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Indice Saggio
0. Abstract
1. Introduction
2. Experimental procedures
3. Results and discussion
4. Conclusions
5. References


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Novità:
L'Alto Adige/Suedtirol dalla caduta dell'Impero Romano all'avvento di Carlo Magno - Damiano Martorelli
Libro disponibile su: La Feltrinelli e Unilibro
   

Lubricated Rolling-Sliding damage
in Powder Metallurgy alloys

by prof.ing. Giovanni Straffelini, dr.ing. Damiano Martorelli
Department of Materials Engineering, University of Trento, Via Mesiano 77, 38122 TRENTO (Italy)
E-mail: Giovanni.Straffelini@ing.unitn.it

3.1 Results and discussion

An examination of the surface of specimen Fe06C after 4.6 × 105 cycles actually revealed the presence of large surface pits. Figure 3.1 shows a top view of the specimen surface with the presence of a pit. In addition, the presence of surface plastic flow, which produced an almost complete closure of surface pores can be clearly appreciated. Not all the specimens revealed the presence of such macroscopic damage before 106 cycles. It was found in the as-sintered materials only: after 4.6 × 102 cycles for Fe06C and 6.5 × 105 cycles for AMoO5C.

Surface Pit in material Fe06C
Figure 3.1 Surface pit in materiai Fe06C after 4.6 105 cycles.

Spalling cracks nucleated both at the surface or subsurface regions and propagated parallel to the surface at the Hertzian depth. Crack nucleation in the subsurface regions was favoured by the presence of pores, whereas crack nucleation at the surface was helped by the formation of surface microcracks by plastic shearing at the asperities, as shown in Figure 3.2a and b. Plastic shearing at the asperities was actually present in all the materials, even if in the case hardened materials its intensity was very low. During rolling-sliding, in fact, a mixed lubrication regime was established and the ripetitive metal-to-metal contacts induced a local continuous accumulation of plastic deformation by ratchetting, [8].

Figure 3.2b also shows that material Fe06C underwent surface densification. This type of damage involved both the as-sintered materials. It was simply due to the high contact pressure, which induced a large plastic deformation of the matrix of the porous materials. On the contrary, the surface treated materials did not experienced such a densification, because of their high matrix hardness. Densification in the as-sintered materials had two effects. From one side it increased material strength by reducing porosity. From the other side it helped in the propagation of fatigue cracks since it triggered an oil hydraulic effect. Such effect was actually absent in the materials which did not densify, since in those cases oil was allowed to penetrate the interior of the material through the interconnected porosity.

Formation of surface cracks in material Fe06Ca)
Formation of subsurface cracks in material Fe06Cb)
Figure 3.2 Formation of surface (a) and subsurface (b) cracks in material Fe06C.

All materials were also characterized by the presence of microcracks at the pore edges. Two types of microcracks were highlighted. The materials which underwent surface densification showed the presence of numerous microcracks at the pores located at the boundary between the densified and undensified region. Material AMo-CN did not show the presence of subsurface microcracks, whereas material Fe-CN was characterized by the presence of a similar population of microcracks at a depth between 550 and 1000 pm from the surface only. Figure 3.3 shows an example of subsurface microcracks in material Fe-CN. In addition, the two carbonitrided materials also showed the presence of microcracks at the pores located at the surface which, as it will shown also later, have a different origin from those located in the subsurface regions.

Subsurface microcracks in material Fe-CN
Figure 3.3 Subsurface microcracks in material Fe-CN at a depth of 800 μm

The formation of subsurface microcracks play clearly an important role in the macroscopic spalling observed in the as sintered materials. In fact, the subsurface microcracks may act as nuclei for a macroscopic crack (in this case we have subsurface crack nucleation) or help in the crack propagation stage. The propagating crack, in fact, finds an energeticaily favourable path in following the pores with the pre-existings microcracks [9].

In order to achieve information on the origin of the subsurface microcracks it is convenient to estimate the 'local equivalent stress', σloc, acting at the pore edges. This can be calculated using the following equation:

σloc = σeq × βk/ Φ

where σeq is the von Mises equivalent stress which can be calculated using the Smith and Liu equations [10], Φ is the fraction of load-bearing section (0=0.6 for a porosity of about 10% [11]) and βk is the notch effect coefficient, due to the pores, which depends on the pore shape and matrix microstructure. Following Pohl [12] and Kubicki [13], βk can be set to 1.9 for the as sintered materials and 2.4 for the carbonitrided materials.

The local equivalent stress can be hence compared to the matrix yieid strength, in order to have an estimation of the probability to have a microcrack nucleation at the pore edges [14]. The matrix yield strength σy0 can be obtained from the microhardness values HV, since σy0 = HV / B, where B is a constant which can be set to 4.5 for the as sintered materials and 4.2 for the others [15].

Comparison of equiv. stresses with the matrix yield strength profiles for as sintered materialsa)
Comparison of equiv. stresses with the matrix yield strength profiles for carbonitrided materialsb)
Figure 3.4 Comparison or the local equivalent stresses with the matrix yield strength
profiles for the as-sintered materials (a) and the carbonitrided materials (b).

The results of the calculations are shown in Figure 4a, for the as sintered materials, and 4b, for the carbonitrided materials. The observation of the figures shows that in the as sintered materials the equivalent stress at the pore edges is higher than the matrix yield strength down to a depth of 500-700 ptm. In particular, the difference between the local equivalent stress and the matrix yield strength is maximum at a depth of 120 ptrm, where the fatigue crack actually was found to propagate. In these materials microcracking at the pore edges can thus easily occur and the local applied stress is of such an integrity as to induce the formation and propagation of a fatigue macrocrack, which leads to spalling.

For material AMo-CN, the local equivalent stress is always lower than the matrix yield strength, in agreement with the fact that both microcracking and spalling are not experimentally observed. For material Fe-CN, the local equivalent stress becomes higher than the matrix yield strength only at a depth between about 600 and 1000 gm. This actually explain the observed formation of microcracks at the pore adges at this depth range.

This analysis is thus able to correctly rationalize the formation of microcracks at the pore edges and also the formation of macroscopic damage by spalling in the as sintered materials. It is, however, not able to explain the formation of the microcracks at the surface pores in materials Fe-CN and AMo-CN, since at the surface regions the local equivalent stress is always higher than the matrix yield strength. A different mechanism is thus responsible for this and it can be proposed that these microcracks are formed by asperity-scale fatigue [16]. In porous materials, in fact, surface pores can behave as cracks if the matrix toughness is low [17]. The formation of the surface fatigue cracks is therefore due to the combination of the presence of a surface tensile (friction) stress and a low matrix fracture toughness, because of the surface hardening treatment.

 
     
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