**4. Failure mechanism: post-creep characterization of SPCT samples**

Scanning electron microscopy (SEM) images of fractured SPCT specimens for different conditions are shown in **Figure 10**. Radial cracks can be observed in all the TMT samples (**Figure 10a**–**c**). This is an evidence of the loss of ductility and indicates a brittle fracture, which is a change in rupture ductility in comparison to the conventionally treated sample. Those samples do not show radial cracks (**Figure 10d**). Besides, a higher reduction in thickness is evident in the conventionally treated sample in comparison to the TMT ones, suggesting a ductile fracture behavior.

To clarify the failure mechanisms, the fractured samples were cut and prepared adequately. **Figure 11(a)** and **(b)** shows the SEM images for the TMT samples ausformed at 600°C with a deformation of 20% and ausformed at 900°C with a deformation of 40%. It is worth noting in those images the existence of cavities nearby coarse particles, which are located at the vicinity of PAGBs. The EDS spectrum shown in **Figure 11(c)** allows us to conclude that these particles are M23C6 carbides with M = (Fe, Cr, Mo).

The greater size of the M23C6 carbides at the vicinity of PAGB contributes to the inhomogeneous and localized deformation experienced by the TMT samples at these locations during creep. The local creep concentration close to PAGB would be promoting the nucleation of cavities that lead to the intergranular fracture with the brittle behavior.

**Figure 12** shows different inverse pole figure (IPF) maps for all the samples under study before and after SPCT. It should be pointed out that, contrary to the lath boundaries that are not correctly indexed due to the step size used for the EBSD mapping, the block boundaries before and after SPCT are clearly disclosed. It is observed that the microstructures of the samples exhibit the characteristic lathlike morphology of the martensitic microstructure. However, such morphology is

#### **Figure 10.**

*Scanning electron microscopy images of the SPCT fracture surfaces for samples tested at 700°C with a load of 200 N: (a) G91-TMT 900\_20; (b) G91-TMT 600\_20; (c) G91-TMT 900\_40; and (d) G91-AR [39].*

*Welding - Modern Topics*

and 1020 m−<sup>3</sup>

*and 900°C [41].*

**Figure 8.**

resistance is.

, respectively [61]. It might be concluded that ausforming promotes a

refining of precipitates, up to five times as compared with conventional processing, as well as an increase in number density up to two orders of magnitude. In fact, these number densities and precipitate sizes are very similar compared to those corresponding to oxides present in oxide dispersion-strengthened (ODS) steels [66, 67]. The elevated number density of nanosized MX precipitates has a direct impact on creep response of this material as it can be clearly observed in **Figure 9**. This figure shows characteristic SPCT curves at 200 N, exhibiting the variation of specimen deflection with time. It might be concluded from this figure that introducing an ausforming step improves the δd significantly, and most precisely, the lower the ausforming temperature, the lower the δd is, and, hence, the better the creep

*(a) MX carbonitrides (white arrows) within laths after thermomechanical treatment ausformed (20%) at 900°C; (b) size distribution of MX precipitates in the TMT samples for the two ausforming temperatures: 600* 

**176**

**Figure 9.**

*SPCT curves for all samples tested at 700°C with a load of 200 N [39].*

#### **Figure 11.**

*Scanning electron microscopy images: (a) M23C6 precipitates located at a prior austenite grain boundary in sample G91-TMT 600\_20. The prior austenite grain boundary in this image has been highlighted with a dash line as a guide to the eye; (b) cavities associated with coarse M23C6 precipitates have nucleated at a prior austenite grain boundary in sample G91-TMT 900\_40. Cavities have been pinpointed with arrows and (c) EDS analysis of the particle marked with a red arrow in image (a), close to a cavity [39].*

blurred in samples ausformed at 600°C because of the high deformation accumulated in the austenite during ausforming [38]. After the SPCT, it is observed that the original lath-like morphology has partially disappeared, and it has evolved towards a fine-grained equiaxed ferritic matrix. One might conclude from the microstructural observations made after SPCT that newly formed equiaxed grains are distributed homogeneously in the conventionally treated sample (AR), while these grains are located mainly nearby the prior austenite grain boundaries in the TMT samples, which is consistent with the fact that it is in these samples where the deformation accumulated is larger during creep.

Therefore, taking into account the results shown previously in the SEM micrographs (**Figure 11(a)** and **(b)**), the microstructural degradation would be a combined consequence of the accumulation of dislocations at the low-angle boundaries and the stress concentration close to the coarse M23C6 carbides, which lead to the progressive loss of the lath-like martensitic microstructure, which evolves to an equiaxed ferritic matrix. As it has been discussed above in the case of the TMT samples, the nucleation of cavities takes place close to M23C6 precipitates located at the prior austenite grain boundaries. The coalescence of the cavities formed surrounding the M23C6 carbides would initiate the cracks, and they will propagate along the prior austenite grain boundaries.

The more homogeneous distribution of the M23C6 precipitates in the conventionally treated sample favors the apparition of equiaxed grains in the whole martensitic matrix and develops the nucleation of cavities intragranularly, which provokes the transgranular fracture. Besides, in the TMT samples, the high austenitization temperature produces an enormous prior austenite grain sizes with concomitant large grain boundary surfaces, facilitating an earlier formation of the critical crack length that causes the brittle fracture [68, 69].

**179**

**5. Conclusions**

*post-SPCT microstructures [39].*

**Figure 12.**

an important drop in ductility.

*High-Chromium (9-12Cr) Steels: Creep Enhancement by Conventional Thermomechanical…*

**Effect of austenitization temperature:** compared to the conventional heat treatments, the use of a higher austenitization temperature (1225°C rather than 1040°C), combined with an ausforming processing step at 900°C, allows the increase of the number density of MX precipitates up to three orders of magnitude after the tempering step, which raises the strengthening capability of the MX at 700°C up to 6.5 times. These microstructures have reduced considerably the minimum disk deflection rate and showed greater time to rupture during the SPCT carried out at 700°C. By contrast, such elevated austenitization temperature induces

*Representative inverse pole figure (IPF) maps of the initial and after SPCT microstructures: (a) G91-AR; (b) G91-TMT 900\_20; (c) G91-TMT 600\_20; and (d) G91-TMT 900\_40. Cavities are in white in* 

**Effect of ausforming:** the SPCT was applied to evaluate the creep behavior of G91 steel after different TMT and heat treatments. The minimum disk deflection rate was lower, and the time to rupture was longer for G91 after the TMT than with the conventional G91 heat treatment (AR). The improvement in creep rupture strength is attributed to the fine and homogeneous distribution of MX carbonitrides. The number density and average precipitate size of MX carbonitrides after the TMT are similar to the oxide particles in ODS steels. These latter steels possess high creep strength due to the high number density of oxides distributed in the

*DOI: http://dx.doi.org/10.5772/intechopen.91931*

*High-Chromium (9-12Cr) Steels: Creep Enhancement by Conventional Thermomechanical… DOI: http://dx.doi.org/10.5772/intechopen.91931*

#### **Figure 12.**

*Welding - Modern Topics*

**Figure 11.**

blurred in samples ausformed at 600°C because of the high deformation accumulated in the austenite during ausforming [38]. After the SPCT, it is observed that the original lath-like morphology has partially disappeared, and it has evolved towards a fine-grained equiaxed ferritic matrix. One might conclude from the microstructural observations made after SPCT that newly formed equiaxed grains are distributed homogeneously in the conventionally treated sample (AR), while these grains are located mainly nearby the prior austenite grain boundaries in the TMT samples, which is consistent with the fact that it is in these samples where the deformation

*Scanning electron microscopy images: (a) M23C6 precipitates located at a prior austenite grain boundary in sample G91-TMT 600\_20. The prior austenite grain boundary in this image has been highlighted with a dash line as a guide to the eye; (b) cavities associated with coarse M23C6 precipitates have nucleated at a prior austenite grain boundary in sample G91-TMT 900\_40. Cavities have been pinpointed with arrows and (c) EDS analysis of the particle marked with a red arrow in image (a), close to a cavity [39].*

Therefore, taking into account the results shown previously in the SEM micrographs (**Figure 11(a)** and **(b)**), the microstructural degradation would be a combined consequence of the accumulation of dislocations at the low-angle boundaries and the stress concentration close to the coarse M23C6 carbides, which lead to the progressive loss of the lath-like martensitic microstructure, which evolves to an equiaxed ferritic matrix. As it has been discussed above in the case of the TMT samples, the nucleation of cavities takes place close to M23C6 precipitates located at the prior austenite grain boundaries. The coalescence of the cavities formed surrounding the M23C6 carbides would initiate the cracks, and they will propagate

The more homogeneous distribution of the M23C6 precipitates in the conventionally treated sample favors the apparition of equiaxed grains in the whole martensitic matrix and develops the nucleation of cavities intragranularly, which provokes the transgranular fracture. Besides, in the TMT samples, the high austenitization temperature produces an enormous prior austenite grain sizes with concomitant large grain boundary surfaces, facilitating an earlier formation of the critical crack

**178**

accumulated is larger during creep.

along the prior austenite grain boundaries.

length that causes the brittle fracture [68, 69].

*Representative inverse pole figure (IPF) maps of the initial and after SPCT microstructures: (a) G91-AR; (b) G91-TMT 900\_20; (c) G91-TMT 600\_20; and (d) G91-TMT 900\_40. Cavities are in white in post-SPCT microstructures [39].*
