fatigue mechanics
TRANSCRIPT
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TABLE OF CONTENTS
Symbol……………………………………………………………………………………………. 2 Abbreviations……………………………………………………………………………………. 2 CHAPTER -1 Introduction to Fatigue of Welded Structures………………………...............................
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CHAPTER -2 2.1 Fatigue as a Phenomenon in the Material………………………………………………… 7 2.2 Different phases of the fatigue life…………………………………………………………… 9 CHAPTER- 3 3.1 Why metal parts fail from repeatedly-applied loads……………………………………… 15 3.1.1 What is fatigue loading……………………………………………………………………… 15 3.1.2 How is the fatigue strength of a metal determined………………………………………… 17 3.1.3 Is there any relationship between UTS and fatigue strength…………………………… 18 3.1.4 Why is the surface so important………………………………………………………… 19 3.1.5 Is the endurance limit an exact number…………………………………………………… 20 3.1.6 Do real-world components exhibit the "laboratory" EL…………………………………… 21 3.1.7 Is fatigue loading cumulative…………….……………………………………………… 22 CHAPTER- 4 4.1 Fracture Mechanisms……………………………………………………………………… 24 4.1.1 Ductile Fracture……………………………………………………………………………… 25 4.1.1.1 General Macroscopic Appearance of Ductile Fractures………………………………… 28 4.1.2 Brittle Fracture……………………………………………………………………………… 30 4.1.2.1 Microscopic Aspects of Fracture………………………………………………………… 32 4.1.3 Transgranular Ductile……………………………………………………………………… 34 4.1.4 Transgranular Brittle Fracture……………………………………………………………… 38 4.1.5 Quasi-Cleavage……………………………………………………………………………… 40 4.1.6 Mechanisms of Intergranular Fracture. …………………………………………………… 44 4.1.6.1 Intergranular brittle cracking………………………………………………………………. 46 4.1.6.2 Dimpled intergranular fractures…………………………………………………………….. 47 4.1.6.3 Intergranular fracture surfaces with corrosion……………………………………………. 48 4.1.7 Macroscopic Aspects of Overload Failures……………………………………………… 50
4.1.8 Cracks propagating from a pre-existing stress raiser or notch………………………….. 61
CHAPTER 5 5.1 Fatigue testing – Part 1………………………………………………………………………. 63 5.1.1 S/N curve………………………………………………………………………………………. 65 5.1.2 Palmgren-Miners rule………………………………………………………………………… 66 5.2 Fatigue testing – Part 2………………………………………………………………………. 67 5.2.1 Preparations and measurements ………………………………………………………….. 69 5.2.2 Test results…………………………………………………………………………………… 74 5.3 Crack growth tests – guidelines for test setup and specimen monitoring……………… 75 5.4 Welded Components…………………………………………………………………………. 80 5.5 Fatigue testing Part 3………………………………………………………………………… 85 5.5.1 BS 7608:1993 ………………………………………………………………………………… 86 5.6 Potential modes of failure of welds………………………………………………………….. 95 5.7 Tubular joints…………………………………………………………………………………. 104 5.8 Weldments ……………………………………………………………………………………. 105
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CHAPTER 6 Designing against Fatigue of Structures…………………………………………………………. 109 8.8 Different types of structural fatigue problems………………………………………… 109 8.8 Designing against fatigue………………………………………………………………… 111 8.8 The crack initiation aspect………………………………………………………………… 112 8.8 Material selection…………………………………………………………………………… 113 8.8 Surface treatments………………………………………………………………………… 113 8.8 Detail design for an improved stress distribution……………………………………… 113 8.8 Large-scale design issues………………………………………………………………… 114 8.8 Uncertainties, scatter and safety margins……………………………………………… 114 8.8.1 Uncertainties………………………………………………………………………………… 114 8.8 Scatter and safety factors………………………………………………………………… 115 8.8.1 The fatigue limit and the safety factor……………………………………………………… 115 8.8 Safety factors for finite fatigue life problems under CA loading……………………… 117 8.8 Safety factors for finite fatigue life problems under VA loading……………………… 118 8.8 Safety factors and fatigue crack growth …………………………………………………. 118 8.8 Safety aspects associated with a corrosive environment and low frequency fatigue.. 122 CHAPTER 7 Methods of revealing fatigue cracks………………………………………………………………
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7.1 Dye-penetrant testing ……………………………………………………………………… 125 7.1.1 An example of dye penetrant testing used on bicycle components…………………… 126 7.2 Photoelasticity ………………………………………………………………………………. 128 CHAPTER- 8 8.1 Causes and recognition of fatigue failures ………………………………………………… 129 8.1.1 General Causes of Material Failures:……………………………………………………… 129 8.1.2 Recognition of Fatigue Failure……………………………………………………………… 129 8.2 Design Considerations……………………………………………………………………… 131 8.2.1 Influence of Processing and Metallurgical Factors on Fatigue ………………………… 131 8.2.1.1Processing Factors …………………………………………………………………………. 131 8.2.1.1 Metallurgical Factors ………………………………………………………………………. 133 8.3 Experimental Analysis of Fatigue Life Curves …………………………………………… 134 8.6 Fatigue Crack Growth……………………………………………………………………….. 134 8.7 Real Life-Design and Manufacturing Considerations ……………………………………. 135 8.8 Recommendations for Designs to Avoid Fatigue Failures ………………………………. 135 Annexure 1 Additional Scanning Electron Microscope Images…………………………………………………
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Annexure 2 Metallography/Microstructure Evaluation…………………………………………………………
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Annexure 3 Microscopic characteristics of fatigue fracture…………………………………………………… 147 Macroscopic characteristics of fatigue fracture…………………………………………………… 147 Lack of Deformation………………………………………………………………………………… 148 Beachmarks…………………………………………………………………………………………… 148 Ratchet Marks………………………………………………………………………………………… 149 Similarities between Striations and Beachmarks………………………………………………… 149 Differences between Striations and Beachmarks………………………………………………… 150 Annexure 4 Samples failure……………………………………………………………………………………… 151 REFERENCES……………………………………………………………………………………… 154
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Symbols a crack length, or semi-crack length, or depth of part through crack a0 initial crack length ac final or critical crack depth af final crack length c (semi) crack length of surface crack C constant in Paris equation D diameter da/dN crack growth rate dU/da strain energy release rate E Young‘s modulus G shear modulus K stress intensity factor ΔK = Kmax − Kmin KIc fracture toughness N number of cycles N fatigue life until failure P load r root radius of notch S nominally applied (gross) stress T temperature Ε strain ν Poisson ratio σ local stress in material σa stress amplitude σm mean stress τ shear stress
Abbreviations
AW As-Welded
BS British Standards
bcc Abbreviation for body-centered cubic crystal structure.
CA Constant Amplitude
CP Cathodic Protection
CT Compact Tension
CTOD Crack Tip Opening Displacement
FC Free Corrosion
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fcc Abbreviation for face-centered cubic crystal structure
FEA Finite Element Analysis
FEM Finite Element Method
HAZ Heat-Affected Zone
HB Hardness Brinell
hcp Abbreviation for hexagonal close-packed crystal structure
HS Hot-Spot
LEFM Linear Elastic Fracture Mechanics
NDI Non-Destructive Inspection
SAW Submerged-Arc Welding
SEM Scanning electron microscopy
TEM Transmission electron microscopy
TIG Tungsten Inert Gas
VA Variable Amplitude
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CHAPTER 1
Introduction to Fatigue of Welded Structures
Fatigue failures in metallic structures are a well-known technical problem. Already in
the 19th century several serious fatigue failures were reported and the first laboratory
investigations were carried out. Noteworthy research on fatigue was done by August
Wöhler. He recognized that a single load application, far below the static strength of
a structure, did not do any damage to the structure. But if the same load was
repeated many times it could induce a complete failure. In the 19th century fatigue
was thought to be a mysterious phenomenon in the material because fatigue damage
could not be seen. Failure apparently occurred without any previous warning. In the
20th century, we have learned that repeated load applications can start a fatigue
mechanism in the material leading to nucleation of a small crack, followed by crack
growth, and ultimately to complete failure. The history of engineering structures until
now has been marked by numerous fatigue failures of machinery, moving vehicles,
welded structures, aircraft, etc. From time to time such failures have caused
catastrophic accidents, such as an explosion of a pressure vessel, a collapse of a
bridge, or another complete failure of a large structure. Many fatigue problems did
not reach the headlines of the newspapers but the economic impact of non-
catastrophic fatigue failures has been tremendous. Fatigue of structures is now
generally recognized as a significant problem.
As a result of extensive research and practical experience, much knowledge has
been gained about fatigue of structures and the fatigue mechanism in the material.
Much has been learned from laboratory research. However, accident investigations
have also highly contributed to the present state of the art. Fatigue failures in service
can be most instructive and provide convincing evidence that fatigue may be a
serious problem. The analysis of failures often reveals various weaknesses
contributing to an insufficient fatigue resistance of a structure. This will be illustrated
here by a case history. The front wheel of a heavy motorcycle completely collapsed,
see Figure 1.1a. Ten spokes of the light alloy casting were broken. Examination of
the failure surfaces indicated that fatigue cracks occurred in all spokes, see Figure
1.1b.Why was the fatigue life of this wheel insufficient? A first question of a failure
analysis must be: Was the failure a symptomatic failure or was it an incidental case?
If it is a symptomatic failure, all motorcycles of the same type are in danger and
immediate action is required. However, the failure may be an incidental case for
some special reason applicable to that single motorcycle only: for instance, unusual
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and severe damage of the material surface. In the case of this motorcycle, the same
failure had occurred in several wheels in different countries, although predominantly
in motorcycles of the police. The wheel shown in Figure 1.1 collapsed when a
policeman suddenly had to use the brakes to stop before a railway crossing. He
survived after some heavy shocks.
Fig 1.1a Front wheel, broken spikes, axle part with drum
Fatigue fractures brakes
Fig. 1.1b Collapse of the front wheel of a motorcycle by fatigue of the spokes.
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A structure should be designed and produced in such a way that undesirable fatigue
failures do not occur during the design life of the structure.
A special issue is how to account for environmental effects. Experimental data used
in the predictions are generally obtained under laboratory conditions and relatively
high testing frequencies. However, in service corrosive environments may be present
and the load frequency can be much lower. As an example, think of a welded
structure for a drilling platform in the sea. The environment is salt water, and the
loading rate of water waves is relatively low [1].
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CHAPTER 2
2.1 Fatigue as a Phenomenon in the Material
In a specimen subjected to a cyclic load, a fatigue crack nucleus can be initiated on a
microscopically small scale, followed by crack grows to a macroscopic size, and
finally to specimen failure in the last cycle of the fatigue life.
Microscopic investigations in the beginning of the 20th century have shown that
fatigue crack nuclei start as invisible micro cracks in slip bands. After more
microscopic information on the growth of small cracks became available, it turned out
that nucleation of micro cracks generally occurs very early in the fatigue life.
Indications were obtained that it may take place almost immediately if a cyclic stress
above the fatigue limit is applied. The fatigue limit is the cyclic stress level below
which a fatigue failure does not occur. In spite of early crack nucleation, micro cracks
remain invisible for a considerable part of the total fatigue life. Once cracks become
visible, the remaining fatigue life of a laboratory specimen is usually a small
percentage of the total life. The latter percentage may be much larger for real
structures such as ships, aircraft, etc. After a micro crack has been nucleated, crack
growth can still be a slow and erratic process, due to effects of the microstructures,
e.g. grain boundaries. However, after some micro crack growth has occurred away
from the nucleation site, a more regular growth is observed. This is the beginning of
the real crack growth period. Various steps in the fatigue life are indicated in Figure
2.1. The important point is that the fatigue life until failure consists of two periods: the
crack initiation period and the crack growth period. Corrosive environments can affect
initiation and crack growth, but in a different way for the two periods. It should be
noted here that fatigue prediction methods are different for the two periods. The
stress concentration factor Kt is the important parameter for predictions on crack
initiation. The stress intensity factor K is used for predictions on crack growth [1].
Fig. 2.1 Different phases of the fatigue life and relevant factors.
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Fig. 2.2. Different scenarios of fatigue crack growth.
The crack initiation period includes the initial micro crack growth. Because the
growth rate is still low, the initiation period may cover a significant part of the fatigue
life. This is illustrated by the generalized picture of crack growth curves presented in
Figure 2.2. which schematically shows the crack growth development as a function
of the percentage of the fatigue life consumed (= n/N), with n as the number of
fatigue cycles and N as the fatigue life until failure. Complete failure corresponds to
n/N = 1 = 100%. There are three curves in Figure 2.2, all of them in agreement with
crack initiation in the very beginning of the fatigue life, however, with different values
of the initial crack length. The lower curve corresponds to micro crack initiation at a
―perfect‖ surface of the material.
The middle curve represents crack initiation from an inclusion.
The upper curve is associated with a crack starting from a material defect which
should not have been present, such as defects in a welded join.
Figure 2.2 illustrates some interesting aspects:
The vertical crack length scale is a logarithmic scale, ranging from 0.1
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nanometer (nm) to 1 meter (1 nanometer = 10−9 m = 4·10−8 inch). Micro
cracks starting from a perfect free surface can have a sub-micron crack length
(<1 μm = 10−6 m). However, cracks nucleated at an inclusion will start with a
size similar to the size of the inclusion. The size can still be in the sub-
millimeter range. Only cracks starting from macro defects can have a
detectable macro crack length immediately.
The two lower crack growth curves illustrate that the major part of the fatigue
life is spent with a crack size below 1 mm, i.e. with a practically invisible crack
size.
Dotted lines in Figure 2.2. indicate the possibility that cracks do not always
grow until failure. It implies that there must have been barriers in the material
which stopped crack growth [1].
2.2 Different phases of the fatigue life
Fatigue fractures have a characteristic appearance which reflects the initiation site
and the progressive development of the crack front, culminating in an area of final
overload fracture.
Initiation site(s).
Progressive development of crack front characterised by beach marks.
Culminating in an area of final fracture.
Fig. 2.3a illustrates fatigue failure in a circular shaft. The initiation site is shown and
the shell-like markings, often referred to as beach markings because of their
resemblance to the ridges left in the sand by retreating waves, are caused by
arrests in the crack front as it propagates through the section. The hatched region
on the opposite side to the initiation site is the final region of ductile fracture.
Sometimes there may be more than one initiation point and two or more cracks
propagate. This produces features as in Fig. 2.3b with the final area of ductile
fracture being a band across the middle. This type of fracture is typical of double
bending where a component is cyclically strained in one plane or where a second
fatigue crack initiates at the opposite side to a developing crack in a component
subject to reverse bending. Some stress-induced fatigue failures may show multiple
initiation sites from which separate cracks spread towards a common meeting point
within the section [2].
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Fig. 2.3 a,b,c
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Fatigue strength is determined by applying different levels of cyclic stress to
individual test specimens and measuring the number of cycles to failure. Standard
laboratory test use various methods for applying the cyclic load, e.g. rotating bend,
cantilever bend, axial push-pull and torsion. The data are plotted in the form of a
stress-number of cycles to failure (S-N) curve, fig 2.4. Owing to the statistical
nature of the failure, several specimens have to be tested at each stress level.
Some materials, notably low-carbon steels, exhibit a flattening off at a particular
stress level as at (a) in Fig.2.4 which is referred to as the fatigue limit. As a rough
guide, the fatigue limit is usually about 40% of the tensile strength. In principle,
components designed so that the applied stresses do not exceed this level should
not fail in service. The difficulty is a localised stress concentration may be present
or introduced during service which leads to initiation, despite the design stress being
normally below the ‗safe‘ limit. Most materials, however, exhibit a continually falling
curve as at (b) and the usual indicator of fatigue strength is to quote the stress
below which failure will not be expected in less than a given number of cycles
which is referred to as the endurance limit.
Fig 2.4
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Fig 2.5 a,b Sample pictures of fatigue failure
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Fig 2.6 Micro crack at grain boundary during fatigue test.
Fig 2.7 Fatigue failure
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Although fatigue data may be determined for different materials it is the shape of a
component and the level of applied stress which dictate whether a fatigue failure is to
be expected under particular service conditions. Surface condition is also important.
Often complete components or assemblies, e.g. railway bogie frames or aircraft
fuselage will be tested by subjecting them to an accelerated loading spectrum
reproducing what they are likely to experience over their entire service lifetime.
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CHAPTER 3
3.1 Why Metal Parts Fail From Repeatedly-Applied
Loads
Long ago, engineers discovered that if you repeatedly applied and then removed a
nominal load to and from a metal part (known as a ―cyclic load‖), the part would break
after a certain number of load-unload cycles, even when the maximum cyclic stress
level applied was much lower than the UTS, and in fact, much lower than the Yield
Stress. These relationships were first published by A. Z. Wöhler in 1858. They
discovered that as they reduced the magnitude of the cyclic stress, the part would
survive more cycles before breaking. This behaviour became known as “FATIGUE”
because it was originally thought that the metal got “tired”. When you bend a paper clip
back and forth until it breaks, you are demonstrating fatigue behaviour.
Some common questions about metal fatigue are:
What is fatigue loading?
How do you determine the fatigue strength of a material?
Does the strength of a material affect its fatigue properties?
Why is the surface of a part so important?
Is fatigue life an exact number?
Do real-world parts behave the same as laboratory tests?
Are fatigue cycles cumulative?
3.1.1 WHAT IS FATIGUE LOADING?
There are different types of fatigue loading. One type is zero-to-max-to zero, where a
part which is carrying no load is then subjected to a load, and later, the load is
removed, so the part goes back to the no-load condition. An example of this type of
loading is a chain used to haul logs behind a tractor.
Another type of fatigue loading is a varying load superimposed on a constant load. The
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suspension wires in a railroad bridge are an example of this type. The wires have a
constant static tensile load from the weight of the bridge, and an additional tensile load
when a train is on the bridge.
The worst case of fatigue loading is the case known as fully-reversing load. One cycle
of this type of fatigue loading occurs when a tensile stress of some value is applied to
an unloaded part and then released, then a compressive stress of the same value is
applied and released.
Fig 3.1
A rotating shaft with a bending load applied to it is a good example of fully reversing
load. In order to visualize the fully-reversing nature of the load, picture the shaft in a
fixed position (not rotating) but subjected to an applied bending load (as shown in Fig
3.1). The outermost fibres on the shaft surface on the convex side of the deflection
(upper surface in the picture) will be loaded in tension (upper green arrows), and the
fibres on the opposite side will be loaded in compression (lower green arrows). Now,
rotate the shaft 180° in its bearings, with the loads remaining the same. The shaft stress
level is the same, but now the fibres which were loaded in compression before you
rotated it are now loaded in tension, and vice-versa.
To illustrate how damaging fully-reversing load is, take a paper clip, bend it out
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straight, then pick a spot in the middle, and bend the clip 90° back and forth at that
spot (from straight to ―L‖ shaped and back). Because you are plastically-deforming the
metal, you are, by definition, exceeding its yield stress. When you bend it in one
direction, you are applying a high tensile stress to the fibres on one side of the OD, and
a high compressive stress on the fibres on the opposite side. When you bend it the
other way, stresses are reversed (fully reversing fatigue). It will break in about 25
cycles.
The number of cycles that a metal can endure before it breaks is a complex function of
the static and cyclic stress values, the alloy, heat-treatment and surface condition of
the material, the hardness profile of the material, impurities in the material, the type of
load applied, the operating temperature, and several other factors.
3.1.2 HOW IS THE FATIGUE STRENGTH OF A METAL DETERMINED?
The fatigue behaviour of a specific material, heat-treated to a specific strength level, is
determined by a series of laboratory tests on a large number of apparently identical
samples of that specific material.
This picture shows a laboratory fatigue specimen. These laboratory samples are
optimized for fatigue life. They are machined with shape characteristics which maximize
the fatigue life of a metal, and are highly polished to provide the surface characteristics
which enable the best fatigue life.
Fig 3.2
A single test consists of applying a known, constant bending stress to a round sample
of the material, and rotating the sample around the bending stress axis until it fails. As
the sample rotates, the stress applied to any fibre on the outside surface of the sample
varies from maximum-tensile to zero to maximum-compressive and back. The test
mechanism counts the number of rotations (cycles) until the specimen fails. A large
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number of tests is run at each stress level of interest, and the results are statistically
massaged to determine the expected number of cycles to failure at that stress level.
The cyclic stress level of the first set of tests is some large percentage of the Ultimate
Tensile Stress (UTS), which produces
failure in a relatively small number of
cycles. Subsequent tests are run at lower
cyclic stress values until a level is found at
which the samples will survive 10 million
cycles without failure. The cyclic stress
level that the material can sustain for 10
million cycles is called the Endurance
Limit (EL).
In general, steel alloys which are
subjected to a cyclic stress level below the
EL (properly adjusted for the specifics of
the application) will not fail in fatigue. That
property is commonly known as ―infinite
life‖. Most steel alloys exhibit the infinite
life property, but it is interesting to note
that most aluminium alloys as well as steels which have been case-hardened by
carburizing, do not exhibit an infinite-life cyclic stress level (Endurance Limit).
3.1.3 IS THERE ANY RELATIONSHIP BETWEEN UTS AND FATIGUE STRENGTH?
The endurance limit of steel displays some interesting properties. These are shown, in
a general way, in this graph, (Figure 7) and briefly discussed below.
It is a simplistic rule of thumb that, for steels having a UTS less than 160,000 psi, the
endurance limit for the material will be approximately 45 to 50% of the UTS if the
surface of the test specimen is smooth and polished.
That relationship is shown by the line titled ―50%‖. A very small number of special case
materials can maintain that approximate 50% relationship above the 160,000 psi level.
However, the EL of most steels begins to fall away from the 50% line above a UTS of
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about 160,000 psi, as shown by the line titled ―Polished‖.
For example, a specimen of SAE-4340 alloy steel, hardened to 32 Rockwell-C (HRc),
will exhibit a UTS around 150,000 psi and an EL of about 75,000 psi, or 50% of the
UTS. If you change the heat treatment process to achieve a hardness of about 50
HRc, the UTS will be about 260,000 psi, and the EL will be about 85,000 psi, which is
only about 32% of the UTS.
Several other alloys known as ―ultra-high-strength steels‖ (D-6AC, HP-9-4-30, AF-1410,
and some maraging steels) have been demonstrated to have an EL as high as 45% of
UTS at strengths as high as 300,000 psi. Also note that these values are EL numbers
for fully-reversing bending fatigue. EL values for hertzian (contact) stress can be
substantially higher (over 300 ksi).
The line titled ―Notched‖ shows the dramatic reduction in fatigue strength as a result of
the concentration of stress which occurs at sudden changes in cross-sectional area
(sharp corners in grooves, fillets, etc.). The highest EL on that curve is about 25% of the
UTS (at around 160,000 psi).
The surface finish of a material has a dramatic effect on the fatigue life. That fact is
clearly illustrated by the curve titled ―Corroded‖. It mirrors the shape of the ―notched‖
curve, but is much lower. That curve shows that, for a badly corroded surface (fretting,
oxidation, galvanic, etc.) the endurance limit of the material starts at around 20 ksi for
materials of 40 ksi UTS (50%), increases to about 25 ksi for materials between 140
and 200 ksi UTS, then decreases back toward 20 ksi as the material UTS increases
above 200 ksi.
3.1.4 WHY IS THE SURFACE SO IMPORTANT?
Fatigue failures almost always begin at the surface of a material. The reasons are that
the most highly-stresses fibres are located at the surface (bending fatigue)
the intergranular flaws which precipitate tension failure are more frequently found
at the surface.
Suppose that a particular specimen is being fatigue tested (as described above). Now
suppose the fatigue test is halted after 20 to 25% of the expected life of the specimen
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and a small thickness of material is machined off the outer surface of the specimen,
and the surface condition is restored to its original state. Now the fatigue test is
resumed at the same stress level as before. The life of the part will be considerably
longer than expected. If that process is repeated several times, the life of the part may
be extended by several hundred percent, limited only by the available cross section of
the specimen. That proves fatigue failures originate at the surface of a component.
3.1.5 IS THE ENDURANCE LIMIT AN EXACT NUMBER?
It is important to remember that the Endurance Limit of a material is not an absolute nor
fully repeatable number. In fact, several apparently identical samples, cut from adjacent
sections in one bar of steel, will produce different EL values (as well as different UTS
and YS) when tested, as illustrated by the S-N diagram below. Each of those three
properties (UTS, YS, EL) is determined statistically, calculated from the (varying)
results of a large number of apparently identical tests done on a population of
apparently identical samples.
The plot below shows the results of a battery of fatigue tests on a specific material.
The tests at each stress level form statistical clusters, as shown. A curve is fitted
through the clusters of points, as shown below. The curve which is fitted through these
clusters, known as an ―S-N Diagram‖ (Stress vs. Number), represents the statistical
behaviour of the fatigue properties of that specific material at that specific strength
level. The red points in the chart represent the cyclic stress for each test and the
number of cycles at which the specimen broke. The blue points represent the stress
levels and number of cycles applied to specimens which did not fail. This diagram
clearly demonstrates the statistical nature of metal fatigue failure.
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Fig 3.4
3.1.6 DO REAL-WORLD COMPONENTS EXHIBIT THE “LABORATORY” EL?
Unfortunate experience has taught engineers that the value of the Endurance Limit
found in laboratory tests of polished, optimized samples does not really apply to real-
world components.
Because the EL values are statistical in nature, and determined on optimized, laboratory
samples, good design practice requires the determination of the actual EL will be for
each specific application, known as the Application-Specific Endurance Limit (AS EL).
In order to design for satisfactory fatigue life (prior to testing actual components), good
practice requires that the ―laboratory‖ Endurance Limit value be reduced by several
adjustment factors. These reductions are necessary to account for:
(a) the differences between the application and the testing environments, and
(b) the known statistical variations of the material.
This procedure is to insure that both the known and the unpredictable factors in the
application (including surface condition, actual load, actual temperature, tolerances,
impurities, alloy variations, heat-treatment variations, stress concentrations, etc. etc.
etc.) will not reduce the life of a part below the required value.
An accepted contemporary practice to estimate the maximum fatigue loading which a
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specific design can survive is the Marin method, in which the laboratory test-
determined EL of the particular material (tested on optimized samples) is adjusted to
estimate the maximum cyclic stress a particular part can survive (the ASEL).
This adjustment of the EL is the result of six fractional factors. Each of these six factors
is calculated from known data which describe the influence of a specific condition on
fatigue life.
Those factors are:
1. Surface Condition (ka): such as: polished, ground, machined, as-forged,
corroded, etc. Surface is perhaps the most important influence on fatigue life;
2. Size (kb): This factor accounts for changes which occur when the actual size of
the part or the cross-section differs from that of the test specimens;
3. Load (kc): This factor accounts for differences in loading (bending, axial,
torsional) between the actual part and the test specimens;
4. Temperature (kd): This factor accounts for reductions in fatigue life which occur
when the operating temperature of the part differs from room temperature (the testing
temperature);
5. Reliability (ke): This factor accounts for the scatter of test data. For example, an
8% standard deviation in the test data requires a ke value of 0.868 for 95% reliability,
and 0.753 for 99.9% reliability.
6. Miscellaneous (kf): This factor accounts for reductions from all other effects,
including residual stresses, corrosion, plating, metal spraying, fretting, and others.
These six fractional factors are applied to the laboratory value of the material endurance
limit to determine the allowable cyclic stress for an actual part:
Real-World Allowable Cyclic Stress = ka * kb * kc * kd * ke * kf * EL
3.1.7 IS FATIGUE LOADING CUMULATIVE?
It is important to realize that fatigue cycles are accumulative. Suppose a part which has
been in service is removed and tested for cracks by a certified aircraft inspection
station, a place where it is more likely that the subtleties of Magnaflux inspection are
well-understood. Suppose the part passes the inspection, (i.e., no cracks are found)
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and the owner of the shaft puts it on the ―good used parts‖ shelf.
Later, someone comes along looking for a bargain on such a part, and purchases this
―inspected‖ part. The fact that the part has passed the inspection only proves that there
are no detectable cracks RIGHT NOW. It gives no indication at all as to how many
cycles remain until a crack forms. A part which has just passed a Magnaflux inspection
could crack in the next 100 cycles of operation and fail in the next 10000 cycles (which
at 2000 RPM, isn‘t very long) [2].
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CHAPTER 4
4.1 Fracture Mechanisms
In very general terms, when crack size (a) in a stressed part reaches a critical size
(acr), the fracture process occurs almost instantaneously with complete and sudden
separation of the part. This stage of final fracture is referred to as overload fracture.
However, it must be recognized that many overload failures occur after subcritical (a ˂
acr) crack growth from various progressive damage mechanisms such as fatigue or
environmentally assisted cracking. Overload cracking can be categorized into three
general types of mechanisms:
● Brittle overload from cleavage (i.e., transgranular brittle cracking)
● Ductile overload failures that involve the fracture mechanisms of ductile tearing
and/ or micro void formation caused by transgranular slip
● Stress rupture (sometimes referred to as de-cohesive rupture), when a pair of free
surfaces are created from a pre-existing grain boundary or second-phase
boundary
Ductile and brittle cracking are the two main types of overload fracture, while stress
rupture includes various types of mechanisms that may be either brittle or ductile. For
example, a brittle stress-rupture failure may occur from grain boundary embrittlement,
while a ductile stress rupture failure may occur from grain-boundary slip (i.e., time-
dependent creep deformation in polycrystalline metals). When examining fracture
surfaces, it is important to obtain an overall perspective from both macroscopic and
microscopic study. Examination beyond the fracture surface also provides information.
For example, visual inspection of a fractured component may indicate events prior to
fracture initiation, such as a shape change indicating prior deformation. Metallographic
examination of material removed far from the fracture surface also can provide
information regarding the penultimate microstructure, including the presence of cold
work (bent annealing twins, deformation bands, and/or grain shape change), evidence
of rapid loading and/or low-temperature service (deformation twins), and so forth.
These types of investigative methods are also important in the analysis of fractures.
In general, identifying the cause and corrective action of a fracture benefits by the
careful documentation of various macroscopic and microscopic observations. Typically,
observation begins with a visual examination of a fracture surface, where general
26
features and surface roughness can be revealed under favourable lighting.
Examination at low magnification (about 15 X or less) can also reveal features
regarding the nature of the fracture path. Metallographic and fractographic techniques
then can be used to reveal microscopic features. The failed piece may be properly
sectioned for preparation of metallographic samples and examination under an optical
microscope. Alternatively or in addition, an electron microscope—typically a scanning
electron microscope (SEM), or a transmission electron microscope (TEM)—with higher
magnifications and depth of field may be used to make direct examination of the raw
fracture surface.
Table 4.1 Maximum shear stress as a function of states of stress
Both the macro- and micro scale appearances of fracture-surface features can tell a
story of how and sometimes why fracture occurred in terms of the following
information:
● Crack initiation site and crack propagation direction
● Mechanism of cracking and the path of fracture
● Load conditions (tension, bending, shear, monotonic, or cyclic)
● Environment
● Geometric constraints that influenced crack initiation and/or crack propagation
● Fabrication imperfections that influenced crack initiation and/or crack propagation
4.1.1 Ductile Fracture
The mechanisms and appearances of ductile fracture are best introduced with the
simple example of an unnotched bar subjected to conventional tension testing at quasi-
static strain rates (˂0.1 s-1). In general, tension test specimens of a ductile material
27
have a visible region of necking (Fig. 4.1), while a brittle specimen results in fracture
with little or no visible evidence of any necking (Fig. 4.2). Necking is a region of strain
localization that forms when the increase in stress due to decrease in the cross-
sectional area of the specimen becomes greater than the increase in the load-carrying
ability of the metal due to strain. Necking generally occurs at the point of maximum
load on the engineering stress-strain curve. Prior to the onset of necking, strain is
uniform along the gage length, while plastic deformation becomes concentrated in the
necked region of the tension specimen. The size of the neck and the extent to which it
is visible depends primarily on strain hardening and strain-rate hardening (when
temperatures are below 0.4 Tm of the metallic material, where Tm is the material
melting point on the Kelvin scale). The resulting necked region is, in effect, a mild
notch, which introduces a complex state of stress that has a large tensile-hydrostatic
component. This tensile-hydrostatic (or triaxial) stress is highest in the centre of the
specimen, where micro voids occur from the tensile separation of the ductile matrix
from harder second-phase particles or inclusions (which are present in most
commercial alloys). This central region of the fracture surface also is typically flat (at
the macro scale), which indicates separation from a tension stress state. Thus, even
though the ductile fracture involves deformation, the microscopic mechanism of crack
initiation involves the brittle like effect of tensile separation in the region of a
Fig. 4.1 Cup-and-cone fracture of a low-carbon steel bar under tension.
28
Fig. 4.2 Brittle fracture of a smooth (unnotched) tensile test specimen.
Deformation-induced notch. The micro voids also grow and connect to ultimately form
and fracture near the central region of the specimen. However, as the central crack
grows, the material in the outer annulus deforms by stresses along the shear plane
(45° to the direction of the tensile load). This results in distinctive shear lips of ductile
fracture and the classic cup-and-cone profile of the fracture surface. One piece has a
cone like profile, while the other piece (shown in Fig. 4.1) has the surface profile of a
cup. The ratio of the area of the flat-face region to the area of the shear lip usually
increases with section thickness. The area of the shear lip also depends on the extent
of necking.
The amount of necking depends on the extent of strain and strain-rate hardening,
which in turn are influenced by factors such as temperature and material condition.
Lowering the temperature below room temperature generally increases the strain-
hardening exponent, n. This increases the strain to neck formation. In contrast, an ideal
plastic material (in which no strain hardening occurs) becomes unstable in tension and
begins to neck as soon as yielding occurs. Most metals and alloys (not heavily cold
worked) undergo strain hardening, which tends to increase the load-carrying capacity
of the specimen as deformation increases. Commercial engineering materials also
typically contain inclusions, second-phase particles, and other constituents. These
microscopic constituents can influence the process of fracture nucleation and crack
growth. In an ideal material containing neither inclusions nor second phases, ductile
fracture would be expected to occur by slip and possibly twinning, resulting in complete
29
reduction in area. Alternately, cleavage across a grain on a single plane would be
expected to result in a smooth fracture surface. Such results are sometimes observed
in high-purity single-crystal specimens, but are seldom seen in commercial engineering
materials.
Fracture by uninterrupted plastic deformation is a special type of plastic fracture. For
metals that do not work harden much, the metal under tension would be drawn down
almost to a chisel edge or a point before breaking apart (see Fig. 4.3a). This special
circumstance can hardly be termed ―fracture‖ in a normal sense, as there is no fracture
surface at all. It is usually referred to as ―rupture‖ because the process arises from
prolonged shear on slip planes within the worked region of the crystals, which finally at
one point shear apart. It should also be noted that strain rate and adiabatic deformation
also play an important role in this type of behaviour. That is, high temperature and very
slow strain rates can result in extensive uninterrupted deformation.
Another form of uninterrupted plastic deformation, which may or may not result in a
―chisel point‖ type fracture, is the shearing of a single crystal. Deformation of a single
crystal is governed by the critical resolved stress and slip occurs in an active slip plane
and slides in a specific direction. Figure 4.3(b) shows a copper-aluminium single crystal
that has gone through such a prolonged extension. Final separation of the specimen
eventually occurs by ―shearing off‖ at one of the slip planes. Whether a single crystal
will fracture by shearing off or by drawing down to a chisel point depends on the slip
system of a particular crystal.
4.1.1.1 General Macroscopic Appearance of Ductile Fractures. For ductile
fracture, macroscopic features include:
● A relatively large amount of plastic deformation precedes the fracture.
● Shear lips are usually observed at the fracture termination areas.
● The fracture surface may appear to be fibrous or may have a matte or silky
texture, depending on the material.
● The cross section at the fracture is usually reduced by necking.
● Crack growth is slow.
30
Fig. 4.3 Examples of uninterrupted shear failure.
(a) Polycrystalline aluminium bars pulled at 600 °C (1110°F).
(b) Extended copper-aluminium single crystal.
Ductile fractures often progress as single cracks, without many separated pieces or
substantial crack branching at the fracture location. The region of crack initiation
typically has a dull fibrous appearance that is indicative of cracking by micro void
coalescence. The crack profiles adjacent to the fracture are consistent with tearing.
The fracture surface may have radial markings, chevrons, and/or shear lips, depending
on the specimen geometry and material condition.
Depending on the state of stress and geometric constraints on macroscopic ductility,
the fracture of a ductile material may occur by plane strain, plane stress, or mixed
mode. In general, these variations in fracture profiles are related to fracture toughness,
which depends on section thickness (B) and the crack size (a) of a pre-existing
discontinuity such as a crack or notch. Figure 4.4 is a schematic illustration of this for
an inherently ductile material with varying section thickness. Plane-strain fracture is
characterized by a flat surface perpendicular to the applied load. Plane stress fracture
occurs when shear strain becomes the operative mode of deformation and fracture (as
maximum stresses occur along the shear plane from the basic principles of continuum
mechanics). In plane-stress cracking, the fracture profile is characterized by shear lips,
which are at about a 45° oblique angle to the maximum stress direction (although this
angle may vary depending on material condition and loading condition).
The classic cup-and-cone appearance that results from ductile fractures of unnotched
cylindrical tension test specimens is a good example of mixed-mode fracture. In this
31
case, crack initiation near the specimen centre occurs from fracture under triaxial
tension and thus has a flat fracture surface normal to the applied load. When fracture
reaches the region near the outer surface, deformation by slip (shear) becomes
dominant, and the stress state of fracture changes to plane stress. However, shear lips
are not necessarily the definitive characteristic of a ductile fracture. The macro scale
appearance of a fracture is also influenced by geometric constraints and stress-state
conditions. For example, even though a flat centre region of crack initiation is
characteristic of tensile (brittle like) separation, the flat fracture region of a ductile
fracture also has a dull fibrous appearance with small microscopic dimples.
Microscopic dimples are indicative of separation from displacement within a ductile
matrix (or localized regions) of a material.
4.1.2 Brittle Fracture
Materials that do not develop a neck before fracture are generally considered brittle.
When the specimen lacks ductility (due to low temperature, environment, strain rate, or
the material itself), the fracture is brittle and occurs by separation on a plane that is
normal to the direction of the applied load. Due to absence of necking, deformation is
approximately uniform until fracture occurs from complete separation under tension.
However, lack of ductility depends on a number of other factors, such as environment,
strain rate, and the internal state of stress created (influenced by part geometry and
discontinuities in the material). Therefore, lack of ductility is not just due to the material
itself, but is also influenced by complex relationships of stress state, part geometry,
localized deformation, and internal discontinuities.
Conversely, a fracture surface normal to the applied load also is not necessarily a
definitive indication of an inherently brittle material or a brittle mechanism (such as
cleavage or intergranular fracture). This important distinction is indicated in the central
region of the classic ductile fracture in Fig. 4.1. Initial deformation causes a region of
localized strain (i.e., necking), which is essentially a mild notch that results in the
development of hydrostatic tensile stresses in the interior of a tension test specimen.
This ―triaxial‖ state of tension then causes crack initiation (void formation) by tensile
separation around small inclusions, second-phase particles, or discontinuities in the
centre region of the tension bar. In effect, ductile mechanisms lead to a stress state
that causes tensile separation around less ductile.
32
Fig. 4.4. Schematic of variation in fracture behaviour and macro
scale features of fracture surfaces for an inherently ductile material. As
section thickness increases, plane strain conditions develop first along
the centre line and result in a flat fracture surface. With further
increases in section thickness, the flat region spreads to the outside of
the specimen, decreasing the widths of the shear lips.
In general, brittle fracture can be distinguished by these characteristics:
● Little or no visible plastic deformation precedes the fracture.
● The fracture surface is generally flat and perpendicular to the loading direction
and to the component surface.
● The fracture may appear granular or crystalline and is often highly light reflective.
Facets may also be observed, particularly in coarse-grain steels.
● Chevron patterns may be present.
● Rapid crack growth results in catastrophic failure, sometimes accompanied by a
loud noise.
Brittle overload failures, in contrast to ductile overload failures, are characterized by
little or no macroscopic plastic deformation. Brittle fracture initiates and propagates
more readily than ductile fracture or for so-called ―subcritical‖ crack propagation
processes such as fatigue or stress-corrosion cracking (SCC). Because brittle fractures
are characterized by relatively rapid crack growth, the cracking process is sometimes
referred to as being ―unstable‖ or ―critical‖ because the crack propagation leads quickly
to final fracture.
The macroscopic behaviour of brittle fracture is essentially elastic up to the point of
failure. The energy of the failure is principally absorbed by the creation of new
33
surfaces—that is, cracks. For this reason, brittle failures often contain multiple cracks
and separated pieces, which are less common in ductile overload failures. All brittle
fracture mechanisms can exhibit chevron or herringbone patterns that indicate the
fracture origin and direction of rapid fracture progression. Herringbone patterns are
unique microscopic features of brittle fractures. Ductile cracking, which occurs by micro
void coalescence, does not result in a herringbone pattern. On a microscopic scale, the
features and mechanisms of fracture may have components of ductile or brittle crack
propagation, but the macroscopic process of fracture is characterized by little or no
work expended from deformation.
4.1.2.1 Microscopic Aspects of Fracture
Although the examination of a fracture surface begins logically with macroscopic
observations, the microscopic aspects of fracture are also essential in understanding
the causes of fracture. In general, fracture on a micro scale can be distinctively defined
in terms of transgranular brittle fracture (cleavage), transgranular slip (ductile fracture),
or intergranular fractures. These three types of fracture paths result in distinctly
different fracture surfaces, as seen in Fig. 4.5. These distinct microscopic appearances
provide important information on the underlying mechanisms of fracture.
On a microscopic level, most engineering materials (alloys) are polycrystalline solids
that consist of many grains (crystals) with grain boundary regions between the crystals.
Typically, the grain boundaries are stronger than individual grains in properly
processing polycrystalline plastic/elastic solids. The reason for this can be understood
in simple terms. The grain boundaries are disruptions between the crystal lattice of
individual grains, and this disruption provides a source of strengthening by pinning the
movement of dislocations. Thus, a finer-grain alloy imparts more grain-boundary
regions for improved strength. Moreover, because a greater number of arbitrarily
aligned grains are achieved when grain size is reduced, the stressed material has
more opportunity to allow slip and thus improve ductility.
However, the grain boundaries are also a region with many faults, dislocations, and
voids. This relative atomic disarray of the grain boundaries, as compared to the more
regular atomic arrangement of the grain interiors, provides an easy path for diffusion-
related (thermally activated) alterations. For this reason, grain boundaries can be a
preferential region for congregation and segregation of impurities, preferential phase
precipitation, and/or absorption of environmental species. These thermally activated
34
alterations are potential mechanisms for weakening or embrittlement along the grain
boundaries. In addition, grain-boundary regions are weakened when the temperature is
high enough to activate diffusion-induced flow deformation along grain boundaries at
stresses below the yield strength. This onset of time-dependent flow (i.e., creep
deformation) is roughly representative of the viscoelastic deformation that occurs when
the temperature is higher than 0.4 Tm.
This basic overview of the relative strength of grains and grain boundaries provides a
general framework to describe the basic mechanisms of transgranular and
intergranular fracture. Transgranular fracture of a crystalline material can occur by
either a brittle process of cleavage or by the ductile process of micro void formation.
These two mechanisms of transgranular fracture are very distinct in terms of
appearance and have clear causes. Brittle transgranular fracture of a crystalline
material takes place by cleavage along low-index crystallographic planes within grains,
while ductile transgranular fracture occurs when small voids (micro voids) form and
coalesce in the region of fracture. These micro voids leave distinctive concave
depressions called dimples on both surfaces of the fracture.
Intergranular fractures are also very distinctive in appearance (Fig. 4.5a), but the
underlying causes or mechanisms can be complex and varied. Grain boundaries are
weakened in various ways, and the surface of an intergranular fracture does not
necessarily reveal the evidence of the underlying mechanism that leads to fracture
along a grain-boundary region. This illustrates the need for careful analysis of the
overall circumstances that lead to fracture. In many instances, SEM fractography
provides a means to identify the fracture path as intergranular, but it may yield little
other information. Additional important information can be obtained by chemical
analysis in the SEM and by microstructural examination. Use of the SEM for
microstructural examination in addition to optical light examination may be appropriate
depending on the scale of micro constituents present. Important clues on the
underlying cause of intergranular separation also may be revealed by fractography at
assorted magnifications.
35
Fig. 4.5 SEM images of (a) intergranular fracture in ion-nitrided layer of
ductile iron (ASTM 80-55-06), (b) transgranular fracture by cleavage in
ductile iron (ASTM 80-55-06), and (c) ductile fracture with equiaxed dimples
from micro void coalescence around graphite nodules in a ductile iron (ASTM
65-40-10). Picture widths are approximately 0.2 mm (0.008 in.) from original
magnifications of 500X.
4.1.3 Transgranular Ductile Fracture (Transgranular Slip and Micro void
Formation).
In terms of inherent material structure of crystalline materials, the deformation
processes of slip and twinning compete with the brittle fracture process of cleavage.
Cleavage is a brittle process that occurs on the plane of maximum normal stress, while
slip mechanisms are associated with plastic deformation and ductile fracture. At
temperatures lower than 0.4 Tm, plastic deformation occurs by transgranular slip and/or
36
twinning in the crystalline lattice. If other events do not intervene, this deformation
culminates in fracture first by strain localization (necking or shear band formation), and
then final fracture occurs in the volume or region of strain concentration. At
temperatures of ~0.4 Tm or higher, however, deformation can occur by slip and viscous
grain-boundary flow, as the grain boundary regions become weakened at high
temperature. Thus, the predominant fracture path becomes intergranular in the region
of creep (time-dependent) deformation at temperatures of _0.4 Tm or higher.
The overall process of ductile fracture is illustrated by the preceding example of a
ductile fracture in an unnotched tension test specimen (Fig. 4.1). Ductile fractures are
uniquely characterized by micro voids that form in the region of high stress. In an
unnotched tension test bar, micro voids nucleate and grow in the central region, where
the diffuse notch (created by necking) causes separation due to triaxial (hydrostatic)
tension. These voids coalesce and join together to form a microscopic crack. At the
same time, more small cavities are formed and distributed over the remaining section
of the test piece. A typical example is presented in Fig. 4.6, which shows a cross
section of a tensile specimen containing numerous voids at a stage between necking
and final fracture. The fracture process consists of these voids joining on the plane
perpendicular to the loading direction and coalescing into a central crack. This crack
grows until it approaches the outer annulus, where a change in fracture path occurs as
the process approaches final fracture. At some point (depending on ductility), final
fracture occurs along the shear plane and results in shear lip. The amount of shear lip
varies, depending on ductility, strain rate, and temperature.
The surfaces of plastic fractures are characterized by microscopic ―dimples,‖ also
called ―cupules.‖ These dimples represent the numerous concave depressions left on
the opposite fracture faces of the broken specimen. Dimples on fracture surfaces are
observed in many materials, including carbon and alloy steels, austenitic steels, alloys
of aluminium, titanium, and copper, and plastics. It has been suggested that dimples
represent the coalesced voids, and the voids initiate from inclusions or intermetallic
particles.
Dimples also can take different shapes, depending on loading condition (Fig. 4.7).
Round dimples occur from separation under tension, while the dimples have an
elongated parabolic shape when ductile fracture occurs from shear, torsion, and
tearing (or bending). In general, the concavity of the parabola is oriented toward the
direction of relative displacement of the other half of the specimen, and the axis of
symmetry
37
Fig. 4.6 Section through the neck area of a tensile specimen of
copper showing cavities and crack formed at the centre of the specimen
as the result of void coalescence.
38
Fig. 4.7 Schematic of plastic fracture. (a) Normal plastic (formation of
round dimples). (b) Shear plastic (formation of elongated dimples pointing in
the direction of shear on each fracture surface). (c) Tear plastic (formation of
elongated dimples pointing in the direction opposite to the direction of each
propagation. (d) Dimple elongation from out-of-plane shear. I Dimple
elongation from mixed mode of screw sliding with ductile tear (c _ d).
39
is parallel to the direction of propagation of the rupture front (Fig. 4.7). Thus, when the
two opposite surfaces of rupture due to in-plane shear (Fig. 4.7) are examined, the
concavity of these dimples is turned in opposite directions on the opposite faces.
Parabolic dimples from torsional loading are shown in Fig. 4.8. An example of matching
surfaces is shown in Fig. 4.9 for dimple shape and orientation for bending/ ductile
tearing. In the case of bending, there may be some question as to how often elongated
dimples are seen in bending loading (or seen at all), because dimples on typical
compact tension specimens (axial _ bending) appear mostly equiaxed.
4.1.4 Transgranular Brittle Fracture (Cleavage).
Cleavage is a low-energy fracture that propagates along well-defined low-index
crystallographic planes known as cleavage planes. Theoretically, a cleavage fracture
should have perfectly matching faces and should be completely flat and featureless.
However, engineering alloys are polycrystalline and contain grain and sub grain
boundaries, inclusions, dislocations, and other imperfections that affect a propagating
cleavage fracture so that true, featureless cleavage is seldom observed. These
imperfections and changes in crystal lattice orientation, such as possible mismatch of
the low-index planes across grain or sub grain boundaries, produce distinct cleavage
Fig. 4.8 Parabolic shear dimples from torsional fracture of cast Ti-6Al-
4V alloy. Original magnification, 2000 X
40
fracture surface features, such as cleavage steps, river patterns, feather markings,
herringbone patterns, and tongues. Cleavage fracture also occurs in ceramics,
inorganic glasses, and polymeric materials.
The cleavage mode of fracture is controlled by tensile stresses acting normal to a
cleavage plane and is brittle in nature. Its fracture surface, which is caused by
cleavage, appears at low magnification to be bright or granular, owing to reflection of
light from the flat cleavage surfaces. It exhibits a river pattern when examined under an
electron microscope. It occurs in bcc and hcp metals, particularly in irons and steels,
below the ductile-to-brittle transition temperature (DBTT). Cleavage fracture is very
seldom found in fcc metals. The fcc metals (e.g., copper, aluminium, nickel, and
austenitic steels) have a large number of slip systems (12), which is one reason why
they exhibit high ductility. The fcc metals also are more closely packed (i.e., a shorter
distance exists between atoms in the crystal cell) than are bcc metals. This partly
explains why cleavage fracture does not normally occur in the matrix of fcc metals.
However, it is not just a matter of the multiplicity of ways for slip or cleavage to occur.
Two other factors also control the inherent ductile brittle behaviour of crystalline
materials:
Fig. 4.9 Fracture markings on Plexiglas. TEM, matching fracture
surfaces. Note the matching features A, A‘ and B, B‘ on the two fracture
faces. The parabola markings are similar to the plastic dimples observed in a
tear ductile fracture of a metallic material.
41
● Critical shear stress required to initiate slip
● Critical normal stress required to propagate a cleavage crack
As long as there is only one type of atom in the lattice, the shear stress for slip is low.
The presence of foreign atoms raises this stress and, depending on the location of the
foreign atoms (random or periodic), may cause a severe loss in ductility, as is the case
for the ordered intermetallic alloys. Thus, the critical stresses required for slip and
cleavage also determine whether or not cleavage occurs.
The cleavage process in bcc and hcp metals occurs by separation normal to
crystallographic planes of high atomic density. Microscopic examination of a fracture
surface from cleavage typically reveals distinctive ―river lines‖ indicative of propagation
by fracture along nearly parallel sets of cleavage planes. The direction of crack
propagation is indicated by the ―flow‖ of the river lines as marked by an arrow in Fig.
4.9.
4.1.5 Quasi-Cleavage.
Totally brittle fracture in metals at the microscopic level (―ideal cleavage‖ or ―pure
cleavage‖) occurs only under certain well-defined conditions (primarily when the
component is in single-crystal form and has a limited number of slip systems) and is
correctly described as ―cleavage fracture.‖ More commonly in metals, the fracture
surface contains varying fractions of transgranular cleavage and evidence of plastic
deformation by slip. Grains oriented favourably with respect to the axis of loading may
slip and exhibit ductile behaviour, whereas those oriented unfavourably cannot slip and
will exhibit transgranular brittle behaviour.
When both transgranular fracture processes operate intimately together, the fracture
process is termed ―quasi-cleavage.‖ The dividing line between cleavage and quasi-
cleavage is somewhat arbitrary. The term quasi-cleavage applies when significant
dimple rupture and/or tear ridges accompany the cleavage morphology.
The fracture surface is typically dominated by cleavage, but there are usually small
patches of micro void coalescence present or thin ribbons of micro void coalescence
contained in the fracture surface. As the patches increase, the fracture surface is more
accurately described as (micro scale) mixed cleavage and micro void coalescence.
Another term is ―cleavage with ductile tear ridges.‖
Quasi-cleavage should not be confused with the decohesion along certain
42
crystallographic planes that can occur by shear, by sliding off (plastic shear), or by
separation along weak, still poorly defined interfaces. This type of decohesion has
been referred to as glide-plane decohesion. Quasi-cleavage fractures also should not
be confused with those in which cleavage appears in brittle second phases with the
characteristic dimples of micro void coalescence appearing in the more ductile matrix.
In quasicleavage, there is no apparent boundary between a cleavage facet and a
dimpled area bordering the cleavage facet (Fig. 4.10).
Figure 4.11 is a schematic representation of quasi-cleavage. The occurrence of quasi-
cleavage is usually distinguished by:
● Initiation within facet boundaries—in contrast to fracture by cleavage, which usually
initiates from one edge of the region being cleaved (Fig. 4.12)
● Cleavage steps appearing to blend directly into tear ridges of the adjacent dimpled
areas
Many high-strength engineering metals fracture by quasi-cleavage, which is a mixed
mechanism involving both micro void coalescence and cleavage. When tested under
embrittling conditions, such as those imposed by corrosive mediums or triaxial stress
states, quasi-cleavage can occur in metals that normally are not known to have active
cleavage planes (e.g., austenitic stainless steels, and nickel and aluminium alloys).
One explanation is that facets that exhibit quasi-cleavage features fracture ahead of
the moving crack front; then, as the stress increases, the cleavage facet extends by
tearing into the matrix around it by micro void coalescence.
Quasi-cleavage fracture surfaces appear in steels from:
sudden or impact loading,
low temperature,
high levels of constraint (ambient temperature),
in heavily cold worked parts (ambient temperature).
Quasi cleavage, or cleavage in complex microstructures, is more difficult to identify
than the cleavage found in low-carbon steel made up of ferrite and pearlite. When
identification is uncertain, it is essential to relate the fracture features to the
microstructure, including the prior austenite grain size, the martensite plate size, and
the distribution, size, spacing, and volume fraction of fine carbide particles precipitated
43
during tempering.
With a few exceptions, intergranular fractures are macroscopically brittle with little or no
mechanical work expended as part of the fracture process. However, the micro
mechanisms of intergranular fracture may be brittle or ductile, depending on how the
grain-boundary regions are weakened or embrittled. For example, grain boundaries
may become embrittled by a film of a brittle phase or by the segregation of an impurity
in the boundary region. In this case, the mechanism of intergranular cracking may be
brittle by cleavage in the brittle phases that congregate in the grain boundaries.
Conversely, intergranular fracture may also occur from ductile micro mechanisms (slip)
that involve localized formation of ductile micro voids in the region near the grain
boundaries. For example, voids along the grain boundaries may form at a particle in
Fig. 4.10 Effect of quasi-cleavage—mixed cleavage and micro void
coalescence—on the fracture surface appearance of 17-PH stainless steel.
TEM p-c replica, 4900X
44
Fig. 4.11 Fracture model showing a cleavage step blending with a tear
ridge in a quasi-cleavage fracture surface. At top left is the lower surface of a
fracture, showing a step at the lower left and a ridge at the upper right. At
right and at bottom are sections through the fractured member, showing
profiles of both the upper and the lower fracture surfaces.
The grain boundary or in a precipitate free zone (PFZ) adjacent to the grain boundary.
This feature is sometimes referred to as dimpled intergranular fracture.
The interpretation of intergranular fracture is more complex than the distinct
mechanisms of transgranular fracture (i.e., fatigue, cleavage, or ductile with micro void
coalescence). However, the appearance of intergranular fracture is relatively easy to
recognize, and the causes are fairly limited. The presence of intergranular fracture
(especially in the region of crack initiation) also is often helpful in narrowing the
potential cause for failure. Some common circumstances that have been known to
induce intergranular cracking have been classified into four general categories:
● Presence of grain-boundary precipitates
● Thermal treatment or exposure that causes segregation of certain impurities to
the grain boundaries without an observable second phase
● Stresses applied at elevated (creep-regime) Temperatures
45
Fig. 4.12 Cleavage in a large second-phase particle on a fracture
surface of A-286 steel
● Environmental assisted alteration or weakening of the grain boundaries by
various mechanisms such as hydrogen embrittlement, liquid-metal embrittlement,
solid metal embrittlement, oxidation or reduction potentially in the grain
boundaries, radiation embrittlement, and SCC
Large grain size also plays a role in causing a change from transgranular to
intergranular cracking and can enhance any of the above mechanisms.
4.1.6 Mechanisms of Intergranular Fracture.
On an atomic scale, crack growth occurs by any one or a combination of the following:
● Tensile separation of atoms (decohesion)
● Shear movement of atoms (dislocation egress or insertion)
● Removal or addition of atoms by dissolution or diffusion
All of these processes can occur preferentially along the grain boundary by various
phenomena, such as:
● Segregation of embrittling elements to the grain boundaries
● More rapid diffusion of elements along grain boundaries than along grain interiors
● More rapid nucleation and growth of precipitates in grain boundaries than in grain
46
interiors
● Greater adsorption of environmental species in the grain-boundary regions
These basic mechanisms of intergranular cracking are more varied than those of
transgranular fracture. However, except for conditions of creep stress rupture at
elevated temperatures, intergranular fracture is not common in properly processed
materials in a benign environment. There also are a fairly limited number of
circumstances of intergranular fracture associated with improper processing of a
material and/or some aggressive service environment. Some specific situations
include:
● High-carbon steels with a pearlitic microstructure
● Segregated phosphorus and cementite at prior-austenite grain boundaries in the
high carbon-case microstructures of carburized steels
● Stress-relief cracking
● Grain-boundary carbide films due to eutectoid divorcement in low-carbon steels
● Grain-boundary hypereutecoid cementite in carburized or hypereutectoid steels
● Iron nitride grain-boundary films in nitrided steels
● Temper embrittlement in heat-treated steels due to segregation of phosphorus,
antimony, arsenic, or tin
● Embrittlement of copper due to the precipitation of a high density of cuprous oxide
particles at the grain boundaries
● Embrittlement of steel due to the precipitation of MnS particles at the grain
boundaries as a result of overheating
Fig. 4.13 Quasi-cleavage in the surface of an impact fracture in a
47
specimen of 4340 steel. The same area is shown in both SEM fractographs,
but at different magnifications. The small cleavage facets in martensite
platelets contain river patterns and are separated by tear ridges. Shallow
dimples, marked by arrowheads, are also visible. Direction of crack
propagation is from bottom to top in each fractograph. The specimen was
heat treated at 845°C (1550°F) for 1 h, oil quenched, and tempered at 425°C
(800°F) for 1 h. Fracture was by Charpy impact at _196°C (_321°F). (a)
1650X. (b) 4140X
● Grain-boundary carbide precipitation in stainless steels (sensitization)
● Improperly precipitation-hardened alloys, resulting in coarse grain-boundary
precipitates and a denuded region (PFZ)
● Embrittlement of molybdenum by interstitials (carbon, nitrogen, oxygen)
● Embrittlement of copper by antimony
● Reduction of Cu2O in tough-pitch copper by hydrogen
● Hydrogen embrittlement by grain-boundary absorption of hydrogen
● Stress-corrosion cracking (sometimes intergranular, but also transgranular)
● Liquid metal induced embrittlement (LMIE), for example, mercury in brass, lithium
in type 304 stainless steel
● Solid metal induced embrittlement (SMIE)
In all these cases, SEM fractography can provide the means to identify the fracture
path. However, it cannot yield sufficient information on the underlying causes (or
mechanisms) of intergranular fracture. Thus, additional important information may be
needed in terms of chemical analysis or fractographic examination at assorted
magnifications. The following sections briefly describe appearances for three general
categories of intergranular fracture:
● Intergranular brittle cracking
● Dimpled intergranular fracture
● Intergranular fracture surfaces with corrosion products
4.1.6.1 Intergranular brittle cracking typically has a relatively ―clean‖ fracture
surface with the faceted appearance of cracking along grain contours. The general
appearance of intergranular brittle fracture may include:
48
● Brittle second-phase particles and/or films in grain boundaries
● Fracture where no film is visible and, due to impurities, atom segregation at the
grain boundary
● Environmentally induced embrittlement where there is neither a grain-boundary
precipitate nor solute segregation
Grain-boundary segregation of elements (such as oxygen, sulphur, phosphorus,
selenium, arsenic, tin, antimony, and tellurium) is known to produce intergranular brittle
fractures. Studies of the effects of such impurities in pure iron have been greatly aided
by the development of Auger electron spectroscopy. In the case of brittle grain-
boundary films, it is not necessary for the film to cover the grain boundaries completely;
discontinuous films are sufficient. Some common examples of intergranular
embrittlement by films or segregants include:
● Grain-boundary carbide films in steels
● Iron nitride grain-boundary films in nitrided steels
● Temper embrittlement of alloy steels by segregation of phosphorus, antimony,
arsenic, or tin
● Grain-boundary carbide precipitation in austenitic stainless steels (sensitization)
● Embrittlement of molybdenum by oxygen, nitrogen, or carbon
● Embrittlement of copper by antimony
Grain-boundary strengthening is characteristic of intergranular fractures caused by
embrittlement. Intergranular brittle fracture can usually be easily recognized, but
determining the primary cause of the fracture may be difficult. Fractographic
examinations can readily identify the presence of large fractions of second-phase
particles on grain boundaries. Unfortunately, the segregation of a layer a few atoms
thick of some element or compound that produces intergranular fracture often cannot
be detected by fractography.
4.1.6.2 Dimpled intergranular fractures result in low macroscopic ductility (i.e.,
the grains separate rather than deform) and a fracture surface that reveals microscopic
dimples at higher magnifications (typically on the order of 1000 to 5000X). For
example, a fracture surface from stress rupture of a nickel-base alloy is shown in Fig.
49
4.14 at two levels of magnification. The higher-magnification image reveals a dimpled
topology on the grain facets.
Another example is shown in Fig. 4.15 for a high-purity aluminium-copper precipitation
hardened alloy with a coarse grain structure. In this example (with the coarse grain
size), ductility is limited, and the yield strength in the local region of the grain boundary
becomes lower than the matrix; thus, fracture tends to develop first within the grain-
boundary zone by micro void coalescence. The two levels of magnification provide a
useful combination of images, one demonstrating the intergranular fracture path and
the other revealing the microscopic mechanism of micro void coalescence. Other
circumstances of dimpled intergranular fracture include:
● Uniform void nucleation aided by the formation of methane bubbles at the grain
boundaries (Ref 4.7)
● Void nucleation at precipitates in the grain boundaries of precipitation-hardening
alloys (such as Al-Mg-Zn alloys) with large precipitates on grain boundaries (Ref
4.8) and wide PFZs
● Voids aided by impurities that adsorb strongly on the grain-boundary surface
● Stress-relief cracking of chromium-molybdenum steels (Ref 4.7)
4.1.6.3 Intergranular fracture surfaces with corrosion products can provide
some evidence of cause. For example, corrosion products are
Fig. 4.14 SEM image of the fracture surface of a nickel-base alloy
50
(Inconel 751, annealed and aged) after stress rupture (730 °C, or 1350 °F;
380 Mpa, or 55 ksi; 125 h). (a) Low-magnification view, with picture width
shown at approximately 0.35 mm (0.0138 in.) from original magnification of
250X. (b) High-magnification view, with picture width shown at approximately
0.1mm (0.004 in.) from original magnification of 1000X.
Fig. 4.15 SEM fractographs of the tensile test fracture surface of a high-
purity, coarse-grained Al-4.2Cu alloy with (a) intergranular facets at low
magnification (10X) and (b) uniform dimples on one facet at higher
magnification (67X). The microstructure indicated alloy depletion at the grain
boundaries.
Frequently observed on the separated grain facets of fractures of intergranular stress-
corrosion cracking (IGSCC) (Fig. 4.15). Intergranular fractures from stress rupture may
also have oxidation products. For example, Fig. 4.16 is an example of turbine blade
failure from a combustion gas environment. The original reference reported it as an
IGSCC failure, when in all likelihood the fracture is one of creep rupture with oxidation
products on the surface. Oxides are commonly observed on the fracture surface of
creep fractures.
This illustrates the importance of using all available information sources, such as stress
analysis and surface chemical analysis. Depending on the environment, cracks from
hydrogen embrittlement may also reveal corrosion products on a fracture surface.
Analysis of the cause of fracture in metal parts and components that have been
exposed to corrosive environments is often difficult because of interactions of fracture
mechanisms or because fractures generated by different mechanisms have similar
51
appearances.
4.1.7 Macroscopic Aspects of Overload Failures
In terms of macroscopic behaviour, overload cracking is either ductile or brittle, but the
entire fracture may occur from different combination:
1. Totally ductile
2. Totally brittle
Fig. 4.15 Intergranular fracture of 201-T6 cast aluminium after SCC
testing. (a) Optical micrograph, Keller‘s etch, approximately 75X. (b) SEM
image of fracture surface.
Fig. 4.16 Surface from fractured U-700 turbine blade. (a) Region with
transgranular and intergranular fracture feature. (b) Debris on intergranular
52
facets may be indicative of oxidation at high temperature after creep
cracking.
4. Initially ductile, then brittle
5. Mixed mode (ductile and brittle)
In the last two cases (4 and 5), the ductile appearance may not be directly visible at the
macro scale. Initially ductile fractures (case 4) are usually associated with rising-load
ductile tearing, or the initial ductility may be inferred by transverse strain at the crack
tip. The size of the plastic zone may be micro scale in this case. Mixed-mode ductile
and brittle cracking (case 5) would be inferred due to the presence of an intimate
mixture of cleavage and micro void coalescence at the micro scale (quasi-cleavage) or
by the presence of shear lips at the macro scale.
On a macroscopic scale, ductile and brittle fractures often are simply determined by
visual examination for evidence of whether the part is deformed excessively, not at all,
or somewhere in between. In the case of a tensile bar, for example, one might judge
the degree of brittleness of that material by its stress-strain curve (i.e., by the amount
of elongation or reduction in area before final fracture). More often, however, the macro
scale appearance is insufficient to convey the full story about a fracture. The
microscopic mechanisms and appearances of fracture also are needed for thorough
understanding and testing of hypotheses in a case study of a fracture. Uncovering both
the macro- and micro scale mechanisms allows the source and cause of a given failure
to be identified, and thus a course of corrective actions can be reached more reliably.
Surface roughness and optical reflectivity also provide qualitative clues to events
associated with crack propagation. For example, a dull/ matte surface indicates micro
scale ductile fracture, while a shiny, highly reflective surface indicates brittle cracking
by cleavage or intergranular fracture. In addition, when intergranular fracture occurs in
coarse-grain materials, individual equiaxed grains have a distinctive rock-candy
appearance that may be visible with a hand lens. In terms of documenting surface
conditions, one major problem with optical (light) macroscopic or microscopic
examination of fracture surfaces is its inability to obtain favourable focus over the entire
surface if the magnification exceeds 5 to 10X. Therefore, SEM also has become a
standard metallographic tool in failure analysis.
Surface roughness provides clues as to whether the material is high strength
(smoother) or low strength (rougher) and whether fracture occurred as a result of cyclic
53
loading. The surfaces from fatigue crack growth are typically smoother than monotonic
overload fracture areas. The monotonic overload fracture of a high strength quenched
and tempered steel is significantly smoother overall than is the overload fracture of a
pearlitic steel or annealed copper. Also, fracture surface roughness increases as a
crack propagates, so the roughest area on the fracture surface is usually the last to fail.
Fracture surface roughness and the likelihood of crack bifurcation also increase with
magnitude of the applied load and depend on the toughness of the material. Brittle
failures often contain multiple cracks and separated pieces, while ductile overload
failures often progress as single cracks, without many separated pieces or substantial
crack branching at the fracture location.
Macroscopic features typically help identify the fracture initiation site and crack
propagation direction. For example, crack branching and T junctions (Fig. 4.17) can
indicate the direction of crack propagation and location of crack initiation. Similar
techniques also apply when brittle materials fracture into multiple pieces (Fig. 4.18).
The orientation of the fracture surface, the location of crack initiation site(s), and the
crack propagation direction should correlate with the internal state of stress created by
the external loads and component geometry. In general,
Fig. 4.17 General features to locate origin from crack path branching (a)
and sequencing of cracking (b) by the T-junction procedure, where fracture A
54
precedes and arrests fracture B
Fracture initiates in the region where local stress (as determined by the external
loading conditions, part geometry, and/or macroscopic and microscopic regions of
stress concentration) exceeds the local strength of the material. This includes micro
scale discontinuities (such as an inclusion or forging seam) and macroscopic stress
concentrations (such as a geometric notch or other change in cross section).
The fracture surface orientation relative to the component geometry may also exclude
some loading conditions (axial, bending, torsion, monotonic versus cyclic) as causative
factors. For example, crack initiation is not expected along the centre line of a
component loaded in bending or torsion, even if a significant material imperfection is
present at that location, because
Fig. 4.18 Characteristics of crack direction and branching in fractures of
brittle materials from (a) impact, (b) bending, (c) torsion, and (d) internal
pressure
55
a shear stress at this location in bending, but in a homogeneous material, it is too small
to initiate fracture. That might not be the case for a laminated structure loaded in
bending.) Likewise, the profile of a fracture surface can indicate the direction of crack
growth. For example, the region of plane-strain fracture indicates the direction of
fracture in a shear overload fracture of annealed iron sheet (Fig. 4.19).
Under the right conditions, fracture surfaces may also have radial marks and chevrons,
which are macroscopic surface features that indicate the region of crack initiation and
propagation direction. They are common and dominant macroscopic features of the
fracture of wrought metallic materials, but are often absent or poorly defined in
castings. The ―V‖ of a chevron points back to the initiation site, and a sequence of ―V‖s
across the fracture surface indicates the crack propagation direction. The appearance
of chevrons or radial marks near the crack origin depends in part on whether the crack-
growth velocity at the surface is greater or less than that below the surface. If crack-
growth velocity is at a maximum at the surface, radial marks have a fan-shape
appearance (Fig. 4.20). If crack growth rate is greatest below the surface, the result is
chevron patterns (Fig. 4.21).
In rectangular sections, specimen dimensions can affect the appearance of radial
markings and chevron patterns. For example, the macro scale fracture appearances of
unnotched sections are shown schematically in Fig. 4.22 for sections with various
width-to-thickness (w/t) ratios. The w/t ratio influences the ability of the sample to
maintain a unidirectional state of stress during tension. In a thick section (top), strain in
the width direction is constrained and thus tends to a condition of plane-strain (mode I)
fracture. In this case, a large portion of the fracture surface
Fig. 4.19 Fracture surface showing a localized zone of plane-strain
fracture (left) from shear overload failure of annealed Armco iron sheet at
_196°C (_321°F). The configuration indicates that the fracture propagated
56
from left to right in this view. Light fractograph, 5X
is comprised of radial markings or chevron patterns indicative of rapid, unstable
cracking. At higher w/t ratios, the radial zone is suppressed in favour of a larger shear-
lip zone. In very thin sections (bottom), plane-stress conditions apply, and the fracture
surface is comprised almost entirely of a shear lip outside the fibrous zone of crack
initiation. Figure 4.23 is a schematic of radial marks and chevrons when fracture
initiates from surface notches.
If conditions are right, radial markings associated with rapid or unstable crack
propagation may also appear on the fracture surface of bar sections. The extent of
radial markings depends on ductility, as seen in Fig. 4.24 from tension testing of
unnotched bars after at different temperatures. As temperature decreases, ductility and
Fig. 4.20 Radial marks typical of crack propagation that is fastest at the surface
(if propagation is uninfluenced by part or specimen configuration)
Fig. 4.21 Chevron patterns typical when crack propagation is fastest below the
surface. It is also observed in fracture of parts having a thickness much smaller
than the length or width (see middle illustration in Fig. 4.22).
57
Fig. 4.22 Typical fracture appearances for unnotched prismatic tension test
sections
Fig. 4.23 Schematic of typical fracture appearances for edge- and side-notched
rectangular tension test sections
58
Fig. 4.24 Fracture surfaces of unnotched AISI 4340 steel specimens (heat
treated to a hardness of 35 HRC) after tension testing at three different
temperatures. (a) A shear lip surrounding a fibrous region is visible in the
specimen tested at 160°C (320°F). (b) At a lower test temperature (90°C, or
195°F), a radial fracture zone formed around the fibrous region, which formed
first; also, the shear lip here is smaller. (c) No fibrous region formed in the
specimen tested at _80°C (_110°F). Instead, fracture formed a radial zone that
extends nearly to the specimen surface and terminates in a narrow shear lip.
Original magnifications all at 15X
the extent of shear-lip formation are reduced. The fracture surface of radial marks is
visually distinct from the fibrous region near the centre on an unnotched bar (Fig.
4.24b). The radial markings (sometimes called a radial shear, star, or rosette) are
perpendicular to the crack front. With an unnotched tension-test bar, radial marks point
to crack initiation in the central region of the bar (Fig. 4.25a). This is not true with
notched bars (Fig. 4.25b), where crack initiation occurs near the root of the notch, and
where the radial zone points to the region of fast final fracture near the central region of
the notched bar. The extent of radial marking depends on the degree of ductility, as
shown in Fig. 4.26 from tension testing of notched 4340 bar at various temperatures.
As an illustration of environmental effects, tension-overload fractures of a notched
quenched and tempered 4340 steel at room temperature are presented in Fig. 4.27,
which shows fracture surfaces from sustained-load cracking after hydrogen charging.
Note that the progressive decrease in fracture stress from Fig. 4.27(a) to (d) is related
to a progressive increase in the size of the fibrous zone in the outer region by the
notch.
59
Fig. 4.25 General fracture-surface regions from ductile fracture of an
unnotched (a) and notched (b) tension test bar. (a) Radial zones on an
unnotched point to the region of crack initiation near the centre of the specimen.
(b) In the notched tensile specimen, the fibrous zone surrounds the radial zone,
because fracture initiates near the root of the notch (and completely around the
specimens in this idealized case without additional stress raisers). Fast final
fracture occurs in the centre.
60
Fig. 4.26 Overload fracture in notched AISI 4340 steel specimens (35 HRC)
from tension testing at three different temperatures. (a) The surface of the
specimen tested at _40°C (_40°F) shows only fibrous marks. (b) The specimen
tested at 80°C (_110°F) has a fibrous zone that surrounds a radial zone, which
is off-centre because of non-symmetrical crack propagation. (c) The specimen
tested at _155°C (_245°F) has a small annular zone of fibrous fracture, with
prominent radial marks in the central region of final fast fracture. All at ~17X
61
Fig. 4.27 Effect of sustained loading with hydrogen charging on the
fracture-surface characteristics of notched specimens of quenched and
tempered AISI 4340 steel tension tested at room temperature. (a) The
specimen with a relatively small fibrous zone at the right edge was broken by
tension overload with notched tensile strength of 2005 Mpa (291 ksi). The
specimen in (a) was not charged with hydrogen, while the three other
specimens were charged with hydrogen and then subjected to sustained
loading as follows: (b) Broke in 1.65 h under a stress of 1380 Mpa (200 ksi). (c)
Broke after 5.35 h under a stress of 1035 Mpa (150 ksi). (d) Broke after 5.5 h
under a stress of 690 Mpa (100 ksi). Notch radius was 0.025 mm (0.001 in.).
Microstructure of all four specimens was tempered martensite. All at ~8X
62
4.1.8 Cracks propagating from a pre-existing stress raiser or notch may
propagate totally in plane stress with net section yielding, totally in plane strain, or
there may be a fracture transition as the crack propagates. In some instances, buckling
may also occur. In terms of crack initiation, the likelihood of brittle crack initiation at the
free surface is not high, unless the material is inherently brittle. Initial cracking is
typically by a ductile mechanism (with three types of appearances: a tear zone, micro
void coalescence, or a tear zone followed by micro void coalescence), but further
cracking may change to cleavage or quasi-cleavage as the crack reaches some (small)
critical length, depending on the temperature, loading rate, and grain size. This is
common in bcc ferrous materials.
In steels, three types of crack initiation are found, depending, in part, on the
temperature.
In the first case, crack-tip blunting followed by ductile crack propagation by tearing
initiates at the free surface.
In the second, quasi-cleavage fracture and/or ductile micro void coalescence,
fracture initiates at the location of maximum constraint.
In the less common third case, the material has sufficient toughness and sufficient
crack blunting occurs that ductile fracture occurs on a shear plane at the crack tip.
If blunting occurs, the peak stress is reduced, and the stress falls off more
gradually behind the notch.
Figure 4.28 shows schematically the differences in cracking behaviour in ductile alloys
that exhibit a ductile-brittle transition with temperature. The presence of the stretch
zone can be used to quantitatively estimate the magnitude of the nominal stress and
the fracture toughness [3].
63
Fig. 4.28 Schematic of the brittle-to-ductile fracture transition. The relative area
on the fracture surface of the three micro scale fracture mechanisms (stretch
zone, dimple zone, and cleavage zone) are indicated.
64
CHAPTER 5
5.1 Fatigue testing – Part 1
Fatigue as a specific failure mechanism has been recognised since the early part of the
nineteenth century but it was the development of rail travel that resulted in a major
increase of interest in this type of fracture. The premature failure of wagon axles led to
Wohler in Germany investigating fatigue failure under rotating loading. This led to the
design of the first standardised test – a reversing stress rotating specimen, illustrated in
Fig. 5.1.
Fig.5.1. Wohler rotating fatigue test
There are many mechanisms that can lead to failure but fatigue is perhaps one of the
most devious since it can lead to a catastrophic failure with little or no warning – one
well known example being the failure of the Comet aircraft in the 1950s.
Failure can occur at a fluctuating load well below the yield point of the metal and below
the allowable static design stress. The number of cycles at which failure occurs may
vary from a couple of hundreds to millions. There will be little or no deformation at
failure and the fracture has a characteristic surface, as shown in Fig.5.2.
65
Fig.5.2. Typical fatigue crack fracture surface
The surface is smooth and shows concentric rings, known as beach marks, which
radiate from the origin; these beach marks becoming coarser as the crack propagation
rate increases. Viewing the surface on a scanning electron microscope at high
magnification shows each cycle of stress causes a single ripple. The component finally
fails by a ductile or brittle overload.
Fatigue cracks generally start at changes in section or notches where the stress is
raised locally and, as a general rule, the sharper the notch the shorter the fatigue life –
one reason why cracks are so damaging.
An unwelded ferritic steel component exhibits an endurance limit – a stress below
which fatigue cracking will not initiate and failure will therefore not occur. This is not the
case with most non-ferrous metals or with welded joints – these have no clearly
defined endurance limit.
The reason for this is that in arc welded joints there is an ‗intrusion‘ – a small defect at
the toe of the weld, perhaps only some 0.1mm deep. Provided that the applied stress is
sufficiently large a crack will begin to propagate within an extremely short period of
time. The endurance limit for a welded joint is therefore dependent on the intrusion size
66
that does not result in crack propagation at the applied stress range. In the case of a
welded joint, therefore, a fatigue limit – a ‗safe life‘ is specified, often the stress to
cause failure at 2x106 or 10 7 cycles.
During fatigue the stress may alternate about zero, may vary from zero to a maximum
or may vary about some value above – or below – zero.
To quantify the effect of these varying stresses fatigue testing is carried out by applying
a particular stress range and this is continued until the test piece fails. The number of
cycles to failure is recorded and the test then repeated at a variety of different stress
ranges.
5.1.1 S/N curve
This enables an S/N curve, a graph of the applied stress range, S, against N, the
number of cycles to failure, to be plotted as illustrated in Fig.5.3. This graph shows the
results of testing a plain specimen and a welded component. The endurance limit of
the plain specimen is shown as the horizontal line – if the stress is below this line the
test piece will last for an infinite number of cycles. The curve for the welded sample,
however, continues to trend down to a point where the stress range is insufficient to
cause a crack to propagate from the intrusion.
By testing a series of identical specimens it is possible to develop S/N curves. In
service however, there will be variations in stress range and frequency. The direction of
the load may vary; the environment and the shape of the component will all affect the
fatigue life, as explained later in this article. When designing a test to determine service
performance it is therefore necessary to simulate as closely as possible these
conditions if an accurate life is to be determined.
67
Fig.5.3. S/N curves for welded and unwelded specimens
5.1.2 Palmgren-Miners rule
In order to enable the fatigue life to be calculated when the stress range varies in this
random manner, the Palmgren-Miners cumulative damage rule is used. This rule states
that, if the life at a given stress is N and the number of cycles that the component has
experienced is a smaller number, n, and then the fatigue life that has been used up is
n/N. If the number of cycles at the various stress ranges are then added together – n 1
/N 1 + n 2 /N 2 + n 3 /N 3 + n 4 /N 4 etc. – the fatigue life is used up when the sum is of
all these ratios is 1. Although this does not give a precise estimate of fatigue life,
Miners rule was generally regarded as being safe [1]. This method, however, has now
been superseded with the far more accurate approach detailed in the British Standard
BS 7608.
The design of a welded joint has a dominant effect on fatigue life. It is therefore
necessary to ensure that a structure that will experience fatigue loading in the individual
joints has adequate strength. The commonest method for determining fatigue life is to
refer to S/N curves that have been produced for the relevant weld designs.
5.2 Fatigue testing – Part 2
Hydraulic digitally-controlled testing machines are used to subject test specimens
to repetitive loading. Although it is possible to simulate variable amplitude loading,
most tests are carried out using constant amplitude loading. The testing is carried
out under constant amplitude loading at various stress levels. The machine is
usually in the load control mode, and it is recommended that the specimen
stresses are monitored by strain gauges in addition to the information provided by
the machine‘s control panel. The fatigue test specimens should usually have
geometry and a loading mode that are representative for in-service condition.
Smaller specimens are tested in standard servo-hydraulic testing machines,
whereas full-scale testing of larger structural parts (e.g. tubular joints, large
girders) are tested in specially built frames with adaptable hydraulic cylinders. A
cruciform fillet-welded joint subjected to axial loading is shown in Figure 5.4. The
specimen has a width of 60 mm and a thickness of 25 mm. The specimen is
mounted directly between the two grips of an AMSLER hydraulic testing machine
with a performance of 250 kN dynamic loading. It is the upper piston that inflicts
the loading by vertical displacement. The main plate of the welded detail is in a
vertical position and the specimen is subjected to an axial force perpendicular to
the welding direction. The electrical cables seen on the photo pertain to
unidirectional strain gauges in the loading direction located 10 mm from the weld
toe.
Figure 5.5 shows a tubular joint with an X configuration in the same type of testing
machine. The diameter of the chord is 320 mm, whereas the diameter of the
branches is 250 mm. The wall thicknesses of the tubes are 16 and 12 mm
respectively. Due to the size of the joint, it is mounted in a frame on the table of the
testing machine. The upper piston of the machine introduces a vertical force on the
chord, with associated bending in the two branches through vertical movement of
the chord. The lower photo also shows all the strain gauges that are attached to
the tube surfaces to control the strain distribution in vicinity of the welded
intersection between the cord and the branches. This type of joint geometry gives
data for establishing the T-class S-N curve. The geometrical stress range is used
as the key parameter to determine the fatigue life for this case.
69
Figure 5.4. Cruciform fillet-welded joint in a hydraulic testing machine
Figure 5.5. Tubular joint with an X geometry subjected to in-plane bending of the braches. Upper part: setup in testing rig. Lower part: detail of branch-chord intersection at the crown point with strain gauges
70
5.2.1 Preparations and measurements
Fatigue testing includes all the phases such as test planning, the preparation of the
specimens and the statistical analyses of the results. It is important to keep track of
all the parameters that have an influence on the test results. This is necessary to
define a homogeny specimen population as a basis for the obtained design curve.
If a population involves test specimens with large differences in these important
parameters, the compiled test results will exhibit large scatter that can be difficult
to explain and cope with. The practical results will be that high-quality joints will be
penalized because they have been merged with joints of lower quality. For the
populations defining categories in current rules and regulations, this is often the
case and improvement in this area will lead to more accurate predictions due to
reduced scatter. Thus, high-quality joints will get the fatigue life predictions they
deserve.
The following background information should be gathered before the test:
- Steel type, chemical composition, mechanical characteristics,
- Welding procedure, method, electrodes, number and sequence of passes, heat
input,
- Manufacturing sequence,
- Global specimen geometry and local weld toe geometry,
- Axial or angular distortion,
- Non-destructive testing (NDT),
- Estimate of residual stresses in the specimens,
- Rolling direction of plates with regard to applied loading,
- Microscopy of the heat-affected zone (HAZ).
Most of this information is straightforward to obtain. The chemical composition and
mechanical characteristic for a C-Mn steel is given in Table 1. The data are typical
for medium-strength steels with nominal yield stress close to 350 Mpa. This is one
of the most widespread steel types used for welding. Both the fillet joint in Figure
5.4 and the tubular joint in Figure 5.5 are made of this steel. The welding has been
carried out by shielded metal arc welding (SMAW). Regarding the given
information for mechanical parameters, hardness measurements should be carried
out after the welding has been completed. This gives important information
regarding the base metal, weld deposit, and the HAZ. Figure 5.6 and Table 5.2
71
show typical results from hardness measurements carried out on the material in
Table 5.1 after welding had been carried out.
Table 5.1. Chemical composition in % and mechanical characteristics for a C-Mn
steel
Table 5.2. Hardness measurement
Figure 5.6. Results for hardness measurements for cruciform plated joint (C-Mn steel in Table 5.1)
72
The local toe geometry parameters are measured by applying replica material on
the weld toe and inserting cuts of replica cross sections into a profile projector with
a typical magnification of 10. The replica cuts must be made transverse to the weld
seam direction. Table 5.3 shows the results from two different test series with
specimens. Series 1 has low toe angle (mean value 30 degrees) and larger radius
(mean value 2.7 mm). Series 2 has a mean angle of 58 degrees and a toe radius
of 0.75 mm only. The fatigue testing actually showed that the mean fatigue life of
series 1 was twice as long as for series 2. If local toe geometry had not been
measured, this difference in fatigue life would have been difficult to understand and
explain.
Table 5.3. Statistical results from local toe geometry measurements of a fillet weld
Non-destructive testing should be carried out at the same level as is typical for
details entering into service conditions. Specimens with flaws that would have
been rejected for in-service purposes should, of course, also be excluded from the
test series. The most frequent method is magnetic particle inspection and no crack
detection is the only acceptable result. Estimating the actual residual stresses is
probably the most challenging part of the preparation work. In general, small
specimens will have lesser residual stresses than larger specimens. If the
specimens are properly stress relieved, the residual stresses may be set to zero. If
not, they can be measured by boring a hole and measuring the change in stresses
in vicinity of the hole by train gauges. This will reveal the residual stress gradient.
The problem is that the residual stresses vary over the volume of the test
specimen. Finally, metallographic study of the HAZ will reveal any abnormal
phases or grain size.
73
During the test the following should be measured and registered:
– loading mode, load range and nominal stress range,
– the applied R-ratio for the applied stresses,
– load frequency,
– stresses by strain gauges monitoring,
– crack growth measurements by an electrical method,
– beach marking of the fatigue surface by temporarily decreasing the stress range,
– number of cycles to final failure.
The loading mode is usually uni-axial loading or bending. One should be aware of
the fact that most S-N data are compiled without taking into account the difference
between these two modes; it is only the maximum stress range that is assumed to
matter. Most specimens are subjected to a positive load ratio to avoid problems
with buckling. Typical values of R are in the range of 0.1 to 0.3. The loading
frequency is assumed to play a minor role in air environment, and most tests are
accelerated so that the test program will not consume too much time. A typical
frequency of 6-10 Hz is often applied. However, special care should be taken for
the environmental condition (e.g. corrosive seawater.) In such cases there is an
interaction between the fatigue damage process related to the oscillating stresses
and the time-dependant electro-chemical reaction at the crack front. Therefore,
test frequency becomes important. The load frequency should usually not exceed
1 Hz. Detailed knowledge of the strain distributions can be obtained by the use of
electrical resistance strain gauges as was shown in Figure 5.5. These gauges
operate on the principle that the resistance of a wire changes when its length
changes due to stretching or compressing. Hence, when bonded to the surface of
the specimen, the local strain at the surface of the material is captured. Some test
specimens should be equipped with strain gauges to reveal the actual strain in
vicinity of the weld toe during the test. Based on these measurements, the strain
concentration can be determined and also the related stress concentration. The
stress should account for the gross stress concentration due to the global
geometry of the joint, but not the local effect caused by the weld profile. The gauge
distance from the weld toe should be large enough to avoid the local effect of the
toe itself, typically 10-15 mm. The strain measurements are normally compared to
stresses obtained by FEA analysis. FEA obtained stresses were shown in Figure
74
4.7. Local strain distribution for a tubular joint is shown in Figure 5.8. The
distribution is obtained with a rack of uni-axial gauges with 0.3 mm length and 2
mm spacing. As can be seen, the strains measured at the chord are higher than
the measurements at the branches.
Figure 5.7 Results from a local FEA model for a cruciform fillet-welded joint
Figure 5.8. Strain measurements at brace and chord at welded intersection
(Strain gauges length 0.3 mm spacing 2 mm, see bottom of Figure 5.5)
75
5.2.2 Test results
Traditional fatigue life testing produced one result only: the number of cycles to
failure. For small test specimens the failure is defined when a final fracture
separates the specimen into two parts. For larger joints the failure state is a
question of definition. It can be defined as a fracture separating the joint into two
pieces or by a through-thickness crack. For the cruciform joint in Figure 5.4, these
two time stages will coincide, whereas for the tubular joint in Figure 5.5, the joint
will still have integrity when a through-thickness crack has appeared. Typical life
data are shown in Figure 5.9 for small-scale testing with base material. The
material is a high strength weld able steel for mooring chains. For these small
specimens, the failure occurred when the crack was less than 1 mm. As can be
seen, the 18 tests are carried out at three stress levels, six tests at each level. The
tests were carried out in seawater without any protection against corrosion and the
load frequency was 1 Hz. Some results in air are added for comparison. The mean
and the design curves are drawn. When this steel is used in a larger structural
item, it must be born in mind that the life data in Figure 5.9 corresponds to early
cracking of the item and to not final failure. Therefore, if the curves in Figure 5.9
are used directly as life data for a structural item, it will be overly pessimistic.
Figure 5.9. Results from S-N testing with small specimens of high-strength steel in
seawater
76
5.3 Crack growth tests – guidelines for test setup and specimen monitoring
Crack growth testing is based on the hypothesis that it is the stress intensity factor
range (SIFR, ΔK) that governs the crack rate. The SIFR uniquely determines the
local severe stress field ahead of the sharp crack front under linear elastic
conditions. Crack growth tests are usually carried out on standard test specimens
for which the SIF (stress intensity factor) can be determined with great accuracy.
One typical test specimen is the compact tension (CT) specimen shown in Figure
5.10.
Figure 5.10. Compact tension specimen for crack growth measurements
The specimen is fabricated with specified dimensions (W and H) that are given
once the thickness T of the specimen is chosen. As can be seen, the specimen
has a sharp prefabricated notch from which the fatigue crack will grow. Definition
of the loading is shown in Figure 5.11. The load F varies with constant amplitude
between its maximum and minimum value. The corresponding SIF can be
calculated and the SIFR is defined as the difference between them. The evolution
of the SIFR with time is shown in Figure 5.12. As can be seen, there is a slight
increase in ΔK for constant ΔF. This increase is due to the increase in crack
length.
77
Figure 5.11. Definition of loading and SIFR
Figure 5.12. SIFR as a function of time
The crack depth is measured as a function of applied number of cycles, as shown
on the left-hand side of Figure 5.13. The measurements can be made optically or
by an ACPD technique. At chosen stages, the crack growth rate Δa/ΔN in m/cycle
is calculated and plotted against the actual range of the stress intensity factor ΔK.
This range has to be based on the mean crack size during the observation period
78
ΔN. The procedure is shown more detailed in Figure 5.13 where one point in the
log a/ΔN-log ΔK diagram is determined to the right on the figure. By following the
same procedure, at several stages along the crack path, the crack growth rate
parameters can be obtained as shown in Figure 5.14 for a log scale.
Figure 5.13. Transforming measurements crack growth rates as functions of ΔK
Figure 5.14. Crack growth rate as function of ΔK for a log-log scale
79
As can be seen from Figure 5.14, the results fall into three different regions
depending on the magnitude of the ΔK. At low ΔK there is a threshold value ΔK0
below which a crack will stop to propagate. Furthermore, there is an almost linear
relation between da/dN and ΔK for a log-log scale in an intermediate region. It is in
this region that the results obey the Paris law. The slope m and parameter C
defined by the intercept with the da/dN axis are obtained. Again, we see from the
figure that the linear relation for a log-log scale is an approximation; deviations do
occur. Hence, a linear regression analysis has to be carried out in this region in the
same way as for the S-N data. At higher values of ΔK, the maximum value of K is
close to the critical value of what the material at the crack front can sustain. The
consequence will be unstable fracture. As a result it is the SIFR that governs the
growth rate, components with various geometries, but made of same material, may
have different a-N curves, but when they are transformed into a da/dN-ΔK
diagram, as in Figure 5.14, the curves will coincide. This is true provided the tested
components have the same loading ratio R and the same environment. The great
benefit is then that the curve in Figure 5.14 can be applied to predict the crack
growth rate for various types of crack in any component if we are able to calculate
the SIFR for the crack and component geometry in question. Hence, the da/dN-ΔK
diagram and related equations are of great importance.
Figure 5.15 shows the test setup and the result for crack growth tests carried out in
seawater. The CT specimens with 25 mm thickness are taken from high-strength
steel in a mooring chain. Steel bars are forged and flash welded to obtain chain
links from this type of steel. The results to the right in the figure pertain to a test in
air, but other specimens were submerged in the small seawater basin as shown to
the left in the figure. The tests were carried out with an applied SIFR in the range
between 10 and 30 Mpam0.5. As can be seen, this gives growth rates from 10-5 to
10-4 mm/cycle. To reveal the threshold value, the testing has to be carried out at
smaller SIFR. Typical values for ΔK0 are found between 2 and 8 Mpam0.5,
dependent on the environment. The threshold value can be defined when the rate
da/dN is less than 10-7 mm/cycle.
80
Figure 5.15. Test setup and growth rate results with high-strength steel
Figure 4.16 shows the final fracture of the CT specimen. This will occur when the
either the net ligament ahead of the crack front is too small (global criterion) or if
the crack front has too high local stresses (local criterion). In the latter case, the
final fracture will be of the brittle type [4].
Figure 5.16. Final fracture of a CT specimen after fatigue crack growth
81
5.4 Welded Components
A welded joint exhibited no clearly established fatigue limit as in an unwelded
component. In service, few structures experience purely static loads and that most
will be subjected to some fluctuations in applied stresses and may therefore be
regarded as being fatigue loaded. Motorway gantries, for example, are buffeted by
the slipstream from large lorries and offshore oilrigs by wave action. Process
pressure vessels will experience pressure fluctuations and may also be thermally
cycled.
If these loads are not accounted for in the design, fatigue failure may occur in as
few as a couple of tens of cycles or several million and the result may be
catastrophic when it does.
Fatigue failures can occur in both welded and unwelded components, the failure
usually initiating at any changes in cross section – a machined groove, a ring
machined onto a bar or at a weld. The sharper the notch the greater will be its
effect on fatigue life.
The effect of a change in section is illustrated in Fig. 5.17, where it can be seen
that the stress is locally raised at the weld toe. The illustration shows a bead-on-
plate run but a full penetration weld will show the same behaviour.
Fig.5.17. Stress concentrating effect of a change in thickness
82
In addition, misalignment and/or distortion of the joint will cause the applied stress to be
further increased, perhaps by introducing bending in the component, further reducing the
expected fatigue life. A poorly shaped weld cap with a sharp transition between the weld
and the parent metal will also have an adverse effect on fatigue performance.
In addition to these geometrical features affecting fatigue life there is also the small
intrusion at the weld toe, as illustrated in Fig.5.18. In an unwelded component the bulk
of the fatigue life is spent in initiating the fatigue crack with a smaller proportion spent in
the crack propagating through the structure. In a welded component the bulk of the
fatigue life is spent in propagating a crack. The consequences of this difference in
behaviour are illustrated in Fig.5.19.
Fig.5.18. Weld toe intrusion
83
Fig.5.19. Effect of stress concentration on fatigue life
This shows that this small intrusion reduces the fatigue life of a fillet welded joint by a
factor of perhaps 10 compared with that of an unwelded item and some eight times that
of a sample with a machined hole. The other consequence is that fatigue cracks in
welded joints almost always initiate at the toe of a weld, either face or root.
It may be thought that the use of a higher strength material will be of benefit in
increasing fatigue life. The rate of crack propagation, however, is determined by Young‘s
Modulus – a measure of the elastic behaviour of the metal – and not simply by tensile
strength.
Alloying or heat treatment to increase the strength of a metal has very little effect on
Young‘s Modulus and therefore very little effect on crack propagation rates. Since the
bulk of a welded component‘s life is spent in propagating a crack, strength has little or
no influence on the fatigue life of a welded item. There is thus no benefit to be gained by
using high strength alloys if the design is fatigue limited. This is illustrated in Fig.5.20
which shows the benefits of increasing the ultimate tensile strength of steel if the
component is unwelded or only machined but how little effect this has on the life of a
welded item.
84
Fig.5.20. Effect of increase in tensile strength on fatigue life
One additional feature in welded joints that set them apart from unwelded or machined
items is the presence of residual tensile stress.
In a welded component there will be stresses introduced into the structure by, for
example, assembly stress. These stresses are long range reaction stresses and from a
fatigue point of view have little effect on fatigue life.
Of far greater significance with respect to fatigue are the short range stresses
introduced into the structure by the expansion and contraction of material close to and
within the welded joint. Whilst the actual level of residual stress will be affected by such
factors as tensile strength, joint type and size and by run size and sequence, the peak
residual stress may be regarded as being of yield point magnitude. The implications of
this are that it is the stress range that determines fatigue life and not the magnitude of
the nominal applied stress.
Even if the applied stress range is wholly compressive and there is apparently no
fluctuating tensile stress to cause a crack to form and grow, the effect of welding residual
stress is to make the structure susceptible to fatigue failure. This is illustrated in
Fig.5.21, where it can be seen that, irrespective of the applied stress, the effective
stress range is up to the level of residual stress at the welded joint.
85
Fig.5.21. Effect of residual stress on stress range
It would seem reasonable, therefore, that a post-weld stress relief treatment would be of
benefit to the fatigue life by reducing the residual stresses to low levels. This is only true,
however, where the applied stress range is partly or wholly compressive. If the applied
stress range is all tensile, research has shown that as-welded and stress relieved
components have almost identical fatigue performances with only a marginal
improvement in the stress relieved joints.
This is the result of the bulk of the fatigue life of a welded joint being spent in crack
propagation where propagation rates are only marginally affected by mean stress. It may
be difficult therefore to justify the cost of stress relief if the only criterion is that of
improving fatigue life [1].
86
5.5 Fatigue testing Part 3
Fig.5.22. Examples of joint classification from BS 7608
87
A welded joint behaves in a radically different way from an unwelded item, even if this
item contains a significant stress raiser. For both welded and bolted steel structures it
has been established that the fatigue life is normally governed by the fatigue behaviour
of the joints, including both main and secondary joints. The best fatigue behaviour will
be obtained by ensuring that the structure is so detailed and constructed that stress
concentrations are kept to a minimum and that where possible the elements are able to
deform in their intended ways without introducing secondary deformations and stresses
due to local restraints. Stresses may also be reduced by increasing the thickness of
parent metal which should improve fatigue life although with some types of joint fatigue
strength tends to decrease with increasing thickness. The best joint performance is
achieved by avoiding joint eccentricity and misalignment and welds near free edges and
by other controls over the quality of the joints. Performance is adversely affected by
concentrations of stress at holes, openings and re-entrant corners [1].
5.5.1 BS 7608:1993
This British Standard gives recommendations for methods for the fatigue design and
assessment of parts of steel structures which are subject to repeated fluctuations of
stress. Table 1 to Table 12 corresponds to the following basic types of details:
plain material;
lapped or spliced, riveted or bolted joints;
fasteners;
continuous longitudinal welded attachments;
other welded attachments;
transverse butt welds in plates;
transverse butt welds in sections and tubes;
load-carrying fillet and T-butt joints;
slotted connections and penetrations through stressed members;
details relating to tubular members;
seam welds in vessels;
branch connections to vessels.
Each classified detail is illustrated and given a type number. Table 1 to Table 12 also
give various associated criteria and the diagrams illustrate the geometrical features and
88
potential crack locations which determine the class of each detail. This information is
used to assist with initial selection of the appropriate type number. A detail should only
be designated a particular classification if it conforms in every respect to the tabulated
criteria appropriate to its type number. Class A is generally inappropriate for structural
work and the special inspection standards relevant to classes B and C cannot normally
be achieved in the vicinity of welds in structural work. In BS 7608 each joint type is
assigned a classification letter. For example, a plate butt weld with cap and root ground
flush is class ‗C‘, an undressed plate butt weld class ‗D‘ and a fillet weld class ‗F‘ (
Fig.5.22). In the case of members or elements connected at their ends by fillet welds or
partial penetration butt welds and flanges with shear connectors, the crack initiation may
occur either in the parent metal or in the weld throat; both possibilities should be
checked by taking into account the appropriate classification and stress range. For other
details, the classifications given in Table 1 to Table 12 cover crack initiation at any
possible location in the detail [5].
89
90
91
92
93
94
95
96
5.6 Potential modes of failure of welds
Some potential modes of failure are given below:-
97
98
99
100
101
102
Fig.5.23. Effects of joint classification on fatigue life
For each classification a fatigue curve has been developed and from these curves the
103
design life can be predicted as shown in fig 5.23. This is obviously an over-simplification
of what can be a very complicated task –the forces acting on a joint arising from
changes in temperature, changes in internal or external pressure, vibration, externally
applied fluctuating loads etc. can be complex and difficult to determine.
Whilst the joint design has a major effect on design life and is the basis for calculating
service performance, the weld quality also has a decisive effect – any fatigue analysis
assumes that the welds are of an acceptable quality and comply with the inspection
acceptance standards. However, in practice it is not always possible to guarantee a
‗perfect‘ weld and cracks, lack of fusion, slag entrapment and other planar defects may
be present, reducing the fatigue life, perhaps catastrophically.
Other less obvious features will also have an adverse effect. Excessive cap height or a
poorly shaped weld bead will raise the stress locally and reduce the design life;
misalignment may cause local bending with a similar effect. Good welding practices,
adherence to approved procedures and competent and experienced staff will all help in
mitigating these problems.
In some applications an as-welded joint will not have a sufficient design life and some
method of improving the fatigue performance needs to be found. There are a number of
options available. The first and perhaps simplest is to move the weld from the area of
highest stress range, the next is to thicken up the component or increase the weld size.
Note that, using a higher strength alloy will not improve the fatigue life.
Local spot heating to induce compressive stresses at the weld toes will also help,
although this needs very accurate positioning of the heated area and very careful control
of the temperature if an improvement is to be seen and the strength of the metal is not to
be affected. For these reasons, spot heating for fatigue improvements has been virtually
discontinued.
Hammer peening with a round nosed tool or needle gun peening gives very good results
although the noise produced may prevent their use. Shot peening can also be used to
introduce compressive stresses at weld toes with equally good results. Compressive
stresses can be induced in a component by overstressing – a pressure test of a pressure
vessel is a good example of this – where local plastic deformation at stress raisers
induces a compressive stress when the load is released. This technique needs to be
approached with some care as it may cause permanent deformation and/or any defects
104
to extend in an unstable manner resulting in failure.
Although the next techniques described are not as beneficial as hammer peening of the
weld toes they have the advantage of being more consistent and easier to control. The
techniques rely upon dressing the weld toes to improve the shape and remove the
intrusion. The dressing may be carried out using a TIG or plasma-TIG torch which melts
the region of the weld toe, providing a smooth blend between the weld face and the
parent metal.
Alternatively the toe may be dressed by the careful use of a disc grinder but for best
results the toe should be machined with a fine rotary burr as shown in Fig.5.24 and
5.25. Great care needs to be exercised to ensure that the operator does not remove too
much metal and reduce the component below its minimum design thickness and that the
machining marks are parallel to the axis of the main stress. Ideally the dressing should
remove no more than 0.5m depth of material, sufficient to give a smooth blend and
remove the toe intrusion [1]. The results of these improvement techniques are
summarised in Fig.5.26.
Fig.5.24. Grinding tools Fig.5.25. Burr machining of weld toes
105
Fig.5.26. Improvement in fillet weld fatigue life
5.7 Tubular joints
One group of joints that needs special attention is welded joints between hollow
sections such as cylindrical or rectangular pipes. One simple example is the T-joint
between two pipes as depictured in Figure 5.27. It consists of a chord member with
diameter D and wall thickness TC and a brace member with diameter d and thickness
TB. The principal loading is the axial force and the in-plane bending moment acting on
the branch. For these kinds of joints the stress concentration factor can be very high
near the intersections mainly due to secondary plate bending in the pipe walls near the
welds. Details are shown to the right in Figure 6 with the stress distribution through the
chord thickness and the maximum surface stress at various distances from the potential
crack site at the weld toe. The section shown is defined as the crown point of the joint.
There is a similar phenomenon on the branch side. The problem is that the stress
concentration can be very different for different loading modes and thickness ratios,
although these loading modes may give the same nominal stresses (elementary beam
stresses) away from the critical intersection. Hence, for these types of joints, the
nominal stresses are not an appropriate key to the fatigue life given by an S-N curve. It
is much more logical to use the stress concentration at the weld toe, ignoring the part of
the concentration created by the notch effect of the weld toe itself. The main argument
for this choice is that whereas the geometrical stresses may differ significantly from one
joint to another due to changes in loading mode and gross geometry, the local stress
106
concentration due to the weld notch will remain the same. Hence, the weld notch effect
is not explicitly taken account of, but will be inherent in the S-N curve. The so-called
geometrical stress concentration or hot-spot stress has also gained popularity for plated
joints, as a strategy to reduce scatter when presenting fatigue results [4].
Fig 5.27. Tubular T-joint with in-plane loading. Geometrical stress (fully-drawn line) and local stress (dotted line) at the weld toe on the chord crown point
5.8 Weldments
Many welding processes are available that can produce satisfactory joints in steels. The
selection of a particular process is based on many factors that include the thickness and
size of the parts to be joined, the position of the weldment, the desired properties and
appearance of the finished weldment, the particular application, and the cost of
fabrication as well as other factors. No single process can be used to produce
satisfactory weldments for all steels, thicknesses, and positions. The most suitable
process is the one that produces the desired properties in the final product at the lowest
possible cost.
In arc welding, which is the most widely used welding process for structural steels; filler
metal is melted and used to fill a weld groove. The arc-welding process, welding
procedure, and joint geometry influence weld penetration and admixture of the filler
metal with the base metal. Because of this admixture, the chemical composition of the
base metal can have significant influence on the microstructural and mechanical
107
properties of the weld metal. This influence is significant; especially for electro slag and
electro gas welds, because they are high heat-input single-weld-pass processes. The
final properties of the weld metal depend on many factors, especially the composition of
the weld metal and the conditions governing its solidification and subsequent cooling.
Because the heat flow in the weld metal is highly directional toward the adjacent cooler
metal, the weld metal develops distinctly columnar grains. Furthermore, the rapid
cooling of the weld metal may not allow sufficient time for diffusion of the chemical
constituents, resulting in microstructural heterogeneities. This segregation and the
directional solidification of the weld metal may result in weld-metal properties having
pronounced directionality.
For the arc welding processes, the maximum temperature of the weld metal is above
the melting temperature of the base metal joined. This temperature decreases as the
distance from the weld increases. Thus, partial melting of the base metal occurs at the
weld-metal-base-metal interface, and microstructural changes occur in the base metal in
the immediate vicinity of the weld, forming a heat-affected zone. The size of this zone is
determined by the rate of heating, the volume and temperature of the weld metal, and
the rate of cooling of the weld metal and surrounding base metal. These factors as well
as the composition and microstructure of the base metal determine the grain size, the
grain-size gradient, the microstructure, and therefore the fracture toughness of the heat
affected zone. Because of the high temperatures and the large variations in temperature
gradient and cooling rate, adjacent regions in the heat-affected zone can exhibit large
differences in microstructure and properties. In general, for carbon and low-alloy steels,
the closer the distance to the weld, the coarser the microstructure.
Coarse-grain regions adjacent to the weld interface generally exhibit the poorest
toughness. Illustrations of intermediate heat-affected-zone microstructure associated
with bead-on-plate deposits on hot-rolled plain-carbon steels and on a low-alloy
quenched-and-tempered steel are shown in Figures 5.28 and 5.29, respectively.
Both welds were made on plates of the same thickness, using identical heat-input
conditions; but, the nominal carbon contents were slightly different. They were 0.20% for
the carbon steel and 0.16% for the low-alloy quenched-and-tempered steel. The heat-
affected zone in a single-pass weld forms under the influence of a single thermal cycle.
The temperature and temperature distribution from the weld metal into the base metal in
a direction perpendicular to the weld is essentially identical at different locations along
the weld groove of a simple butt joint for two constant-thickness plates. Consequently,
the various microstructural regions in the heat-affected zone can be continuous.
108
However, the weld metal in a multipass weld is built up by the deposition of successive
weld beads. The structure and properties of deposited weld beads and existing heat-
affected zones are usually altered by the heating effects of subsequent weld-bead
deposits. The heat from subsequent weld passes may refine the grain size of the
deposited weld metal and existing heat-affected zone, may change the columnar
structure of the weld metal to an equiaxed structure, and may temper the microstructure
of the existing heat-affected zone. In multipass welds, unlike single-pass welds, the
heat-affected zone regions that exhibit relatively low toughness occur intermittently
adjacent to the weld interface. In either case, these lower-toughness regions are
surrounded by heat-affected zone regions of higher toughness[6].
Figure 5.28 Microstructures typical of the weld metal and the heat-affected base metal in a mild-steel weld.
109
Figure 5.29 Microstructures typical of the weld metal and the heat-affected base metal in a quenched-and-tempered low-alloy steel weld.
110
CHAPTER 6
Designing against Fatigue of Structures 8.8 Different types of structural fatigue problems The question about how to define problems of designing a structure against fatigue is
obviously associated with the goals to be achieved. In principle it implies that
satisfactory fatigue properties of a structure should be obtained, but it depends on the
type of structures which fatigue properties should be explored. For the present
discussion three categories are considered:
Structures for which fatigue failures are unacceptable.
Structures in which fatigue cracks may occur after a sufficient lifetime but without
the risk of a complete failure.
Structures for which crack initiation and crack growth until a complete failure are
acceptable, but for which a reasonable lifetime is still desirable.
Rotating blades of turboprop engines, wind turbines and compressors are examples in
the first category. Many components of various engines are also in this category with a
crankshaft as a well-known case. A fatigue failure in such components would be a kind
of a disaster. The fatigue limit of the structure is the important fatigue property and high-
cycle fatigue is an important issue. However, fatigue failures may also be unacceptable
in pressure vessels for which the number of pressurization cycles is not very large, e.g.
not exceeding 105. If all cycles have practically the same load range, the relevant fatigue
property is the crack initiation life under CA loading.
A variety of structures can also occur in the second category. Obviously the crack
initiation life is again of interest, and it should be large enough for a satisfactory lifetime
in service. If a complete failure is unacceptable, a reliable inspection procedure is
indispensable. This applies to aircraft structures, and it can also be applicable to several
welded structures. As a consequence both the crack initiation life and the crack growth
life are of interest. Moreover, fatigue under VA amplitude loading may also be a relevant
condition.
The third category includes various utilities for which final failure simply implies that it
must be replaced by a new one. Various housekeeping articles are in this category, e.g.
washing-machines, vacuum cleaners, but not stairs. Bicycles are another typical
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example in which fatigue failures do occur. The fatigue property is the fatigue life until
failure with lifetime as an economical criterion. Data on crack initiation life and crack
growth properties are not required, but again both CA and VA load histories can be
significant. The three categories of structure have been defined because within each
category similar fatigue properties should be predicted. The literature on fatigue
prediction problems is quite diverse. The world of building steel bridges and the world of
manufacturing wind turbines are two different cultures, but still with similar fatigue
problems. As an example, in both worlds load spectra are consisting of a combination of
deterministic loads and random loads.
In practice the designer who is faced with fatigue endurance problems, must also
consider other durability issues, such as: maintenance, inspections, repairs,
replacements, service conditions with implications for corrosion, wear and tear. They
are all a matter of concern dealing with the structure as an object that should be in
function for a long time. Anyway, the possibility of fatigue crack initiation is a relevant
problem because it can have a large economic impact. Designing against fatigue crack
initiation is one of the responsibilities of the designer of the structure. Following
information is required:
information about the structure,
analysis and fatigue data, and
last but not least, the load spectrum.
In the literature it is sometimes suggested that our fatigue problems are solved if an
accurate prediction model would be available. This is misleading. The present physical
understanding about fatigue damage accumulation is reasonably well developed in a
qualitative sense. And just because of this understanding, it must be accepted that
accurate quantitative predictions on fatigue lives are illusory.
Problems of fatigue life and crack growth prediction for notched elements indicate how
estimates of the fatigue limit of notched elements could be obtained. Unfortunately,
similar prediction procedures are not applicable to fatigue of joints. Empirical data of
joints must be available to arrive at estimated of the fatigue limit. Predictions on the
fatigue life under VA loading are even more complicated. The Miner rule is physically
rather primitive. The rule starts from the idea that damage can be characterized by a
single damage parameter which essentially disagrees with the present knowledge about
fatigue damage accumulation. At best, the Miner rule gives some weighted indication of
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the load spectrum, but not of the severity of the load spectrum. The Miner rule fully
breaks down in comparisons of load spectra severities. When the Miner rule is used to
obtain some rough indication of the fatigue life under VA loading, one should realize that
the prediction is an extrapolation of S-N data which by itself have already a limited
reliability. The situation appears to be more convenient for predictions of fatigue crack
growth. Crack growth prediction for CA loading based on the well-known fracture
mechanics methodology can be reasonably reliable. But the situation is less satisfactory
for fatigue crack growth under VA loading. A major problem is to account for interaction
effects of cycles with different amplitudes. If the interaction effects are ignored,
predictions will probably be conservative, but it can lead to significant under-predictions
for crack growth under steep load spectra.
In general terms, it must be accepted that fatigue predictions are speculative in a way
that the order of magnitude may be instructive, but the predictions should be evaluated
with appreciable judgement. In cases of doubt, the design variables should be
reconsidered to see where weak links are present. Estimates of fatigue properties can
be improved by experiments. Whether this is really necessary depends on safety
margins and costs involved. Detailed stress analysis, fatigue experiments and load
spectrum measurements can improve the significance of predictions. It is possible that a
simple fatigue analysis shows that the occurrence of a fatigue failure problem is very
unlikely, and no further design improvements are necessary. It is also possible that
fatigue failures in service are acceptable because a simple replacement of the failed
element is not expensive and safety is not involved. In such cases, a cost-benefit
analysis can show that efforts to improve the fatigue prediction are not really worthwhile.
But it is also possible that a simple prediction indicates that structural improvements
must be considered, i.e. designing against fatigue. It then is useful to have some idea
about the accuracy of preliminary fatigue life predictions. Several sources of
uncertainties in the prediction technology should be considered, including the strategy of
applying safety factors.
8.8 Designing against fatigue
A designer should know whether he is designing against crack initiation, or for an
acceptable crack growth behaviour, or for both. Moreover, he also should be aware of
the question whether his problem is associated with high-cycle fatigue or low-cycle
fatigue.
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The initiation period is basically a material surface phenomenon, whereas crack growth
is a matter of crack growth resistance of the material as a bulk property. As a
consequence, fatigue related influences are essentially different for the two periods. The
crack initiation period and the fatigue limit are heavily depending on material surface
conditions, whereas most of these conditions are practically irrelevant for the crack
growth period. Understanding of the effects of these variables is essential for designing
against fatigue.
Fig. 6.1 Survey of topics associated with designing against fatigue. 8.8 The crack initiation aspect
It is easily understood from Figure 5.1 that designing against fatigue crack initiation is
concerned with the general layout of a structure, detail design, material selection and
surface treatments. The layout of a structure depends on the purpose of the structure.
But there are various possibilities to obtain an improved load distribution in a structure,
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e.g. by changing local dimensions such as a locally increased thickness to reduce the
stress level around a fatigue critical detail. Another example is associated with
eccentricities which are causing unfavourable secondary bending.
8.8 Material selection
The selection of the material depends on many circumstances, such as static
properties, workshop properties, corrosion resistance, thermal properties, costs, etc. A
material with a higher S0.2 may have a higher Sf for unnotched specimens, but also an
increased notch sensitivity. Similarly, welded joints of a higher strength material usually
do not necessarily have an improved fatigue strength. If a new material is considered for
a structural application, it should be supported by results of service-simulation fatigue
tests on specimens which are representative for fatigue critical details of the structure
under consideration.
8.8 Surface treatments
The designer can specify the quality of the material surface, and also certain surface
treatments. Some typical examples of surface treatments are: fine machining, nitriding
of steel, shot peening, surface rolling, prevention of fretting and corrosion protection. It
has been pointed out that surface treatments are carried out for various purposes:
improvement of fatigue properties, protection against corrosion, improved wear
resistance, restoring poor surface quality, and cosmetic reasons. Surface treatments
can increase the hardness of a surface layer, and thus hamper cyclic micro plasticity. At
the same time, residual compressive stresses restrain the opening of micro cracks in
the surface layer and thus will reduce or even arrest the growth of these cracks. As a
result, the major benefit of surface treatments is on the crack initiation period. They are
important for high-cycle fatigue, and in particular for the fatigue limit.
8.8 Detail design for an improved stress distribution
An essentially different approach is associated with the reduction of stress
concentrations. For fatigue critical notches it generally boils down to increasing root radii
or applying stress relieving grooves if that is possible. Non-circular fillets are rarely
considered. Recently Kt –values were calculated for elliptical fillets and results have
shown significant reductions of the Kt –value. It may be repeated here that various Kt –
values are not always very accurate as a result of older techniques used to determine
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the various graphs in the book. By now FE analysis can produce more accurate Kt –
values as well as stress gradients.
8.8 Large-scale design issues
Detail design is associated with dimensions which are significant for local stress
concentration, e.g. a hole diameter or notch root radii. On a larger scale, the designer is
considering the general concept of the structure. As an example, although perhaps a
somewhat curious one, rather different concepts can be contemplated for designing a
bridge. Another noteworthy example, in various structures joints are present, but the
variety of different joint concepts is also large. Decisions to be made on the type of
structure are generally depending on experience of the industry, and in the industry on
economic implications.
8.8 Uncertainties, scatter and safety margins
The purpose of designing against fatigue is to prevent disasters, and also to avoid non-
fatal incidents in view of unwanted economic consequences. Unfortunately,
uncertainties about the fatigue performance of a structure cannot be solved by accurate
and rational arguments. It implies that some philosophy about safety factors or other
measures should be considered. A solid rational frame work to arrive at safety factors
cannot be formulated. Statistical distribution function are unknown. Information about
scatter of fatigue properties is largely coming from laboratory test series and not from
service experience. The choice of reasonable safety factors is a matter of experience
and engineering judgement. Both economic and safety consequences of the occurrence
of a premature fatigue failure must be considered. In view of limited accuracies of
quantitative fatigue predictions, it must be asked how this situation should be carried on.
The variety of sources for uncertainties is fairly large.
8.8.1 Uncertainties
Three reasons for uncertainties about the prediction of the fatigue performance of a
structure are:
Uncertainties about the load spectrum.
Uncertainties about the fatigue properties of the structure.
Uncertainties about the reliability of predictions.
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Not all these uncertainties can be associated with scatter of some properties. Variations
of the load spectrum include differences between deterministic loads (applied by the
operator) and stochastic loads (random type loads depending on environmental
conditions). With respect to the deterministic loads, all structures of the same type are
not used in exactly the same way. The designer must consider the variability of loads
which should be taken into account (functional loads, manoeuvres). But another part of
the variability of the load spectrum does not depend so much on the operator.
Stochastic loads are relevant to structures operating under a variety of weather
conditions (aircraft, boats, drilling platforms) or moving over various roads (passenger
cars, trucks, coaches). Statistical distribution functions and power density spectra can
be involved for air turbulence, sea-waves and road roughness. Several types of
structures will see a combination of deterministic and stochastic loads. Cranes, bridges
and buildings offer interesting combinations of both types of loads. The second reason
for uncertainties is associated with fatigue properties of the structure. These
uncertainties are of an entirely different nature. Statistical variations are related with
material properties and production quality. The fatigue properties of a material with a
standardized composition may be obtained from data banks, but it cannot be
guaranteed that these properties are always the same. Scatter may occur as a result of
batch to batch differences, but even in a single plate statistical variations are possible.
Moreover, the crack initiation fatigue life is depending on the quality of the production of
components. There are sufficient reasons why components produced during a number
of years cannot be considered to be samples of the same statistical population.
Finally, the third source of uncertainties is associated with the reliability and accuracy of
a prediction model. Estimates can be obtained for S-N curves, fatigue limits and crack
growth.
8.8 Scatter and safety factors
8.8.1 The fatigue limit and the safety factor
An important category of problems of designing against fatigue is associated with high-
cycle fatigue and a flat load spectrum. If fatigue failures are unacceptable, the criterion
is that all load cycles should be below the fatigue limit of the structure. The variables
involved are the maximum load cycles occurring in service load spectra and the fatigue
limit of the structure. They are both affected by uncertainties. It can be tried to obtain an
estimated value of the fatigue limit of the structure, but even if the analysis would be
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supported by experiments, some unknown scatter must be expected. With respect to
the load spectrum uncertainties are involved, not so much as a consequence of scatter,
but due to different utilizations. Under these twofold conditions, a safety factor cannot be
defined with rational arguments. If a fatigue failure of the structure would cause a fatal
accident, relevant experimental efforts should be considered.
Fig. 6.2 Safety margin on load level S1 for required life time N1. A similar margin for the fatigue limit is unrealistic.
A full-scale fatigue test on a representative part of the structure with the step by step
increasing load (see Figure 6.3) can give useful information about the fatigue limit.
Fig. 6.3 Load history in a step test to obtain an approximate fatigue limit with
a single specimen. A small Δsa and a large ΔN-value should be adopted.
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A full-scale CA load tests at the estimated load level of the load spectrum cannot be
recommended. The test should be continued to a very high number of cycles, say > 108.
However, if the load level is just below the unknown fatigue limit, see Figure 6.2, then
failure will not occur. In view of the scatter band of the fatigue limit, an other similar
structure can fail at a fatigue life between 106 and 107 cycles, just above the average
fatigue limit. It implies that information about the safety level remains unknown. In the
high-cycle fatigue regime and for the fatigue limit, scatter of fatigue lives is not the
relevant issue. Scatter of the fatigue strength, and in this case of the fatigue limit, is
crucial. For this reason the step by step increasing test of Figure 6.3 should be
preferred. Of course the number of cycles in each step (ΔN) should be large enough in
order to be in the high-cycle fatigue regime, for instance ΔN = 106 or 2 × 106 cycles. The
fatigue limit Sf obtained with the step-by-step method and also the load spectrum in
service are not free from uncertainties. A safety factor should be adopted. Since
quantitative indications on scatter are lacking, an intelligent guess must be made.
Possible consequences of fatigue failures in service have to be considered. It is
believed that a safety factor of 1.5 can be sufficient in many cases. However, if more
confidence is desirable, more fatigue tests should be carried out. Another approach is to
carry out load history measurements in service to have more information about the load
spectrum.
8.8 Safety factors for finite fatigue life problems under CA loading
Crack initiation cannot be avoided if stress amplitudes above the fatigue limit occur in
the service load spectrum. As a consequence, fatigue crack initiation is possible and a
finite life should be considered. A typical example is represented by a pressure vessel.
A safe approximation of the load spectrum is that the pressure vessel is always loaded
to the same maximum operational pressure. Load spectra of other structures with a flat
load spectrum can be approximated in the same way. A safety factor can now be
defined in two different ways. The factor can be applied to the fatigue life or to the
fatigue strength. If a finite life is envisaged, the natural approach is to think in terms of
endurances which guarantee a sufficient lifetime. If N1 is the required lifetime and N2
the estimated fatigue life, see Figure 6.2, then the safety factor is f N = N2/N1. However,
in terms of the fatigue strength, if S1 is the required fatigue strength and S2 is the
estimated fatigue strength, then the safety factor f S = S2/S1. Adopting the Basquin
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relation (S k ·N = constant), the relation between the two safety factors is f N = (f S) k. If
loads exceeding S1 should not be expected or even be impossible, and then the safety
factor for the fatigue life should be considered. However, if required lifetimes larger that
N1 are of little interest then the safety factor for the stress level is more appropriate. The
size of these safety factors to be adopted depends on the consequences of a fatigue
failure. Obviously larger factors are necessary if fatal accidents are possible; say 1.5 on
the stress level or 6.0 on lifetimes. In such a case, a realistic experimental verification
test must be advised. If the consequences of a final failure are not serious, a smaller
safety factor can be adopted, say 1.2 on the stress level, or 2.5 on the fatigue life. If the
quality of the stress raisers is poor (e.g. in low-quality welds), larger values may be
worthwhile. Engineering judgement and experience from previous structures should be
practiced.
8.8 Safety factors for finite fatigue life problems under VA loading
The VA load case offers an additional uncertainty if compared to the CA load case.
Predictions for a VA load history are affected by the unreliability of the Miner rule. It is
difficult to understand how this might be accounted for by a safety factor. When using
the Miner rule, it appears to be wise to extrapolate S-N data below the fatigue limit. In
cases of doubt, some exploratory service-simulation fatigue tests are much
recommended.
8.8 Safety factors and fatigue crack growth
Safety factors on fatigue crack growth have to be considered if the crack growth period
covers an essential part of the lifetime in service. This can occur when cracks are
initiated at material defects, corrosion pits, or sharp corners with a high stress
concentration. It can also start from unintentional surface damage caused in-service
(nicks, dents, scratches, impact damage, etc.). In welded structures, crack initiation is
possible from weld defects, but also at the edge of the weld toe due to a locally
unfavourable profile. All these situations are undesirable, but they cannot always be
avoided. In view of safety, it may be necessary to consider fatigue lives with a practically
zero crack initiation period. It is kind of a worst case analysis which should be made if
complete failure is unacceptable.
Two different cases can be defined:
Crack growth is accepted, but the occurrence of a complete failure must be
prevented by periodic inspections.
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The crack growth period until failure should be larger than the design lifetime of the
structure because inspections for cracks in service are undesirable or not feasible.
The first case is well-known for aircraft structures for which so-called damage tolerance
requirements are laid down in official airworthiness regulations. It can also be applicable
to nuclear pressure vessels or other structures if fatigue failures are inadmissible and
periodic inspections must be done to detect fatigue cracks before failure occurs. The
problem setting is illustrated in Figure 6.4 by a schematic crack growth curve and a
corresponding curve of the decreasing static strength of the structure caused by the
growing fatigue crack. Failure of the structure is supposed to occur at a critical crack
length, ac. Cracks can be detected at the crack length denoted
Fig. 6.4 Principle of safe crack growth by period inspections. As ad . The period for crack detection covers crack growth from ad to ac, see Figure
6.4. The number of uncertainties is fairly large:
the initial crack length a0,
the final crack length ac,
the crack growth data of the material,
the load spectrum,
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the crack growth prediction model,
the probability of detecting a fatigue crack.
The probability of crack detection depends on the non-destructive techniques adopted.
Questions can be raised whether a surface crack with a length of a few illimetres can
be detected. In general, very small cracks, say 1 mm (0.04 inches) cannot be detected
reliably. Crack detection of invisible cracks, e.g. in joints, must be done with special
inspection techniques.
Secondly, it must also be established how far the crack may grow before the risk of a
large failure is present. The crack must be found within the crack growth range between
the detectable crack size (ad) and the critical crack size (ac). A safety factor should then
be applied to this period to assess the inspection period. In the past, a factor 3 has been
used for transport aircraft, but more recently, the tendency is to use a factor 2.
Obviously, the choice of the safety factor is a matter of judgement, which requires that
all sources of uncertainties are recognized and understood. It should also include the
human factor of the inspection procedure. If a large number of structures must be
inspected, most of which will be free from cracks, an occasionally occurring small crack
might escape detection. Situations of finding cracks in order to prevent dangerous
situations are not confined to aircraft. It also applies to other types of structures if a
fatigue failure cannot be accepted, e.g. for pressure vessels. Operators of large
structures try to combine inspections with periodic maintenance for economic reasons.
Actually, operators prefer structures which do not require inspections.
The size of the initial crack length (a0) must be associated with the size of some initial
defect. This is a difficult issue because the crack growth rate of initially small cracks is
very low. As a consequence, the predicted crack growth life will significantly increase for
a smaller value of a0 (see Table 6.1).
Table 6.1 Illustrative crack growth life predictions for a carbon steel.
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It is more conservative to select a larger a0-value, but which size? The final crack length,
ac, is reached at the moment of failure. It requires that the reduction of the residual
strength of the structure is calculated as a function of the increasing crack length, which
is not a simple calculation because macro plasticity will occur. However, the crack
growth rate in the last part of the crack growth period is relatively high, and assuming a
lower ac will have a small effect on the crack growth period, see again Table 6.1. The
crack growth prediction model is less problematic for a CA load spectrum than for a VA
load spectrum. In case of CA loading, predictions may give reasonably reliable results
provided that K solutions are available. Quite often, K solutions are not available, even
for structural elements with a simple geometry. Small cracks are usually part through
cracks at the material surface. If K-values are not available, they can be calculated with
FE techniques, but it requires expertise on this topic. Predictions on crack growth during
VA loading offer problems due to interaction. Ignoring these effects should be expected
to give a conservative prediction for most load spectra. The basic CA crack growth data
used in the prediction are also subjected to uncertainties. Variations can occur between
nominally similar materials from different producers. Even differences between batches
from the same producer have been found, see Figure 6.5.
Fig. 6.5 Comparison between crack growth lives of sheet specimen of different
producers and different batches
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Small cycles with ΔK < ΔKth can still contribute to crack growth. It was proposed to
extrapolate the da/dN−ΔK function in the Paris regime to low ΔK < ΔKth.
8.8 Safety aspects associated with a corrosive environment and low frequency
fatigue
The effect of corrosion on fatigue depends on the material/environment system.
Unfortunately, most types of steel and aluminium alloys are sensitive to corrosion. It can
imply that these materials are also sensitive to the frequency and wave shape of the
load cycles. Unfortunately, the effect of corrosion fatigue cannot simply be described by
a quantitative model. Experience should indicate how to deal with safety issues
introduced by a corrosive environment. Corrosion can affect both crack initiation and
crack growth. Obviously, the application of safety factors does not preclude the
occurrence of corrosion. Pitting and other local corrosion phenomena can occur in a
corrosive environment, and subsequent crack growth will be activated. It might be
hoped that cracks should not grow at low stress amplitudes, but it would require a high
safety factor (see Figure 5.6 for mild steel).
Fig. 6.6 The effect of environment and load frequency on the S-N curve of
unnotched mild steel specimens
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The best solution is to prevent corrosion at the material surface. Sometimes this is done
by preventing the access of the aggressive environment to a fatigue critical element of a
structure. Corrosion resistant surface layers can be considered also, but experience
should indicate whether this will be successful. Another solution is shot peening of the
material surface. This would not prevent corrosion at the material surface, but the
residual compressive stresses may prevent crack opening and further crack growth. An
example of this application is shot peening of springs used in cars. If water is trapped in
the structure, the consequences of a stagnant water environment may be disastrous.
Trapping of water should be avoided, either by design or sealing of critical locations.
Corrosion fatigue can be problematic for structures used in the open air or in the sea,
e.g. for bridges, cranes, ships, offshore structures, but also for many other structures. In
the open air, rain and fog are causing a moist environment of usually polluted water,
which is an aggressive environment. After fatigue cracks have been initiated, the
corrosive environment can enhance crack growth. On welded joints, accelerated crack
growth has been observed in comparative tests carried out in air and salt water. In salt
water, crack growth could be about three times faster. A safety factor of three applied on
the crack growth life may be reasonable. If fatigue failures in the environment of the
structure would have serious consequences, it might be necessary to support the
fatigue analysis by relevant experimental work. The problem is how a service-simulation
fatigue test should then be carried out in view of corrosion being a time dependent
phenomenon. The frequency of the cyclic loads in service may be low and an exact
simulation can imply an unacceptably long duration of the test. A compromise should be
considered. Certain parts of the load-time history can be simulated faster than the
history in service, while the more damaging load cycles can be applied with the loading
rates relevant for the service load-history. It then should be recognized that the
increasing load part of a cycle is the most important part for fatigue crack increments.
Another interesting alternative to service-simulation tests is to build a few prototypes of
the structure and to test these prototypes in a realistic but severe application in service.
This has been done for cars and trucks, which were tested by severe driving along
selected tracks with rough road conditions. Actually, such tests are not done for fatigue
only. It should show a satisfactory functioning of all parts of a structure under severe
conditions. However, it also can reveal insufficient fatigue properties[1].
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CHAPTER 6
Methods of revealing fatigue cracks
126
7.1 Dye-penetrant testing
Dye-penetrant testing; A method of
examining components to detect surface-
breaking flaws, such as cracks. The technique
is based on the ability of a liquid to be drawn
into a "clean" surface-breaking flaw by
capillary action.
It is essential that the component to be
inspected is thoroughly cleaned to remove all
traces of dirt and grease.
It is then sprayed with a penetrating liquid,
usually a brightly coloured liquid or a
fluorescent dye, which penetrates any surface-
breaking cracks or cavities. The liquid is
allowed to soak into the components surface.
(Fig1 right)
After soaking, the excess liquid penetrant is
wiped from the surface and a developer
applied. The developer is usually a dry white
powder, which draws the penetrant out of any
cracks by reverse capillary action to produce
indications on the surface.(Fig2 right)
These (coloured) indications are broader than
the actual flaw and are therefore more easily
visible.
Fluorescent penetrants are normally used with
a UV lamp to enhance sensitivity.
These systems are often used to check weld
quality during fabrication.
7.1.1 An example of dye penetrant testing used on bicycle components.
This cycle crank arm was returned to the
supplier after a very short time in use. The
owner had seen a crack coming from the square
taper axle attachment and suspected a smaller
crack close to the pedal thread
The suspect areas were sprayed with red
penetrant dye and left to soak. The square hole
location was clearly cracked but the minor region
may be a surface scratch.
Such distinctions are very important in the
performance of engineering components.
After the dye was cleaned off the
component was sprayed with chalk
developer. The crack running from the
square axle drive hole gave a very distinct
red indication at its precise location,
indicating it was clearly cracked. The other
feature showed no red line on development
indicating a surface scratch not an
embyronic crack.
The crank spider arm of this
chainset fractured and
unseated the cyclist in heavy
city traffic.
It was old but well looked after
and cleaned regularly The
growing fatigue crack was
undetected until the dangerous
failure.
The cyclist was concerned if a
similar crack had been
nucleated in the matching plain
crank arm shown alongside.
The equivalent area was
sprayed with red penetrating
dye which was then left to soak
into any cracks or fissures in
the component.
After several minutes of
soaking the dye was cleaned
off the components surface.
The pre-soaked and cleaned
area was then sprayed with
developer spray which is
basically chalk powder in a
volatile carrier. Any defects
present show up as the red dye
is pulled out of any cracks or
fissures in the component.
It was concluded that none
were present. The only red
marks were from dye that had
been retained in the stamped
product identification marks.
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7.2 PHOTOELASTICITY
A method of examining transparent polymer models of structures etc. to isolate stress
concentrations and other weak zones. The model is placed between crossed circular polarizing
filters (eg Polaroid sheets) and a force applied. The technique also enables residual stress to
be shown in transparent articles. Stress fields (applied and residual) can be exposed using
models of structures in photosensitive material placed between polarising filters in the crossed
polar position. Here the stresses in a 7 member model bridge truss, centrally loaded and simply
supported are shown. These injection moulded safety spectacles contain residual moulding
stresses shown here using photo elastic viewing techniques [2].
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CHAPTER 8
8.1 CAUSES AND RECOGNITION OF FATIGUE FAILURES
8.1.1 General Causes of Material Failures:
Design deficiencies
Manufacturing deficiencies
Improper and insufficient maintenance
Operational overstressing
Environmental factors (i.e. heat, corrosion, etc.)
Secondary stresses not considered in the normal operating conditions
Fatigue failures
Improper and insufficient maintenance seems to be one of the most contributing factors
influenced by some improper designs such as areas that are hard to inspect and
maintain and the need for better maintenance procedures. In many circumstances the
true load is difficult to predict resulting in a structure being stressed beyond its normal
capabilities and structural limitations. When a structure is subject to cyclic loads, areas
subject to fatigue failure must be accurately identified. This is often very hard to analyse,
especially in a highly composite structure for which analysis has a high degree of
uncertainty. Thus, in general, experimental structural fatigue testing is frequently
resorted to.
8.1.2 Recognition of Fatigue Failure
Two fatigue zones are evident when investigating a fracture surface due to fatigue, the
fatigue zone and the rupture zone. The fatigue zone is the area of the crack
propagation. The area of final failure is called the rupture or instantaneous zone. In
investigation of a failed specimen, the rupture zone yields the ductility of the material,
the type of loading, and the direction of loading. The relative size of the rupture zone
compared with the fatigue zone relates the degree of overstress applied to the
structure. The amount of overstressing can be determined from the fatigue zone as
follows: highly overstressed if the area of the fatigue zone is very small compared with
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the area of the rupture zone; medium overstress if the size or area of both zones are
nearly equal; low overstress if the area of rupture zone is very small. Figure 8.1
Figure 8.1
Figure 8.2 Typical fatigue zone with identifying marks.
Describes these relations between the fatigue and rupture zones.
The fatigue zone can be described as follows: a smooth rubbed, and velvety
appearance, the presence of waves known as ―clam-shells‖ or ―oyster-shells‖, ―stop
marks‖ and ―beach marks,‖ and the herringbone pattern or granular trace which shows
the origin of the crack.
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In general, stop marks indicate the variations in the rate of crack propagation due to
variations in stress amplitude in a cyclic application varying with time. Figure 7.2 is a
schematic representation of the fatigue zone.
8.2 Design Considerations
Even if careful attention to good design practices is constantly the goal of design
engineers, fatigue problems are sometimes introduced into the structure. Fatigue failures
are often the result of geometrical or strain discontinuities, poor workmanship or
improper manufacture techniques, material defects, and the introduction of residual
stresses that may add to existing service stresses.
Typical factors affecting fatigue include the following: Stress raisers, usually in the form
of a notch or inclusion; most fatigue fractures may be attributed to notch effects,
inclusion fatigue specimens are rare. High strength materials are much more notch-
sensitive than softer alloys. Corrosion is another factor that affects fatigue. Corroded
parts form pits that act like notches. Corrosion also reduces the amount of material
which effectively reduces the strength and increases the actual stress. Decarburization,
the loss of carbon from the surface of the material, is the next factor. Due to bending
and torsion, stresses are highest at the surface; decarburization weakens the surface
by making it softer. Finally, residual stresses which add to the design stress; the
combined effect may easily exceed the limit stress as imposed in the initial design.
8.2.1 Influence of Processing and Metallurgical Factors on Fatigue
A myriad of factors affect the behaviour of a material under fatigue loading. Obvious
factors include the sign, magnitude, and frequency of loading, the geometry and
material strength level of the structure and the ambient service temperature. However,
processing and metallurgical factors are not often considered, but these factors
determine the homogeneity of materials, the sign and distribution of residual stresses,
and the surface finish. Thus, processing and metallurgical factors have an overriding
influence on the performance of a structure.
8.2.1.1 Processing Factors
Stresses are normally highest at the surface of a structure, so it follows that fatigue
usually initiates at the surface. Stress raisers are more likely to be present as a result of
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surface irregularities introduced by the design of the structure or produced in service or
resulting from processing. Processing factors can introduce a detrimental or beneficial
effect into a structure, usually in the form of effect on strength level or residual stress
condition of the surface material. Therefore, the effect of processing on the mechanical
properties of a material, especially the surface of the material, directly affects fatigue
properties. Processing factors that influence the fatigue life of a structure include the
following: the process by which a part is formed, such as die casting; the heat treatment
of a material, such as quenching, which builds up residual stresses and annealing,
which relieves internal stress (see Figure 7.3); case hardening, such as carburization or
nitriding, which increases surface hardness and strength (see Figure 7.4); surface finish,
such as polished smooth by electro polishing; cold working, which increases strength;
also, cladding, plating, chemical conversion coatings, and anodizing.
Figure 8.3 Effect of hardness on the fatigue life of threads rolled before and
after heat treatment.
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Figure 8.4 Bending fatigue test results on sections from crankshafts: endurance
limit versus surface treatment.
8.2.1.1 Metallurgical Factors
Metallurgical factors refers to areas within the material, wither on the surface or in the
core, which adversely affect fatigue properties. These areas may arise from melting
practices or primary or secondary working of the material or may be characteristic of a
particular alloy system. In virtually all instances the detriment to fatigue properties
results from a local stress-raising effect. Therefore, metallurgical factors affecting
fatigue include the following:
surface defects
sub-surface and core defects
inhomogeneity, anisotropy
improper heat treatment
localized overheating
corrosion fatigue
fretting corrosion.
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8.3 Experimental Analysis of Fatigue Life Curves
Failure due to repeated loading is known as fatigue. A small crack, a scratch, or some
other such minor defect causes localized deformation. This deformation leads to a
small crack if one was not initially present. After cyclic loading, that is, loading in the
same way multiple times, the crack grows, and eventually the material fails. A fatigue life
curve is a graphical representation of the cyclic loading. Simply, a fatigue life curve,
also known as an S-N curve is a plot of the stress amplitude versus the number of cycles
the material goes through before it fails. That is, for a certain stress, the material will fail
within a certain number of cycles. Figure 8.5 is an example of a typical fatigue life
curve.
Figure 8.5 Typical Fatigue Life Curve.
8.6 Fatigue Crack Growth
If an engineering component contains a crack, and if a cyclic or repeated load is applied,
the crack is likely to grow slowly with increasing number of load cycles. This process is
known as fatigue crack growth. In a fatigue crack growth experiment, the progress of a
crack growing under a cyclic load is measured, and the results are plotted as a fatigue
crack growth rate curve, da/dN versus K (that is, change in crack length divided by
change in number of cycles to failure versus change in fracture toughness). A typical
fatigue crack growth curve is shown in Figure 8.6.
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Figure 8.6 Crack growth rates obtained from adjacent pairs of a vs. N data
points.
In the simplest form of a fatigue crack growth rate test, a cyclic load is applied that has
fixed maximum and minimum loading levels. The test specimen is usually a plate of
material in which a crack has already been started at the end of a V-bottom machined
slot. In a typical fatigue crack growth experiment, the sample is loaded in a closed-loop
servo hydraulic testing machine and data for crack length, number of cycles to failure,
and fracture toughness is recorded. From this data the mechanical behaviour for a
certain material can be described under fatigue crack growth loading by the fatigue
crack growth rate curve. This sort of experiment is useful for materials that would
undergo high cyclic loading stresses such as an airplane wing or a helicopter rotor [2].
8.7 Real Life-Design and Manufacturing Considerations
The following describes a relationship between factors that shape the S-N curves as
they are influenced by design and manufacturing conditions and the effects of such
conditions on the fatigue properties of materials, components, and structures.
8.8 Recommendations for Designs to Avoid Fatigue Failures
A designer can help to minimize the possibility of fatigue failure by proper design of
structural components. Many fatigue failures may be attributed to lack of sufficient
consideration of design details or a lack of appreciation of engineering principles.
These principles, which are an integral part of good design of structures subject to
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fatigue, are well reported in literature, but this information has been scattered throughout
sources and may be inaccessible to a designer who needs to understand and utilize the
principles. It is good design practice to seek out sources of this information and to
utilize the principles before, during and after the design process.
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Annexure 1
Additional Scanning Electron Microscope Images [2]
Scanning Electron Microscopy
Scanning Electron Microscope (SEM) image of intergranular fracture indicative of hydrogen embrittlement. (Mag: 400X)
Scanning Electron Microscope (SEM) image of fatigue striations indicative of cyclic crack propagation. (Mag: 700X)
Scanning Electron Microscope (SEM) image of ―beach marks‖ indicative of a progressive fatigue failure. The area of fatigue initiation is noted at the arrow. (Mag: 180X)
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Scanning Electron Microscope (SEM) image of a brittle fracture surface in an aluminium casting. The angular particles in the surface are silicon particles that contribute to the brittleness of the material. (Mag: 1,000X)
Scanning Electron Microscope (SEM) image of microbiological activity in a fire protection system piping. (Mag: 700X)
Scanning Electron Microscope (SEM) image of the surface of a casting void in the fracture surface of an aluminium casting. (Mag: 1,000X)
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Annexure 2
Metallography/Microstructure Evaluation
The properties of a material and its performance in a specific application depend on its
microstructure. Our metallographs (light microscopes) are capable of examinations at
magnifications from 15X to 1,000X. Analyses of microstructure and material defects in
cross-sectioned samples determine material properties, flaw characteristics, and defect
mechanisms.
Metallurgical Technologies, Inc. (Mti) has full metallographic preparation capabilities
from sectioning and mounting the specimen through the grinding and polishing stages
to proper selection and etching techniques of the tested material.
View of intergranular stress corrosion cracking (IGSCC) in an Inconel heat exchanger tube. Note that the crack follows the grain boundaries. (Mag: 500X)
View of chloride stress corrosion cracking in a 316 stainless steel chemical processing piping system. Chloride stress corrosion cracking in austenitic stainless steel is characterized by the multi-branched ―lightning bolt‖ transgranular crack pattern. (Mag: 300X)
A cross-section through a seam weld in a 400 series ferritic stainless steel tube. The seam exhibits a wide fusion zone and a large grain size contributing to brittleness of the weldment. (Mag: 25X)
Cyclic Fatigue Cracks Propagated by a Rust Pit (stress corrosion)
Again, many of the high strength steel alloys are susceptible to stress corrosion.
The photos illustrate such a failure. The first picture is a digital photo with an
arrow pointing to the double origin of the fatigue cracks. The second photograph
at 30X magnification shows a third arrow pointing to the juncture of the cracks
propagating from the rust pits. L- 19, H-11, 300M and Aeromet 100, are
particularly susceptible to stress corrosion and must be kept well-oiled.
Microstructure evaluation of the heat-affected zone of a welded stainless steel piping flange etched to reveal the carbide distribution. Fine carbide particles outline the grain boundaries, indicating a "sensitized" condition resulting in susceptibility to intergranular corrosion. (Mag: 600X)
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Slide 1: Typical fatigue failures in steel components.
Slide 2: Striations in an aluminium alloy.
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Slide 3 : Fatigue failures in the Alexander L Kielland platform.
Slide 4 : Fatigue crack initiation at an inclusion in a high strength steel alloy.
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Fractures are analysed using the latest scanning electron microscopy (SEM) and
other metal testing techniques. You receive a comprehensive written report with
photographic documentation showing each stage of the laboratory analysis. We
identify the cause of failure and recommend correction of material processing such
as heat treatment, plating, machining, and/or design to prevent recurrence of the
problem [2].
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Annexure 3 Microscopic characteristics of fatigue fracture
Striations are the most characteristic microscopic evidence of fatigue fracture,
although striations are not always present on fatigue fracture surfaces, as will be seen.
However, each time the crack is opened by a tensile stress of sufficient magnitude,
creating a tiny ridge, or striation, on each of the mating fracture surfaces. If the
maximum cyclic load remains constant, the striations near the fatigue origin are
extremely small and closely spaced; the crack grows at a slow rate because the part is
still quite strong. However, as the crack gradually propagates, the spacing between
striations increases and the crack grows at an increasingly rapid rate because the crack
greatly weakens the section. Eventually, complete final fracture (stage 3) and
separation occur. Unfortunately, striations are not always visible on fatigue fracture
surfaces for a variety of reasons: On very hard or very soft metals. Artifacts caused by
rubbing or other post fracture damage may produce parallel ridges that resemble
striations. Certain lamellar microstructures in metals resemble fatigue striations.
However, careful study in the electron microscope will reveal that the orientation of the
platelets varies randomly from one location to another, whereas true striations are
generally concentric around the origin.
Macroscopic characteristics of fatigue fracture
Information can be learned about a fatigue fracture with only macroscopic examination.
That is, study with the unaided eye and relatively low magnification – up to perhaps 25 to
50 times magnification – is usually the most important single way to study and analyse
fatigue fractures.
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Lack of Deformation
Since initiation of fatigue fracture does not require a high stress, there is usually little or
no deformation in a part or specimen that has fractured by fatigue. If the maximum
stress did not exceed the yield strength (actually the elastic limit), there can be no gross
plastic, or permanent, deformation, although the final rupture region may have some
obvious macroscopic deformation. The typical fatigue fracture that occurs in most load-
bearing parts, which have relatively low-stress, high-cycle loading.
Not only the fracture surface but the entire part should be examined for deformation.
For example, if a unidirectional (one-way) bending fracture is observed, it is useful to
carefully reassemble the pieces to determine if there was gross deformation in the part
prior to fracture. Of course, the origin of the fracture would be on the convex side,
which is the tension side in bending.
As pointed out at the beginning of this section, in a ―true‖ high cycle fatigue fracture,
there will be no deformation in the fatigue region, provided that there has been no post
fracture damage to the fracture surface. If the final rupture region (stage 3) is ductile, the
resulting deformation will prevent close realignment of the fractured pieces; however, if
the final rupture region is a truly brittle fracture, there should be no gross deformation,
except for post fracture damage. A partially ductile/brittle final rupture region probably
will show some degree of deformation.
Beachmarks
―Beachmarks‖ are a unique feature found in many fatigue fractures,
and their presence is a positive means of identifying fatigue
fractures. Beachmarks also have been called ―stop marks,‖ ―arrest
marks,‖ ―clamshell marks,‖ and ―conchoidal marks,‖ all in an attempt to describe their
origin or characteristic appearance. The term ―beachmarks‖ is the most commonly used
term but is not really as descriptive as some of the others.
At any rate, this term is used to describe macroscopically visible marks or ridges that
are characteristic of interruptions in the propagation periods (stage 2) of fatigue
fractures in relatively ductile metals.
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Beachmarks must not be confused with striations, although they frequently are present
on the same fracture surface; there may be many thousands of microscopic striations
between each pair of macroscopic beachmarks.
Ratchet Marks
The term ―ratchet marks‖ is used to describe features that are very useful in
identification of fatigue fractures and in locating and counting the number of fatigue
origins. These marks are essentially perpendicular to the surface from which fatigue
fractures originate. Therefore, in circular, shaft-like parts, the ratchet marks are
essentially radial, pointing toward the centre; in flat parts, such as leaf springs, they
initially are perpendicular to the surface but may curve if the bending is unidirectional.
The ratchet marks are not the origins themselves; each ratchet mark separates two
adjacent fatigue fractures. As the cracks become deeper, the cracks from each origin
tend to grow together and become essentially one fatigue fracture that has numerous
origins. The number of ratchet marks equals or is one less than the number of origins;
thus recognition of the number of ratchet marks is important in determining the number
of origins.
Similarities between Striations and Beachmarks
Both striations and beachmarks identify the position of the tip of the fatigue crack
at a given point in time.
Both striations and beachmarks expand from the fatigue origin or origins, often in
a circular or semicircular fashion.
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Both striations and beachmarks are relatively parallel ridges which do not cross
similar features from another origin.
Some fatigue fracture surfaces have neither striations nor beachmarks. Artifacts, or
false features, can confuse observation of both striations and, beachmarks.
Differences between Striations and Beachmarks
The most obvious difference between striations and beachmarks is size. Striations
are extremely small ridges, visible only with an electron microscope. Beachmarks
are much larger than striations. If they are present, they are normally visible to the
unaided eye.
The other difference between striations and beachmarks, as previously
mentioned, is the factors that cause them. Striations represent the advance of the
crack front by one load application in many ductile metals, whereas beachmarks
locate the position of the crack front when repetitive, fluctuating loading was
stopped for a period of time [2].
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Annexure 4
SAMPLES FAILURE
Stub Axle failure This is the classic reverse bending fatigue of a steel stub axle from a road vehicle. Notice cracks have grown from 8 o‘clock upwards and to a lesser extent from 2 o‘clock downwards. The rough central region is the final ductile rupture.
Bending fatigue fracture This 100 mm diameter steel shaft failed after a long period of service on a large dumper truck. The keyway terminated in a circumferential groove approximately half the depth of the keyway.
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Fatigue cracks in steel cycle frame Although steel has a fatigue endurance limit certain parts of the frame are stressed above this limit and are prone to fatigue cracking. The rear triangle, comprising the chainstays and seatstays are particularly vulnerable.
On the opposite side, the associated cracked point had caused surface corrosion. Notice that the cracks had occurred close to the brazed-on cross member. The heat affected zone associated with brazing and welding usually reduces the endurance limit and raises the chance of fatigue crack initiation.
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A metallurgical investigation of the extent of the heat affected zone is in progress. The manufacturers did not fulfil their ―lifetime‖ guarantee in this instance by suggesting that damaged paintwork caused corrosion, which in turn initiated fatigue.
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REFERENCES
[ 1 ] Jaap Schijve, Fatigue of Structures and Materials, Second Edition, Springer
2009, ISBN-13: 978-1-4020-6807-2
[ 2 ] An Introduction of mechanical Testing Pictorial Basic
CMM NDT Services (www.cmmok.xinwen520.com )
[ 3 ] Alan F. Liu, Mechanics and Mechanisms of Fracture: An Introduction,
ASM International, 2005, ISBN: 0-87170-802-7
[ 4 ] Tom Lassen, Fatigue Life Analyses of Welded Structures, ISTE UK 2006,
ISBN-10: 1-905209-54-1
[ 5 ] BRITISH STANDARD BS 7608:1993, Fatigue design and assessment of
steel structures, ISBN 0 580 21281 5
[ 6 ] John M. Barsom, Fracture and Fatigue Control in Structures: Applications
of Fracture Mechanics, ASTM 1999, ISBN 0-8031-2082-6