Effect of Metallurgical Variables on Fatigue

The fatigue properties of metals are quite structure-sensitive. However, at the present time there are only a limited number of ways in which the fatigue properties can be improved by metallurgical means. By far the greatest improvements in fatigue performance result from design changes, which reduce stress concentration and from the intelligent use of beneficial compressive residual stress, rather than from a change in material. Nevertheless, there are certain metallurgical factors, which must be considered to ensure the best fatigue performance from a particular metal or alloy.

Fatigue tests designed to measure the effect of some metallurgical variable, such as special heat treatments, on fatigue performance are usually made with smooth, polished specimens under completely reversed stress conditions. It is usually assumed that any changes in fatigue properties due to metallurgical factors will also occur to about the same extent under more complex fatigue conditions, as with notched specimens under combined stresses.

Fatigue properties are frequently correlated with tensile properties. In general, the fatigue limit of cast and wrought steels is approximately 50 percent of the ultimate tensile strength. The ratio of the fatigue limit (or the fatigue strength at 10⁶ cycles) to the tensile strength is called the fatigue ratio.

Several nonferrous metals such as nickel, copper, and magnesium have a fatigue ratio of about 0.35. While the use of correlations of this type is convenient, it should be clearly understood that these constant factors between fatigue limit and tensile strength are only approximations and hold only for the restricted condition of smooth, polished specimens which have been tested under zero mean stress at room temperature.

For notched fatigue specimens the fatigue ratio for steel will be around 0,20 to 0,30. However, as yield strength is increased by the various strengthening mechanisms, the fatigue limit usually does not increase proportionately. Most high-strength materials are fatigue-limited.

Several parallels can be drawn between the effect of certain metallurgical variables on fatigue properties and the effect of these same variables on tensile properties. The effect of solid-solution alloying additions on the fatigue properties of iron and aluminum parallels nearly exactly their effect on the tensile properties. Gensamer showed that the fatigue limit of a eutectoid steel increased with decreasing isothermal-reaction temperature in the same fashion as did the yield strength and the tensile strength.

However, the greater structure sensitivity of fatigue properties, compared with tensile properties, is shown in tests comparing the fatigue limit of a plain carbon eutectoid steel heat-treated to coarse pearlite and to spheroidite of the same tensile strength. Even though the steel in the two structural conditions had the same tensile strength, the pearlitic structure resulted in a significantly lower fatigue limit due to the higher notch effects of the carbide lamellae in pearlite.

There is good evidence that high fatigue resistance can be achieved by homogenizing slip deformation so that local concentrations of plastic deformation are avoided. This is in agreement with the observation that fatigue strength is directly proportional to the difficulty of dislocation cross slip.

Materials with high stacking-fault energy permit dislocations to cross slip easily around obstacles, which promotes slip-band formation and large plastic zones at the tips of cracks. Both of these phenomena promote the initiation and propagation of fatigue cracks. In materials with low stacking-fault energy, cross slip is difficult and dislocations are constrained to move in a more planar fashion. This limits local concentrations of plastic deformation and suppresses fatigue damage.

While the concept has been useful in understanding fatigue mechanisms, the ability to control fatigue strength by altering stacking-fault energy has practical limitations. A more promising approach to increasing fatigue strength appears to be the control of microstructure through thermomechanical processing to promote homogeneous slip with many small regions of plastic deformation as opposed to a smaller number of regions of extensive slip.

The dependence of fatigue life on grain size varies also depending on the deformation mode. Grain size has its greatest effect on fatigue life in the low-stress, high-cycle regime in which stage 1 cracking predominates. In high stacking-fault-energy materials (such as aluminum and copper) cell structures develop readily and these control the stage 1 crack propagation. Thus, the dislocation cell structure masks the influence of grain size, and fatigue life at constant stress is insensitive to grain size. However, in a low slacking-fault-energy material (such as alpha brass) the absence of cell structure because of planar slip causes the grain boundaries to control the rate of cracking. In this case, fatigue life is proportional to grain diameter.

In general, quenched and tempered microstructures result in the optimum fatigue properties in heat-treated low-alloy steels. However, at a hardness level above about Rc 40, a bainitic structure produced by austempering results in better fatigue properties than a quenched and tempered structure with the same hardness. Electron micrographs indicate that the poor performance of the quenched and tempered structure is the result of the stress-concentration effects of the thin carbide films that are formed during the tempering of martensite.

For quenched and tempered steels the fatigue limit increases with decreasing tempering temperature up to a hardness of Rc 45 to Rc 55, depend on the steel. The fatigue properties at high hardness levels are extremely sensitive to surface preparation, residual stresses, and inclusions. The presence of only a trace of decarburization on the surface may drastically reduce the fatigue properties. Only a small amount of non-martensitic transformation products can cause an appreciable reduction in the fatigue limit. The influence of small amounts of retained austenite on the fatigue properties of quenched and tempered steels has not been well established.

The results indicate that below a tensile strength of about 200,000 psi (~1400 MPa) the fatigue limits of quenched and tempered low-alloy steels of different chemical composition are about equivalent when the steels are tempered to the same tensile strength. This generalization holds for fatigue properties determined in the longitudinal direction of wrought products. However, tests have shown that the fatigue limit in the transverse direction of steel forcing may be only 60 to 70 percent of the longitudinal fatigue limit. It has been established that practically all the fatigue failures in transverse specimens start at nonmetalic inclusions.

Nearly complete elimination of inclusions by vacuum melting produces a considerable increase in the transverse fatigue limit. The low transverse fatigue limit in steels containing inclusions is generally attributed to stress concentration at the inclusions, which can be quite high if an elongated inclusion stringer is oriented transverse to the principal tensile stress.

However, the fact that nearly complete elimination of inclusions by vacuum melting still results in appreciable anisotropy of the fatigue limit indicates that other factors may be important. Further investigations of this subject have shown that appreciable changes in the transverse fatigue limit which cannot be correlated with changes in the type, number, or size of inclusions are produced by different deoxidation practices. Transverse fatigue properties appear to be one of the most structure-sensitive engineering properties.

The existence of a fatigue limit in certain materials, especially iron and, titanium alloys, has been shown to depend on the presence of interstitial elements. The S-N curve for a pure metal will be a monotonic function with N increasing as stress decreases. The introduction of a solute element raises the yield strength and since it is more difficult to initiate a slip band, the S-N curve is shifted upward and to the right. If the alloy has suitable interstitial content so it undergoes strain aging, there is an additional strengthening mechanism.

Since strain aging will not be a strong function of applied stress, there will be some limiting stress at which a balance occurs between fatigue damage and localized strengthening due to strain aging. With enhanced strain aging, brought about by higher interstitial content or elevated temperature the fatigue limit is raised and the break in the curve occurs at a lower number of cycles. In quenched and tempered steels, which do not normally exhibit strain aging in the tension test, the existence of a pronounced fatigue limit presumably is the result of localized strain aging at the tip of the crack.