TR-91-5

Air Force Office of Scientific Research

PANEL DISCUSSION

ON

ORIGIN OF IMPERFECTIONS

1 November 1991 Holiday Inn, Dayton Ohio

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Panel Discussion on Origin of Imperfections

6. AUTHOR(S)

George Sendeckyj Spencer Wu

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Air Force Office of Scientific Research Building 410 Boiling AFB, DC 20332-6448

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Air Force Office of Scientific Research Building 410 Boiling AFB, DC. 20332-6448

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ASIAC-TR-91-5

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13. ABSTRACT (Maximum 200 words)

This report is a surrmary of a panel discussion on Origin of Imperfections, and the presentations that were included. This meeting was sponsored by AFOSR and was held at Wright-Patterson AFB, Ohio on 9 October 1991.

The panel discussion had two major objectives. The first objective was to present current research activities under the AFOSR Initiative on the "Origin of Imperfections." The second objective was to obtain a consensus on science issues in this critical research area and to identify topics for future research efforts.

14. SUBJECT TERMS Acoustic Microscopy, Imperfections, Acoustic,

Emission, Non-destructive Evaluation, NDE, Defects, Failure, Materials, Composites, Micro-plasticity, Micro-Fracture, Micro-

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FOREWORD

The Air Force Office of Scientific Research (AFOSR) sponsored two contractor

meetings on the "Mechanics of Materials" and "Structural Dynamics" at the Holiday

Inn in Fairborn, Ohio during the week of 7 October 1991. The meetings were hosted

by Mr. Robert M. Bader, Chief of the Structures Division, Flight Dynamics

Directorate, Wright Laboratory at Wright-Patterson Air Force Base. The technical

sessions were organized and coordinated by Dr. Spencer Wu, AFOSR Program

Manager for Aerospace Sciences. A panel discussion on the Origin of Imperfections

was conducted in conjunction with these meetings on Wednesday, 9 October. This

panel session was chaired by Dr. George Sendeckyj, of the Structural Integrity Branch

of the Structures Division.

This report includes a summary of the panel discussion and the six technical

presentations that were part of the program.

The administrative arrangements for the AFOSR meetings and panel

discussions were conducted by the Aerospace Structures Information and Analysis

Center (ASIAC), which is operated for the Flight Dynamics Directorate, Wright

Laboratory by CSA Engineering, Inc. The efforts of Mr. Gordon R. Negaard, ASIAC

Technical Manager, and all the ASIAC staff for coordinating and arranging the

meeting details are greatly acknowledged.

li

TABLE OF CONTENTS

Section Page

I Panel Discussion Summary - Dr. George Scndeckyj, Wright- Laboratory 1

1. Introduction 2

2. Summary of Formal Presentations 3

2.1 Introductory Comments by Sprncer Wu, AFOSR/NA 3 2.2 "The Microscopic Origin of Inelastic Behavior" 3 2.3 "Quantitative Non-Destructive Evaluation of Fatigue- Induced

Surface Damage by Line-Focus" 3 2.4 "Origins of Imperfections in Composites" 5 2.5 "Examination of Precursors to Fracture" 5 2.6 "Hysteresis and Incremental Collapse of Structures, A Paradigm

for the Strength of Metals" 6 3. Summary of General Discussion 7 4. Conclusion 10

II Technical Presentations: 11

Dr. Spencer T. Wu, AFOSR 12

Prof. J. D. Achenbach, Northwestern 17

Prof. D. E. Chimenti, John Hopkins University 33

Prof. J. T. Dickinson, Washington State University 46

Prof. S. A. Guralnick, IIT 78

Dr. Harold Weinstock, Boiling AFOSR 125

in

SECTION I

ORIGIN OF IMPERFECTIONS PANEL

DISCUSSION SUMMARY

Dr. George P. Sendeckyj

Aerospace Research Scientist

Fatigue, Fracture & Reliability Group

Structural Integrity Branch

Structures Division

Flight Dynamics Directorate

Wright Laboratory

United States Air Force

Wright-Patterson Air Force Base, Ohio

1. INTRODUCTION

A panel discussion on the "Origin of Imperfections" was held in Dayton OH on 9

October 1991 as part of the Air Force Office of Scientific Research (AFOSR) Contractors

Review. The discussion consisted of brief introductory remarks by Dr Spencer Wu,

AFOSR/NA, five (5) formal presentations by panel members, and a general discussion on the

origin of imperfections. The presentation material is reproduced in the Appendix, with some

minor editing to correct obvious typographical errors.

To put the panel discussion in proper perspective, it is necessary to define what is

meant by an imperfection. With this in mind, an imperfection is defined as a shortcoming,

defect, fault or blemish that has a detrimental effect on the performance of a structural

material system for a particular application. Imperfections can occur on different scales in

structural materials. At the macro or structural scale, the nature of imperfections in materials

is well understood and their effect on structural performance can be predicted with a high

degree of accuracy. Our understanding of the nature of imperfections and their consequences

on structural performance is relatively poor at the mini- and micro-scales, and almost

non-existent at the nano-scalc.

The Aerospace Sciences Directorate of the Air Force Office of Scientific Research has

a small research initiative to develop a fundamental understanding of the origin of

imperfections. The objectives of the initiative are to develop experimental tools for observing

the origin of imperfections at the micro- and nano-scale and their evolution to the macro-scale

in structural materials, and to formulate theoretical models of the damage evolution process.

The panel discussion had two major objectives. The first objective was to present

current research results obtained under the AFOSR Initiative on the "Origin of Imperfections."

The second objective was to obtain a consensus on science issues in this critical research area

and to identify topics for future research activities.

2. SUMMARY OF FORMAL PRESENTATIONS

2.1 Introductory Comments by Dr Spencer Wu, AFOSR/NA.

Dr Spencer Wu, AFOSR/NA, opened the panel discussion with a short introductory

presentation. He introduced the panel members, indicated the objectives and funding level of

the AFOSR Initiative on the Origin of Imperfections and stated the objectives to be achieved

during the panel discussion.

2.2 "The Microscopic Origin of Inelastic Behavior" by Dr Harold Weinstock,

AFOSR/NE

Dr Weinstock presented some research results on the use of the Barkhausen effect to

study the onset of hysteresis and changes in surface residual compressive stresses in

ferromagnetic materials. He became interested in this area when he saw some results on the

Kaiser effect in acoustic emission monitoring. He decided to determine whether the Kaiser

effect also occurs in the magnetic field induced in ferromagnetic materials under mechanical

loading. He performed some experiments and observed that induced magnetization undergoes

a jump at approximately 60% of elastic limit strain. This jump has been interpreted as an

indication of the occurrence micro-plasticity, which can be interpreted as the onset of

hysteresis. He is pursuing further research in this area. A practical application of the

Barkhausen effect is in inspection of surface hardened steel structures, which rely on residual

compressive surface stresses to achieve required performance.

2.3 "Quantitative Non-Destructive Evaluation of Fatigue-Induced Surface Damage by

Line-Focus Acoustic Microscopy" by J. D. Achenbach, M. E. Fine, J. O. Kim and S. M.

McGuire, Northwestern University

Dr Achenbach presented results of an experimental investigation of surface

micro-cracking in SiC whisker reinforced aluminum alloys using scanning acoustic

microscopy. He indicated that Dr Fine, one of his colleagues, had performed some tedious

experiments on the evolution of surface micro-cracking in SiC whisker reinforced aluminum

alloys. He presented some of those result on the surface crack length distributions as a

function of number of fatigue cycles for whisker reinforced aluminum composites by way of

an introduction of their current research on the application of line-locus acoustic microscopy

to damage detection and quantification. He described the operation of the scanning acoustic

microscope and discussed how data on the stale of near surface imperfections are generated.

He presented some results on the variation of surface wave speed with anisolropy and number

of fatigue cycles. He indicated that the missing link is the correlation of surface wave speed

data with the actual damage state. He believes that this correlation can only be obtained by

analyzing wave propagation in a half-space with a random distribution of surface cracks. He

and his colleagues are working this problem.

2.4 "Origins of Imperfections in Composites" by D. E. Chimenti, Johns Hopkins

University

Dr Chimenti discussed the application of scanning electron thermal microscopy

(SETM) to the detection of imperfections in composite materials. The SETM process consists

of inducing thermal waves in a sample inside a scanning electron microscope (SEM). The

thermal waves are due to localized periodic heating of the surface and they can be used to

image subsurface features. The SETM has high spatial resolution that is limited by the elastic

wave length. He also described their acoustic microscope. Finally, he presented the goals of

the funded research work.

2.5 "Examination of Precursors to Fracture" by J. T. Dickinson, Washington State

University

Dr Dickenson discussed his research on the examination of precursors to fracture. By

way of a general introduction, he made some general comments on the failure process and

sources of pre-existing and grown-in defects. He also discussed some work on the simulation

of the damage evolution process. After the introductory comments, he described his research

work on fracto-emission during the failure process. He presented results on the emission of

argon during tensile fracture of argon sorbed polystyrene. The results indicate that argon

emission correlates with craze volume evolution during the fracture process. He also

presented results on electron emission during fracture of polycarbonate, alkali emission from

NaCl, and 02 emission during microcracking in single crystal MgO. The results indicate that

there are many different precursors to failure that have not been fully exploited by the

research community. Dr Dickinson concluded his presentation with the following list of

research issues:

a. Detection and characterization of defects

b. Description and kinetics of evolution of damage

c. Materials design to control degradation

d. Statistical physics concepts (fractals, self-organized criticality, scaling phenomena)

e. Experimental advances

i. sensitivity to differences (i.e., perfect versus flawed, defects as a

function of time)

ii. surface versus interior phenomena

iii. probes in hostile environments (temperature, atmoshere/chemical, radiation)

f. Modeling efforts (global, atomic, etc.)

2.6 "Hysteresis and Incremental Collapse of Structures, A Paradigm for the Fatigue

Strength of Metals" by S. A. Guralnick and T. Erber, Illinois Institute of Technology

Dr Guralnick presented results of a research program with the objectives to (a)

correlate fatigue and hysteresis by recognizing the processes which begin with

micro-imperfections, progress through macro-imperfections and culminate in fatigue rupture

and (b) correlate acoustic emission phenomena with the inception of the damage growth

process which ultimately results in fatigue rupture. He began the presentation discussing

possible defects in metal crystals and polycryslals, and possible micromechanical responses of

metals to cyclically-varying loads. He then presented various experimental evidence to

support the concept that microplaslicity is a precursor to fatigue crack initiation, and outlined

two scenarios for the progression of the fatigue process. His model of the fatigue processes

consists of the following steps:

a. Inception of microplaslicity

b. Cooperative organization of zones of microplasticity

c. Damage growth and accumulation leading to micro-fracture

d. Macro-crack initiation

e. Crack propagation and organization leading to rupture of the material.

Dr Guralnick outlined a number of diagnostic techniques for studying the damage

evolution process in metals. In addition to those discussed by the previous speakers, he

mentioned acoustic emission monitoring as a viable damage documentation technique. He

also presented some acoustic emission results for aluminum. Finally, he indicated directions

for future research that he plans to pursue.

3. SUMMARY OF GENERAL DISCUSSION

The formal presentations described a number of experimental tools for observing

microscopic imperfections in materials and their evolution into macroscopic defects that lead

to structural failure. The experimental tools addressed by the presenters include acoustic

emission monitoring, acoustic microscopy, scanning electron thermal microscopy,

fracto-emission monitoring, and magnetic effect monitoring. They permit the observation of

surface or near surface imperfections and their evolution. Since x-ray techniques can easily

image interior defects in solids, Dr. T. Moran of the Materials Directorate of the Wright

Laboratory was asked to brielly indicate the capabilities of their micro-focus computer-aided

tomography (CAT) scanner. He indicated that their unit has the capability to image defects

and imperfections with dimensions on the order of 250 micro-meters. The CAT scanner has

another unique capability in that it can be used to experimentally determine internal strain

distributions in materials with texture. By imaging the texture in the material under different

loads and superimposing the images, interference patterns related to relative displacement are

formed. The patterns can be analyzed to determine strain distributions. Dr Moran indicated

that their CAT scanner does not possess the required resolution for this purpose, but other

CAT scanners with the required resolution are available within the research community.

After the comments on the CAT scanner, the discussion was opened for general

comments and questions from the audience. Dr Chimcnli was asked how severe the localized

heating could get in the SETM. He indicated that local temperatures can be very high (on the

order of thousands of degrees Centigrade). This led to a discussion of how to determine that

the SETM gives a true indication of the local subsurface features. A procedure for

establishing the electron beam modulation parameters was suggested. This indicated a set of

experiments that should be performed in the SETM.

The discussion then turned to where the new experimental tools were leading the

research community and whether the study of the origin of imperfections was necessary. It

was pointed out that new research tools for documenting imperfections and their evolution

have better resolution than previously available tools and that the resolution continues to be

improved. The question was asked when this will stop. By way of a response, the analogy

between the development of medical and materials diagnostic tools was made. In the early

days of medicine, section of cadavers was used to infer information about the human body.

The counterpart of this is the use of sectioning to study damage in composite materials, which

is tedious. Then the stethoscope was developed as means of diagnosing living organisms.

The conventional ultrasonic techniques are the counterpart of the stethoscope. They provide

much useful information on the nature of the damage evolution process in composite

materials, but at present do not have the required resolution to guide the formulation of

models of the damage evolution process at the lamina level. The x-ray techniques, including

CAT scanning, were developed for the medical profession to provide improve diagnostic

information. They were adapted to damage documentation in composite materials. As a

result, physically realistic models of the damage accumulation process are being developed.

These models are sufficiently accurate for engineering simulation of the damage accumulation

process in composite materials, but they do not provide adequate information on the

imperfections that lead to the damage observed by the x-ray techniques. Thus, better research

tools are required to identify the imperfections that lead to the damage. Such tools are the

acoustic microscope, the scanning electron thermal microscope, and fracto-emission

monitoring. Whether these research tools are adequate to determine the origins of

imperfections is not clear at present, but their use is a step in the right direction. It was the

consensus of the panel and audience that research to develop experimental tools with still

better resolution will improve the understanding of the behavior of materials and should be

pursued. The concern was raised about reaching the point when the observation tools affect

the observations, that is, reaching an uncertainty point. This was not fully addressed during

the discussion.

The question was raised about what kind of defects are detectable with field

instrumentation. This is an important consideration since it addresses the practical problem

faced by the materials user community. The materials user is only interested in when the

material develops defects that threaten safely and affect durability of the structure. He is also

interested in how often the structure must be inspected using the available inspection tools

and how fast can the inspections be performed. It is doubtful that inspection tools with

higher resolution will impact the damage tolerance philosophy and design criteria. It is

expected thai inspection tools with higher resolution will impact the durability design

procedures, since they arc based on the ability to model the growth and find small

imperfections. As a consequence, it is important to be able to correlate the amount of initial

imperfections with the resulting defects and to determine how soon the imperfections evolve

into defects that lead to eventual failure of the structure.

It became obvious during the discussions that the primary reason for studying the

origin of imperfections is to develop the necessary understanding to design better materials

for structural applications. This requires an in-depth understanding of imperfections and how

they evolve into defects, such as cracks, that finally lead to structural failure. This requires

the use of modern experimental tools to observe the origin of imperfections and their

evolution into structural defects. At the very least this research will lead to a better

understanding of the behavior of structural materials. It may also lead to a breakthrough in

materials design once the defect causing imperfections are understood and means for

eliminating them are developed.

Finally, the research on the origin of imperfections is an excellent example of natural

philosophy. The term natural philosophy is used for the scientific process of observing and

developing theories to explain phenomena. In the present case the phenomenon of interest is

the evolution of structural defects from material imperfections. Experimental observations are

necessary to guide the formulation of models and simulations of the damage evolution

process. The models and simulations must capture the salient features of the damage

accumulation process. Moreover, experiments must be designed to determine the accuracy

and soundness of the models and simulations. If successful, this will lead to a better

understanding of the behavior of structural materials. Moreover, this understanding can be

used to design improved structural materials.

4. CONCLUSIONS

The major conclusions of the panel discussion on the origin of imperfections arc:

1. Understanding of imperfections and their evolution into defects that eventually lead

to structural failure is necessary. This will not only lead to a better understanding of material

behavior, but will also provide guidelines for the design of belter structural materials.

2. The research tools being developed for studying the origin of imperfections include

acoustic microscopy, scanning electron thermal microscopy, micro-focus tomography, etc.

With the exception of acoustic emission monitoring and micro-focus tomography, these tools

only permit of observation of the evolution of damage from near surface imperfections.

3. It is unlikely that an understanding of the origin of imperfections will impact the

damage tolerance design philosophy. It will have an eventual impact on the durability design

philosophy.

10

SECTION II

PANEL DISCUSSION

ON

ORIGINS OF IMPERFECTIONS

TECHNICAL PRESENTATIONS

Dr. Spencer T. Wu Program Manager, Aerospace Sciences

Boiling Air Force Base "Introduction and Overview"

Professor J. D. Achenbach Northwestern University

"Quantitative Non-Destructive Evaluation of Fatigue-Induced Surface Damage"

Professor D. E. Chimenti John Hopkins University

"Origins of Imperfections in Composites"

Professor J. T. Dickinson Washington State University

"Examination of Precursors to Fracture"

Professor S. A. Guralnick Illinois Institute of Technology

"Hysteresis and Incremental Collapse of Structures"

Dr. Harold Weinstock Boiling Air Force Base

"The Microscopic Origin of Inelastic Behavior"

11

TECHNICAL PRESENTATION

11 Panel Discussion on Origin of

Imperfection "

BY

Dr. Spencer T. Wu

AFOSR

Boiling Air Force Base

12

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TECHNICAL PRESENTATION

" Quantitative Non-Destructive

Evaluation of Fatigue-Induced

Surface Damage"

BY

Professor J. D. Achenbach

Center for Quality Engineering and Failure Prevention

Northwestern University

17

Material:

Silag SiC (20 vol.% whisker) / Al (2124-T4) whiskers : diameter ~ 1jim, length ~ 10 ^im

rolled (tends to align whiskers at surface => surface anisotropy)

Loading : cyclic (0.5 Hz)

tension (stress range 222 M Pa)

QNDE technique :

line-focus acoustic microscope 225 MHz

V(z) curves at increasing numbers of cycles (N)

surface wave speed vs. angle (anisotropy) for increasing N

surface wave speed vs. N for various angles

18

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20

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21

The Scanning Acoustic Microscope

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At 200 MHz the wavelength in water is 7.5 /mi and the congressional

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information from the acoustic microscope by relating the voltage from the

microscope to the mechanical properties of the insonified material.

22

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23

METAL MATRIX COMPOSITE

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TECHNICAL PRESENTATION

"Origins of Imperfections in

Composites1 f

BY

Professor D. E. Chimenti

John Hopkins University

33

Objectives

• Conduct research leading to the identification and char- acterization of origins of imperfections in advanced com- posite materials.

• Demonstrate optimum characterization methods for the defects to which the imperfections lead.

• Detect and characterize incipient damage in composites that have been mechanically, thermally, and environmen- tally stressed.

• Examine occurrence, development, and growth of imper- fections such as matrix plastic deformation, microcrack- ing, fiber fracture, fiber/matrix debonding, and other po- tential failure modes.

34

Incipient Damage

As-fabricated It has been demonstrated that origins of fu- ture damage are often introduced into the material in the fabrication stage

Mechanical Loading in situ static and cyclic stress to induce the nucleation and development of mechanical damage

Thermal Loading Electron beam energy deposited in the SETM causes large rise in local temperature (~ 103 °C)

Environmental Degradation Electrochemical or corrosive action in MMC's to nucleate imperfection sites leading to damage

35

Approach to Microscopic Characterization

Scanning Electron Thermal Microscopy Performed in an electron microscope and detected by delicately sensing the minute bending moments thermally induced in the sample

Acoustic Microscopy A derivative of an established method, but performed at intermediate frequencies and with an ad- vanced lens design to permit critical study of subsurface material properties

36

Scanning Electron Thermal Microscopy

• Based on thermal wave imaging

Performed in SEM by injecting heat (and electrons) with E-beam

• Modulated E-beam gives periodic heat source

Functions over broad frequency range (10-1000 kHz)

• Thermal energy deposited below the surface (dependent on accelerating potential, target density)

37

Scanning Electron Thermal Microscopy, cont

Extremely high spatial resolution (like SEM)

• Detected by observing thennoelastic signal generated within sample

• Periodic sample deformations detected phase-sensitively by piezoelectric element

• Very large local temperature excursions can be produced with E-beam

38

Scanning electron microscope column HeNe laser

Blanking plates

Position sensitive detector

Function generator

Lock-in amplifier

Spatial filter

Computer

GPIB Interface Drawing not to scale

39

Analysis of SETM signals

Coupled acoustic field equations and thermal diffusion equa-

tion

//V2u +. (A + //,)V(V • u) - p^ = «T(3A + 2//.)VT

v2r--—= —ffx,f , a ar K

A, /u Lame elastic constants

u particle displacement

p mass density

aT coefficient of thermal expansion

a thermal diffusivity

K thermal conductivity

iJ(x,f) power density of the heat source, a tabulated function for a circular source.

40

Analysis of SETM signals

Voltage induced in a piezoelectric element is

V^KC ' (1)

Po incident power

ß electron attenuation coefficient

Rs E-beam attenuation coefficient (a beam energy)

7 thermal stress constant

uj modulation frequency

C heat capacity

41

Acoustic Microscopy

• Can create images of surface or subsurface features

• Sensitive to weak, small scatterers and near-surface elastic property variations

• Large aperture lens permits Rayleigh wave imaging at in- termediate wavelengths (2-10 /mi)

• Spherical focus lens best for detecting, imaging small sam- ple features

Cylindrical focus lens best for quantitative property de- termination

• Can also map surface topography

42

43

Goals for first year

1.1 Phase I: Establishment of initial capability

Task 1: Scanning electron thermal microscopy

• Confirm operation of e-beam modulator on the Center SEM.

Design and install mechanical load stage on SEM to per- form SETM simultaneously with application of mechani- cal load.

• Setup and test and detection electronics and data acqui- sition system.

• Verify spatial resolution and operation over band width of interest using calibration standards.

44

Goals for first year, cont

Task 2: Acoustic microscopy

• Construct scanning acoustic microscope system with com- puterized motion control and data acquisition; system to be capable of 2 /mi spatial resolution.

• Build or procure cylindrical microscope lens for operation with wideband transducer in the frequency range 50 ~ 100 MHz.

• Integrate system components and test operation on ap- propriate standard samples.

1.2 Phase II: Sample selection and defect initiation procedures

• Select specimen composite material systems of interest in consultation with sponsor or sponsor-related organiza- tions.

45

TECHNICAL PRESENTATION

"Examination of Precursors to

Fracture"

BY

Professor J. T. Dickinson

Dept. of Physics & Material Science Program

Washington State University

46

General Comments

Catastrophic failure initiates when the Griffith (defect size) criteria is exceeded either by:

(a) increasing the stress until the critical size equals the size of a preexisting defect, or

at lower stresses:

(b) increasing size of existing defects (subcritical. crack growth) until a critical defect size(causing catastrophic crack growth) is reached

(c) nucleating defects (void formation, pile up of dislocations- crack formation, instabilities-e.g., necking) which then grow to critical size [in ceramics grain size

(d) linkage of defects to produce critical dimensions.

Service loads are typically well below the stress required for catastrophic failure —> lifetime predictions require modelling of defect growth from assumed distributions of initial defects.

In many applications, most failures occur early in the expected life cycle (weak items-infant mortality) or late in the expected life cycle (defect growth-wear out, or induced failure).

Sources of preexisting defects:

inherent defects (impurities, inhomogeneities -- including fluctuations, dislocations)

processing defects

voids

occluded particles (composite fillers, dust in polymers)

design factors (stress concentrations, rivet holes, exposure to environment)

47

Sources of grown-in defects (usually requires preexisting defect such as above)

localized phases (e.g., glassy phase in polyxtal. ceramic- secondary phases usually initiates cracks)

cracks (for materials without secondary phase -- subcritical crack growth)

shear bands (fracture of metallic glass often nucleates at surface step associated with shear band -- Koh cracks in inorganic crystals)

' crazes (fracture of PS usually initiates at crazes)

synergisms with chemical effects (humidity, saltwater, fingerprints etc.)

residual stresses (tempered glass, interfaces)

EXAMPLE: GLASSES

surface flaw (often due to mechanical contact)

flaws: devitrification at inhomogeneities (misfit stresses) atmospheric corrosion (pits)

subcritical crack growth

catastrophic crack growth

Environmentally Induced Fracture of Glass--(Michalske, Freiman, Wiederhorn).

Glass = inorganic substance which is a non-crystalline fluid with extremely high viscosity-for all practical purposes it is a rigid solid

Silica glasses based on SiCh structure built from tetrahedra:

48

Metallic Glasses

surface steps formed by shear banding (concentrate stress)

cavitation within shear bands

devitrification (decohesion -> void formation)

EXAMPLE; POLYCARBONATE

Typical path to failure:

processing defect—e.g., surface flaw, dust particles

initial craze growth (p.m dimensions) [crazes -- thermal/stress activated in the presence of defects void formation + highly oriented bridging fibrils]

craze rupture (usually at craze boundary)

crack often halted by interaction with cold drawing

tearing of cold drawn material

catastrophic crack growth

In craze resistant polymers: shear bands form, followed by tearing, (crazing possible precursor).

The relation between crack velocity, v, and strain energy release rate, G, may be written (after Maugis)

G - 2y = 2y a(T) f(v) [G - 2y is the dissipative part of the energy loss]

7 is the minimum energy required for surface formation

49

a(T), f(v) are functions which describe the temperature and velocity dependence of dissipative processes. At Gc (a critical value of G) crack velocity increases discontinuously.

Distributed Flaws : The Real World

t / '—

\^ \

~\ r ' -> , , N.

X

X —

\

Distribution of flaws under stress -- changes with time.

50

For ith flaw

first region to fail determined by distribution of

damage zones evolve — A zone that fails will tend to fail again.

When repeated failure starts, ~ at critical breaking point <—> damage zones compete - larger cracks grow fastest. Coalescence very effective in localization of failure zone.

a enhanced linkage probable

crack shielding

51

Monte Carlo Simulation of a "Brittle Polymer" Paul Meakin

Assume a bond breaking rate probability for the ith bond to be:

Pi - exp(ki8i2)/kBT

where ki is the force constant for the ith bond and 8j is the elongation of the bond. Sample is stretched in one direction; Bonds are chosen by a Monte Carlo algorithm and tested for bond breaking by testing Pi.

At low forces (mostly small 8i) bond breaking is random. At higher forces, cracks form that have very irregular shapes and small cracks resemble large cracks. Suggests fractal geometry of cracks = 1.27 - 1.31 depending on method.

Figure 9.1 Results from a typical simulation carried out using the two dimensional model of Termonia and Meakin J9J At the start of this simulation the network of nodes and bonds was stretched by 20% in the y direction (without allowing bond breaking to occur). The parameters used to obtain these results were kx = 20, ky = 20 and kBT = 1. Figures a, b and c show three different stages well before failure, close to failure and at failure.

52

W^M

WITHOUT RADIATIONS^!

PTFE WITH ELECTRON BEAM

Extension

Figure 1.10 Breaking characteristics for a typical polymer tested at different temperatures: (a' brittle, (b) ductile, (c) necking and cold-drawing and (d) rubber-like behaviour.

Yield point

Strain \Slrain .softening hardening

Figure 2.5.9 Cross-section of a polymer craze.

Figure 2.5.7 Typical stress-strain curves of polymers.

PA66

Figure 2.5.10 Types of shear bands observed between cross-polarizers.

54

SEM Micrographs of PS Fracture Surfaces

Heavily Crazed

Lightly Crazed

55

FRACTO-EMISSION (FE):

DURING AND FOLLOWING THE DEFORMATION AND FRACTURE OF MATERIALS: THE EMISSION OF ELECTRONS, IONS, NEUTRALS, AND PHOTONS (INCLUDING LONG WAVELENGTH ELECTROMAGNETIC RADIATION - RADIO WAVES).

PROPERTIES WE MEASURE: TYPE OF PARTICLE nwrnKMM) YYniYift/MWM, TIME DEPENDENCE (VS CRACK GROWTH)

INTENSITIES ENERGIES (electrons, ions; photons [ X ]) MASS (neutrals, ions) TIME CORRELATIONS BETWEEN SPECIES

FE DEPENDS ON: MATERIAL LOCUS OF FRACTURE SAMPLE STRENGTH CRACK VELOCITY ENVIRONMENT (in some cases)

56

Mass Scan during Abrasion of Polystyrene (V^ ^ 100

» 80 JO u a 5 60

1 40 »■*

2 20 ■ 41 > <

0 4

V M

lit I -1 I

I I

to*

f*t»t«Mtf»

T—*—i—•—i—' i ' i ' i •—r 0 20 40 60 80 100 120 140 160 180 200 220 240

Detected Ion Mass (amu)

-S 11 •_• * so ***»t

57

Mass 40 During the Tensile Fracture of Argon Sorbed Polystyrene

Time (ms)

Time (ms)

91-461 & 462 Mass 40 (Argon) during the tensile fracture, 461 (top) shows a lot of crazing and the final fracture surface is smooth. 462 (bottom) does not show as much crazing and has a more classic fracture (mirror, mist, hackle zone) appearance. Strain rate is 0.07 mm/s for each.

58

Ar Emission and Craze Volume in Polystyrene

i 11 11111111 ii i

0 200 i i i i I 11 i i i 11111 111 i i i 11 i I

400 600 800 Time (ms)

Craze volume estimate at constant strain rate:

Assume all deviations from elastic behavior due to craze formation.

Load = E (£applied" e craze)

where E is Young's modulus. The strain due to the craze is proportional to the craze volume, where

craze Load _ o

£ ^applied

The Ar emission intensity is proportional to the total craze volume.

59

60

Ar Emission During the Deformation and Fracture of HIPS

600

< c w

2 300 M «n CS

(a) Mass 40

emission onset (1% strain)

2 Time (s)

t fracture

(3% strain)

180-f

(b)Load

i

2 Time (s)

61

Craze Volume Fraction and Mass 40 Signal

units

) ca

led

Sign

als

(arb

itrar

y

o

Craze Volume

\

Mass 40 Signal jf

w 0:

( ) 2 4 Time (s)

v craze ^sample

00 Ccraze <* Ctotal - £uncrazed

06 Ljnitial ' t - L

H \

Etotal: Total strain at constant strain rate.

Ccraze : "Plastic" strain in crazed material.

Cuncrazed : Elastic strain in uncrazed material. •

Linitial : Initial slope of load curve.

L : Load.

62

Bond Energies of Various Bonds of Polycarbonate

CH L H-C-H 0

_oHry_p0lVn-O-^^Ö-oTH CH, 7VH

Polycarbonate Structure

No. Bond Bond energy, krnl./mole

1 C—H 98.7 •i C—CH» 60.0

3 C—C 82.6 4 C=C 147.5 0 C—() 83.5 G C=<) 176.0-179.0 i C=C-H 102.0

S O—If •. 101 5

L. H. Lee, J. Poly. Sei. A 2,2859 (1964) review paper.

63

Mass 20 and Load During the Tensile Fracture of Polycarbonate

§ tftlUtr TÜR •

L *"

Y

10 15 Time (s)

10 15 Time (s)

20

800- (b) Load fl ^**~

g 1

0 ■ ■ i

25

64

Mass 20 and Load During the Tensile Elongation of Polycarbonate

0 2 3 Time (s)

0 2 3 Time (s)

65

Mass 20 and Load During the Elongation of Polycarbonate

50

< a "8 s •■■

©

o T 0

A , . a) Mass 20 V«'UH%f iHfUtt)

* StfW..* ^ i»A^ vAMo

6 9 Time (s)

12 T 15

0 6 9 Time (s)

66

0

Mass 20 and Load During the Tensile Elongation of Polycarbonate

Time (ms)

NOT HciK* ̂

67

A*&L6D

<*•+» l»{o i <-«4U<b'Kt

Timing Light and Load from Fracture of Polycarbonate £ -f Adwj

140 > E <w 3 a 3 o

■o o 1 o

QU

0

(a) timing light

0

"1

23

Time (seconds) 6 btHr*iA*+»i*\

1600

1280 V) G O

! c

T3

§

Time (seconds)

?ö/30* 68

EE and Timing Light Signal from Fracture of Polycarbonate

40

E

3 a «■» a O « 20 ©

'•5 o

Ol

60 90 Time (jxs)

(b) ti n ing light

***

30 60

ISO*1/

90 120 ^ Time (|is)

150

—r 150

%)/3o> 69

Crack Velocity during Tensile Loading of PC Inferred from Light Transmission Measurements

(LED Source)

i i i l i i i i I i i i i l i i i i I i i i i | i i i i | i i

10 12 14 16 18 20 22 24 Time (s)

Geometry of Crack Growth

70

Small Bursts in Total Alkali Emission from NaCl

25 - large bursts off scale v.

< S

u u 9

0.5 1.0

time of fracture

1.5

Time (s)

2.0

fcuds rs

^ <UKct4>

2^

71

EVIDENCE OF 02 EMISSION DUE TO IHCROCRACKINQ IN SINGLE CRYSTAL MflO

s 7.

*!

!l

BRSSION PRIOR TO FRACTURE

A

• • • I

PEAKS A AND RARE SHOWN ON

FASTtR TRIE SCALES BELOW.

TIME (8) T ••*

•/ V

»!

• • % **'...«

Vs

(mS)

,.-.•• •-...-..-... ...../ •••• •■

72 (m8)

Ox PftE-CuftSogS FR.OM ft^Q

• • PoUU.

x20 ♦—J • • —

i •

V <

i

•

5-»

• •

• • • • •

M

3

•• •

• • • v • • v > * • «• >».7 \ • •••«** • • *■

• • •

••• • • • •/ •

• • •

•t • • • # • • • • \ • • •

•

*

t

%

8

Time (scc^

\)v\ poli*/i*A.

0.1

X20

.*

\

Tfi*c (M*^ 0.05 73

Interfaces

(Evans, Ruhle, Kinloch, Mattox)

l .2

I

Ä es B

t

perfect material

1 s

f interface regioji

(interphase) 1 . ■m.»»««»».»««»««»M

interface regioi (interphase) y

2 \^

perfect material

2

interface plane

Defects at or near interfaces:

chemical defects segregation chemical reactions chemical gradients interphases (reaction products)

structural defects facets, steps misfit dislocations lattice dislocations impurities pores precipitates

Chemical Defects often reduce Wad (Work of adhesion (reversible work/area to separate the bond into two free surfaces)

74

Wad =Yi +Y2-Y12

where yl , y2 - free energies of relaxed surfaces; yl2 = energy of

a relaxed interface between the two materials. Wad cannot be measured directly.

In practice use, equilibrium contact angle, 9, measured:

Wad =Yi (1 + cosG )

Fracture Resistance, Interface > Wad (typically 1 J/m2).

Tinterface ~ Metal/ceramic ~ 50-300 J/m2. Due to dissipative processes: plastic deformation of one or both materials. Important to realize that Hnterface usually scales with (provided a strong singularity remains at the crack tip) Wad: i«e., the stronger the adhesion, the more plastic deformation of one or both materials .

Internal Stress: often due to differences in a (thermal exansion).

Phase Transitions in one component.

75

V WD CO E

c/> 08

O c 0> CM o

CM O •pp

0)

o 13

CM ©

• Pp

13 CO

C/} ÖD C o

'S a> 5> Q

• pp W PP 03 CM a> N O 'M

• PP

o M 0>

CM

• PP

3 CO M CO

ts u H

c •pp o

M • • C/5

U T3 •PP CM

.N

c cd

C CO

•PP C/2

PG

c Q PLH i c o

• M

o •pp

OH

C/3

'S •pp

13 •pp

• • C/5

cd ■*—>

o o

a! Q

•pp U C/5 0) Q

M

03

2

M CM

•pp

CO

in

G ©

•

CO VH

PH

•

13 00 •

• • • • w* n <T> ^

<tf C a> 6 o Ö <D

>> •H

13 c •4—> 13 u 00

76

a

»9

$ 9 ©

•5:

C/5

E >

■4—1 o

0) OH

C/5 4—1 o

Q

as C 0) a o c a»

PL.

o

t/3 >

u CM U

0 M

• UN

k c w cu <D

• P4 VH •*-» 3 1/1 •4-» O cd H a) c B vx <D a> H -fi O • u

PH

13

4> XI U ~a3 0) X a, O

G O

•rH

cd

W) S3

o CO *s

o O

o

>/3 so

77

TECHNICAL PRESENTATION

"Hysteresis and Incremental

Collapse of Structures"

A Paradigm for the Fatigue

wStrength of Metals

BY

Professor S. A. Guralnick

Dept. of Civil Engineering

Illinois Institute of Technology

78

OBJECTIVES

To Coxß&iAre FAT/<SI'£F JA/P //rsrejees/s ßy #£cc6wz/AtG T/t*- Pae cesses IA////C* £fG/A/ W/TH Af/cAO'/AIPeAF&Cr/MSj Pxa&xess

fo Caßߣt/iT£ <4COV3T/C £M/SS/0M fiveA/wey/l v//r# 7#e /*/<?e/>r/0// &r Tfre V/?A//?G£:

lessors /*/ /r/)T/&t/e /ä/PFf/zs.

79

/)P/=>/ZO/ICH

To /Aj\/e$T/6tiT£ TM COMPLBTS EMLunOAf or HrsTeßes/Sj f#oM fr/zsr Cyctz To LAST CYCLFJ 3Y A/*/)A/S *&

TO CYC*-/C JLO/7I>5> (&=0 s4*p

80

STRESS AND CYCLIC STRAIN PARAMETERS

AO"= stress range

t

total strain range

81

SiMcrt/ßAL SCALE

.rrv JO

m 10

cm

mm 10 .5

assemJ>/aje

Structure

element

fzolycrystaf

/U — JO -6 criistal

10 ■lO at otox

82

POSSIBLE DEFECTS IN METAL CRYSTALS AND POLYCRYSTALS

CRYSTAL DEFECTS

• POINT DEFECTS ATOMS INSERTED OR SUBSTITUTED IN SOLID SOLUTIONS OR VACANCIES

LINE DEFECTS (DISLOCATIONS) EDGE DISLOCATIONS SCREW DISLOCATIONS DISLOCATION LOOP OR LINE

SURFACE DEFECTS (SURFACES OF SEPARATION) DISLOCATION LOOPS TWIN CRYSTAL BOUNDARIES

POLYCRYSTAL DEFECTS

SURFACE DEFECTS (SURFACES OF SEPARATION) CRYSTAL OR GRAIN BOUNDARIES INTERFACES BETWEEN TWO PHASES (EXOGENOUS AND ENDOGENOUS)

COHESION DEFECTS MICROCRACKS AT CRYSTAL BOUNDARIES CAVITIES BETWEEN CRYSTALS

83

POSSIBLE HICROHECHANICAL RESPONSES OF METALS TO CYCLICALLY-VARYING LOADS

ELASTIC DEFORMATIONS (REVERSIBLE)

• VARIATION IN INTERATOMIC SPACING

PLASTIC DEFORMATIONS (IRREVERSIBLE)

SLIP ALONG CRYSTALOGRAPHIC PLANES OF MAXIMUM ATOMIC DENSITY IN SINGLE CRYSTALS

PLANAR GLIDE IN SINGLE CRYSTALS

• DEFORMATION TWINNING AND/OR KINKING

SLIDING OR SLIPPING BETWEEN GRAIN BOUNDARIES

FRACTURE

INTERGRANNULAR FRACTURE MICROVOID FORMATION MICROVOID COALESCING GRAIN BOUNDARY CRACK AND CAVITY FORMATION DECOHESION BETWEEN CONTIGUOUS GRAINS DUE TO PRESENCE OF IMPURITIES SEPARATION DUE TO INSUFFICIENT NUMBER OF INDEPENDENT SLIP SYSTEMS TO ACCOMMODATE PLASTIC DEFORMATIONS BETWEEN CONTIGUOUS GRAINS

INTRAGRANNULAR FRACTURE CLEAVAGE FRACTURE

TRANSCRYSTALLINE FRACTURE ALONG SPECIFIC CRYSTALOGRAPHIC PLANES

84

INFLUENCE OF NUMBER OF CYCLES OF LOADING UPON MECHANICAL PROPERTIES

OF MILD STEEL IN TENSION (AFTER PUSKAR AND GOLOVIN, 1985)

!

85

MAXIMUM STRESS VS. CYCLIC LIFE CURVES (AFTER LEMAITRE AND CHABOCHE, 1990)

V)

2 b

N¥ (cycles)

86

GOODMAN-GERBER DIAGRAN FOR 'SAFE' RANGE OF STRESS

Ultimate Tensile Strength, S„

-0.50

87

CONVENTIONAL REPRESENTATION OF THE FATIGUE FAILURE PROCESS

(AFTER BANNANTINE, COMER AND HANDROCK, 1990)

Crack Propagation Period

Fatigue Life (log scale)

88

CONSTANT STRESS RANGE CRACK GROWTH VS. NUMBER OF CYCLES

(AFTER ROLFE AND BARSOM, 1977)

2.2

•£ 2.0 o

c

•6 1.6 o o

a*

Sinusoidal Stress-Constant Amplitude Ao-3>Acr2>cr1

£

Ao-3

k

1.2-

1Q

£ hA- 0, o

0

0 o

—r" Ao^ 8

#

<? <9

2 3 4 5 6 Number of Cycles, N*105

8

89

CRACK GROWTH RATE VS. STRESS INTENSITY FACTOR RANGE

(AFTER HERZBERG, 1989)

A K th JLKC

\OQAK

90

KtKCLNI Uh LJLht Tu URACK INITIATION VS. CYCLIC LIFE

(AFTER LAIRD AND SMITH, 1963)

lOOr-

60

PERCENT OF LIFE TO

CRACK INITIATION

60

40

20

10*

:CDO

2.5 N -1/3 f

♦ O ▲

410 STEEL 400fF TEMPER 410 STE€L 850#F TEMPER 2024-T4 ALUMINUM 4130 HARD STEEL POLYCARBONATE PURE ALAND PURE Nl

L L I I0J 10*

1 I 10^ 10* IOl Mr

CYCLIC LIFE

91

WAVY GLIDE IN HIGH STACKING FAULT ENERGY MATERIAL

(AFTER HERZBERG, 1989)

92

STRIATIONS PRODUCED AT VARIOUS ALTERNATING STRESS LEVELS

(AFTER SANDOR, 1972)

\> v

93

A471 STEEL REVEALING LOCALIZED EVIDENCE OF INTERGRANNULAR FAILURE AS A

RESULT OF CYCLIC LOADING (AFTER HERZBERG, 1989)

94

PLASTIC STRAIN-INDUCED SURFACE OFFSETS (A) MONOTONIC LOADING, (B) CYCLIC LOADING LEADING TO EXTRUSIONS AND INSTRUSIONS AND (c) PHOTOMICROGRAPH OF

INTRUSIONS AND EXTRUSIONS ON PREPOLISHED SURFACE

(after Herzberg, 1984)

(a) (6)

95

HETALLOGRAPHIC SECTION REVEALING ONSET OF MACROCRACKING

(AFTER HERZBERG, 1989)

96

PROGRESSION OF THE FATIGUE PROCESS (AFTER PUSKAR AND GOLOVIN, 1985)

1. DISLOCATION SEGMENTS BLOCKED BY PRECIPITATES OR BY ATOMS OF ADMIXTURES. LOCAL MOVEMENTS ARE REVERSIBLE, (0«X«Xr )

LE 2. DISLOCATION SEGMENTS CAN BE DISPLACED BY SHORT

SEGMENTS AND LOCALIZED MICROPLASTICITY OCCURS. (0CE

K a ' V 3. ACTIVITATION OF DISLOCATION SOURCES RESULTING

IN FINE TRANSITION SLIP LINES. (<T > (Tn ) LC

4. DISLOCATION DENSITY INCREASES AND THEIR DISTRIBUTION CHANGES FROM A STATISTICALLY RANDOM DISTRIBUTION THROUGH THE OCCURRENCE OF CLUSTERS AND TANGLES VIA AN IMPERFECT CELL STRUCTURE TO A PERFECT CELL SUBSTRUCTURE.

5. PERMANENT SLIP LINES AND BANDS FORM. CRACKS BEGIN TO NUCLEATE IN POINTS OF INTRUSIONS.

6. FRACTURES PROPAGATE FIRST IN A CRYSTALLOGRAPHIC AND LATER IN A NON-CRYSTALLOGRAPHIC WAY.

7. FRACTURES ORGANIZE INTO NETWORKS AND COMPLETE RUPTURE ENSURES. (CT < 0" )'

97

PROGRESSION OF THE FATIGUE PROCESS (AFTER LEMAITRE AND CHABOCHE, 1990)

1. INITIATION STAGE - CYCLIC ELASTICITY, NEARLY COMPLETELY

REVERSIBLE BEHAVIOR AND 'INFINITE' LIFE. (O < & < 0" )

2. ACCOMMODATION STAGE (OR NUCLEATION) - CYCLIC PLASTIC MICRODEFORMATIONS

OCCUR WHICH, IN TURN, BLOCK FURTHER SLIP BY VIRTUE OF MULTIPLICATION OF DISLOCATION NODES (<T*Q±Jd*ic)

1 limit '

3. INITIATION PHASE OF MICROCRACKS (STAGE 1) - DISLOCATION CLIMBS ACCOMPANIED BY FORMATION

OF VOIDS FORMATION OF PERMANENT SLIP BANDS AND DECOHESION

- INTRUSION-EXTRUSION MECHANISMS

4. GROWTH PHASE OF MICROCRACKS (STAGE 2) - ORIENTATION OF MICROCRACKS PERPENDICULAR

TO PLANE OF MAXIMUM PRINCIPAL STRESS - PROGRESSION OF MICROCRACKS THROUGH

SUCCESSIVE GRAINS OR ALONG GRAIN BOUNDARIES. MACROSCOPIC CRACK INITIATION TAKES PLACE.

5. GROWTH PHASE OF MACROCRACKS - CYCLIC OPENING AND CLOSING OF CRACKS RESULTS

IN ALTERNATING PLASTIC SLIPS AT THE CRACK TIP, WHICH IN DIFFERENT CRYSTALLOGRAPHIC PLANES FORM A RIDGE OF CLEAVAGE AT EACH GROWTH OF THE CRACK.

- CRACKS GROW UNTIL A CRITICAL SIZE IS REACHED AT WHICH PART BECOMES UNSTABLE.

- CRACKS PROPAGATE RAPIDLY UNTIL COMPLETE 98 RUPTURE OCCURS (<T = <T ) .

STAGES IN THE PROGRESSION LEADING TO RUPTURE

s s t c a a g 1 e e

} '

S c a 1 e

100 M

mm

10 mm

99

s t a g e

-2 10 M

'Cooperative Organization'

of Zones of Microplasticity

Damage Growth and

Accumulation, Micro-Fracture

2

i I

i >

1 .01M

Inception of

Microplasticity

Macro Crack Initiation 2-

t ,

" '

1 Evolution of

Stresses and Strains

Crack Propagation and

'Organization' 3

i

Material Rup ture a - Oc

/ Initial Conditi ans

>

\ Cyclic Loadin g

/NCBPTfO/J OF Aftc&OpLAST/C/TYs Pj /A/CSPT/M op AfAC&OCÜACKWGj F

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lous- cycle *> hijk- cyc/e.

Ji

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100

POSSIBLE DIAGNOSTIC TECHNIQUES

1. ANALYSIS OF HYSTERESIS EVOLUTION

2. MAGNETIC EFFECTS (IN STEEL)

3. ACOUSTIC TRANSMISSION AND HARMONIC GENERATION

4. ACOUSTIC EMISSION

5. ACOUSTIC SCATTERING (MACRO-CRACK GROWTH MONITORING)

6. MICRO-MECHANICAL PHYSICAL TESTING IN CONJUNCTION WITH: A. SCANNING TUNNELING MICROSCOPE B. SCANNING ELECTRON MICROSCOPE c. TRANSMISSION ELECTRON MICROSCOPE D. POSITRON ANNIHILATION

7. PHOTOSTIMULATED EXO-ELECTRON EMISSION

8. SURFACE EXAMINATION A. OPTICAL HOLOGRAPHY

9. EDDY CURRENT PROBES (MICRO-CRACK MONITORING)

101

RESPONSE OF A STRAIN-SOFTENING MATERIAL TO CYCLIC APPLICATIONS

OF LOAD UNDER STRAIN CONTROL

CO CO

73 -25-E

-50-E

-75-5

-100- I i i i i i i i i i i i i i i i i i i

-0.02 -0.01 0.00 0.01 Strain (in/in)

0.02

102

TYPICAL HYSTERESIS LOOP SHOWING ELASTIC-PLASTIC STRAIN DIVISION

103

STRESS VS. STRAIN DIAGRAM FOR RIMMED AISI 1018 UNANNEALED STEEL

80-r-

60-

w w <u u

.4-) w

40-

20-

0 I I I M I I I I I | I I

0.00 0.01

I I I I I I I I M I I I M I I M I I

0.02 0.03 Strain (in/in)

104

CYCLIC AND MONOTONIC STRESS-STRAIN CURVES

W CO

OT _20-

-0.02 -0.01 0.00 0.01 Strain (in/in)

0.02

105

HYSTERESIS LOSS PER CYCLE VS. NUMBER OF CYCLES

(A €/2 - 0.006 in/in)

0.8

£0.6 I p, 2

o O0.4 H

a w w

.3

.3 0.2 w <u u

to

0.0

■ B ■ ■ • » » • Q «

»»"» Specimen 127 »•♦•* Specimen 224

ii mi ii 11) i) i u ii 11 in ii ii ii |i ii ii ii ii )i ii mi ii 11 ilium 500 1000 1500 2000

Cycles (N) 2500 3000

106

HYSTERESIS LOSS PER CYCLE VS. NUMBER OF CYCLES FOR A

SPECIMEN SUBJECTED TO OVERLOADS

0.5 -r—

0.4

I a S -

Ü

0.3 -

u a0.2 01 w

w

QJ 0.1 U a>

0.0 I I IlI I I I M ' I I I I I I I I llI I I I I l I I '' I

1000 2000 3000 Cycles (N)

1111111111111'' . 4000 5000

107

AVERAGE HYSTERESIS LOSS PER CYCLE VS. STRAIN AMPLITUDE

99999 0.0055(Ar/2(0.015 ooooo 0.0022<Ac/2<0.0055 *±±±f 0(Ac/2(0.0022

i i i i i i i i i i i i i i i i i i i i i 0.004 0.008 0.012

Strain Amplitude (in/in) 0.016

108

HYSTERESIS LOSS PER CYCLE VS. NUMBER OF CYCLES

We,

CO

CYCLES, A/

109

ZONES IN THE EVOLUTION OF HYSTERESIS LOSS PER CYCLE

Co

2

CrcLeSjN

110

DISSIPATED ENERGY OF PLASTIC DEFORMATION UNTIL FAILURE AND ITS ACTIVE PART,

TYPE 316 STAINLESS STEEL AT ROOM TEMPERTURE (AFTER LEMAITRE AND CHABOCHE, 1990)

E .E •a •a ***

u.

0.1 10a

—r— 10J

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it

to /ai'Jur-e ** COVSt arvC

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JLoj Cyc/es & Failure 'N<

112

C(/MvL/ir/\/e DAM/MB AA/D Sm/j/A/

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113

*

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Failure tjut/l occur* uj^ettj

Frc**> £*S, / CLrvct £j

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NFI Nfz Afcß

Mtuuilit tM15MUN IN ALUMINUM (AFTER GRAHAM AND MORRIS, 1975)

1. ACOUSTIC EMISSION BEHAVIOR OF VARIOUS MATERIALS IN DIFFERENT CONDITIONS IS MARKEDLY DIFFERENT

2. AT LEAST FOUR DIFFERENT TYPES OF EMISSION WERE OBSERVED

3. PRIMARY SOURCE OF ACOUSTIC EMISSION IS BRITTLE FRACTURE OF INTERMETALLICS

4. SECONDARY SOURCE OF ACOUSTIC EMISSION IS BRITTLE FRACTURE OF MATRIX MATERIAL

5. TERTIARY SOURCES OF ACOUSTIC EMISSIONS ARE: - GRAIN BOUNDARY SEPARATION - DELAMINATION ALONG GRAIN BOUNDARIES - MATRIX CLEAVAGE

6. GROSS PLASTIC DEFORMATION OF THE ALUMINUM MATRIX YIELDS VERY LITTLE MEASURABLE ACOUSTIC EMISSION

115

THE FOUR DISTINCT TYPES OF ACOUSTIC EMISSION OBSERVED

FOR INDIVIDUAL ACOUSTIC EMISSION EVENTS (AFTER GRAHAM AND MORRIS, 1975)

> Q

•o

LU Q

a.

T—i—i—i—\—i—i—i—r

(a) TYPE WN

(b)TYPE II

(c) TYPE I

A (d)TYPE III

0 1 FREQUENCY (MHz)

116

MU5I LlKbLY bUUKttb PKUUUCING IHt FOUR TYPES OF ACOUSTIC EMMISSION

(AFTER GRAHAM AND MORRIS, 1975)

TYPE FREQUENCY CONTENT

UN WHITE NOISE

II HIGH FREQUENCY I Low FREQUENCY

III PEAKED

LIKELY SOURCE

MOSTLY BRITTLE FRACTURE OF INTERMETALLICS; SOME BRITTLE MATRIX FRACTURE

SAME AS FOR TYPE WN DELAMINATION ALONG GRAIN BOUNDARIES

SOURCE UNDETERMINED

117

PRELIMINARY CORRELATION OF METALLURGICAL

STATE WITH ACOUSTIC EMISSION RATE

SCALE STAGE ACOUSTIC EMISSION

• INCEPTION OF MlCROPLASTICITY

OlyU 1 KAISER EFFECT OCCURS

• COOPERATIVE ORGANIZATION OF MLCROPLASTICITY

10// 1 - 2 NON- VANISHING COUNTS, FELICITY EFFECT

. DAMAGE GROWTH AND ASSUMULATION, HLCRO FRACTURE

100 jti 2

• MACRO CRACK INITIATION

MM 2 - 3

• MACRO CRACK PROPAGATION AND ORGANIZATION

10 MM 3

t

RUPTURE > 10 MM

118

ULTIMATE TB/S/LB 5ru//or//? ai-- 77.7z AU

Y/ELD STfätfGT//(0.2% WFJEr], <Ty 70.77h/

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MODULUS cp £iAS7/c/rr; E* 2?}6M fa

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119

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122

UlKtUIONS FOR FUTURE RESEARCH

1. CORRELATION OF ACOUSTIC EMISSION DATA WITH ACTUAL METALLURGICAL STATES USING SEM OBSERVATIONS

2. OBSERVATION OF FATIGUE PROGRESSION ON A MICRO- MECHANICAL SCALE USING A MICRO TESTER IN CONJUNCTION WITH THE TEM

3. CORRELATION OF HYSTERESIS DATA WITH ACTUAL METALLURGICAL STATES USING SEM OBSERVATIONS

4. CORRELATION OF MAGNETIC FIELD EVOLUTION WITH PROGRESSION OF FATIGUE PROCESS (IN COOPERATION WITH NRL)

5. INVESTIGATION OF INCEPTION OF MICROPLASTICITY BY HYSTERESIS MEASUREMENTS WITH STM (IN COOPERATION WITH C. TEAGUE OF NIST)

123

REFERENCES

Anon, 1978, "Fatigue and Microstructure", Paper presented at the 1978 ASM Materials Science Seminar. American Society for Metals. Metals Park, Ohio.

Bannantine, J.A., Comer, J.J., and Handrock, J.L., 1990. Fundamentals of Metal Fatigue Analysis", Prentice-Hall. Inc..

Englewood Cliffs, New Jersey.

Hertzberg, Richard W., 1989, "Deformation and Fracture Mechanics of Engineering Materials", Third Edition, John Wilev and Sons. New York, N.Y. —

Lemaitre, J. and Chaboche, J., 1990, "Mechanics of Solid Materials , Cambridge University Press. New York, N.Y.

Parker, A.P., 1981, "The Mechanics of Fracture and Fatigue", E. and F.N. Spon, in association with Methuen. Inc.. New York,

Puskar, A., and Golovin, S.A., 1985, "Fatigue in Materials: Cumulative Damage Processes", Elsevier Science Publishing Co.. Inc.. New York, N.Y.

Sandor, B.I., 1972, "Fundamentals of Cyclic Stress and Strain", J_he University of Wisconsin Press. Madison. Wisconsin.

Suresh, S., 1991, "Fatigue of Materials", Cambridge University Press. New York, N.Y.

124

TECHNICAL PRESENTATION

11 The Microscopic Origin of

Inelastic Behavior"

BY

Dr. Harold Weinstock

AFOSR

Boiling Air Force Base

125

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128

Schema'/o or cxferimevTdl ^r^ngcm^f MaghefVc field--fret e.vv/irorwy\e.nf Ha$v\efic -field coil in -transverse. Jirecf/or)

PREAMPJ:

SQUID

I -^LIQUID

1 HELIUM OEWAR

J f ^-GRAOIOMETER [8_J COILS

'TRANSVERSE' MAGNETIZATION

r~-MICRO-OERSTED REGION

12.7 M DIAMETER HELIWOLTZ COILS

129

Stress + Mawefic fasponze, fo Sfra'm for a Sfee/ oar

REVERSIBLE REGIME

STRAIN

130

if-

I «■

TO

</)

3)

131

Stress a*\J Magnetic Response t» Sfr<aiy\ for a 5>feeJ ßar

PLASTIC FLOW REGIME

• i

t \ t

ill A<t

ii

Q < O

STRAIN

132

^0

5

•5 VA

' —

V)

1 < I : <

cat <

«^ < • i <

4B

» 1

' 5" of k 3. ■B

•.

• i •

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m/uuftw^i •|B| «t)tS{UI3 9RIM9V 133

FERROMAGNETIC SAMPLES IN AN APPLIED FIELD

BARKHAUSEN NOISE

n • V • \ K

• «...

+ j

AA - ::::

•' :••; •:'.!

-.;:; • ^V I: 4 -

. * V ;;;! \ * •

:::: \: !v I : .::. /■■

• :: .:: \ • ■\i ;■

'... * ::: ■ j

Polycryst aline iron samples

Field ramped (Increasing from right to left)

Appearance of Barkhausen jumps corresponds to the transition from reversible to irreversible magnetiza- tion, i.e. the onset of hysteresis.

Repeated field cycles at threshold remove jumps and extend the range of reversible magnetization.

"Threshold of Barkhausen emission and onset of hysteresis in iron,"

H. Weinstock, T. Erber, and M. Nisenoff," Physical Review B, 31

153541553 (1985). 134