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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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1 November 1991
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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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— 500 cycles
— 1000 cycles
2 3 4 5 6 7 CRACK LENGTH, MICRONS
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The Weibull probability density function as a function of crack length. A: Silag SiC (20 vol. pet. whisker)/Al (2124-T4), Heat #A6104Z-2A, at. 100, 500 and
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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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The purpose of the present work is to extract quantitative
information from the acoustic microscope by relating the voltage from the
microscope to the mechanical properties of the insonified material.
22
Transducer
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23
METAL MATRIX COMPOSITE
Silicon Coated Carbon Fibers in Titanium
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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
sConventional «J ** ™f ^ fic/*t''onsA//b; Complete Kufture
lous- cycle *> hijk- cyc/e.
Ji
Cyc L/C
10« ~*rf
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
—T— 104 NF (cycles)
it
to /ai'Jur-e ** COVSt arvC
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STRAW AMPLWDB />A/Z> /Jt/Ai8B/L
loy Cycles to Failvm,
*^~~~ tiAUd(%)
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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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/
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^
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€, > eL > 63 > £v > es > c
113
*
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Failure tjut/l occur* uj^ettj
Frc**> £*S, / CLrvct £j
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doe £> Ci/mictfd'^e <J<tm<*je *AJ;// O*CUV* «^te^
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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/
Pß0F0£T/0N4L lw/7, cy U^Asi
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MODULUS cp £iAS7/c/rr; E* 2?}6M fa
E/WU/?/)A/C6 //M/r ? S * A<=r~ 3Z.ttA*i
119
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120
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Sr^ESS #/i//6£, Acr^ ksi 121
CcWCL t/S/OtfS
By SEPAPAT/MG //rSTEKES/S /ATO
Two COMPONENTS — 0ue DOE TO
MFA7MG AMD TAE LOTAEP IA/AIC#
CAUSES DAMAGP-- /r /J PJSS/ME
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DAMAGE /S /)PPL/C*SLE. COA/I/£ASEIY
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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
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STRAIN
132
^0
5
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1 < I : <
cat <
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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