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The mechanical behavior of the fully austenitic matrix of high-chromium cast steel (HCCS) alloy is determined by external compression stress applied at 300 and 700 C. The microstructure is roughly characterized toward both optical and scanning electron microscopy analyses. Dilatometry is used during heating from room temperature up to austenitization to study the solid-state phase transformations, precipitation, and dissolution reactions. Two various strengthening phenomena (precipitation hardening and stress-induced bainite transformation) and one softening mechanism (dynamic recovery) are highlighted from compression tests. The influence of the temperature and the carbide type on the mechanical behavior of the HCCS material is also enhanced. Cracks observed on grain boundary primary carbides allow establishing a rough damage model. The crack initiation within the HCCS alloy is strongly dependent on the temperature, the externally applied stress, and the matrix strength and composition.

Numerous studies have been undertaken to understand the work roll behavior during hot rolling in order to prevent or to better control both surface and internal damage. These works are related to both experimental and modeling studies, and different damage phenomena are involved, such as banding failure, thermal fatigue, fire cracking, wear, etc. (Ref 10). In addition, residual stress distribution within the work roll in service is also taken into account in order to reach accurate simulations (Ref 11).

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The studied alloy is an HCCS grade originating from the shell material of a work roll obtained from a vertical spin casting process and dedicated to the roughing stands of an HSM. The work roll has an outer diameter of 1200 mm and a shell depth of 80 mm. The average chemical composition of the HCCS alloy as obtained from optical emission spectroscopy is given in Table 1.

Microstructure characterization is carried out at different stages of the work using various techniques such as optical microscopy (OM), scanning electron microscopy (SEM), and Vickers hardness measurements.

The CT300-B samples exhibit bainite with oriented sheaves in their matrix, and either martensite or some retained austenite in addition (Fig. 8a). The crack path in the CT300-B clearly looks more complex and more extended than the previous one observed in CT300-A, due to the branching that occurs within the cracks (Fig. 8b). But the primary cracks remain parallel to the stress direction. Similar to CT300-A, evidence of numerous intragranular secondary carbides is enhanced in the optical micrograph after OPS preparation (Fig. 8c). In addition, the corrosion pits observed on the CT300-B are probably due to the infiltration of the nital etching into the large cracks that exist within the primary carbides (Fig. 8a). PFZs still exist on the CT300 samples (Fig. 7c and 8c).

While all the compressed samples seem to exhibit preferential oriented bainitic sheaves in various amounts and sizes, the stress-free SF300 sample contains large quantities of sheave-like bainite structures with haphazard oriented directions inside the grains (Fig. 10a). In addition, the grain boundary carbides do not exhibit cracks as no external stress was applied. Widespread secondary carbides are present inside the grains with PFZs around them (Fig. 10b), such an observation being similar to the previous ones made on the compressed samples.

Primary carbides are made of Cr-rich M7C3 carbides and Mo-rich M2C carbides, which are located at grain boundaries. But only the M7C3 carbides, which represent the majority of the grain boundary carbides, contain more or less cracks that are related to the total elongation achieved under the external compression stress applied (Fig. 7b, 8b, 9b). In addition, the maximum elongation under compression at 300 C seems to be related to the progress of the internal damage, as the CT300-A samples exhibit a total elongation lower than that achieved on the CT300-B samples.

Focusing on the crack path, it is made of primary carbides containing cracks parallel to the external compression stress, which correspond to the beginning of the internal damage process (CT300-A samples, Fig. 7b). With increasing strain, the crack path changes to a more complex oriented one (CT300-B samples, Fig. 8b). Therefore, it could be assumed that the primary cracks parallel to the external stress correspond to the initiation of the internal failure mechanism, which is in good agreement with Dünckelmeyer et al. (Ref 42). They found that, for similar carbides, primary cracking is always parallel to the loading direction. The crack occurs when the normal stress of the carbide induced by the plastic deformation of the surrounding matrix is exceeded. This can be achieved either from dislocation pileups on the grain boundary carbides (Ref 43, 44), or due to micro-stresses resulting from differences between the thermal and the elastic properties of the matrix on the one hand, and of the eutectic carbides on the other (Ref 45).

In addition, Margolin et al. (Ref 46) in an earlier study that focuses on brittle fracture mechanisms set criteria for crack nucleation in the form of micro-discontinuities within the material under external stress. In most cases, the crack preferentially nucleates on a grain boundary carbide because the strength of such a phase is lower than the strength of the carbide/matrix interface. The micro-crack occurs in a carbide when the local stress at the head of the dislocation pileup exceeds the strength of the carbide, while remaining lower than the more resistant carbide/matrix bound at the same time. The same authors also found that the crack nuclei change from carbide micro-cracking to delamination at the carbide/matrix interface, only if higher stresses are reached for the same loading. Then a pore is formed at the carbide/matrix bound due to the existence of anti-parallel dislocation pileups within the interface, due to the higher local stresses. Moreover, Kroon and Faleskog (Ref 44) in a more recent work on ferritic steels established that the cleavage fracture is often initiated on brittle carbide as a result of a fiber loading mechanism in which the stress levels in the carbides are raised while the surrounding ferrite undergoes plastic deformation.

The other cracks especially observed on the CT300-B samples, which are oriented in directions different from that of the external stress, are probably ascribed to the progress of the internal damage process with the micro-plastic strain evolution, which allows the occurrence of other critical shear directions within the carbides.

In conclusion, the primary grain boundary M7C3 carbides serve as preferential crack initiation sites under external compression stress. The cracking of the carbides occurs due to their brittleness as they do not allow plastic deformation contrary to the surrounding matrix.

Undissolved tertiary carbides significantly strengthen the HCCS alloy under compression stress at 300 C, in such a way that the cracking that was initiated on grain boundary primary carbides is delayed. The secondary carbides probably increase the matrix hardness but have little influence on the strengthening mechanism itself, regardless of the test temperature. Primary carbides promote crack initiation under external stresses.

In conclusion, although high plastic deformation occurs within the grains associated with significant total elongation, only a limited shearing phenomenon occurs on grain boundaries, leading to short cracks within primary carbides. The final result is a smaller fracture strain at 700 than at 300 C.

When the compression stress is applied at 300 C on an austenitic matrix containing undissolved tertiary carbides, high work hardening rates can be expected, as these ultrafine and coherent carbides may pin the dislocations while at the same time impeding their glide toward the grain boundaries. Undissolved tertiary carbides strengthen the matrix in such a way that increased compression stresses are necessary to yield sufficient micro-plastic deformation and new dislocation generation within the grains. In addition, the SIBT which occurs will also account for the strengthening phenomenon. When dislocations are not pinned by intragranular carbides, then the dislocation glide may be possible in the direction of the grain boundaries where primary carbides are located. The presence of PFZs close to the grain boundaries may also promote dislocation gliding within these regions. Subsequent pileup of the dislocations on the primary carbides leads to a shearing process which in turn yields sharp micro-cracks as these species are known to be very brittle. The crack initiation within the grain boundary carbides will then follow the direction of the externally applied stress.

When the compression stress is applied at 300 C on a supersaturated austenitic matrix where tertiary carbides have been dissolved, the matrix is softer and it can undergo significant plastic strain under limited externally applied stress. The increase in the compression stress leads to an increase in the dislocation density within the grains. Due to the weak presence of fine precipitates having high coherency with the matrix, such as undissolved tertiary carbides, most of the newly generated dislocations are free to move toward grain boundaries where they will pileup on primary carbides. The crack initiation process will then occur in the same manner as for the HCCS material containing undissolved tertiary carbides, with sharp micro-cracks parallel to the externally applied stress. Similarly, SIBT may also occur under stress. Nevertheless, the applied stress necessary to initiate cracks within the grain boundary carbides will be many times lower than that of the same material, which contains undissolved tertiary carbides. Moreover, the crack path inside the softer material will quickly change to a more complex and branched one. Such a complex crack path will contain, in addition to the large cracks parallel to the direction of the external stress, other micro-cracks which are not parallel to the external stress. 2ff7e9595c