翻译英文版(刘宏健)

Fatigue of structures and materials in the 20th century

and the state of the art

J. Schijve

Delft University of Technology, Faculty of Aerospace Engineering, Kluyverweg 1, 2629HS Delft, The Netherlands Received 30 October 2002; received in revised form 22 January 2003; accepted 4 February 2003

Abstract

The paper surveys the historical development of scientific and engineering knowledge about fatigue of materials and structures in the 20th century. This includes fatigue as a material phenomenon, prediction models for fatigue properties of structures, and load spectra. The review leads to an inventory of the present state of the art. Some final remarks follow in an epilogue.2003 Elsevier Science Ltd. All rights reserved.

Keywords: Fatigue mechanism; Fatigue properties; Prediction; Load spectra; History

1. Introduction

An evaluation of fatigue of structures and materials in the 20th century raises the question what happened in the 19th century? The answer is that fatigue of structures became evident as a by-product of the industrial revolution in the 19th century. In some more detail, it was recognized as a fracture phenomenon occurring after a large numbers of load cycles where a single load of the same magnitude would not do any harm. Fatigue failures were frequently associated with steam engines, locomotives and pumps. In the 19th century, it was considered to be mysterious that a fatigue fracture did not show visible plastic deformation. Systematic fatigue tests were done at a few laboratories, notably by August Wo¨hler. It was recognized that small radii in the geometry of the structure should be avoided. Fatigue was considered to be an engineering problem, but the fatigue phenomenon occurring in the material was still largely in the dark. Some people thought that fatigue implied a change from a fibrous to a crystalline, brittle structure in view of the absence of visible plastic deformation. A fundamental step regarding fatigue as a material problem was made in the

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beginning of the 20th century by Ewing and Humfrey in 1903 [1]. They carried out a microscopic investigation which showed that fatigue crack nuclei start as microcracks in slip bands. Much more evidence about fatigue as a material phenomenon was going to follow in the 20th century.

Fatigue as a technical problem became evident around the middle of the 19th century. About 100 years later, in the middle of the 20th century, the development of fatigue problems were reviewed in two historical papers by Peterson in 1950 [2] and Timoshenko in 1954

[3]

. Both authors were already well-known for important

publications. Peterson reviewed the discussion on fatigue problems during meetings of the Institution of Mechanical Engineers at Birmingham held just before 1850. He also mentioned historical ideas about fatigue as a material phenomenon and the microscopic studies carried out by Gough and co-workers and others around 1930. Crack initiation occurred in slip bands and (quoting Peterson) ―one or more of these minute sources starts to spread and this develops into a gross crack which, in general, meanders through the grains in zigzag fashion in an average direction normal to the direction of tensile stresses. It should be remembered, however, that although the fractured surface generally follows a normal stress field, the microscopic source of failure is due to shear‖. Peterson also refers to the concept of the ‗endurance limit‘, as already defined by Wo¨hler. In this paper the endurance limit is generally referred to as the fatigue limit which is an important material property for various engineering predictions on fatigue.

Timoshenko in his review discussed the significance of stress distributions and emphasized stress concentrations around notches. According to Timoshenko, the importance was recognized by design engineers around the end of the 19th century, and the knowledge was further refined in the beginning of the 20th century. Timoshenko referred to the significance of theoretical stress analysis employing complex variables (Kolosov, Inglis, Mushkelisvili, Savin and others). But he considered experimental studies on stress distributions and stress concentrations to be of prime importance. He mentioned several developments on strain measurements, basically by using mechanical displacement meters, strain gauges and photo-elastic models. A famous book published in 1950 was Handbook of Experimental Stress Analysis by Hete´nyi [4]. Timoshenko thought that great progress had been made. He also raised the question ―how does a high, localized stress weaken a machine part in

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service? This important question can be satisfactorily answered only on the basis of an experimental investigation‖.

The above re´sume´ of developments before 1950 now seems to be ‗old stuff‘, primarily because substantial improvements of our present knowledge about fatigue occurred in the second half of the 20th century. The improvements became possible due to the development of essentially new experimental facilities, computers and numerical stress analysis. However, some basic concepts remained, such as that fatigue in metallic materials is due to cyclic slip, and stress concentrations contribute to a reduced fatigue endurance. One other characteristic issue of a more philosophical nature also remained, the question of whether fatigue is a material problem or an engineering problem, or both in some integrated way? The present paper primarily covers developments in the second half of the previous century. It is not the purpose to summarize all noteworthy happenings in a historical sequence, also because informative reviews about the history of ‗fatigue‘ have been presented in the last decades of the 20th century, e.g. by Mann [5], Schu¨ tz [6], Smith [7] and others. Moreover, collections of significant publications have been compiled [8,9]. The emphasis in this paper will be on how the present knowledge was acquired. The development of fatigue problems of structures and materials in the 20th century was fundamentally affected by milestone happenings, important discoveries, and various concepts of understanding fatigue phenomena. Furthermore, the approach to solving fatigue problems and the philosophy on the significance of fatigue problems is of great interest.

The efforts spent on fatigue investigations in the 20th century is tremendous, as illustrated by numerous publications. John Mann [10] published books with references to fatigue. Later he continued this work to arrive at about 100 000 references in the 20th century compared to less than 100 in the 19th century. The large number of publications raises an obvious question. Is the problem so difficult and complex, or were we not clever enough to eliminate fatigue problems of our industrial products? Various conferences on fatigue of structures and materials are already planned for the forthcoming years of the 21st century implying that the fatigue problem is apparently not yet fully solved. If the problem still exists after 100 years in the previous century, there is something to be explained.

In a recent textbook [11] the author has used the picture shown in Fig. 1 to survey

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prediction problems associated with fatigue properties of structures. The predictions are the output of a number of procedures and Fig. 1 presents the scenario of the various aspects involved.

The input problems occur in three categories: (i) design work, (ii) basic information used for the predictions, and (iii) fatigue load spectra to which the structure is subjected. Each of the categories contains a number of separate problems, which again can be subdivided into specific aspects, e.g. ‗joints‘ cover welded joints, bolted joints, riveted joints, adhesively bonded joints.

Fig. 1 illustrates that the full problem can be very complex depending on the structural design, type of material, production variables, load spectra and environment. Prediction models are presented in the literature and software is commercially available. The prediction of the fatigue performance of a structure is the result of many steps of the procedures adopted, and in general a number of plausible assumptions is involved. It implies that the accuracy of the final result can be limited, the more so if statistical variables also have to be considered. The reliability of the prediction should be carefully evaluated, which requires a profound judgement, and also socalled engineering judgement, experience and intuition. It has persistently been emphasized in Ref. [11] that physical understanding of the fatigue phenomena is essential for the evaluation of fatigue predictions. A designer cannot

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simply rely on the validity of equations. Behind an equation is a physical model and the question is whether the model is physically relevant for the problem considered. This implies that each topic in Fig. 1 should also be a relevant subject for research, and the number of variables which can affect the fatigue behavior of a structure is large. Without some satisfactory understanding of aspects involved, predictions on fatigue become inconceivable. In this paper, it will be summarized how the understanding in the previous century has been improved, sometimes as a qualitative concept, and in other cases also quantitatively. It should already be said here that qualitative understanding can be very important, even if a strictly quantitative analysis is not yet possible. The major topics discussed in the following sections are associated with: (i) material fatigue as a physical phenomenon (Section 2), (ii) the S-N curve and the fatigue limit (Section 3), (iii) prediction of fatigue properties (Section 4), and (iv) fatigueload spectra in service (Section 5). These topics are first discussed to see the development of the knowledge about fatigue of structures and materials in the 20th century. Afterwards, the text covers an evaluation of the present understanding also in relation to the engineering significance (Section 6). The paper is concluded with some general remarks about the present state of the art and expectations for the 21st century (Section 7).

2. Fatigue of materials as a physical phenomenon

2.1. Fatigue crack initiation

As said before, fatigue damage in steel in the 19th century was associated with a mysterious crystallizing of a fibrous structure. It was not yet defined in physical terms. In the first half of the 20th century, cyclic slip was considered to be essential for microcrack initiation. Cracks, even microcracks, imply decohesion in the material and should thus be considered to be damage. But is cyclic slip also damage, and what about cyclic strain hardening in slip bands? In the thirties, Gough [12] postulated that fatigue crack initiation is a consequence of exceeding the limit of local strain hardening. The idea was adopted by Orowan in 1939 [13] who argued that the local exhaustion of ductility leads to a localized increase of the stress and ultimately to cracking. This concept was used in 1953 by Head [14] in a model for obtaining an equation for fatigue crack growth.

An important question about the ductility exhaustion theory is how cracking

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occurs on an atomic level. Stroh [15] analyzed the stress field around a piled-up group of dislocations. According to him, the local stress can become sufficiently high to cause local cleavage. However, it was difficult to see why high local stresses can not be relaxed near the material surface by plastic deformation in a basically ductile material. The ductility exhausting theory did not become a credible crack initiation model, the more so since the detection of striations in the late 1950s [16,17] indicated that crack extension occurred in a cycle-by-cycle sequence, and not in jumps after intervals of cycles required for an increasing strain-hardening mechanism.

In the 1950s, the knowledge of dislocations had beenwell developed. Cyclic slip was associated with cyclic dislocation movements. It is not surprising that people tried to explain the initiation and crack growth in terms of creating crevices in the material or intrusions into the material surface as a result of some specific dislocation mobilities. Interesting dislocation models were proposed in the 1950s, noteworthy by Cottrell and Hull, based on intersecting slip systems [18], and by Mott, based on generation of vacancies [19]. Microscopic observations were made to see whether the proposed models for crack initiation and crack growth were in agreement with a model. Several papers of historical interest were collected in 1957 [20] and 1959 [21] respectively. The microscopic work of Forsyth [22] on extrusions and intrusions in slip bands should be mentioned, see Fig. 2. Similar figures have been used by several authors to discuss basic aspects of the fatigue crack initiation process. Three fundamental aspects are: the significance of the free material surface, the irreversibility of cyclic slip, and environmental effects on microcrack initiation. Microcracks usually start at the free surface of the material,1 also in unnotched specimens with a nominally homogeneous stress distribution tested under cyclic tension. The restraint on cyclic slip is lower than inside the material because of the free surface at one side of the surface material. Furthermore, microcracks start more easily in slip bands with slip displacements normal to the material surface [23] which seems to be logical when looking at Fig. 2. It still remains to be questioned why cyclic slip is not reversible. Already in the 1950s, it was understood that there are two reasons for non-reversibility. One argument is that (cyclic) strain hardening occurs which implies that not all dislocations return to their original position. Another important aspect is the interaction with the environment. A slip step at the free surface implies that fresh material is exposed to the environment. In a non-inert

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environment, most technical materials are rapidly covered with a thin oxide layer, or some chemisorption of foreign atoms of the environment occurs. An exact reversibility of slip is then prevented. A valid and important conclusion is that fatigue crack initiation is a surface phenomenon.

In the 1950s, microscopical investigations were still made with the optical microscope. It implies that crack nucleation is observed on the surface where it indeed occurs. As soon as cracks are growing into the material away from the free surface, only the ends of the crack front can be observed at that free surface. It is questionable whether that information is representative for the

growth process inside the material, a problem sometimes overlooked. Microscopic observations on crack growth inside the material require that cross-sections of a specimen are made.

Several investigations employing sectioning were made in the 1950s and before. These showed that in most materials fatigue cracks are growing transcrystalline. Although the fatigue fractures looked rather flat as viewed by the unaided eye, it turned out that the crack growth path under the microscope could be rather irregular depending on the type of material. In materials with a low stacking fault energy (e.g. Cu- and Ni-alloys), cross slip is difficult and as a result cyclic slip bands are narrow and straight. Crack growth on a microscale occurs in straight segments along these bands. In materials with a high stacking fault energy (e.g. Al-alloys) cross slip is easy. Moreover, in the Al crystal lattice there are many slip systems which can easily be activated. As a consequence, slip lines are wider and can be rather wavy. Crack

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growth on a micro scale does not suggest that it occurs along crystallographic planes. As a result, fatigue on a microscale can be significantly different for different materials. The behavior is structure-sensitive, depending on the crystal structure (fcc, bcc, or hexagonal), elastic anisotropy of the crystalline structure, grain size, texture, and dislocation obstacles (e.g. pearlite bands in steel, precipitated zones in Al-alloys, twins, etc.). An extensive survey of the material fatigue phenomenon was recently presented in a book by Suresh [24]. 2.2. Fractographic observations

The description of the fatigue mechanism in different materials was studied in the

1950s and in the following decades. A significant experimental milestone was the introduction of the electron microscope (EM), originally the transmission electron microscope (TEM) in the 1950s, and later the scanning electron microscope (SEM) in the 1970s. Microscopic investigations in the TEM are more laborious than in the SEM because either a replica of the fracture surface must be made, or a thin foil of the material. The thin foil technique is destructive and does not show the fatigue fracture surface, but information on the material structure can be obtained, such as forming of subgrains under cyclic loading. The thin foil technique requires a good deal of experimental expertise.

Investigations of fatigue fracture surfaces in the SEM are now a rather well standardized experimental option, which can indicate where the fatigue fracture started, and in which directions it was growing.2 A fundamental observation was made with the electron microscope around 1960. Fractographic pictures revealed striations which could be correlated with individual load cycles. By mixing of small and large load cycles in a fatigue test the occurrence of one striation per load cycle was proven by Ryder [17]. An example is shown in Fig. 3. The striations are supposed to be remainders of microplastic deformations at the crack tip, but the mechanism can be different for different materials. Several models for forming striations were proposed in the literature, two early ones in 1967 by Pelloux and Laird, respectively

[25]

. Because of microplasticity at the crack tip and the crack extension mechanism in

a cycle, it should be expected that the profile of striations depends on the type of material. Terms such as ductile and brittle striations were adopted [22]. Striations could not be observed in all materials, at least not equally clearly. Moreover, the

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visibility of striations also depends on the severity of load cycles. At very low stress amplitudes it may be difficult to see striations although fractographic indications were obtained which showed that crack growth still occurred in a kind of a cycle-by-cycle sequence [26].

Striations have also shown that the crack front is not simply a single straight line as usually assumed in fracture mechanics analysis. Noteworthy observations on this problem were made by Bowles in the late 1970s [27,28] who developed a vacuum infiltration method to obtain a plastic casting of the entire crack. The casting could then be studied in the electron microscope. An example is shown in Fig. 4 which illustrates that the crack front is indeed a curved line and the crack tip is rounded. Macroscopic shear lips, see Fig. 5, were well known for aluminium alloys from the early 1960s [29], but they were also observed on fatigue cracks in other materials

[30–32]

. The width of the shear lips increased for faster fatigue crack growth, and

finally a full transition from a tensile mode fatigue crack to a shear mode fatigue crack

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can occur. The shear lips are a surface phenomenon because crack growth in the

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shear mode is not so constrained in the thickness direction. Shear lips are a macroscopic deviation from a mode-I crack assumed in a fracture mechanics analysis.

Fatigue cracks in thick sections can be largely in the tensile mode (mode I) because shear lips are then relatively small. However, the topography of the tensile mode area observed in the electron microscope indicates a more or less tortuous surface although it looks rather flat if viewed with the unaided eye. Large magnifications clearly show that the fracture surface on a microlevel is not at all a nicely flat area. It is a rather irregular surface going up and down in some random way depending on the microstructure of the material. It has also been shown for aluminium alloys that the roughness of the fracture surface depends on the environment [33]. An inert environment increased the surface roughness whereas an aggressive environment (salt water) promoted a more smooth fracture surface. Similarly, shear lips were narrower in an aggressive environment and wider in an inert environment. These trends were associated with the idea that an aggressive environment stimulates tensile-decohesion at the crack tip, whereas an inert environment promotes shear decohesion. It should be understood that the crack extension in a cycle (i.e. the crack growth rate) depends on the crack growth resistance of the material, but also on the crack driving force which is different if deviations of the pure mode I crack geometry are present, e.g. shear lips and fracture tortuosity.

2.3. More about fatigue crack growth

In the 1950s, many investigators mentioned how early in the fatigue life they

could observe microcracks. Since then it was clear that the fatigue life under cyclic loading consisted of two phases, the crack initiation life followed by a crack growth period until failure. This can be represented in a block diagram, see Fig. 6. The crack initiation period may cover a large percentage of the fatigue life under high-cycle fatigue, i.e. under stress amplitudes just above the fatigue limit. But for larger stress amplitudes the crack growth period can be a substantial portion of the fatigue life. A special problem involved is how to define the transition from the initiation period to the crack growth period.

It was in the early 1960s that the stress intensity factor was introduced for the

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correlation between the crack growth rate, da/dN, and the stress intensity factor range,

K. The first paper was published in 1961 by Paris, Gomez and Anderson [34], and it

turned out to be a milestone publication. They adopted the K-value from the analysis of the stress field around the tip of a crack as proposed by Irwin [35] in 1957, another milestone of the application of fracture mechanics. The well-known general equation in polar coordinates for the stress distribution around the crack tip is:

ijkf(ij)（）

2rwith K as the stress intensity factor and the polar coordinates r and q (see Fig. 7). Eq. (1) is an asymptotic solution which is valid for small values of r only, i.e. r＜＜a with ‗a‘ as the crack length. The stress intensity factor is given by:

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ksa （2）

with b as the geometry factor. The results of the crack growth tests of Paris et al. were expressed in terms of da/dN as a function of K on a double log scale,3 see Fig. 8a, which shows a linear relation between log(da/dN) and log(K). Many more crack growth tests carried out later indicated the same trend which led to the well-known Paris equation:

dadNCkM （3）

with C and m as experimentally obtained constants. The equation is a formal description of results of a fatigue crack growth experiment. At the same time, it must be recognized that fatigue crack growth is subjected to physical laws. In general terms, something is driving the crack extension mechanism which is called the crackdriving force. This force is associated with the K-value .The stress intensity factor is related to the strain energy release rate, i.e. the strain energy in the material which is available for producing crack extension. The relation to be found in textbooks is:

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dUKdaE2 （4）

With E=E(Young‘s modulus) for plane stress, and E/(1-v2) for plane strain (v= Poisson‘s ratio). The strain energy looks like a characteristic variable for energy balances. The material response (da/dN) is characterized in Eq. (3), but the experimental constants C and m are not easily associated with physical properties of the material. However, the crack growth rate obtained is representing the crack growth resistance of the material.

Already by the 1960s it was clear that the correlation of da/dN and K depends on the stress ratio R. This could be expected because an increased mean stress for a

constant S should give a faster crack growth while the R-value is also increased. Furthermore, results of crack growth tests indicated systematic deviations of the Paris equation at relatively high and low K-values. It has led to the definition of three regions in da/dN-K graphs, regions I, II and III respectively, see Fig. 8b. Obvious questions are associated with the vertical asymptotes at the lower K boundary of region I and the upper K boundary of region III. The latter boundary appears to be logical because if Kmax exceeds the fracture toughness (either Kc or KIc) a quasi-static failure will occur and fatigue crack growth is no longer possible. It still should be recognized that the Kmax value causing specimen failure in the last cycle of a fatigue crack growth experiment may well be different from Kc or Kic measured in a fracture toughness test.

From the point of view of fracture mechanics, the occurrence of a lower boundary in region I is not so obvious. As long as a K-value can be defined for the tip of a crack, a singular stress field should be present and micro-plasticity at the tip of the crack should occur. So, why should the crack not grow any more; for which physical reason should there be a threshold K-value (Kth).

New ideas on Kth were associated with observations on so-called small cracks. These cracks occur as microcracks in the beginning of the fatigue life starting at the material surface or just subsurface. The first relevant paper was published by Pearson

[36]

in 1975 who observed that small surface cracks were growing much faster than

large macro cracks at nominally similar Kvalues. It was confirmed in several

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investigations that microcracks could grow at low K-values, whereas macrocracks did not grow at these low K-values where K ＜Kth. Illustrative data of Wanhill are shown in Fig. 9 [37]. The small-crack problem became a well recognized subject for further research. Various crack growth barriers offered by the material structure (e.g. grain boundaries, pearlite in steel, phase boundaries in general) could be significant for microcracks [38] whereas they were less relevant to macrocrack growth. As a result, considerable scatter was observed in microcrack growth rates, see Fig. 9. Moreover, the barriers affecting microcrack growth could be quite different for different materials. Although proposals for fracture mechanics predictions of the growth of microcracks were presented in the literature, the publications about this issue in the last decades of the previous century were not always convincing. Actually, it should be recognized that the K-concept for such small cracks in a crystalline material becomes questionable. The plastic zone is a slip band and its size is not small compared to the crack length of the microcrack. They may even have a similar size.

Another question about the _Kth concept applies to macrocracks. Why do large

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cracks stop growing if K ＜Kth? A formal answer to this question is because the crack driving force does not exceed the crack growth resistance of the material. At low K-values the crack driving force is low which affects the crack front microgeometry, the crack front becomes more tortuous, and also the crack closure mechanism is changing [39]. It may then occur that the crack driving force is just no longer capable of producing further crack growth.

A concept to be discussed here is the occurrence of crack closure, and more specifically plasticity induced crack closure. In the late sixties [40,41], Elber observed that the tip of a growing fatigue crack in an Al-alloy sheet specimen (2024-T3) could be closed at a positive stress (tensile stress).4 Crack opening turned out to be a non-linear function of the applied stress, see Fig. 10. During loading from S 0 to Sop the crack opening displacement (COD) is a non-linear function of the applied stress. For S ＞Sop the behavior is linear with a

slope corresponding to the specimen compliance with a fully opened crack. The same non-linear response was observed during unloading. During the non-linear behavior the crack is partly or fully closed due to plastic deformation left in the wake of the growing crack. Elber argued that a load cycle is only effective in driving the growth of a fatigue crack if the crack tip is fully open. He defined the effective S and K as:

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SeffSmaxSopandkeffSeffa （5）

(b is the geometry factor). He then assumed that the crack growth rate is a function of

Keff only.

dadNfkeff （6）

Elber found that the crack opening stress level depends on the stress ratio for which he proposed the relation:

USeffSfR0.50.4Rfor2024T3A1alloy（7）

This relation is an empirical result. Moreover, Elber proposed that the relation should be independent of the crack length. The Elber approach was carried on in later investigations, partly because it was attractive to present crack growth data of a material for various R-values by just one single curve according to Eq. (6). It turned out that the relation in Eq. (7) could be significantly different for other materials which is not surprising because the cyclic plastic behavior depends on the type of material. In the 1980s, the crack closure concept was much welcomed by investigators on crack growth models for fatigue under VA loading [42]

3. The S-N curve and the fatigue limit

3.1. Aspects of the S-N curve

Wo¨hler had already carried out experiments to obtain S-N curves in the 19th century. For a long time such curves were labeled as a Wo¨hler curve instead of the now more frequently used term S-N curve. In the 20th century numerous fatigue tests were carried out to produce large numbers of S-N curves. In the beginning, rotating beam tests on unnotched specimens were popular because of the more simple experimental facilities available in the early decades. The significance of testing notched specimens was recognized, especially by engineers. Fatigue-testing machines for loading in tension, torsion and bending were available before 1940. The excitation of cyclic loads occurred by mechanical or hydraulic systems. High

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frequencies were obtained with resonance machines. Fatigue became more and more recognized as a problem to be considered for various small and large industrial products in view of economical reasons. Apart from basic research on unnotched specimens, many test series were also carried out on notched specimens. The tests were performed with a constant mean stress and a constant stress amplitude (referred to as constant-amplitude tests, or CA tests). The S-N curves should give useful information about the notch and size effect on the fatigue life and the fatigue limit. Initially, the fatigue life N was plotted on a logarithmic scale in the horizontal direction, and the stress amplitude on a linear scale in the vertical direction. Many curves can be found in the literature and collections of these curves have been published as data banks while commercial software also contains this type of information.

For low stress amplitudes, the S-N curve exhibited a lower limit which implies that fatigue failures did not occur after high numbers of load cycles, see Fig. 11. The horizontal asymptote of the S-N curve is called the fatigue limit (in some publications the name endurance limit is used). The fatigue limit is of practical interest for many structures which are subjected to millions of load cycles in service while fatigue failures are unacceptable. The fatigue limit is considered in more detail later.

At the upper side of the S-N curve (large stress amplitudes) another horizontal asymptote appears to be present. If failure did not occur in the first cycle, the fatigue life could be several hundreds of cycles. Such fatigue tests were not easily carried out on older fatigue machines because adjusting the correct load amplitude required too many cycles. A real breakthrough for fatigue testing equipment occurred in the 1950s and 1960s when closed-loop fatigue machines were introduced employing a feedback signal from the specimen to monitor the load on the specimen. With this technique, the fatigue load could be adjusted by a computer-controlled system. Furthermore, if S-N curves were plotted on a double-logarithmic scale, the curves became approximately linear (the Basquin relation).

The interest for short fatigue lives is relevant for structures with a load spectrum of small numbers of severe load cycles only (e.g. high-pressure vessels). In the 1960s, this has led to -N curves. Instead of applying a stress amplitude to a specimen, a constant strain amplitude is maintained in the critical section of the specimen. The

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problem area was designated as ‗low cycle fatigue‘ (Fig. 11) which actually implies that macro plastic deformation occurs in every cycle. It turned out that the -N curve in the low-cycle regime, again plotted on a doublelog scale, was a linear function:

aNconstantC （8）



The equation is known as the Coffin–Manson relation [43,44].

At lower stress amplitudes macro plastic deformation does not occur and the fatigue phenomenon is labeled as ‗high-cycle fatigue‘, see Fig. 11. It is assumed that cyclic deformations on a macro scale are still elastic.

Questions to be considered with respect to the S-N curve are: (i) Is the fatigue phenomenon under highcycle and low-cycle fatigue still a similar phenomenon, and (ii) what is the physical meaning of the fatigue limit? With respect to the first question, it should be pointed out that crack initiation and crack growth are both significant for the fatigue life. The crack initiation is easily affected by a number of material surface conditions (surface roughness, surface damage, surface treatments, soft surface layers, surface residual stress). These conditions are important for high-cycle fatigue and the fatigue limit, but they are generally overruled during

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low-cycle fatigue where plastic deformations at the material surface will occur anyway. Moreover, macro crack growth under low-cycle fatigue is rather limited because small cracks will already induce complete failure at the high stress levels. This implies that low-cycle fatigue is important for special problems where large plastic deformations cannot be avoided. Low-cycle fatigue problems are also often associated with hightemperature applications. However, during high-cycle fatigue significant macrocrack growth occurs. At still lower stress amplitudes the physical meaning and the engineering significance of the fatigue limit have to be considered. 3.2. The fatigue limit

The formal definition of the fatigue limit (see footnote 5) appears to be rather obvious. It is the stress amplitude for which the fatigue life becomes infinite in view of the asymptotic character of the S-N curve. However, fatigue tests must be terminated after a long testing time. If it occurs after 107 cycles (see papers in ref.

[45]

), the nofailure stress level need not be a fatigue limit. Actually, it was also labeled

as an endurance limit associated with a certain high number of cycles.

From an engineering point of view, it appears to be more logical to define the fatigue limit as the highest stress amplitude for which failure does not occur after high numbers of load cycles. This definition should cover situations where fatigue failures are unacceptable, e.g. in various cases of machinery. Design stress levels must then remain below the fatigue limit which emphasizes the significance of the fatigue limit as a material property.

References

[1] Ewing JA, Humfrey JCW. The fracture of metals under repeated alternations of stress. Phil Trans Roy Soc 1903;A200:241–50.

[2] Peterson RE. Discussion of a century ago concerning the nature of fatigue, and review of some of the subsequent researches concerning the mechanism of fatigue. In: ASTM Bulletin, 1. American Society for Testing and Materials; 1950. p. 50–6. [3] Timoshenko S. Stress concentration in the history of strength of materials. The William M. Murray Lecture. Proc Soc Exp Stress Analysis (SESA) 1954;12:1–12. [4] Hete´nyi M. Handbook of experimental stress analysis. New York: John Wiley; 1950.

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[5] Mann JY. Aircraft fatigue—with particular emphasis on Australian operations and research. In: Proc 12th Symp of the Int Committee on Aeronautical Fatigue (ICAF). Centre d‘Essais Ae´ronautique de Toulouse; 1983.

[6] Schu¨ tz W. A history of fatigue. Engng Fract Mech 1996;54:263–300. [7] Smith RA. Fatigue crack growth, 30 years of progress. Pergamon; 1986.

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