Evaluation of the possibility of improving the durability of tools made of X153CrMoV12 steel used in the extrusion of a clay band in ceramic roof tile production

Within the investigations of the mechanical properties, hardness measurements of three samples were conducted. The measurement results have been collected and presented in Figure 3. The hardness of sample 960/450 equaled 485 ± 7 HV1 (≈ 48 HRC). Also determined was the material hardness in the matrix areas devoid of large primary carbide precipitations, which equaled 414 ± 23 HV0.1 (≈ 41 HRC). For sample 1060/450, the hardness of steel NC11LV after the applied heat treatment equaled much less, i.e., 796 ± 10 HV1 (≈ 64 HRC), whereas the hardness of the matrix was 598 ± 28 HV0.1 (≈ 55 HRC). Sample 1020/200 (cut out of the material of the tool [plate]) was also characterized at very high level of hardness—about 805 HV1 ± 14 HV1 (≈ 65 HRC), with the matrix hardness reaching 700 HV1 (≈ 61 HRC). A higher matrix hardness after the classic heat treatment results from a lower degree of martensite decomposition. Clear differences were, in turn, observed in the case of the matrix of the material tempered at the same temperature but with a different austenitization temperature. This should be connected with a different degree of saturation by the alloy elements. In comparison, the commercial wear-resistant Hardox 600 steel, commonly used for machine parts in the ceramics industry, demonstrates hardness at the level of 590 HV1.

Fig. 3.

The impact strength was also determined in these studies. The Charpy V-notch values (CVN) obtained have been collected and are presented in Table 3. The results for tool steel NC11LV are presented in comparison with the results for wearresistant steel Hardox 600. For all the samples, the impact strength measurement was conducted at two temperatures, 50°C and 200°C, where the first value corresponded to the working temperature of the tool under industrial conditions, and the second value corresponded to the extreme temperatures present in the process of producing ceramic roof tiles during drying.

As we can observe, tool steel NC11LV under different states of heat treatment, the impact strength undergoes a slight change. Only in the case of sample 1020/200 can we see that the CVN coefficient is minimally lower than in the case of the other two samples, and it equals 2.8 J at the working temperature. We can conclude from the tests performed that the impact strength of steel NC11LV demonstrates stability at 200°C.

The impact strength test results unequivocally show that tool steel NC11LV hardened to the value of about 60 HRC (hardness Rockwell C) exhibits low ductility, which can favor a tendency for cracking. This is caused mainly by the presence in the structure of hard and undeformable carbides, which lower the hardness of tool steel NC11LV and at the same time significantly lower its impact strength.

As a comparison of the crack resistance of wear-resistant materials, Table 3 also includes steel Hardox 600. This steel is characterized by an impact energy equaling 21.6 J. In the case of application for plates forming a clay band, no large risk exists for critical impact loads of the examined steel NC11LV. In extreme cases, there is, however, a possibility of the plate’s breaking in the area of the notch, which is formed as a result of deep scratches in the working surfaces. Examples of such damage to the tool made of NC11LV are shown in Figure 4.

Fig. 4.

In most cases, the forming tools are exposed to mostly continuous loads resulting from the pressure of the extruded mass onto the working surface of the plate.

In the analysis of the abrasive wear of NC11LV steel, the material loss was determined based on the traces of wear obtained in the conducted ballon-disc test. An Al₂O₃ ball was used as the clamp element. The test temperature was assumed to be 50°C, which was in agreement with the actual temperature present on the clay band-tool contact in the band extrusion process during the production of ceramic roof tiles. Figure 5 presents the test results for NC11LV steel in three heat treatment variants. The results presented illustrate the effects of the ball-on-disc test on the length of 300 m together with the magnified three-track areas. We can observe that the highest wear occurs for sample 1060/450, which was quenched at 1060°C and tempered at 450°C. The mean depth of the wear tracks for this sample equals 22.8 μm. Slightly better results were obtained in the case of sample 960/450, which was quenched from a lower temperature (960°C) and also tempered at 450°C. The averaged depth of the wear track from the three examined areas for sample 960/450 equals 13.7 μm. The lowest wear in the ball-on-disc test is exhibited by sample 1020/200 after the classic heat treatment. In the case of sample 1020/200, the mean track depth is at the level of 6.7 μm. The cause of the highest wear of sample 1060/450 is the presence of numerous spallings, which were observed by means of further microscopic observations. The spallings were created especially in the presence of primary carbides. The effect of this is excessive material wear caused by spalling of the material, which, at further stages, constitutes an abradant, contributing to intensification of the abrasion wear. The effect of the carbide particles’ spallings on the increase of wear has also been confirmed by other authors [26].

Fig. 5.

In the analyzed areas of all the samples examined, we can see different levels of wear. This is connected with a non-uniform distribution of the carbide precipitates, which, in the friction process, undergo spalling in a random manner.

Figure 6 presents the track profiles obtained during the ball-on-disc tests for the materials analyzed. On the profiles, we can clearly see differences in the track depths. The track depth for sample 1020/200 is almost three times smaller than in the case of sample 1060/450 and nearly twice as low as for sample 1060/450. Large discrepancies can be seen in the case of the track widths. The track width of the sample that had worn the most is over twice as high as that in the other cases. The highest track width of sample 1060/450 additionally confirms that, in the case of this sample, material spalling took place in the extreme areas of contact between the ball and the steel disk, which intensified the material wear. Phenomena leading to the formation of numerous spallings on the surface of this sample were also observed in the microscopic tests, which are presented in subsequent sections of the article.

Fig. 6.

The most important parameters obtained in the ball-on-disc test for the analyzed materials have been compiled in Table 4. This illustrates the big differences between the wear of the particular samples.

In ball-on-disc test, the values of the friction coefficient were also determined from the test for the whole measurement line 300 m, shown in Figure 7.

Fig. 7.

On the basis of the obtained friction coefficient courses, we can observe some differences between the sample that underwent the smallest wear and the other two samples. The friction coefficient values for samples 960/450 and 1060/450 in the initial phase have a tendency to increase, and stabilization of the values can be seen after 100 m. The final value of the friction coefficient for these samples stabilizes within a range of 0.5–0.55. In the case of sample 1020/200, we can also see, just like at the first stage of the ball-on-disc test, that the friction coefficient increases for the first 50 m, and next oscillates around 0.35.

Figure 8 presents the microstructure of steel for sample 960/450. The samples for the microscopic tests were collected from the material’s core. The microstructure of the material examined was constituted of precipitates of primary carbides type M₇C₃, elongated in the direction of plastic deformation in the martensitic matrix. For an additional identification of the phases formed in the microstructure of steel NC11LV, element distributions with the EDS method were carried out for the examined area of sample 960/450. The results of this analysis have been presented in Figure 9. The element distribution obtained for the tested material shows that the chemical composition of the carbides includes chromium or vanadium. The presence of chromium in the precipitates excludes the possibility of the formation of the interstitial carbides with a simple, closed-packet structure: MC and M₂C type, which demonstrate very good melting points. The co-existence of chromium and vanadium points to the formation of the interstitial carbides with a complex hexagonal, closepacked structure of the M₇C₃ type. In this case, vanadium, similarly to molybdenum, replaces the chromium atoms, forming a multi-component carbide (Cr,V,Mo)₇C₃. It was established that even a 4% addition of vanadium affects the type of the formed carbide. At the same time, no presence of separate MC carbides was observed [27]. Analytical tests were performed for the other samples, obtaining similar element distributions.

Fig. 8.

Fig. 9.

The microstructure of sample 1060/450 showed a similar character (Figure 10). A difference can be noticed in the form of tempered martensite, which, in this case, has a more crypto-acicular character. It was observed in the microscopic examinations that, together with the application of a higher austenitizing temperature (1060°C), the number of carbides was significantly lower compared to the other samples, which were quenched from lower temperatures. In the case of sample 1020/200 (Figure 11), the matrix was constituted of low-temperature-tempered martensite. In its background, one could observe primary carbide precipitates elongated in the direction of plastic deformation.

Fig. 10.

Fig. 11.

The microscopic examinations were also conducted in the sub-surface area in order to determine the effect of the mechanical operation on the changes in the operational surface layer. For the sub-surface layer of sample 960/450, in the light microscopic image, no significant changes caused by the mechanical operation were observed in the structure (Figure 12). In the sub-surface area, we can see only slight plastic deformations along the friction direction. The carbide bands in all the tests were arranged perpendicular to the examined surface.

Fig. 12.

In the case of sample 1060/450, in the subsurface area, in the microscopic images, it is possible to see the presence of cracks propagating perpendicular to the surface the friction track (Figure 13). They were formed as a result of the mechanical operation resulting from the pressure of the ball (Hertzian stress) constituting a countersurface in the tribological tests. Cracks were observed only in the area of the friction track.

Fig. 13.

The increased material brittleness is connected with the very high hardness of steel obtained after the heat treatment. In this case, no significant changes were observed (apart from the cracks along the friction track) in the sub-surface area in the light microscopy image. Austenitizing at a higher temperature should result in a higher residual austenite content. However, the retained austenite decomposes during tempering at 450°C. For this reason, the increased hardness of 1060^°C-quenched steel should not be associated with the residual austenite content. In addition, its presence should contribute to a reduction in hardness.

In the analysis of the light microscopy images of sample 1020/200 (Figure 14) in the sub-surface area, we can see slight plastic deformations in the longitudinal direction to that of the friction test. We can observe a local loss of continuity between the carbide precipitates and the matrix, although without macroscopic material decohesion.

Fig. 14.

The investigations were also extended by an analysis of the sub-surface area with the use of scanning electron microscopy methods. Detailed analyses were performed especially of sample 1060/450, which demonstrated cracks visible in the light microscopy image (Figure 13). For this sample, it was established that the developing crack propagated further within the matrix, but the cracking could also have occurred in the case of the carbide precipitates. The cracks initiated in the matrix propagate further along it. The presence of carbides favors a change in the direction and separation of the developing crack, causing a reduction of the crack propagation speed (Figure 15). When the carbides were within the friction area, the weakened carbides underwent local cracking, which would favor spalling at the further stages. This translates to a higher wear of this sample, despite the clearly higher hardness compared to the other samples. The authors of the study [28] stated that cracks are formed as a result of cooperation of tensile and shear stresses working on the surface, and the crack itself is dependent on their orientation in respect to the abraded surface and the abrasion direction.

Fig. 15.

The dissolution of carbides in austenite depends on temperature, and the properties of tool steel after hardening are dependent on the amount of the alloy elements dissolved in the austenite. The SEM images obtained from the material austenitized at a higher temperature (sample 1060/450), which proves that the higher austenitization temperature enabled the initiation of dissolution (partial melting) of the carbides contained in the microstructure (Figure 16). This resulted in an increase of the solution strengthening of the formed martensite, which translated to a much higher hardness of the matrix and thus the whole material. A side effect was a simultaneous increase of material brittleness. The increase in brittleness was conducive to easier spalling of the surface during tribological interactions. This ultimately translated into greater wear, despite the higher hardness of the steel under these conditions.

Fig. 16.

The data in the literature show that a temperature of 1120°C is sufficient for the dissolution of M₇C₃-type carbides. During the heating of steel until tempering, the M₇C₃ carbides are dissolved first, while the thermally more stable M₆C and MC carbides hinder the grain growth. It has been established that, in the case of chromium carbides Cr₇C₃, the melting point is within a range of 900°C to 950°C [29]. Vanadium increases the melting point by about 90°C.

The data in the literature demonstrate that, during prolonged annealing, M₇C₃-type carbides can also undergo transformation into M₂₃C₆ carbides by way of diffusion. Wieczerzak et al. [18] stated that the transformation takes place at high temperatures owing to the diffusion of iron into the matrix and chromium from the matrix to the peripheral regions. This is attributed to the fact that the peripheral regions are closer to the matrix, with the result that the diffusion occurs easier. In consequence, the transformation takes place mostly in the peripheral regions of the M₇C₃ type carbides [30]. In the meantime, Kondrat’ev et al. [17] suggest that the transformation of M₇C₃ into M₂₃C₆ begins from the core (the center) of the M₇C₃ carbides. Regardless of the mechanism, however, it has been established that during the 100 h of annealing at 1050°C, a complete transformation of M₇C₃ to M₂₃C₆ can take place [31]. The conducted microscopic SEM observations of the sample in state 1060/450 did not show any clear inhomogeneity, which would point to the occurrence of a transformation (Figure 17).

Fig. 17.

The microscopic SEM analysis performed in the direction transverse to the friction direction of the remaining sample (Figure 18) demonstrated the formation of surface irregularities, which point to the occurrence of the ridging phenomenon.

Fig. 18.

The visible traces of spallings in the magnified image are the effect of the high hardness of the carbides, whereas the detached fractions are smaller than in the case of samples 960/450 and 1060/450, which can be observed in the measurements of the material loss in the ball-on-disc test.