**3.2 Effect of austenitization temperature**

The effect of austenitization temperature on the temper microstructure of G91 steel is analyzed in this section. **Figure 4** schematically illustrates the two alternative processing routes considered:


#### **Figure 3.**

*(a) Resulting hierarchy microstructure achieved by conventional heat treatment; (b) and (c) SEM micrographs of the as-received state; (d) and (e) TEM micrographs. Arrow heads point out the location of the M23C6 carbides on lath boundaries and MX carbonitrides within the laths [37].*

**171**

**Figure 4.**

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

For the sake of comparison, **Figure 4** also includes the industrial manufacturing conditions for G91 steel named as-received (AR) condition. The goal of exploring the effect of austenitization temperature on the microstructure is to enhance the precipitation of nanoparticles during the subsequent tempering stage indicated in **Figure 4**. As it was mentioned above, the main cause for creep softening in conventional G91 is due to the recovery of the martensitic lath microstructure because of mechanisms, such as the dislocation movement, controlled by diffusion [12, 27]. The dislocation pinning by nanosized MX precipitates can delay this phenomenon, since they present an enhanced ripening resistance [8, 52–54]. The goal of undergoing such elevated temperatures in the HAT treatment as compared to conventional austenitization heat treatment (AR treatment) is to dissolve all the primary carbides in the microstructure and drive to solid solution all the potential carbide former elements. Therefore, the martensite formed after quenching from such elevated austenitization temperature keeps in solid solution most of the precursor elements of MX carbides (M = Nb, V; X = C, N) that might precipitate during the subsequent tempering. It is important to consider that the austenitization temperature has to be high enough to eliminate as much as possible the primary carbides formed during the casting process, but lower than the delta ferrite formation temperature, in order to avoid the detrimental effect of this phase from a long-term creep property point of view. Computational thermodynamic calculations by means of Thermocalc® determine the optimum austenitization temperature in 1225°C (**Figure 5**).

The interest of TMT relies on the role that austenite deformation has on refining the martensitic microstructure [55, 56]. Depending on the deformation temperature, several are the mechanisms that affect the austenite microstructure, and hence, that could be transferred to the martensite upon quenching. If deformation temperature is above the non-recrystallization temperature, the freshly formed austenite microstructure will present a significantly reduced grain size that would induce the concomitant martensitic microstructural refinement. Similarly, by applying plastic deformation to the austenite at temperatures below the nonrecrystallization temperature, which is the so-called ausforming processing [57],

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

*Thermomechanical treatments investigated in this study [40].*

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

**Figure 4.** *Thermomechanical treatments investigated in this study [40].*

For the sake of comparison, **Figure 4** also includes the industrial manufacturing conditions for G91 steel named as-received (AR) condition. The goal of exploring the effect of austenitization temperature on the microstructure is to enhance the precipitation of nanoparticles during the subsequent tempering stage indicated in **Figure 4**. As it was mentioned above, the main cause for creep softening in conventional G91 is due to the recovery of the martensitic lath microstructure because of mechanisms, such as the dislocation movement, controlled by diffusion [12, 27]. The dislocation pinning by nanosized MX precipitates can delay this phenomenon, since they present an enhanced ripening resistance [8, 52–54]. The goal of undergoing such elevated temperatures in the HAT treatment as compared to conventional austenitization heat treatment (AR treatment) is to dissolve all the primary carbides in the microstructure and drive to solid solution all the potential carbide former elements. Therefore, the martensite formed after quenching from such elevated austenitization temperature keeps in solid solution most of the precursor elements of MX carbides (M = Nb, V; X = C, N) that might precipitate during the subsequent tempering.

It is important to consider that the austenitization temperature has to be high enough to eliminate as much as possible the primary carbides formed during the casting process, but lower than the delta ferrite formation temperature, in order to avoid the detrimental effect of this phase from a long-term creep property point of view. Computational thermodynamic calculations by means of Thermocalc® determine the optimum austenitization temperature in 1225°C (**Figure 5**).

The interest of TMT relies on the role that austenite deformation has on refining the martensitic microstructure [55, 56]. Depending on the deformation temperature, several are the mechanisms that affect the austenite microstructure, and hence, that could be transferred to the martensite upon quenching. If deformation temperature is above the non-recrystallization temperature, the freshly formed austenite microstructure will present a significantly reduced grain size that would induce the concomitant martensitic microstructural refinement. Similarly, by applying plastic deformation to the austenite at temperatures below the nonrecrystallization temperature, which is the so-called ausforming processing [57],

*Welding - Modern Topics*

parallel close-packed plane relationship in the Kurdjumov-Sachs (K-S) orientation relationship, denoted as "CP group" recently. In general, a packet is partitioned into several blocks, each of which contains laths of a single variant of the K-S relationship. Blocks and packets are mostly surrounded by high-angle boundaries, whereas

The microstructure resulting from conventional industrial heat treatment consists of tempered martensite, which presents elongated subgrains with an average size of 0.25–0.5 μ m (**Figure 3**). Two types of precipitates, M23C6 carbides rich in chromium and MX carbonitrides rich in V or Nb, are present in the microstructure. The size of M23C6 carbides is around 100–200 nm, and they are precipitated on subgrain boundaries and prior austenitic grain boundaries. The size of MX carbonitrides is much smaller than M23C6 carbides, 20–50 nm, and they are in the matrix [51]. The purpose of this work is to produce a dispersion of nanosized precipitates by a controlled TMT, bearing in mind that a high number density of fine MX precipitates (Nb-MX and V-MX) should display superior high-temperature performance.

The effect of austenitization temperature on the temper microstructure of G91 steel is analyzed in this section. **Figure 4** schematically illustrates the two alterna-

• High austenitization temperature (HAT): In order to achieve an almost complete solid solution in austenite of most of the potential MX precipitate formers, the austenitization condition set will imply an elevated temperature.

• Thermomechanical treatment: The combined effect of the elevated austenitization temperature and a subsequent deformation will be studied with the aim of optimizing the MX-nanoprecipitate distribution during tempering of the

*(a) Resulting hierarchy microstructure achieved by conventional heat treatment; (b) and (c) SEM micrographs of the as-received state; (d) and (e) TEM micrographs. Arrow heads point out the location of the M23C6*

*carbides on lath boundaries and MX carbonitrides within the laths [37].*

lath boundaries inside a block are of low-angle type.

**3.2 Effect of austenitization temperature**

tive processing routes considered:

martensitic microstructure.

**170**

**Figure 3.**

#### **Figure 5.**

*Temperature evolution of phase mole fraction in G91 calculated by Thermocalc® [38].*

an austenitic microstructure with a high population of deformation bands will be formed. This would directly induce the preferential formation of some specific martensitic variants upon austenite transformation (martensite variant selection), leading to the development of strong transformation texture.

**Figure 6** illustrates the IPF maps, SEM and TEM micrographs after HAT and TMT processing routes, and the reference (AR) condition. The first conclusion obtained is the coarsening of the block size (white arrows in **Figure 6**) in HAT and TMT conditions as compared with AR condition, because of the high austenitization temperature. Block widths of 2.7 ± 0.2 μm for AR condition were obtained; meanwhile, values of 4.12 ± 0.37 μm for HAT and 3.21 ± 0.27 μm for TMT were measured. The coarser the prior austenite grain, the coarser the block size. However, it is worth noting that finer block size is observed after TMT than with HAT, which is consistent with the fact that thermomechanical processing increases the low-angle substructure and decreases the block size of as-quenched martensite.

The dislocation density after HAT and TMT was measured by XRD [40]. The results show a dislocation density of (14 ± 0.1) × 1014 m−<sup>2</sup> and (28 ± 0.1) × 1014 m−<sup>2</sup> after austenitization and ausforming, respectively. One might conclude from these results that the dislocation density in the as-quenched martensite after the TMT is substantially increased as compared with conventional treatment. A similar effect of the ausforming on the dislocation density was reported by other authors [58, 59]. Finally, TEM examination of the microstructure allowed us to determine the lath width of the martensitic microstructure. Values of 360 ± 35 nm for AR condition, 350 ± 20 nm for HAT condition, and 318 ± 32 nm for TMT condition were obtained, which are significantly finer than those reported after conventional treatments, i.e., lath size ranging from 300 to 500 nm [60].

The distribution of M23C6 precipitates in the tempered martensitic microstructure is also worth analyzing. **Figure 6** illustrates the distribution of M23C6 carbides after AR, HAT, and TMT processing routes. Coarse and closely spaced M23C6 carbides, about 70 to 500 nm, were observed. The number density and average particle size of these carbides were determined by studying several SEM micrographs to determine values of 6.19 × 1019 m−<sup>3</sup> and 141 ± 3 nm for AR condition, 8.24 × 1019 m−<sup>3</sup>

**173**

fracture.

**Figure 6.**

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

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

and 124 ± 3 nm for HAT condition, and 4.11 × 1019 m−<sup>3</sup>

12 ± 1 nm with a number density of 7.20 × 1021 m−<sup>3</sup>

size of 25 ± 5 nm with a number density of 8.14 × 1019 m−<sup>3</sup>

after conventional heat treatment [35].

*thermomechanical and heat [40].*

a number density of 1.86 × 1022 m−<sup>3</sup>

steel. These values are very similar to those reported by Klueh et al. for the steel

*Martensite matrix, M23C6 precipitate, and MX-nanoprecipitate distributions after the different* 

On the contrary, the finely dispersed MX nanoprecipitates present inside the martensitic laths and associated with dislocations are also observed in **Figure 6**. Therefore, this result suggests the role of dislocations as potential nucleation sites for MX nanoprecipitates. Hence, the importance of ausforming in generating a homogeneous distribution of nanosized MX particles in the microstructure might be also foreseen. These spherical MX nanoprecipitates had a mean particle size of

cases, significantly smaller than those measured after AR condition, i.e., particle

after HAT and TMT are smaller, and the number density higher, than measurements reported in the literature after conventional heat treatments [61].

**Figure 7** shows the disk deflection versus time curves obtained for the three conditions studied (AR, HAT, and TMT) at 700°C with a load of 200 N. The curves exhibit the three stages of creep that were described in previous sections. The first stage corresponds to the loading region where the spherical indenter loads the sample, and the mode of deformation is by bending. The second stage is characterized by a decrease in deflection rate and corresponds to the steady-state region with a minimum disk deflection rate. Finally, the third stage consists of an acceleration of disk deflection and fracture region. In the secondary and tertiary stages, stretching is the prominent deformation mode. Once a crack propagates to a critical length, the sample is no longer in balance, leading to an increase in deflection rate and to a reduction in the structure stiffness in the tertiary region until the final

and 143 ± 5 nm for TMT

for HAT steel and 9 ± 1 nm with

. The size values obtained

for TMT steel. The MX precipitates are, in both

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

**Figure 6.**

*Welding - Modern Topics*

**Figure 5.**

an austenitic microstructure with a high population of deformation bands will be formed. This would directly induce the preferential formation of some specific martensitic variants upon austenite transformation (martensite variant selection),

**Figure 6** illustrates the IPF maps, SEM and TEM micrographs after HAT and TMT processing routes, and the reference (AR) condition. The first conclusion obtained is the coarsening of the block size (white arrows in **Figure 6**) in HAT and TMT conditions as compared with AR condition, because of the high austenitization temperature. Block widths of 2.7 ± 0.2 μm for AR condition were obtained; meanwhile, values of 4.12 ± 0.37 μm for HAT and 3.21 ± 0.27 μm for TMT were measured. The coarser the prior austenite grain, the coarser the block size. However, it is worth noting that finer block size is observed after TMT than with HAT, which is consistent with the fact that thermomechanical processing increases the low-angle

The dislocation density after HAT and TMT was measured by XRD [40]. The

after austenitization and ausforming, respectively. One might conclude from these results that the dislocation density in the as-quenched martensite after the TMT is substantially increased as compared with conventional treatment. A similar effect of the ausforming on the dislocation density was reported by other authors [58, 59]. Finally, TEM examination of the microstructure allowed us to determine the lath width of the martensitic microstructure. Values of 360 ± 35 nm for AR condition, 350 ± 20 nm for HAT condition, and 318 ± 32 nm for TMT condition were obtained, which are significantly finer than those reported after conventional

The distribution of M23C6 precipitates in the tempered martensitic microstructure is also worth analyzing. **Figure 6** illustrates the distribution of M23C6 carbides after AR, HAT, and TMT processing routes. Coarse and closely spaced M23C6 carbides, about 70 to 500 nm, were observed. The number density and average particle size of these carbides were determined by studying several SEM micrographs to

and (28 ± 0.1) × 1014 m−<sup>2</sup>

and 141 ± 3 nm for AR condition, 8.24 × 1019 m−<sup>3</sup>

leading to the development of strong transformation texture.

*Temperature evolution of phase mole fraction in G91 calculated by Thermocalc® [38].*

substructure and decreases the block size of as-quenched martensite.

results show a dislocation density of (14 ± 0.1) × 1014 m−<sup>2</sup>

treatments, i.e., lath size ranging from 300 to 500 nm [60].

determine values of 6.19 × 1019 m−<sup>3</sup>

**172**

*Martensite matrix, M23C6 precipitate, and MX-nanoprecipitate distributions after the different thermomechanical and heat [40].*

and 124 ± 3 nm for HAT condition, and 4.11 × 1019 m−<sup>3</sup> and 143 ± 5 nm for TMT steel. These values are very similar to those reported by Klueh et al. for the steel after conventional heat treatment [35].

On the contrary, the finely dispersed MX nanoprecipitates present inside the martensitic laths and associated with dislocations are also observed in **Figure 6**. Therefore, this result suggests the role of dislocations as potential nucleation sites for MX nanoprecipitates. Hence, the importance of ausforming in generating a homogeneous distribution of nanosized MX particles in the microstructure might be also foreseen. These spherical MX nanoprecipitates had a mean particle size of 12 ± 1 nm with a number density of 7.20 × 1021 m−<sup>3</sup> for HAT steel and 9 ± 1 nm with a number density of 1.86 × 1022 m−<sup>3</sup> for TMT steel. The MX precipitates are, in both cases, significantly smaller than those measured after AR condition, i.e., particle size of 25 ± 5 nm with a number density of 8.14 × 1019 m−<sup>3</sup> . The size values obtained after HAT and TMT are smaller, and the number density higher, than measurements reported in the literature after conventional heat treatments [61].

**Figure 7** shows the disk deflection versus time curves obtained for the three conditions studied (AR, HAT, and TMT) at 700°C with a load of 200 N. The curves exhibit the three stages of creep that were described in previous sections. The first stage corresponds to the loading region where the spherical indenter loads the sample, and the mode of deformation is by bending. The second stage is characterized by a decrease in deflection rate and corresponds to the steady-state region with a minimum disk deflection rate. Finally, the third stage consists of an acceleration of disk deflection and fracture region. In the secondary and tertiary stages, stretching is the prominent deformation mode. Once a crack propagates to a critical length, the sample is no longer in balance, leading to an increase in deflection rate and to a reduction in the structure stiffness in the tertiary region until the final fracture.

**Figure 7.**

*(a) SPCT curves measured for the samples after the different thermomechanical and heat treatments and the creep fracture micrographs for the (a) AR, (b) HAT, and (c) TMT [40].*

As indicated above, the minimum disk deflection rate (δd ) is an important parameter that can be evaluated by SPCT. The evolution of disk deflection rate with the applied load might be described by an equivalent expression to the conventional Norton's power law for creep, which is similar to the expression used in Eq. (1):

$$
\delta\_{\mathbf{d}} = \mathbf{A} \cdot \mathbf{F}^{\mathbf{n}} \tag{2}
$$

**175**

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

increase the dislocation density in the resulting martensitic microstructure. On the other hand, Takahashi et al. [64] reported recently the formation of Nb-cottrell atmospheres in low-carbon Nb-microalloyed steels. The authors explained that this mechanism is based on the fact that segregation energy of Nb to edge dislocation core was almost the same as the energy for grain boundary segregation. Besides, the large attractive interaction between Nb and dislocation core was due to its large atomic size. Therefore, such interaction between Nb atoms and dislocations retards the recovery of dislocation at high temperatures and, hence, stabilizes the microstructure at high temperatures. It might be expected that Nb presents the same behavior in the studied steel, preventing recovery after ausforming and promoting the fine and

homogeneous MX carbonitride precipitation during tempering accordingly.

Log ρd = 9.28480 +

This expression is valid only when the martensite start temperature (T) is

Extracting the data of the martensite start temperature from a previous work [38], the estimation of the dislocation density obtained after the different ausforming conditions can be estimated. In this sense, ausformed samples at 600°C present a martensite start temperature of 338°C; introducing this value in Eq. (3), a disloca-

900°C with a martensite start temperature of 374°C, the dislocation density calcu-

measured by X-ray diffraction, which demonstrate that the ausforming increases

During the final stage (tempering), MX carbonitrides and M23C6 carbides precipitate, and the recovery of dislocations takes place. Because of the higher dislocation density of ausformed samples, the number density of finer MX increases, and these precipitates are found homogeneously distributed within laths, as it can be

The number density of MX precipitates (N) was determined through the direct measurements of spacing (λ) between MX carbonitrides from several TEM micro-

**Figure 8(b)** shows the size distribution of the precipitates in ausformed material. In the material ausformed at 600 and 900°C, the average size of MX carbonitrides was 5.6 nm and 7.4 nm, respectively. The number density of MX carbonitrides

material ausformed at 900°C. On the other hand, the reported values of the size and number density of MX carbonitrides after the conventional processing were 30 nm

for the material ausformed at 600°C and 6.4 × 1022 m−<sup>3</sup>

ing condition studied [41]. Values of (2.8 ± 0.1) × 1015 m−<sup>2</sup>

allows to estimate the dislocation density (ρd):

between the range 297 and 647°C.

tion density of 5.97 × 1015 m−<sup>2</sup>

graphs as indicated by Eq. (4):

was 9.39 × 1022 m−<sup>3</sup>

the dislocation density in the martensite.

seen in **Figure 8(a)** pointed out by white arrows.

lated is 4.62 × 1015 m−<sup>2</sup>

In this work the role of ausforming temperature by selecting 600 and 900°C, at a constant deformation of 20% (**Figure 4**), is explored. As mentioned above, the dislocation densities were estimated by XRD in fresh martensite after each ausform-

were obtained for the ausforming at 600 and 900°C, respectively. These results show that the lower the ausforming temperature, the higher the dislocation density introduced in austenite is, which might be due to the fact that some of the dislocations in fresh martensite are inherited from deformed austenite as it was mentioned above. On the other hand, Bhadeshia and Takahashi reported [65] an expression that

> \_ 6880 T −

\_1780360

T2

is calculated. Similarly, for the material ausformed at

N = 1/ λ<sup>3</sup> (4)

for the

. These results are in the same order of magnitude than those

and (1.9 ± 0.1) × 1015 m−<sup>2</sup>

(3)

that might be transferred to martensite (011)M planes. Therefore, ausforming might

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

where A is a temperature-dependent constant, F is the force applied on the specimen, and *n* is the force exponent. One might conclude, therefore, from **Figure 7** that the creep strength has significantly improved after the TMT condition. The time to rupture was 2.5 and 1.24 times greater than AR condition, from 38 to 95 h and 48 h for the TMT and HAT, respectively. The δd was 2.9 μm·h<sup>−</sup><sup>1</sup> for the TMT sample, while for the HAT sample, it was 3.7 μm·h<sup>−</sup><sup>1</sup> . These minimum disk deflection rates were significantly slower than the minimum disk deflection rate measured for the G91 in the AR condition, which was 9.5 μm·h<sup>−</sup><sup>1</sup> .

The results obtained suggest that the increase in the number density of MX precipitates enhances the strengthening capability at high temperature, since they are able to pin more effectively the dislocations. Hence minimum creep rate is reduced and the onset of tertiary creep is retarded. The differences in minimum disk deflection rate and time to rupture between the sample after TMT and HAT support the importance of ausforming on improving creep resistance.

#### **3.3 Effect of ausforming**

The next stage in the TMT after austenitization is the ausforming as shown in **Figure 4**. The effect of ausforming on low-carbon lath martensitic microstructure has been already described by Miyamoto et al. [62]. The authors reported that martensite variants with habit planes that are nearly parallel to the close-packed primary and secondary slip planes in austenite transform preferentially, i.e., martensite habit planes such as (575)γ that are nearly parallel to (111)γ and (−111) γ in asutenite [63]. Since strain is accumulated preferentially in (111)γ and (−111) γ slip planes during ausforming, this results in an increasing number of dislocation

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

that might be transferred to martensite (011)M planes. Therefore, ausforming might increase the dislocation density in the resulting martensitic microstructure.

On the other hand, Takahashi et al. [64] reported recently the formation of Nb-cottrell atmospheres in low-carbon Nb-microalloyed steels. The authors explained that this mechanism is based on the fact that segregation energy of Nb to edge dislocation core was almost the same as the energy for grain boundary segregation. Besides, the large attractive interaction between Nb and dislocation core was due to its large atomic size. Therefore, such interaction between Nb atoms and dislocations retards the recovery of dislocation at high temperatures and, hence, stabilizes the microstructure at high temperatures. It might be expected that Nb presents the same behavior in the studied steel, preventing recovery after ausforming and promoting the fine and homogeneous MX carbonitride precipitation during tempering accordingly.

In this work the role of ausforming temperature by selecting 600 and 900°C, at a constant deformation of 20% (**Figure 4**), is explored. As mentioned above, the dislocation densities were estimated by XRD in fresh martensite after each ausforming condition studied [41]. Values of (2.8 ± 0.1) × 1015 m−<sup>2</sup> and (1.9 ± 0.1) × 1015 m−<sup>2</sup> were obtained for the ausforming at 600 and 900°C, respectively. These results show that the lower the ausforming temperature, the higher the dislocation density introduced in austenite is, which might be due to the fact that some of the dislocations in fresh martensite are inherited from deformed austenite as it was mentioned above.

On the other hand, Bhadeshia and Takahashi reported [65] an expression that allows to estimate the dislocation density (ρd): \_1780360

$$\begin{array}{l} \text{location density (}\rho\_{\rm d}\text{):}\\\\ \text{Log }\rho\_{\rm d} = \text{9.28480} + \frac{6880}{\text{T}} - \frac{1780360}{\text{T}^2} \end{array} \tag{3}$$

This expression is valid only when the martensite start temperature (T) is between the range 297 and 647°C.

Extracting the data of the martensite start temperature from a previous work [38], the estimation of the dislocation density obtained after the different ausforming conditions can be estimated. In this sense, ausformed samples at 600°C present a martensite start temperature of 338°C; introducing this value in Eq. (3), a dislocation density of 5.97 × 1015 m−<sup>2</sup> is calculated. Similarly, for the material ausformed at 900°C with a martensite start temperature of 374°C, the dislocation density calculated is 4.62 × 1015 m−<sup>2</sup> . These results are in the same order of magnitude than those measured by X-ray diffraction, which demonstrate that the ausforming increases the dislocation density in the martensite.

During the final stage (tempering), MX carbonitrides and M23C6 carbides precipitate, and the recovery of dislocations takes place. Because of the higher dislocation density of ausformed samples, the number density of finer MX increases, and these precipitates are found homogeneously distributed within laths, as it can be seen in **Figure 8(a)** pointed out by white arrows.

The number density of MX precipitates (N) was determined through the direct measurements of spacing (λ) between MX carbonitrides from several TEM micrographs as indicated by Eq. (4):

$$\mathbf{N} = \mathbf{1}/\lambda^3 \tag{4}$$

**Figure 8(b)** shows the size distribution of the precipitates in ausformed material. In the material ausformed at 600 and 900°C, the average size of MX carbonitrides was 5.6 nm and 7.4 nm, respectively. The number density of MX carbonitrides was 9.39 × 1022 m−<sup>3</sup> for the material ausformed at 600°C and 6.4 × 1022 m−<sup>3</sup> for the material ausformed at 900°C. On the other hand, the reported values of the size and number density of MX carbonitrides after the conventional processing were 30 nm

*Welding - Modern Topics*

**Figure 7.**

As indicated above, the minimum disk deflection rate (δd ) is an important parameter that can be evaluated by SPCT. The evolution of disk deflection rate with the applied load might be described by an equivalent expression to the conventional Norton's power law for creep, which is similar to the expression used in Eq. (1):

*(a) SPCT curves measured for the samples after the different thermomechanical and heat treatments and the* 

where A is a temperature-dependent constant, F is the force applied on the specimen, and *n* is the force exponent. One might conclude, therefore, from **Figure 7** that the creep strength has significantly improved after the TMT condition. The time to rupture was 2.5 and 1.24 times greater than AR condition, from 38 to 95 h and 48 h for the TMT and HAT, respectively. The δd was 2.9 μm·h<sup>−</sup><sup>1</sup>

deflection rates were significantly slower than the minimum disk deflection rate

The results obtained suggest that the increase in the number density of MX precipitates enhances the strengthening capability at high temperature, since they are able to pin more effectively the dislocations. Hence minimum creep rate is reduced and the onset of tertiary creep is retarded. The differences in minimum disk deflection rate and time to rupture between the sample after TMT and HAT support the

The next stage in the TMT after austenitization is the ausforming as shown in **Figure 4**. The effect of ausforming on low-carbon lath martensitic microstructure has been already described by Miyamoto et al. [62]. The authors reported that martensite variants with habit planes that are nearly parallel to the close-packed primary and secondary slip planes in austenite transform preferentially, i.e., martensite habit planes such as (575)γ that are nearly parallel to (111)γ and (−111) γ in asutenite [63]. Since strain is accumulated preferentially in (111)γ and (−111) γ slip planes during ausforming, this results in an increasing number of dislocation

the TMT sample, while for the HAT sample, it was 3.7 μm·h<sup>−</sup><sup>1</sup>

*creep fracture micrographs for the (a) AR, (b) HAT, and (c) TMT [40].*

importance of ausforming on improving creep resistance.

**3.3 Effect of ausforming**

measured for the G91 in the AR condition, which was 9.5 μm·h<sup>−</sup><sup>1</sup>

δd = A · F<sup>n</sup> (2)

for

. These minimum disk

.

**174**

#### **Figure 8.**

*(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 and 900°C [41].*

and 1020 m−<sup>3</sup> , 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 resistance is.

**177**

**Figure 10.**

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

**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

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

**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

*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].*

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

carbides with M = (Fe, Cr, Mo).

brittle behavior.

**Figure 9.** *SPCT curves for all samples tested at 700°C with a load of 200 N [39].*

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