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Direct observations have been made of the emission and distribution of dislocations at the crack tip of propagating cracks in various metals during tensile and cyclic deformation in an electron microscope. The results show a number of new findings regarding the dislocation behavior near the crack tip and its relationship to the propagation of cracks. Depending on the loading geometry, it was possible to observe the propagation of shear cracks of Mode III type as well as tensile cracks of Mode I type. Some of the observations were made using a video recording system. For shear cracks of Mode III type, the crack tip generated screw dislocations on planes which were coplanar with the cracks. The dislocations were generally in the form of an inverse pileup but a dislocation-free zone was present near the crack tip. For tensile cracks of Mode I type, edge dislocations were generated at the crack tip on planes which were inclined to crack. As the dislocations were emitted, the crack tip was blunted and crack propagation ceased. Under cyclic loading conditions, some of the dislocations near the crack tip reversed their direction and returned to the crack tip and disappeared as the stress was reduced. In order to explain the physical origin of the dislocation-free zone near the crack tip, the dislocation theory of fracture was extended and the dependence of the stress intensity factor on the number of dislocations in the plastic zone was calculated. It was found that the dislocation-free zone was expected if there was a critical stress intensity factor required for the generation of dislocations at the crack tip. The magnitudes of the critical stress intensity factor for dislocation generations were estimated for various metals from the dislocation theory and these values were compared with the values determined from the electron microscope fracture experiments. The problem of a linear array of dislocations on a plane which is inclined to the crack plane was also considered for Mode III geometry.
Several studies have examined the FCG behavior of 6-4 brass (α-β brass). Murakami et al. [8] performed FCG experiments using several 6-4 brasses with different values of d to study the effects of d and R in the low-FCG velocity range. They reported that in the high FCG region, multiple slip systems operated simultaneously ahead of the fatigue crack tip before fatigue crack growth, indicating that the slip surface separation mechanism operated. They also reported striations on the fracture surfaces. In the low-FCG region, fatigue cracks were developed by a single slip system because the slip systems that could operate were limited. The fatigue fracture surface exhibited morphologies, such as cleavage, grain boundary facets, and striped patterns. Therefore, in the low-FCG region, the influence of d and R on the FCG velocity becomes significant. According to Sugeta et al. [9], in the low-FCG velocity range, owing to the effects of grain boundaries and repeated strain hardening, the slip system that had been operating until it was suppressed, and other slip systems began to operate, leading to zigzag FCG behavior. The above observations were similar to those for 7-3 brass [6]. Lin et al. [10] reported that the effect of f on the FCG properties of 6-4 brass materials was affected by R. It was also shown that the effect of specimen thickness on the FCG behavior was affected by the corrosive environment [11].
As aforementioned, several studies have been conducted on the fatigue behavior of Cu and Cu alloys (pure Cu, 7-3 brass, 6-4 brass). However, many of these studies are concerned with the FCG behavior of long through-cracks, whereas there seems to be a lack of examples of the FCG behavior of short surface fatigue cracks.
The specimen shown in Figure 4c was used for the FCG experiment. The observation area was narrowed to facilitate the measurement of the short surface fatigue crack. That is, a constriction with a depth of 0.25 mm and a radius of 12 mm was provided in the parallel portion of the smooth specimen. Additionally, it was designed so that short surface cracks would occur there. The stress concentration factor of the constricted part was 1.05.
Hatanaka et al. [29] reported that in Cu and 7-3 brass (α-brass), slips generated within the crystal grains accumulate at the grain boundaries, causing fatigue cracks. Similar results have been reported for Cu and Cu alloys [5,7]. Therefore, in free-cutting brass, the fatigue crack generation mechanism is different from that of other Cu alloys because Bi or Pb that are added to improve the free-cutting property cause cracks.
Regarding the FCG behavior in the low ΔK region, the difference between Bi,A and Bi,B materials was considered. As shown in Figure 1, the α/β-phase boundaries in the Bi,B materials are more strongly bonded (in comparison to the Bi,A materials) and intermesh with each other. In the low ΔK region of the Bi,B material, fatigue cracks tend to avoid the strong α/β-phase boundaries (Figure 11d) and propagate inside the β phase (Figure 11c). Consequently, the FCG velocity of the Bi,B material is lower than that of the Bi,A material.
A detailed observation of Figure 12b revealed that the da/dN values of Bi,B, and Pb were approximately the same. In addition, in the region below M = 4 MPam1/2, the da/dN of the Bi,A material was faster than that of the Bi,B, and Pb materials, indicating that the FCG resistance of the Bi,B, and Pb materials was higher than that of the Bi,A materials. This is because, in the Bi,B materials, fatigue cracks propagate more frequently within the β phase than in the Bi,A materials (Figure 11b,c). The hardness of the β phase (HV 187, body-centered cubic lattice) was higher than that of the α phase (HV 151, face-centered cubic lattice), and their crystal structures were different. Therefore, it is assumed that the FCG resistance of the β phase is higher than that of the α phase. 2ff7e9595c