviations from the mirror 

size relationship occur at large mirror sizes but also at very small sizes, the 

latter due to internal stress effects, e.g., from thermal expansion anisotropy of 

grains in ceramics. 

8 Show at least one photo with arrows or lines marking the mirror size. 

This simple step will help clear up a lot of the doubt and confusion about what 

an investigator has actually measured. 

9. Report how the mirrors were measured. 

This simple last step is often overlooked or ignored, and the reader is left 

wondering exactly what had been done. The fractographer should report the 

microscope used, whether interpretation was made while looking through the 

microscope or from photos, and approximately what magnifications were used. 

The direction the mirror radii were measured should be documented. The 

approximate shape of the mirrors (semicircular, circular, or elliptical) should be 

noted. It should also be noted whether the mirrors were an appreciable fraction 

of the size of the cross section or not. Lastly, and most importantly, the 

judgment criterion used should be reported. 

10. Use the stress at the origin site (additional steps to determine A). 

If the specimen was broken in controlled conditions where the stress 

distribution was known (e.g., beams, rods, or plates in flexural loadings) correct 

the stress for location in specimens with stress gradients. 

No correction is needed if a part was stressed in uniform tension. On the other 

hand, many parts or laboratory specimens do have stress gradients. The general 

principal that should be followed is that the mirror formation is guided by the 

D-25
 
 

 

 

 

 

 

 

 

 

 

y

 Fractography of Ceramics and Glasses 

stresses that acted on and in the immediate vicinity of the flaw origin. While 

this may seem obvious, it is probable that some analysts in the past have 

erroneously used nominal stresses in a specimen or component rather than the 

actual stress that was acting upon the mirror region in a body. In contrast, 

some researchers have correlated the stress at every site along the mirror 

periphery with the mirror radius at that periphery site, but this complex process 

is not practical on a routine basis. 

Example 1: Flexural strength test specimens 

The nominal flexural strength is the maximum stress at the outer fiber or tensile 

surface of the bar. The stress correction for locations beneath the surface is 

y h 

y′

(a) (b) 

(c) 

y 

y′

Figure D.14 The flexural stress gradient is shown in (a) and (b) for a 

rectangular beam. (c) shows how the origin location may be measured from a 

rod neutral axis. 

D-26
 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Appendix D Guidelines for Fracture Mirrors 

simple if the material is linearly elastic. The stress decreases from a maximum 

at the outer tensile surface to zero at the middle “neutral axis” which 

corresponds to the centroid in a square, rectangular, or circular cross-section 

beam specimen as shown in Figure D.14. At the origin, 

σ = (2y/h) σmax (D.5) 

where y is the distance from the centroid and h is the beam thickness or rod 

diameter. Alternatively, in terms of the distance from the tensile surface, y’: 

σ = (1 – 2y’/h) σmax. (D.6) 

In similar fashion, if a fracture origin is not directly under the middle loading 

point in three-point bending, or if the origin is outside the inner gage length in 

four-point bending, then the stress should be correspondingly adjusted to the 

stress that actually acted on the origin site. Corrections for breaks outside the 

gage section or away from the middle loading point are also linear. The stress 

is reduced in proportion to the distance from the fracture plane to the outer 

loading points. This correction requires knowledge of the loading point 

locations. This is an important reason why loading points should be marked on 

any test piece prior to its fracture. For three-point loading: 

σ =(2x/L) σmax (D.7) 

where x is the distance from the outer loading point to the fracture plane and L 

is the total three-point span. For four-point flexure, 

σ = (x/a) σmax (D-8) 

where x is the distance from an outer loading point to the fracture plane, and a 

is the distance from the outer loading point to an inner loading point. If an 

origin in a flexural beam is subsurface and outside the inner span, then both 

corrections should be applied. If the origin is on the tensile surface, within the 

inner span length, and the mirror size is small and semicircular, then no stress 

correction is needed. 

Example 2: Ring-on-ring biaxial strength test specimens. 

If the origin is on the tensile surface in the inner loading circle, and the mirror 

size is small and semicircular, then no stress correction is needed. Fracture 

mirrors from volume flaws located beneath the tensile surface should be 

corrected in the same manner as above for the uniaxial bend test methods. 

D-27
 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Corrections for mirrors located outside the inner loading ring also should be 

corrected for location. The corrections are not as simple in this case, however. 

In the annulus outside the inner loading ring, the radial and circumferential 

(hoop) stresses decrease out to the specimen rim but at different rates.21,

22 

Radial stresses decrease faster, and are zero at the outer rim, and are therefore 

less likely to induce fracture. Fracture planes and mirrors will usually be 

perpendicular to the hoop stresses outside the inner ring, so use the hoop stress 

in the analysis. 

On the other hand, hoop stresses decrease more slowly and are non-zero at the 

outer circumference of the specimen. So although the outer hoop stresses are 

only 10 % to 50 % of the maximum stress in the inner circle (depending upon 

the disk and fixture geometries, and especially the disk thickness), large 

grinding or handling flaws on the specimen rim can cause fracture. 

The same comments apply to piston on three-balls, ball on three balls, or even 

pressure loaded plates in bending. Consult appropriate references as needed 

for the stress distributions in these configurations. 

Example 3: Component failure analyses 

If the mirror constant for the material is known and a mirror size is measured, 

the stress at the origin site can be calculated. The peak stress or the nominal 

stress in the part may be different, however. So as a rudimentary example, 

imagine a bend bar that broke from an internal origin site, but at an unknown 

stress. The fractographer uses the mirror size and the mirror constant to 

estimate the stress at the origin site. The nominal strength of the beam (the 

maximum outer fiber stress) or the “flexural strength” is greater, however. In 

this case, the stress adjustment should use σ, A, h, x, y or y’ to compute the 

greater nominal stress, σmax. 

11. Evaluate the Fracture Mirror Constants 

Once a set of matching mirror radii and fracture stresses is compiled, it is 

customary to plot the data on either a graph of log stress versus log mirror size 

(Figure D-7a) or linear stress versus inverse square root of mirror size (Figure 

D.7b). A linear regression analysis is then performed and a mirror constant 

calculated from the regressed line. The mirror constants should be reported as 

either MPa√m (or ksi√in if the older units are in use). Both are widely found 

in the literature. Each is discussed in turn and it is this author’s conclusion that 

the stress versus inverse square root of size procedure is better. 

D-28
 
 

 

 

 

 

 

 

 

 

 

 

Appendix D Guidelines for Fracture Mirrors 

(a) (b) 

Figure D.7 (duplicate ) Plot of applied stress, σa, versus 1/√R, the preferred 

regression method. (a) shows the trend for parts with no residual stress. (b) 

shows the effects of residual stresses. Compressive residual stresses move the 

line up with a positive intercept σr, but with the same slope. Tensile residual 

stresses shift the data downwards. 

11a. Plot σ versus 1/√R. (Preferred) 

Use linear regression methods to obtain A in accordance with equation D.9 

(same as D.2) with a zero intercept. Use units of MPa for stress and 1/√m for 

the graph axes. In a typical strength test experiment in a laboratory, applied 

stress, σa, and mirror radius, R, are independently measured. It is customary to 

regress σa on R and this procedure shall be followed in this Guide. A is the 

slope of the regression line: 

A (D-9) σ a = R 

Use some judgment in the regression analysis since fracture mirror data usually 

has moderate scatter. If the data does not appear to fit a trend that has a zero 

intercept, regress the data with a non-zero intercept as shown in Figure D.7b. 

Again use some judgment in the interpretation, since a strict linear regression 

fit may produce implausible outcomes, particularly if the data is collected over 

a limited range of mirror sizes and stresses. Report the intercept if it deviates 

significantly (> 10 MPa) from zero. Investigate possible residual stresses or 

specimen size or shape issues in such cases. 

Consistent units should be used. That is to say, if the stress axis is MN/m2

 or 

MPa, then the abscissa (horizontal axis) should be 1/√R where radius is in units 

D-29
 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

of meters. If the mirror is measured in mm or μm, then appropriate conversion 

factors should be added, but this is not trivial and can cause confusion, since 

the square root of a conversion factor of 1000 (e.g., meters to mm) is an odd 

value. (If the stress units are psi or ksi, the mirror radii should be measured in 

inches.) 

The mirror constant as a slope is easily visualized. Deviations from the trend 

usually cause a nonzero intercept, which may be conveniently interpreted as an 

effective residual stress. If residual stresses σr are present in addition to the 

externally applied stress, σa, then the net stress acting on the origin site is: σa + 

σr. The fracture and the mirror markings respond to the actual net stress, σnet: 

A (D.10) 

σ net = (σ a r + σ ) = R 

A

σ a = −σ r (D.11) 

R 

An intercept below the origin corresponds to a net tensile residual stress. A 

positive intercept corresponds to residual compressive stress since the usual 

sign convention is for compressive stresses to have a negative sign. 

Some caution is advised since residual stresses are often nonuniform. The 

residual stress estimated from the intercept is an effective residual stress, which 

in reality may vary in magnitude through the mirror region. If the mirror is in a 

heat-strengthened or tempered piece (where stress may be constant along the 

surface, but change dramatically through the thickness) the mirrors should only 

be measured along the surface or just underneath to avoid the cusp. Residual 

stresses from an indentation or impact site are very local to the origin and may 

have very little effect on a mirror size. 

Although most researchers have felt that the regressed lines should go through 

the origin in annealed test pieces, there is evidence by J. Quinn23 that a small 

but measurable intercept may exist in even annealed materials. The intrinsic 

intercept was evaluated as 10 MPa (1,500 psi) for glass, a value that 

interestingly concurs with Orr’s24 estimate of the minimum stress necessary for 

branching in glass. 

b. Plot log σ versus log R. (Alternative method) 

Use linear regression methods to fit the data in accordance with equation D.12 

(slope set at – 1⁄2) as shown in Figure D.15a. Use units of MPa for stress and 

meters for the graph axes. This will simplify the computation of the mirror 

D-30
 
 

 

 

 

 

 

 

 

 

 

 

 

 

Appendix D Guidelines for Fracture Mirrors 

log 

σa -1⁄2 

(a) 

σ = A / √1 

large 

mirrors 

R = 1 log R Log R = 0 

log
	 slope < - 1⁄2

σa
	 -1⁄2 

(b) 

σ = A / √1 

σr 

R = 1 log R Log R = 0 

Figure D.15 Plot of log σa versus log R for residual stress-free parts (a) and 

parts with residual stress (b). Compressive residual stresses move the locus 

upwards, but with a different slope and intercept. Tensile residual stresses 

move the loci below the baseline curve (not shown). 

constants in units of MPa√m. The mirror constant A corresponds to the stress 

that would create a mirror of size = 1 meter. 

1 log σ a = − log R + log A (D.12) 2 

where σa is the stress at the origin site, A is the mirror constant and R is the 

mirror radius in the direction of constant stress. Use some judgment in the 

regression analysis since fracture mirror data frequently has moderate scatter. 

If the data does not appear to fit a trend with slope -1⁄2, then regress the data 

with equation D.13: 

log σ a = m log R + log A' (D.13) 

D-31
 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

where σa is the stress at the origin estimated from known applied stresses. R is 

the mirror radius in the direction of constant stress. A’ is a modified mirror 

constant. Report the slope m and the alternative constant A’. Again use some 

judgment in the interpretation, since a strict linear regression fit may produce 

implausible outcomes, particularly if the data is collected over a limited range 

of mirror sizes and stresses. Investigate possible residual stresses or specimen 

size or shape issues if m deviates significantly from the value –1⁄2. 

If stresses are in units of MN/m2 (MPa) and the mirror size is measured in 

meters, then the mirror constant A has units of MN/m1.5 or MPa√m. If the 

mirror size is 1 m, then log R = 0. Then log σ = log A and hence, σ = A. Thus, 

the mirror constant A corresponds to the value of stress that would create a 

mirror of size 1 m. 

Use consistent units with this approach for the same reasons mentioned in 

section 11a. If the stresses are in MPa, then the abscissa (horizontal axis) 

should be with radii in units of meters or micrometers. (If the stress units are 

psi or ksi, the mirror radii should be in inches.) 

Since most actual mirrors that are measured are usually much smaller than unit 

size, it is apparent from Figure D.15 that the mirror constant (or the stress at R 

= 1) lies somewhat beyond the range of data usually collected. This method of 

showing the results and calculating a mirror constant was common in the older 

technical literature and is occasionally still found today. Deviations from the 

linear relationship on a log - log plot occur when residual stresses are present 

but unaccounted for (since the plotted stress may not be the true stress at the 

origin), or when the mirror size becomes large relative to the component size, 

or when there are stress gradients. The residual stress deviations cause the line 

to have a slope other than -1⁄2 as shown in Figure D.15b. Attempts to compute 

the residual stresses may then be made by guessing values of the residual 

stresses σr, replotting the data with a vertical axis of (σa - σr) , and checking the 

goodness of fit of a line of slope -1⁄2. This is a fairly cumbersome process and 

the alternative procedure in section 11a and Figure D.7b may be simpler and 

more effective. 

If a single mirror is measured, and the externally applied stress, σa, and the 

mirror constant A are both known, then: 

(σ R = A (D.14) a + σ r )

and A

σ = R −σ a (D.15) r 

D-32
 
 

 

 

 

 

 

 

 

 

 

 

 

Appendix D Guidelines for Fracture Mirrors 

Mecholsky and colleagues25,26 have shown excellent examples how residual 

surface stresses in tempered or clad glasses may be estimated from equation 

D.8 and fracture mechanics analysis. 

11c. Comparison of the two curve fitting and regression approaches 

The merits of the two plotting - regression schemes vary. They put different 

weights on large and small mirror measurements. In one case, the mirror 

constant is a slope of a line, in the other it is an intercept at R = 1, a rather large 

mirror size not likely to be realized in practice. 

Some researchers may have evaluated test pieces with residual stresses, and 

then force-fitted regression lines through the data with zero intercepts in the 

former scheme, or lines of slope -1⁄2 through the data in the latter scheme. This 

probably contributed to variability in published mirror constants. 

Figure D-16 shows both graph types for a common set of data from the author’s 

work on the effects of machining on the strength of silicon nitride.6,7 A large 

number of rod- or bar-shaped specimens were tested. They had various grinding treatments. The mirror constants from the two analyses are quite similar, 

although there is a small difference (9%) between the constants for the rods and 

bars. 

(a) (b) 

Figure D-16 Fracture minor data for silicon nitride.6,7 The round blue 

symbols are for 60 rods and the square red symbols are for 80 bars. 

D-33
 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Data from flexurally-loaded specimens frequently deviate from the trends when 

mirror sizes are large and are a significant fraction (50 % or greater) of the 

cross section size. Upward deviations from the log stress - log radius graphs 

have been noted in a number of studies (e.g., Shand,20,27 Orr24). The 

experimentally measured mirrors are larger than they otherwise would be in a 

uniform tension stress field. Regression lines chase the upward deviations from 

the trend and dramatically alter the estimate of the mirror constant. On the 

other hand, with the σ versus 1/√R graph, large or oversized mirror data points 

are closer to the origin and have less influence on the regression line, and hence 

have less effect upon the slope. 

Regression analysis with the σ versus 1/√R approach minimizes the deviations 

of σ from the fitted line. Regression analysis for the log σ versus log R 

approach minimizes deviations of log σ from the fitted line. 

In the 1950s and 1960s many researchers plotted log stress versus log radius 

probably because they were not confident of the theoretical justification for the 

-1⁄2 slope. They let the exponent vary on the log-log plots and discussed the 

differences, if any, relative to -1⁄2 power. Differences were usually due to 

residual stresses or overly-large mirrors relative to the specimen size. 

Gradually the case for the σ√R relationship solidified and more researchers 

began to plot stress versus 1/√R. 

In summary, the linear stress versus inverse square root radius approach is 

superior and is recommended by this Guide. Analyses are simple and intuitive. 

The mirror constant is the slope of the regressed line. Residual stresses may be 

interpreted from non-zero intercepts. The uncertainty of the slope can be 

estimated from routine analyses available in many statistical software packages. 

Data deviations due to large mirrors in flexure specimens have less effect upon 

the regression process and the mirror constant estimates. The alternative 

method (log stress – log radius) may be useful in some cases for displaying data 

with an unusually large range of mirror sizes and stresses.28 

12. Mirrors sizes should be collected over a broad range of sizes and 

fracture stresses if possible. Data from different specimen types and sizes 

may be combined. 

This is a fairly obvious conclusion in light of the discussion in the previous 

paragraph. Superb examples are shown by Kerper and Scuderi28, for 

borosilicate glass rods with diameters that varied by a factor of ten, and by 

Mecholsky and Rice29 for various sized fused silica rods, disks, and fibers. 

D-34
 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Appendix D Guidelines for Fracture Mirrors 

Ideally, data from many small specimens could be complemented by testing of 

a few large specimens. Another common procedure is to anneal or fine 

grind/polish some specimens to obtain high strengths, but abrade or damage 

others to obtain low strengths. Sometimes the mode of loading can be changed 

to alter the fracture stress. For example, some studies have used mirrors with 

large four-point and small three-point flexure specimens. Some specimens may 

be tested in inert conditions and others tested in conditions conducive to slow 

crack growth. 

D.4 Some Final thoughts 

The goal of these guidelines is to bring some consistency to procedures used to 

measure fracture mirrors. They were adopted as ASTM Standard C1678 in 

2007. This should facilitate improved data bases and better estimates of failure 

stresses. These Guidelines were prepared on the basis of the author’s own 

experiences as well as a careful review of sixty years of literature as summarized in reference 1. One is struck by the conclusion, which nearly all writers 

have reached, that measurement of the mirror sizes requires subjective 

interpretation. The perception of the observer and the type of equipment are 

important factors. Although advanced microscopy and software tools hold 

considerable promise in the future, it is unlikely that a simple definitive 

criterion (such as a set level of surface roughness) will emerge. Despite this, 

most students of the technique have concluded that consistent readings are 

possible between observers. The quotation from Johnson and Holloway3 at the 

start of this Appendix is one example. Another is from Mecholsky and 

Freiman19 who said: 

“While one might think initially that the measurements of a mirror 

boundary using a microscope is quite a qualitative operation and would 

vary from observer to observer; in fact, experiments performed over a 

number of years by a large number of investigators have shown that the 

values of mirror constants obtained in different laboratories are quite nearly 

the same.” 

They then listed some values for a few glasses and ceramics that did in fact 

vary as much as 20 % to 30 %. For example, the soda-lime glass values varied 

by 23 %. Hopefully, adoption of the guidelines in this Guide and ASTM C 

1678 will improve the consistency of future data to within 10%. 

One positive conclusion is that fracture mechanics principles do seem to 

account for the observed shape variations in mirrors.4,5 (A caveat is that the 

fracture mechanics analyses have been based on static loading, whereas the 

D-35
 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

mirror boundaries are formed by a very dynamic crack traveling at or near 

terminal velocity.) A logical conclusion of this finding is that equation 1 is 

over simplistic, since it does not account for stress gradients or geometric 

effects. Nevertheless, it does have the proper functional form for the stress 

intensity of a crack in a far-field tension stress. For a small mirror in a uniform 

tensile stress field, equation 1 is completely justifiable. 

Another interesting finding from the literature review is how few authors have 

shown good fracture mirror photos. Even fewer have marked them. Perhaps 

the authors were not sure or were hesitant to show an interpretation for fear of 

criticism. Hopefully, step 8 of this Guide (which requires marked photos) and 

the examples shown in this Guide will help future authors improve their 

reporting. One is also struck by the fact that nearly all the mirrors shown in the 

literature, even in the classical papers, are not exactly semicircular or circular, 

despite all the schematics that imply that they are. Fractographers should not 

be alarmed if their mirrors are not perfect. 

Many of the steps in this Guide have already been proposed. Shand recommended that stresses be corrected for the origin location,20 that radii be 

measured beneath the surface to avoid surface effects,18 that low-angle grazing 

illumination be used.20 He also warned about deviations from the trends if the 

mirror sizes were too large relative to the component thickness.20 Shand also 

said that mist could not be discerned in glass ceramics.20 Morrell et al. 30 

agonized over the interpretation of mirrors in Y-TZP zirconia (such as shown in 

Figures D.4 – D.6), but settled on a set of specific criteria. Optical microscopy 

with a stereo optical microscope at a fixed magnification was used, with 

grazing incidence illumination. No reflective coating was applied to the 

surface. The specimen sides were masked to block transmitted-light scatter. 

Matching fracture halves were mounted together to aid the interpretation. The 

best set of recommendations predating this Guide were crafted by Mecholsky 

and Freiman.

19 Six of their recommendations match steps in this Guide: (a) 

optical microscopy is preferred over scanning electron microscopy whenever 

possible, (b) suitable magnifications should be used, (c) mirror boundary arcs 

should be projected to the outer surface to complete a circular arc to eliminate 

the surface cusps, (d) lighting should be varied to obtain optimum contrast, (e) 

radii should be measured in directions of constant stress and not into gradients, 

and (f) trend deviations occur for mirrors that are large relative to part 

thickness. 

D-36
 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Appendix D Guidelines for Fracture Mirrors 

Appendix D References 

1.	 G. D. Quinn, “Guidelines for Measuring Fracture Mirrors,” pp. 163 – 190 in: 

Fractography of Glasses and Ceramics, V, eds., J. R. Varner, G. D. Quinn, M. 

Wightman, American Ceramic Society, Westerville, OH, 2007. 

2.	 ASTM C 1678, “Standard Practice for Fractographic Analysis of Fracture Mirror 

Sizes in Ceramics and Glasses,” Annual Book of Standards, Vol. 15.01 ASTM Int. 

West Conshohocken, PA, 2007. 

3. J. W. Johnson and D. G. Holloway, “On the Shape and Size of the Fracture Zones 

on Glass Fracture Surfaces,” Phil. Mag., 14 (1966) 731 - 743. 

4. H. P. Kirchner and J. W. Kirchner, “Fracture Mechanics of Fracture Mirrors,” J. 

Am. Ceram. Soc., 62 [3-4] (1979) 198 - 202. 

5. H. P. Kirchner and J. C. Conway, Jr., “Criteria for Crack Branching in Cylindrical 

Rods: I, Tension and II, Flexure,” J. Am. Ceram. Soc., 70 [6] (1987) 413 - 425. 

6. G. D. Quinn, L. K, Ives, and S. Jahanmir, “On the Fractographic Analysis of 

Machining Cracks in Ground Ceramics: A Case Study on Silicon Nitride,” SP 996, 

NIST, Gaithersburg, MD, May, 2003. 

7. G. D. Quinn, L. K. Ives, S. Jahanmir, and P. Koshy, “Fractographic Analysis of 

Machining Cracks in Silicon Nitride Rods and Bars,” pp. 343 - 365 in 

Fractography of Glasses and Ceramics IV, Ceramic Transactions, Vol. 122, eds., 

J. R. Varner and G. D. Quinn, American Ceramic Society, Westerville, OH, 2001. 

8. G. D. Quinn, J. Eichler, U. Eisele, and J. Rödel, “Fracture Mirrors in a Nanoscale 

3Y-TZP,” J. Am. Ceram. Soc., 87 [3] (2004) 513 - 516. 

9. W. H. Duckworth, D. K. Shetty, and A. R. Rosenfeld, “Influence of Stress 

Gradients on the Relationship Between Fracture Stress and Mirror Size for Float 

Glass,” Glass Technol., 24 [5] (1983) 263 - 273. 

10. D. M. Kuluwansa, L. C. Jensen, S. C. Langford, and J. T. Dickinson, “Scanning 

Tunneling Microscope Observations of the Mirror Region of Silicate Glass 

Fracture Surfaces,” J. Mater. Res., 9 [2] (1994) 476 – 485. 

11. D. Hull, “Influence of Stress Intensity and Crack Speed on Fracture Surface 

Topography: Mirror to Mist Transition,” J. Mat. Sci., 31 (1996) 1829 - 1841. 

12. D. Hull, Chapter 5, in Fractography, Observing, Measuring and Interpreting 

Fracture Surface Topography, Cambridge Univ. Press., Cambridge, 1999. 

13. C. Wünsche, E. Rädlein, and G.H. Frischat, “Morphology of Silica and 

Borosilicate Glass Fracture Surfaces by Atomic Force Microscopy,” Glastech. Ber. 

Glass Sci. Technol., 72 (1999) [2] 9 - 54. 

14. E. Gölz, “Übermikroskopische Feinstructuren an Glasbruchflächen,” Zeitschrift für 

Physik, 120 (1943) 773 - 777. 

15. E. K. Beauchamp, “Mechanisms of Hackle Formation and Crack Branching,” pp. 

409 - 446 in Fractography of Glasses and Ceramics III, eds., Varner, Fréchette, 

Quinn, Cer. Trans., Vol. 64, 1996, American Ceramic Society, Westerville, OH. 

16. E. K. Beauchamp, “Fracture Branching and Dicing in Stressed Glass,” Sandia 

Research Report SC-RR-70-766, January 1971, Sandia National Laboratory, 

Albuquerque, NM. 

D-37
 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

17.	 R. J. Castilone, G. S. Glaesemann, and T. A. Hanson, “Relationship Between 

Mirror Dimensions and Failure Stress for Optical Fibers,” pp. 11 - 20 in Optical 

Fiber and Fiber Component Mechanical Reliability and Testing II, eds., M. J. 

Matthewson and C. R. Kurkjian, Proc. SPIE, 4639 (2002). 

18. E. B. Shand, “Strength of Glass – The Griffith Method Revised,” J. Amer. Ceram. 

Soc., 48 [1] (1965) 43 - 48. 

19. J. J. Mecholsky and S. W. Freiman, “Determination of Fracture Mechanics 

Parameters Through Fractographic Analysis of Ceramics,” pp. 136 -150 in 

Fracture Mechanics Applied to Brittle Materials, ASTM STP 678, ed., S. W. 

Freiman, ASTM Int., West Conshohocken, PA, 1979. 

20. E. B. Shand, “Breaking Stress of Glass Determined from Dimensions of Fracture 

Mirrors,” J. Am. Ceram. Soc., 42 [10] (1959) 474 - 477. 

21. H. Fessler and D. C. Fricker, “A Theoretical Analysis of the Ring-On-Ring 

Loading Disk Test,” J. Am. Ceram. Soc., 67 [9] (1984) 582 - 588. 

22.	 J. A. Salem and L. Powers, “Guidelines for the Testing of Plates,” Cer. Eng. Sci. 

Proc., 24 [4] (2003) 357-364. 

23. J. B. Quinn, “Extrapolation of Fracture Mirror and Crack-Branch Sizes to Large 

Dimensions in Biaxial Strength Tests of Glass,” J. Am. Ceram. Soc., 82 [8] (1999) 

2126 - 2132. 

24. L. Orr, “Practical Analysis of Fractures in Glass Windows,” Materials Research 

and Standards, 12 [1] (1972) 21 - 23, 47. 

25. J. C. Conway, Jr. and J. J. Mecholsky, Jr., “Use of Crack Branching Data for 

Measuring Near-Surface Residual Stresses in Tempered Glass,” J. Am. Ceram. 

Soc., 72 [9] (1989) 1584 - 87. 

26. J. J. Mecholsky and M. G. Drexhage, “Comparisons of Optical and Fractographic 

Measurements of Residual Stress in Compressively Clad Glass Rods,” J. Am. 

Ceram. Soc., 63 [5-6] (1980) 347 - 349. 

27. E. B. Shand, “Breaking Stresses of Glass Determined from Fracture Surfaces,” The 

Glass Industry, April (1967) 190 - 194. 

28. M. J. Kerper and T. G. Scuderi, “Relation of Fracture Stress to the Fracture Pattern 

for Glass Rods of Various Diameters,” Bull. Am. Ceram. Soc., 45 [12] (1966) 1065 

- 1066. 

29. J. J. Mecholsky, Jr. and R. W. Rice, “Fractographic Analysis of Biaxial Failure in 

Ceramics,” pp. 185 - 193 in Fractography of Ceramic and Metal Failures, eds., J. 

J. Mecholsky, Jr. and S. R. Powell, ASTM STP 827 American Society for Testing 

and Materials, West Conshohocken, PA, 1984. 

30. R. Morrell, L. Byrne, and M. Murray, “Fractography of Ceramic Femoral Heads,” 

pp. 253 - 266 in Fractography of Glasses and Ceramics, IV, Ceramic 

Transactions, Vol. 122, eds., J. R. Varner and G. D. Quinn, American Ceramic 

Society, Westerville, OH, 2001. 

D-38
 
 

 

 

 

 

 

 

 

 

 

 

Index 

INDEX 

A 

Abrasion	 6-39, 6-66 – 6.69 

Acid, aqua regia	 3-26 

Acid storage tank	 9-6 

Agglomerate, see origins 

Alumina	 4-18, 4-32, 4-45, 5-19, 5-68, 5-71, 6-7, 

6-8, 6-11, 6-12, 6-14, 6-23, 6-30, 6-70, 

6-77 

10-30 –10-37, 10-43 – 10-45 

Alumina, β	 6-16 

Alumina furnace plate	 4-32 

Alumina crown	 4-45, 10-30 – 10-37, 10-43 – 10-45 

Alumina single crystal, see sapphire 

Alumina, whisker reinforced	 7-71 

Aluminum oxynitride	 5-34, 5-67, 5-74, 6-8, 6-13, 6-75, 10-16 – 

10-17 

Aqua regia	 3-26 

Arrest lines	 5-58 - 5-60 

Atomic Force Microscope	 3-60 

Automobile window	 4-31, 4-50, 7-7 

B 

Ball bearings	 3-68, 6-39 

Ball mill	 6-60 

Baseline microstructural flaws	 6-75 

Batman mirror	 8-19 

Bend bar	 4-14 – 4-18, 10-19 – 10-25 

Bevels, see chamfers 

Biaxial disk 3-23, 4-5, 4-9, 4-19 – 4-21, 5-72, 6-29, 

6-31, 7-4 

Biaxial stress 1-3, 1-4, 4-19 – 4-21 

Bifurcation 

Black light 3-55 

Blanchard grinding 6-40 

Bleach 3-6 

Blisters, see origins 

Blood 5-63 

Boron carbide 4-50, 7-61 

Bottles 4-7, 4-38, 7-3, 10-14, 10-15 

Boundary phases, ceramics 5-78, 5-79 

I-1 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Branching 4-3 – 4-10, 7-7 – 7-9 

Branching angles 4-6 – 4-9 

Branching constant 4-9, 7-8, C-1 

Branching distance 4-9, 7-7 - 7-9 

Branching constants 4-9, Appendix C 

Brick, Roman 4-41 

Bubbles 6-7 – 6-11, 6-78, 6-79 

C 

C crack 3-68, 6-35 

Camera 3-7 – 3-11, 3-14 – 3-16 

Camera, single lens reflex 3-8 

Camera stand 3-7 

Camera mount, F-mount or C-mount 3-14 

Cantilever curl 4-15 – 4-17, 5-2 

Capacitor, multilayer 3-23 

Cathedral mirror 8-9, 8-10, 8-16, 8-24 

Cavitation, scarp 5-61, 5-62 7-53 

Cell phone 3-10, 3-11 

Cellulose acetate 3-36 - 3-38, 3-42 

Center heated plate 4-32, 4-33 

Chain of events 1-10 

Chamfers 6-16, 6-23, 6-40, 6-53 

Characteristic strength 7-72 

Chatter marks, see origins 

Chill checks, see origins 

Chips, chipping, see edge chips 

Clay 3-5, 6-70, 6-71 

Cleaning coatings off 3-26, 3-71 

Cleaning specimens 3-5, 3-6, 3-71, 5-1 

Cleaning, dental plaque 3-6 

Cleaning, green dye 3-29 

Cleaning, removal of gold coatings 3-26 

Cleaner, ultrasonic 3-71 

Cleavage 5-3, 8-1, 8-2, 8-5 

Cleavage step hackle 8-8 

Coatings 3-26 – 3.29 

Combination flaw 6-73, 6-74 

Component fracture patterns 4-22, 4-44, 4-45 

Composites 6-20, 7-24, 7-72, 9-1 – 9-11 

Compositional inhomogeneity, see origins 

Compound optical microscope 3-32 – 3-35 

I-2 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Index 

Compression curl 4-15 – 4-18 

Compression fracture 4-51 – 4-54 

Conchoidal fracture 5-3 

Cone cracks, see Hertzian cone crack 

Confocal microscope 3-62, 3-63 

Contact cracks 3-68, 4-30, 4-37, 6-34– 6-39, 10-6, 10-9, 

10-10, 10-31, 10-46 - 10-48 

Contaminants, see also clay 6-70 - 6-72 

Conversion factors 7-1 

Cooking ware 6-63, 6-66, 6-68 

Cords, see origins 

Corner hackle 10-36 

Corrosion flaws 4-49, 6-28 

Crack branching 4-3 – 4-10, 7-7 – 7-9, 7-24 

Crack bridging 7-24, 7-55 

Crack front shape 5-46 

Crack healing 4-28, 7-43 

Crack, processing 4-10, 6-54 - 6-58 

Cracks, intersecting 4-26, 4-27 

Crack velocity 7-46 – 7-53 

Crack velocity, terminal 5-6, 5-7, 7-46 – 7-49 

Crazing 6-78 

Creep fracture 5-77 – 5-81 

Critical strain energy release rate, GIc 7-22 

Crossing cracks 4-27 

Crowns, see dental 

Cusps 6-76 

D 

Damage wave 5-6 

Dandruff 6-69 

Dark field illumination 3-20, 3-24, 3-32, 8-5, 8-6 

Debonding 10-20, 10-21 

Debris on fracture surface 5-83, 5-84 

Deep Reactive Ion Etching (DRIE) 6-76, 8-23, 8-24, 10-25 – 10-27 

Deer 4-23 

Defect, see also flaws and origins 6-1 

Delaminations 6-57 

Delta markings 5-3, 5-36 

Dental ceramics, crowns, bridges 3-38, 3-40, 3-41, 10-30 – 10-45 

Alumina 3-41, 4-45, 6-55, 6-58, 6-87, 10-30–10-37, 

10-43 – 10-45 

I-3 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Bovine dentin 4-15 

Bridge, lithium disilicate 6-79 

Bridge, zirconia 6-55, 6-78 

Cerestore alumina-magnesia spinel 10-31 – 10.33 

Composites, filled and unfilled 6-20, 6-61, 6-69, 9-6 – 9-8 

Denture material 4-43 

Dicor 10-30 

Empress II glass ceramic 5-34 10-37 – 10-42 

Human enamel 4-4 

Lithium disilicate glass ceramic 4-41, 6-60, 6-71, 6-79, 10-37 – 10-42 

Porcelain 4-3, 4-43, 6-10 

Resin, PEG with ACP filler 3-23 

Resin infiltrated ceramic, Enamic 6-60 

Replicas 3-38 – 3-41 

Veneers, glazes 5-33, 5-34, 6-78, 6-79 

Zirconia, see also zirconia 4-43, 6-55, 6-58 

Dental plaque 3-6, 6-71, 6-78 

Devitrification stones, see origins 

Differential Interference Contrast D-12, D-13 

Digital camera, photos 3-3 – 3-11, 3-14 – 3-16 

Digital image formats 3-8, 3-9 

Digital image processing 3-17 

Digital microscope 3-19, 3-20 

Digs 6-31, 6-33, 6-34 

Discussion stereo microscope 3-18, 3-19 

Disk specimen, see biaxial disk 

Dome, IR 8-7, 8-21, 10-16 – 10-17 

Dome, see sapphire domes 

Double cantilever beam specimen 8-3 

Double torsion specimen 5-36, 5-49, 8-3 

Drawings 4-24 – 4-26 

Dunt crack 6-58 

Dust 6-69 

Dye staining 3-28, 3-29, 4-48 

Dye penetration 3-66 – 3-70, 4-48, 6-51 

Dynamic fatigue (variable rate strength test) 7-52 

E 

E-glass composite 9-4, 9-6 

Edges 6-1, 6-2, 6-40, 6-53 

Edge chips 4-40 – 4-44, 10-36, 10-37, 10-43 – 10-45 

Elastic waves 5-6, 5-47 – 5-58, 7-46 – 7-50 

I-4 
 

 

 

 

 

 

 

 

 

 

 

 

Index 

Electrical insulators 3-37, 6-8, 6-37 

Electronic part, multilayer capacitor 3-23 

Electronic part, thick film resistor 5-65 

Electronic part, varistor 6-16 

Epoxy, replica 3-40, 3-42, 3-45 

Electron backscatter diffraction 3-55 

Extrinsic flaws 6-4 

F 

Failure analysis 1-6 – 1-8 

Far field stress 7-21 

Fatigue crack 3-24, 5-81 – 5-84 

Fiber composite 9-1 – 9-6 

Fiberglass- epoxy composite 9-4, 9-5 

Fiber reinforcement 7-24, 9-1, 9-2 

File formats, digital images 3-8, 3-9 

Finite element analysis 1-5, 4-46, 7-18 

Firing cracks 4-10, 6-54 – 6-56, 6-58 

Flaw 6-1 

Flaw bluntness 7-20, 7-39 – 7-41 

Flaw linkage 6-50, 6-73, 6-74, 6-80 – 6-82, 7-44, 7-45 

Flaw management 6-75, 7-75 

Flaw shielding 7-44, 7-45 

Flaw size 6-3, 6-75, 7-18, 7-19 

Flaw size distribution 7-73, 7-74 

Flaw truncation 6-77, 7-42, 7-44, 7-45 

Flexural strength 4-14 – 4-18, 5-20 

Flexural stress gradient D-18, D-19, D-26, D-27 

Flow lines of stress 7-20, 7-21, 7-44 

Fluorescent dye penetration 3-67 – 3-68 

Focus stacking 3-17 

Ford Rotor 10-1 – 10-7 

Ford Turbocharger 6-56 

Fractal analysis 7-65 – 7-67 

Fractography, definition 1-1 

Fractography, Laws of 11-1 – 11-3 

Fracture, conchoidal 5-3 

Fracture energy 7-18, 7-22 

Fracture map 5-80, 5-81 

Fracture mirror, definition 5-3 

Fracture mirror 5-3 – 5-30, Appendices C and D, 

see also mirror 

I-5 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Fractographic montages 5-84, 5-85, 6-84 – 6-86 

Fracture origin, definition 4-1 

Fracture origins, see origins 

Fracture toughness 7-21 – 7-25, 7-40 

Fracture toughness, single crystal 8-2, 8-3 

Fragmentation 4-10 – 4-12, 4-19, 7-2 – 7-7 

Furnace plate 4-32, 10-18, 10-19 

G 

Gas turbine rotor 10-1 – 10-7 

Glass 

Borosilicate, Pyrex 7-16 

Borosilicate crown 4-5, 4-21, 6-31, 6-59, 7-4, 7-34, 7-61 

Fused silica 5-2, 5-4, 5-5, 6-6, D-3, D-22 

Fused silica, sintered 10-20 – 10-22 

Soda lime 4-23, 6-18, 6-57 – 6-64, 7-3, 7-5, 7-6, 

7-9 

Glass bottles, see bottles 

Glass ceramic 6-7, 6-11, 6-60 

Glass ceramic, dental 3-29, 6-60 

Glass disk, see biaxial disk 

Glass rod, flexure 1-1, 4-15, D-3, D-20, D-22, D-26 

Glass tube 4-36 

Glass windows, see windows 

Glazes, veneers 5-33, 5-34 

Gloves, cotton 3-2 

Glue, specimen reconstruction 4-2, 4-3 

Glue chips 5-63 - 5-66 

Grazing illumination 3-12, 3-22 

Green pen, dyes 3-28, 3-29, 10-42 

Griffith flaw 6-1, 7-18, 7-19 

Griffith criterion, equation 7-18, 7-19 

Grinding cracks, see origins, grinding cracks 

Grinding damage, glass disks 4-20, 4-21 

Gull wings 5-14, 5-49, 5-50, 7-46, 7-47 

Gull wing mirrors 8-9, 8-15, 8-18 

Gun Barrel 10-7 – 10-10 

H 

Hackle 5-10 – 5-16, 5-31 - 5-44 

cleavage step 8-8 

corner 5-43 

I-6 
 

 

 

 

 

 

 

 

 

 

 

 

Index 

definition 5-12 

coarse microstructural 5-31, 5-32 

grinding crack 6-40 – 6-52 

fine microstructural 5-36, 5-41 

mist hackle 5-12 

shear hackle 5-42 

step hackle 5-43, 5-44 

twist hackle 5-36 – 5-41 

velocity hackle, definition 5-12 

wake hackle 5-33 – 5-36, 5-49 

Hailstone 7-4 

Hair 3-23, 6-69 

Halo, slow crack growth 5-71 

Hand magnifier 3-2 

Handling damage 6-29, 6-36, 6-39, 8-7 

Healed cracks 4-28 

Heat sleeks 6-63 

Hertzian cone cracks 6-34 – 6-40, 10-10, 10-31, 10-47, 10-48 

High speed photography 3-64 

Hinge fracture 4-34 – 4-36 

History of brittle materials fractography 1-8, 1-9 

Holders 3-4, 3-5 

Hoop stress 4-33, 4-37 – 4-38 

Hopf bifurcation 4-32 

Human enamel 4-4, 4-46 

Human skin contamination 6-70 

Hybrid flaws 6-73, 6-74 

Hydrogen peroxide 3-29 

I 

Illumination 3-12, 3-20 – 3-26 

Impact 4-29, 4-35, 4-49 

Impact origins, see origins impact 

Impact, tertiary Wallner lines 5-53 

Inclusion, see origins-inclusion 

Incomplete bonding 6-16 

Indentation fracture mechanics 7-25, 7-61 – 7-65 

Insect 3-37, 6-69 

Insulator, electric 6-8 

Intrinsic flaws 6-3, 6-4 

Intergranular fracture 5-17, 5-66 - 5-70, 5-75, 5-79 

Intersecting cracks 4-26, 4-27 

I-7 
 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Invisible cracks 4-28 

IR dome 8-7, 8-21, 10-16, 10-17 

J 

Jeweler’s Loupe 3-3, 3-4 

Joint 5-3 

JPEG format 3-8, 3-9 

K 

Kitchenware 6-63, 6-66, 6-68 

Knoop indentation crack, see origins 

K-V diagram (slow crack growth) 7-51 – 7-53 

L 

Laboratory fractures 4-11, 4-44 

Lambda lines 5-58 

Lamp fracture 4-11, 7-7 

Lances, see also hackle 5-16, 5-36, 8-13, 8-17 

Law of normal crack propagation 4-3 

Lead zirconium titanate (PZT) 5-35, 6-57 

Leader crack 4-35, 4-36 

Lint 6-8, 6-69 

Lithic fractures 5-3, 5-62 

Lithic fractures - Obsidian 5-35, 5-49, 5-62 

Litigation 1-8 

Longitudinal grinding 6-40, 6-41 

Loupes 3-3, 3-4 

Low-angle grazing illumination 3-12, 3-22, 5-19, 6-45, 8-5 

M 

Machine gun barrel 10-7 – 10-10 

Machine shops 10-23 – 10-25 

Machining cracks, see origins 

Machinist loupe 3-3, 3-4 

Macrofractography camera stand 3-7 

Magnesium fluoride 5-71 

Magnesium oxide 8-1 

Magnesium aluminate spinel, see spinel 

Magnifying glass 3-2 

Mandelbrot relationship 7-66 

Margin damage, dental crowns 10-32 – 10-37, 10-43 – 10-45 

Maximum distortion energy criterion 1-5 

I-8 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Index 

Mechanical engineers 1-4 

Mechanical overload 4-47 

Medicinal vial 4-37, 6-39, 10-14, 10-16 

Microcracking 7-24 

Microelectromechanical structures (MEMS) 8-23 – 8-24, 10-25 – 10-27 

Microflaw pocket 6-80 – 6-82 

Micrometer, damage from 6-36 

Micrometer, stage 3-30 

Microscope 

Atomic force 3-60 

Compound optical 3-32 – 3-35 

Confocal 3-62 

Discussion stereo optical 3-19 

Environmental scanning electron 3-59 

Field emission scanning electron 3-58 

Scanning electron 3-43 – 3-57 

Stereo optical 3-11 – 3-16 

Transmission electron 3-59 

Mirror, fracture 

Alumina bend bar 5-19 

Borosilicate crown glass 5-22, 5-26 

Bent 5-22 

Glass rod 1-1, 3-61, 3-63, 5-2, 5-4, 5-6, 5-24 

Medicinal vial 4-37 

Multiple 5-20, 5-21 

Silicon carbide 5-18 

Silicon nitride 5-20, 5-27 

Tempered glass 5-29, 5-30 

Zirconia 3-27, 5-17, 5-19 

Mirrors, single crystals 8-2, 8-6, 8-8, 8-9 

Batman 8-19 

Cathedral 8-9, 8-10, 8-16 

Gull wings 8-9, 8-15, 8-18 

Mirror constants 5-13, 7-10–7-16, 7-40, 7-41, 

Appendices C and D 

Mirror cusp D-21 – D-24 

Mirror, fibers 9-5 

Mirror size, radius 7-10 – 7-16, Appendix D 

Mirror size, residual stresses 7-69, 7-70, D-20 – D-23 

Missile domes 

Missile radomes 10-20 – 10-23 

Mist 5-9 – 5-16 

I-9 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Mist hackle, definition 5-12 

Mohr’s circle 1-5 

Montages 5-84, 5-85, 6-84 – 6-87 

N 

Near surface 6-3 

Neutral axis, bending 4-14 – 4-16 

Nickel sulfide 6-18 

Nomarski illumination 5.62, D-12, D-13 

Nomenclature 1-8, 1-9, 5-1, 5-3, 6-1 

O 

Obsidian 5-14, 5-35, 5-49, 5-62 

Optical comparator 3-3 

Optical profilometer 3-60 – 3-62 

Origin 6-1 

Abrasion tracks 6-66 – 6-48 

Agglomerate 3-51, 6-6, 6-13, 9-1 

Ball mills 6-5, 6-69 

Baseline microstructure 6-75 

Blank pressing flaws 6-11 

Blisters, glass 6-61 

Bubbles 6-7, 6-9, 6-10, 6-11 

Chatter marks, cracks 5-29, 6-65 – 6-68 

Chill checks 4-28, 4-48, 6-64, 6-65 

Chips 6-53 

Contact cracks 3-42, 3-68, 6-34 – 6-39 

Cords, glass 6-61 

Compositional inhomogeneity 3-51, 6-21, 6-22 

Cusps, geometric irregularities 6-76 

Dandruff 6-69 

Delaminations 6-58 

Devitrification stones, inclusions 6-61 

Dunt cracks 6-57 

Etch groove 10-26, 10-27 

Etch pits 8-24 

Fatigue crack 3-24 

Feces 6-69 

Fiber 6-8 

Fiber clump 9-1 

Firing cracks 6-54 – 6-56, 6-58 

Geometric sharp points 6-76 

I-10 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Index 

Grain boundary 6-25, 6-26, 6-75 

Grinding cracks, see origins, machining 

Hair 3-23, 6-8 

Handling damage 6-29, 6-36, 6-39, 8-7 

Heat sleeks 6-63 

Hybrid 6-73, 6-74 

Impact, blunt 4-29, 4-30, 6-34 - 6-39 

Impact, sharp 4-29, 4-30, 6-33, 8-22 

Inclusion 3-52 – 3-54, 6-17 – 6-20, 6-61, 9-8 

Insect 6-69 

Knoop indentation crack 3-29, 3-68, 5-32, 5-71, 5-73, 6-59, 6-60, 

7-61, 7-63, 8-2 – 8-4, 10-12 – 10-14, D-3 

Large grain 6-22 – 6-24 

Lint 6-8, 6-69 

Machining cracks 6-6, 6-28 – 6-36, 10-5, 10-6, 10-16–10-18, 

10-20, 10-21, 10-23 

Herring bone 6-51 

Orthogonal cracks 6-41 – 6-43 

Parallel cracks 6-41 – 6-46, 6-48 

V-grinding cracks 6-42, 6-46, 6-47 

Zipper cracks 6-44 – 6-46 

Machining crack skin zone 6-46 

Microflaw pockets 6-80 – 6-82 

Nickel sulfide, glass 6-18 

Pecks 6-33, 6-63 

Pits 6-28 

Polishing surface flaws 3-65, 5-72, 6-29 – 6-32, 8-21, 10-48 

Pore 6-4, 6-8, 6-13 

Porous region 6-13, 6-14, 10-12, 10-13 

Porous seam 6-14, 9-3 

Pressing flaws 6-11 

Processing cracks 4-10, 6-14, 6-54 – 6-58, 6-83 

Quartz grains, Quartz inversion 6-24, 6-57 

Scratch 6-6, 6-29 – 6-32, 8-7 

Seeds, glass 6-61 

Silicon vein 6-69 

Skin 6-70 

Sleeks 6-65 

Stones, glass 6-61 

Striae, glass 6-61 

Surface void 6-27 

I-11 
 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Vickers indentation crack 5-22, 6-59, 6-33, 7-25, 7-25, 7-59, 7-61 – 

7-65 

Whisker clump 9-1 

Origin, fracture definition 4-1 

Orthogonal machining crack, see origins 

Oxidation 4-49, 6-28 

P 

Parallel machining crack, see origins 

Parfocal	 3-12, 3-20 

Parting	 5-3 

Patriot missile	 10-20 

Pattern recognition	 1-3 

Pecks	 6-33, 6-63 

Penny shape cracks	 7-26, 7-29, 7-40 

Pharmaceutical vials and syringes	 4-37, 5-62, 6-39, 10-14, 10-15 

Phase instability	 6-69 

Photos, Photography	 3-7 – 3-11 

Pits, see origins 

Plaque, dental	 3-6 

Plasma asher	 8-27 – 8-29 

Plane strain	 1-5 

Plane stress	 1-5 

Plaque, dental	 3-6, 6-71 

Plate fracture patterns	 4-28 – 4-30 

Plate fracture patterns, thermal	 4-32, 4-33 

Plate, whiteware	 3-42 

Polariscope	 3-70, 3-71 

Polarizer	 3-25, 3-26, 8-5, 8-26 – 8-29, 

Polished microstructural section	 6-5, 6-75, 6-82, 6-83 

Polishing surface flaws, see origins 

Polycrystalline alumina (PCA)	 10-17 

Porcelain	 3-37, 3-42, 3-66, 4-41, 6-8, 6-10, 6-24, 

6-37, 6-52 

Porcelain electrical insulator	 3-37, 6-8, 6-37 

Polyvinylchloride (PVC), replicas	 3-38 

Polyvinylsiloxane (PVS) replicas	 3-38, 3-41 

Pore, see origins 

Porous region, see origins 

Porous seam, see origins 

Pressing flaws	 6-11 

Pressure flaking	 5-63 

I-12 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Index 

Pressure vessel 4-7, 4-33 – 4-38 

Processing cracks, see origins 

Proof test, radome 10-20 

PZT, see Lead zirconium titanate 

Q 

Quartz grains, Quartz inversion 6-24, 6-57 

Quasi static loading 6-35 

R 

R-curve 5-73, 6-48, 7-23, 7-25, 7-55 – 7-60 

Radome, fused silica 10-20 – 10-22 

Rayleigh wave velocity 5-5 

RAW format 3-8 

Reliability 10-1 – 10-7 

Renegade abrasive grit 6-48, 6-49 

Replicas 3-35 – 3-42, 6-82 

Residual stress 4-49 – 4-50 

Compression induced cracking 5-63 

Estimates of magnitude 7-68– 7-70, D-30 – D-34 

Grinding 6-40, 7-68 

Heat strengthened glass 5-28, 5-29 

Indentation 6-59, 6-60, 7-63 – 7-65, 7-70 

Mirror shapes, effect on 5-28, D-18 – D-24 

Stable crack extension from 5-73 

Tempered glass 4-28 – 4-31, 4-49, 5-30 

Resin-matrix composite, dental 6-20, 6-21, 6-69, 9-6 – 9-8 

Retroreflective flaws 6-69 

Rib marks 5-45, 5-58 

Rings of Saturn 6-7, 7-20, 7-39, 7-40 

Ripples 5-45, 5-56 

River deltas 5-3, 5-36 

Rock fracture, see lithic fracture 

Roman brick 4-41 

Rotor, gas turbine 10-1 – 10-7 

Round robin, flexural strength 10-23 – 10-25 

Round robin, fractography 10-28 – 10-30 

Round robin, machining 10-23 – 10-25 

S 

Saliva 5-63
 

Sapphire, single crystal alumina 3-65, 6-31, 8-4, 8-7, 8-19, 8-21, 8-22,
 

I-13 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

8-25 - 8-29, 10-46 – 10-48 

Sapphire dome 3-65, 4-25, 5-62, 6-31, 8-7, 8-21, 8-22 

Sapphire scarp 5-62 

Sapphire tubes 8-27 – 8-29 

Scanner 3-71 

Scanning electron microscope 3-43 – 3-57 

Scarps 5-61, 5-62, 7-53, 8-20, 10-14 

Scratch, see also origin 6-6, 6-30 – 6-32 

Scratch/Dig specification 6-31, 6-33 

Secondary fracture, edge chips 4-44 

Secondary origins 4-3, 4-18, 4-44 

Secondary Wallner line 5-50, 5-53 

Seeds, see origins 

Self-toughened ceramics 9-2 

Shark’s teeth 6-42 – 6-46 

Shear hackle 5-42 

Shrink fit 10-7 – 10-9 

Sialon 6-18, 6-30, 7-72, 7-73 

Side wall damage 6-52 

Sierra scarp 5-61, 5-62 

Silica, fused 5-2, 5-4, 5-5, 6-6, 10-20 – 10-22, 

D-3, D-22 

Silicon 6-76, 7-33, 8-2, 8-3, 8-6, 8-20, 8-23, 8-27 

Silicon carbide 4-4, 5-18, 5-40, 5-62, 5-75, 6-5, 6-8, 6-13, 

6-16, 6-21, 6-23, 6-37, 6-53, 6-59, 6-69, 

6-70, 6-83, 7-32, 9-2, 9-3, 

10-9 – 10-15, 10-18, 10-19, D-5 

Silicon carbide, CVD 5-40, 9-3, 10-25 – 10-27 

Silicon carbide, fibers 9-1 – 9-3, 9-6 

Silicon carbide machine gun barrel 10-7 – 10-10 

Silicon carbide furnace plate 10-18, 10-19 

Silicon nitride 3-51, 3-52, 3-53, 4-9, 4-15, 5-17, 5-21, 

5-27, 5-32, 5-68, 5-73, 5-76, 5-77, 5-79, 

5-80, 5-81, 6-6, 6-8, 6-13, 6-16 – 6-18, 

6-21 – 6-23, 6-27, 6-28, 6-30, 6-32, 6-36, 

6-47 – 6-49, 6-51, 6-53, 6-54, 6-56, 6-60, 

7-58, 

9-2, 9-3, 10-1 – 10-7, D-6, D-33 

Single crystals 8-1 – 8-32 

Alumina, see sapphire 

Spinel 8-9 - 8-18, 8-25 

Silicon 6-76, 7-33, 8-2, 8-3, 8-6, 8-20, 8-24 

I-14 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Index 

Zirconia 8-1 

Sketches 4-24 – 4.26 

Sleeks 6-65 

Slow crack growth 4-51, 5-69 – 5-76, 7-23, 7-25, 7-51 – 

7-54, 10-1 – 10-5, 10-11 – 10-14 

Slow crack growth exponent, N 7-51, 7-52 

Soda Lime silica, see glass 

Spinel, magnesium aluminum oxide 6-25, 6-26, 8-9 - 8-18, 8-25 

Spray-dried particles 6-5, 6-15, 6-16, 6-55 

Stable crack extension 5-69 - 5-76, 7-42, 7-43, 7-51 – 7-60, 10-11 

Staining 3-26– 3-29, 3-66 – 3-70 

Standards 

ASTM C 1322, fractography 6-73, 6-84, 7-28, 7-32, 8-9, 10-28
 

ASTM C 1421, fracture toughness 7-28, 7-31
 

ASTM C 1678, fracture mirrors C-2, D-1
 

ISO 14704, flexural strength 10-23
 

ISO 18756, fracture toughness 7-31
 

MIL STD 1942 (MR), flexural strength 10-23
 

MIL HDBK 790, fractography 10-28
 

Standard Reference Material 2100, KIc 7-25 

Stage micrometer 3-30 

Stage, traversing 3-31, 3-34 

Static fatigue 7-51, 7-52, 10-11 – 10-14 

Steel, fatigue crack 3-24 

Step hackle 5-43, 5-44 

Stereoptical microscope 3-1, 3-11 – 3-16 

Stereoscope 3-57 

Stones, glass, see origins 

Streaks 5-61, 5-62 

Stress 1-4, 7-17 – 7-19 

Stress concentration 7-17, 7-18, 7-20, 7-39 

Stress, biaxial 1-4 

Stress corrosion 10-13 

Stress, far field 7-21 

Stress, flexural 4-14 – 4-17 

Stress intensity 7-17 - 7-29 

Stress intensity shape factor 7-21 – 7-38 

Stress intensity factors, Newman-Raju 7-30 – 7-36, 7-53 

Stress rupture 7-52, 10-3, 10-11 – 10-14 

Stress state 1-4 

Stress, triaxial 1-4 

Stress uniaxial 1-4 

I-15 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 Fractography of Ceramics and Glasses 

Stress wave fractography 

Striae, see also origins 

Striations, fatigue cracks 

Striations, grinding 

Subcavitation hackle 

3-64
 

6-61
 

5-81, 5-82
 

6-6, 6-41, 6-43
 

5-61, 5-62, 5-72
 

Surface crack – fracture mechanics 7-27 – 7-38, 7-53, 7-54
 

Surface crack in flexure, see origins Knoop 6-59
 

Surface finish 

Surface grinding 

Surface void, see origins 

Syringes 

T 

Teeth 

Telescoping denture 

Tempered glass, heat strengthened glass 

Tension strength 

Terminal velocity 

Tertiary Wallner lines 

Thermal crack 

Thermal fracture 

Thermal shock 

Thermal stresses 

Thermocouple bead origin 

Theta strength specimen 

Time dependent fracture 

TIFF format 

Torsion 

Transformation toughening 

Transgranular fracture 

Transillumination 

Transmission electron microscope 

Transmitted illumination 

Transverse grinding 

Triaxial stress 

Tube fracture 

Tungsten inclusion 

Turbocharger, silicon nitride 

Twinning 

Twist hackle 

6-40
 

6-41 – 6-52
 

6-39
 

4-4, 4-46
 

6-58
 

3-10, 4-11, 4-29 – 4-31, 6-18, 6-66, 6-68, 

7-5 - 7-9
 

4-13, 10-25 – 10-27, D-5
 

5-6, 5-7
 

5-53 – 5-57
 

6-57
 

4-32, 4-33, 4-47, 4-48, 10-18,
 

10-46 – 10-48
 

4-47, 4-48, 6-58, 8-27
 

6-58, 8-7, 8-27, 10-18, 10-46 – 10-48
 

6-69
 

8-23, 8-24
 

4-51
 

3-8, 3-9
 

4-7, 4-18, 4-39
 

7-24, 7-25, 7-57
 

5-17, 5-66, 5-67, 5-69, 5-71, 10-11, 10-12
 

3-24, 3-25, 4-3, 4-4, 8-5, 10-27
 

3-59, 5-79
 

3-24, 3-25, 4-3, 4-4, 8-5, 9-8,10-27
 

6-41, 6-44 – 6-49, 6-52
 

1-4
 

4-36, 8-27 – 8-29, 10-7– 10-9
 

6-18, 6-19
 

6-56
 

3-64, 4-54, 8-5, 8-25 - 8-30, 10-46 – 10-48
 

5-36 – 5-42
 

I-16 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

V 

Index 

U 

Ultrasonic cleaner 3-71 

Ultrasonic fractography 3-64, 5-56, 5-57 

Uniaxial stress 1-4 

USB microscope 3-19 

VAMAS 10-28 

Vanadium inclusion 6-17 

Velocity hackle 5-11, 5-12, 5-31 

Velocity, terminal 5-6, 5-7 

Veneer, dental 6-36, 6-71, 6-77 – 6-79, 6-87 

Vials, medicinal 4-37, 5-62, 6-39,10-14, 10-15 

Vicinal illumination 3-22, 3-23 

Vickers indentation flaws, see origins 

Video camera 3-15 

V machining crack, see origins 

Von Mises failure criterion 1-5 

W 

Wake hackle 5-33 – 5-36, 5-49, 9-3, 10-16, 10-17, 

10-33 

10-35, 10-36, 10-45 

Wallner lines 5-39– 5-58, 7-46 – 7-50, 8-8 

Gull wings 5-14, 5-49, 5-50 

Lambda 5-58 

Primary 5-45 – 5-47 

Secondary 5-50 – 5-52 

Tertiary 5-53 – 5-55, 8.22 

Watch bracelet 6-56 

Watchmaker’s loupe 3-3 

Water hammer 4-34, 4-35 

Wave velocity, see elastic waves 

Weibull distribution 6-86, 7-71 – 7.76 

Weibull modulus 7-59, 7-71 – 7-76 

Whisker lances, single crystal 8-9, 8-13, 8-17 

Whisker reinforcement 7-24, 9-1, 9-2 

Wing cracks 4-53 

Window failure 4-23 – 4-33 

Window failure, glue chips 5-66 

Window patterns 4-28 – 4-33 

I-17 
 

 

 

 

 

 

X 

 Fractography of Ceramics and Glasses 

Window, thermal fracture 4-32, 7-2 

Witness marks 4-22, 4-49, 6-33, 6-34, 6-39 

X-ray topography 3-65, 8-5, 8-7 

X-Y stage, traversing 3-31, 3-34 

Z 

Zinc oxide, varistor	 6-16 

Zipper crack, see origins, machining 

Zirconia, cubic single crystal	 8-1 

Zirconia, magnesia stabilized	 5-31, 6-25, 6-26, 6-81, 6-82, 7-23, 7-57, 

7-58 

Zirconia, yttria stabilized	 5-17, 5-19, 6-8, 6-13, 6-18, 6-56, 6-58, 

6-72, 7-14, 7-15, 7-16, 7-60, D-7 – D-9 

I-18