**2. Brief remarks on the evolution of 9Cr FM steel for the power-generation industry**

As it has been reviewed by Klueh in his seminal work on high-chromium FM steels [7], the design and production of 9-12Cr FM steels began in 1912 when Krupp and Mannesmann produced a 12 wt. % Cr steel containing 2–5 wt. % Mo. This type of steel was used for steam turbine blades, and it is still in use under the designation of X22CrMoV12. The 2¼Cr-1Mo bainitic steel grade normally known as ASTM Grade 221 (with nominal composition of Fe-2.25Cr-1.0 Mo-0.3Si-0.45Mn-0.12C) was firstly introduced in fossil fuel power plants in the 1940s and is nowadays widely used. The 9Cr-1Mo FM steel grade (known as Grade 9) is a natural evolution from Grade 22, seeking a better corrosion resistance and, hence, increasing the chromium addition. These two steel grades are the reference steels for heat-resistant application in power plants. Since then, the steady need of pushing up the operating conditions in conventional fossil-fired power-generation systems led to the development of several "generations" of steels with improved elevated-temperature strengths. The evolution of steel compositions (**Figure 1**), which began with G22 and G9 (zeroth generation) with 100,000 h creep rupture strengths at 600°C of about 40 MPa, has allowed for increased operating steam temperatures and pressures [1–3, 8–12]. Three generations of steels have been introduced since the introduction of G22 and G9, and a fourth generation is in development.

The strategy adopted for improved corrosion and oxidation resistance for elevated-temperature operating conditions was the addition of carbide formers such as vanadium and niobium to add precipitate strengthening. Hence, the zeroth generation containing mainly 9-12Cr evolved to the 12Cr-MoV steels introduced in the power plants in the mid-1960s for thin- and thick-walled power station

**165**

strength [13, 14].

**Figure 1.**

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

components. Their creep strength is based on solution hardening and on the precipitation of M23C6 carbides. These steels have been applied successfully in power stations over several decades [10]. These steels had increased 105

structure (tempered martensite and δ-ferrite). These steels have 105

The second generation, developed in the late 1970s, is based on the modified 9Cr-1Mo, designated as G91 and HCM12 (see **Table 1**), which were developed for manufacturing of pipes and vessels for fast breeder reactors [10]. In this steel class, C, Nb, and V contents were optimized, N (0.03–0.05 wt. %) was added, and the maximum operating temperature increased to 593°C. The new steels have a duplex

strengths at 600°C of about 100 MPa. Of these latter steels, G91 has been used most extensively in the power-generation industry in all new power plants with operational temperatures up to 600°C [7]. The responsible mechanism for this substantial increment of creep strength as compared with 12Cr-MoV steels is the formation of thermally stable V and Nb carbonitrides. Besides lowering the Cr content down to 9 wt. %, tempered martensite microstructure also contributes to the higher creep

The Japanese steel development program led by Nippon Steel achieved the development of the P92 steel (NF616). This steel grade, designated as Grade 92, presents a further increases in stress rupture by the addition of 0.003 wt. % B and 1.8 wt. % W and reducing the Mo content from 1 to 0.5 wt. % [15, 16]. The addition of B ensures thermally stable M23(C,B)6 precipitates, whereas the higher W content leads to a higher amount of precipitated Laves phase [17, 18]. Grade 92, firstly introduced in the 1990s along with equivalent steel such as E911, fulfills the niche of steam

Finally, the goal for the next steel generation being developed at present is pushing the limit of operation temperature above 650°C. This so-called fourth generation differs from the previous ones mainly by the addition of 3.0 wt. % Co

h creep rupture strengths at 140 MPa.

strengths at 600°C of up to 60 MPa (**Table 1**).

*Flowchart showing the evolution of 9-12Cr FM steels [7].*

operational temperature of 620°C for 104

h rupture

h rupture

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

<sup>1</sup> Grade 22 and the other commercial steels are given designations by ASTM (e.g., Grade 9 is 9Cr-1Mo, and Grade 91 is modified 9Cr-1Mo). The steels are further distinguished as T22 or T91 for tubing, P22 and P91 for piping, F22 and F91 for forgings, etc. The "G" designation will mainly be used here.

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


**Figure 1.**

*Welding - Modern Topics*

improved greatly.

Grade 221

**power-generation industry**

superheater and reheater tubing in boilers, turbine valve chest, rotors and casings, main steam and reheat pipework, generator rotors, and bolts used for high-temperature applications. The boiler components are limited by corrosion and creep. Pipework also suffers creep as well as weld cracking and thermal fatigue. Turbine components are subjected to creep and fatigue (both thermal and mechanical). Therefore, the development of improved structural materials to increase in thermal efficiency has been the driving force to develop new generations of 9-12Cr ferritic/martensitic (FM) steels [1–3]. The most relevant in-use properties that heat-resistant steels employed to manufacture components in power plants should fulfill are good mechanical properties, fabricability, corrosion resistance, and creep strength. As indicated above, creep strength has been the most studied and has led to innumerable research activities, aiming at improving the creep strength in 9-12Cr FM steel developments [4–6]. The disadvantage of these steels is their loss of strength beyond 600°C, so they need to be optimized to guarantee their use in the future power plants. In this chapter one of the most promising ideas described is applying a thermomechanical treatment (TMT) instead of a conventional treatment. The main contribution of the TMT is the ausforming, which, as other authors have reported, allows increasing considerably the number density of the thermally stable precipitates, i.e., MX nanoprecipitates. Consequently, the creep strength has

**2. Brief remarks on the evolution of 9Cr FM steel for the** 

As it has been reviewed by Klueh in his seminal work on high-chromium FM steels [7], the design and production of 9-12Cr FM steels began in 1912 when Krupp and Mannesmann produced a 12 wt. % Cr steel containing 2–5 wt. % Mo. This type of steel was used for steam turbine blades, and it is still in use under the designation of X22CrMoV12. The 2¼Cr-1Mo bainitic steel grade normally known as ASTM

was firstly introduced in fossil fuel power plants in the 1940s and is nowadays widely used. The 9Cr-1Mo FM steel grade (known as Grade 9) is a natural evolution from Grade 22, seeking a better corrosion resistance and, hence, increasing the chromium addition. These two steel grades are the reference steels for heat-resistant application in power plants. Since then, the steady need of pushing up the operating conditions in conventional fossil-fired power-generation systems led to the development of several "generations" of steels with improved elevated-temperature strengths. The evolution of steel compositions (**Figure 1**), which began with G22 and G9 (zeroth generation) with 100,000 h creep rupture strengths at 600°C of about 40 MPa, has allowed for increased operating steam temperatures and pressures [1–3, 8–12]. Three generations of steels have been introduced since the

introduction of G22 and G9, and a fourth generation is in development.

The strategy adopted for improved corrosion and oxidation resistance for elevated-temperature operating conditions was the addition of carbide formers such as vanadium and niobium to add precipitate strengthening. Hence, the zeroth generation containing mainly 9-12Cr evolved to the 12Cr-MoV steels introduced in the power plants in the mid-1960s for thin- and thick-walled power station

<sup>1</sup> Grade 22 and the other commercial steels are given designations by ASTM (e.g., Grade 9 is 9Cr-1Mo, and Grade 91 is modified 9Cr-1Mo). The steels are further distinguished as T22 or T91 for tubing, P22 and P91 for piping, F22 and F91 for forgings, etc. The "G" designation will mainly be used here.

(with nominal composition of Fe-2.25Cr-1.0 Mo-0.3Si-0.45Mn-0.12C)

**164**

*Flowchart showing the evolution of 9-12Cr FM steels [7].*

components. Their creep strength is based on solution hardening and on the precipitation of M23C6 carbides. These steels have been applied successfully in power stations over several decades [10]. These steels had increased 105 h rupture strengths at 600°C of up to 60 MPa (**Table 1**).

The second generation, developed in the late 1970s, is based on the modified 9Cr-1Mo, designated as G91 and HCM12 (see **Table 1**), which were developed for manufacturing of pipes and vessels for fast breeder reactors [10]. In this steel class, C, Nb, and V contents were optimized, N (0.03–0.05 wt. %) was added, and the maximum operating temperature increased to 593°C. The new steels have a duplex structure (tempered martensite and δ-ferrite). These steels have 105 h rupture strengths at 600°C of about 100 MPa. Of these latter steels, G91 has been used most extensively in the power-generation industry in all new power plants with operational temperatures up to 600°C [7]. The responsible mechanism for this substantial increment of creep strength as compared with 12Cr-MoV steels is the formation of thermally stable V and Nb carbonitrides. Besides lowering the Cr content down to 9 wt. %, tempered martensite microstructure also contributes to the higher creep strength [13, 14].

The Japanese steel development program led by Nippon Steel achieved the development of the P92 steel (NF616). This steel grade, designated as Grade 92, presents a further increases in stress rupture by the addition of 0.003 wt. % B and 1.8 wt. % W and reducing the Mo content from 1 to 0.5 wt. % [15, 16]. The addition of B ensures thermally stable M23(C,B)6 precipitates, whereas the higher W content leads to a higher amount of precipitated Laves phase [17, 18]. Grade 92, firstly introduced in the 1990s along with equivalent steel such as E911, fulfills the niche of steam operational temperature of 620°C for 104 h creep rupture strengths at 140 MPa.

Finally, the goal for the next steel generation being developed at present is pushing the limit of operation temperature above 650°C. This so-called fourth generation differs from the previous ones mainly by the addition of 3.0 wt. % Co


#### **Table 1.**

*Chemical composition and creep rupture strength at 600°C of the 9-12Cr heat-resistant steels from 1950 to 2005 [19].*

as an austenite stabilizer because of the adverse effect of nickel on creep. They have projected 105 h creep rupture strengths at 600°C of 180 MPa [7]. In these steels with about 0.1 wt. % carbon, molybdenum has been further reduced or eliminated, and tungsten (2.6–3.0 wt. %) has been increased compared to third-generation compositions. In **Table 1**, an overview of the historical development of the 9-12Cr heat-resistant steels from 1950 to 2005 is shown.

#### **2.1 Creep deformation**

Creep deformation is a thermally activated process, and the rate of deformation (creep rate) is extremely temperature sensitive. In metals, creep deformation becomes important at temperatures greater than about 0.4TM, where TM is the absolute melting temperature [20]. In the case of 9-12Cr FM steels, this temperature is approximately 450°C. Clearly, power plant materials operate in the temperature regime where creep process is significant. The creep properties of the material used limit the operating temperature of many power plant components, such as the turbines. Development of materials with an increased creep resistance is central to the use of power plants with higher steam temperatures.

Creep deformation can occur by a variety of different mechanisms. The mechanism that dominates depends on the stress and temperature conditions as well as the microstructure of the material.

In the case of power plant steels, the stress levels are relatively high, and the temperatures (compared with melting point) are relatively low. In the case of creep deformation, it is controlled primarily by dislocation movement and the thermal energy available for dislocations to overcome obstacles. A deformation mechanism map gives information about which mechanism will dominate for a particular set

**167**

**Figure 2.**

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

of conditions. Such a diagram for a G91 steel is shown in **Figure 2**. For the exposure conditions, for this material, a power law creep (dislocation creep) is expected to

Power law creep involves the movement of dislocations, and the creep rate is a result of the balance between work hardening and recovery. Work hardening results in an increase in the dislocation density, while recovery leads to a reduction in the dislocation density. If the dislocation density remains constant, then the creep rate

*A* (σ*<sup>a</sup>* + σ*<sup>l</sup>* + σ*<sup>u</sup>* + σ*g*)

In this equation σ*a* is the applied stress, σ*l*, σ*u*, and σ*g* are the internal stresses due to solution, precipitation, and grain boundary hardening, respectively. The most effective method for reducing the creep rate is therefore to form a suitable distribu-

<sup>4</sup> (1)

ε = ̇

tion of particles which are also able to act as barriers to dislocation motion. A study of the possible creep mechanisms suggests microstructures would be expected to have good creep resistance under conditions used in power plant. In general, creep-resistant alloys are based on a matrix which is a solid solution. The presence of misfitting solute atoms in solid solution makes the passage of dislocations through the matrix more difficult. However, the majority of the creep resistance, at least in the early stages of service, is derived from precipitate particles. Ideally these particles should be small, and they should be widely and homogeneously distributed in large numbers through the matrix. The particles need to be stable at operating temperatures for which the alloy is designated, and they should be resistant to coarsening, as this will reduce their effectiveness as strengtheners. In general, excessive work hardening and very fine grain sizes, which provide strengthening at ambient temperatures, are considered detrimental in high-temperature alloys. This is because both of them provide easy diffusion paths and therefore

*Deformation mechanism map (D-MAP) of T91 steel calculated from experimental reported values [22, 23].* 

*The red line indicates the experimental conditions considered in this work.*

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

is given by Norton's law [21]:

lead to an increase in the creep rate.

dominate.

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

of conditions. Such a diagram for a G91 steel is shown in **Figure 2**. For the exposure conditions, for this material, a power law creep (dislocation creep) is expected to dominate.

Power law creep involves the movement of dislocations, and the creep rate is a result of the balance between work hardening and recovery. Work hardening results in an increase in the dislocation density, while recovery leads to a reduction in the dislocation density. If the dislocation density remains constant, then the creep rate is given by Norton's law [21]:

$$
\dot{\varepsilon} = \mathcal{A} \left( \sigma\_d + \sigma\_l + \sigma\_u + \sigma\_\mathcal{g} \right)^4 \tag{1}
$$

In this equation σ*a* is the applied stress, σ*l*, σ*u*, and σ*g* are the internal stresses due to solution, precipitation, and grain boundary hardening, respectively. The most effective method for reducing the creep rate is therefore to form a suitable distribution of particles which are also able to act as barriers to dislocation motion.

A study of the possible creep mechanisms suggests microstructures would be expected to have good creep resistance under conditions used in power plant. In general, creep-resistant alloys are based on a matrix which is a solid solution. The presence of misfitting solute atoms in solid solution makes the passage of dislocations through the matrix more difficult. However, the majority of the creep resistance, at least in the early stages of service, is derived from precipitate particles. Ideally these particles should be small, and they should be widely and homogeneously distributed in large numbers through the matrix. The particles need to be stable at operating temperatures for which the alloy is designated, and they should be resistant to coarsening, as this will reduce their effectiveness as strengtheners. In general, excessive work hardening and very fine grain sizes, which provide strengthening at ambient temperatures, are considered detrimental in high-temperature alloys. This is because both of them provide easy diffusion paths and therefore lead to an increase in the creep rate.

#### **Figure 2.**

*Welding - Modern Topics*

as an austenite stabilizer because of the adverse effect of nickel on creep. They have

*Chemical composition and creep rupture strength at 600°C of the 9-12Cr heat-resistant steels from 1950 to* 

with about 0.1 wt. % carbon, molybdenum has been further reduced or eliminated, and tungsten (2.6–3.0 wt. %) has been increased compared to third-generation compositions. In **Table 1**, an overview of the historical development of the 9-12Cr

Creep deformation is a thermally activated process, and the rate of deformation (creep rate) is extremely temperature sensitive. In metals, creep deformation becomes important at temperatures greater than about 0.4TM, where TM is the absolute melting temperature [20]. In the case of 9-12Cr FM steels, this temperature is approximately 450°C. Clearly, power plant materials operate in the temperature regime where creep process is significant. The creep properties of the material used limit the operating temperature of many power plant components, such as the turbines. Development of materials with an increased creep resistance is central to

Creep deformation can occur by a variety of different mechanisms. The mechanism that dominates depends on the stress and temperature conditions as well as the

In the case of power plant steels, the stress levels are relatively high, and the temperatures (compared with melting point) are relatively low. In the case of creep deformation, it is controlled primarily by dislocation movement and the thermal energy available for dislocations to overcome obstacles. A deformation mechanism map gives information about which mechanism will dominate for a particular set

heat-resistant steels from 1950 to 2005 is shown.

the use of power plants with higher steam temperatures.

h creep rupture strengths at 600°C of 180 MPa [7]. In these steels

**166**

projected 105

**Table 1.**

*2005 [19].*

**2.1 Creep deformation**

microstructure of the material.

*Deformation mechanism map (D-MAP) of T91 steel calculated from experimental reported values [22, 23]. The red line indicates the experimental conditions considered in this work.*

## **2.2 Creep and microstructural evolution**

The new environmental regulations and commercial needs of the industry are the driving force for the development of new heat-resistant steels that push forward the operational limits of current steels. In this framework, the high-Cr FM steels applied as structural materials in fossil-fired and in nuclear power plants need to implement the operating temperatures above 650°C [1–3, 8–12]. The mechanism responsible for creep strengthening in these steels is the solid-solution and dispersion strengthening.

In the particular case of the so-called 9Cr FM steels, the creep degradation is a consequence of the thermal evolution of their hierarchal martensitic microstructure constituted by prior austenite grains, martensitic packets, blocks, and laths [24]. The microstructural degradation during creep consists of the coarsening of the lath structure [12, 24]. Such coarsening is governed by the subgrain boundary formation and evolution inside the laths, which can be prevented at high temperatures, and virtually frozen, by the dispersion of proper precipitates. The precipitates pin boundary migration and dislocation motion, slowing down the degradation of the martensitic microstructure and hence reducing creep rates [25, 26].

There are two main actors for the microstructural stability driven by precipitation in 9Cr FM steels: The first one is the coarse M23C6 carbides located mainly at the grain boundaries either from the prior austenitic grains or from the blocks or martensite lath boundaries. The second one is the V- and Nb-rich MX carbonitrides. Contrary to M23C6, those MX precipitates are homogeneously distributed within martensite laths. Therefore, the ideal situation would consist of reducing the presence of the M23C6 carbides to the minimum since their fast coarsening induces crack formation at the particle-matrix interface and promoting the formation of MX carbonitrides (nanometric in size), since they will delay the lath coarsening as mentioned above; it has been studied extensively [27–29].

Thermomechanical processing of 9Cr FM steels has been revealed as a promising tool to promote a high number density of MX carbonitrides [30–38]. TMT involves different steps that need to be optimized to produce the most favorable precipitate microstructure for elevated-temperature strength.

#### **2.3 Creep tests**

The creep behavior of a material may be characterized by a number of different parameters which can be measured by performing the appropriate creep test. For metallic materials most creep tests are conducted in uniaxial tension with a dumbbell-shaped specimen similar to that used for tensile testing. The tests are carried out at a constant temperature and under either a constant load or stress. Applying a constant stress is more useful if the test is being employed to provide information about a creep mechanism.

The conventional treatments (AR) and TMT considered in this work were carried out on 10 mm in length and 5 mm in diameter cylindrical samples using a DIL 805A/D plastodilatometer (TA instruments) as described elsewhere [39, 40]. Due to the limited amount of material available after the TMT is carried out in the plastodilatometer, the creep properties were investigated by means of the small punch creep test (SPCT) performed at 700°C as it has been previously reported [41, 42]. The SPCT samples were cut transversally, from cylindrical specimens, with a thickness of 600 μm and a diameter of 8 mm. Then, the disks were ground on both sides down to a final thickness of 500 μm. In the setup of the SPCT, the lower and upper dice are connected via a thread to ensure the clamping of the sample. The load is applied by a ceramic punch ball which is in contact with the sample. A plunger rod is used to transmit the dead weight load to

**169**

current power plants.

dispersion-strengthened steels.

**3.1 Microstructure after conventional heat treatment**

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

different regions similarly to conventional strain vs. time creep curves obtained from uniaxial testing. However, the failure in SPCTs occurs away from the load line with cracks propagating in a circumferential direction due to membrane stretching. Therefore, the first part of the disk deflection vs. time curve corresponds to the loading region where the spherical indenter loads on a very small contact area of the sample are. Since the stresses will be higher than the yield stress of the material, local plasticity and an initial large deformation are produced. This large deformation is accumulated in a short period of time. The second stage corresponds to the steadystate region, which coincides with most of the sample life, where the disk deflection rate reached almost a minimum. Finally, the third stage consists of an acceleration of disk deflection and fracture region. The interpretation of this behavior is that 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. Another explanation might be due to the localized necking without crack presence. The deformation mechanism in the tertiary region is a mixture among accumulation of creep damage, geometric softening, and crack growth effect.

**3. Effect of thermomechanical control processing on microstructure**

steels present an upper service temperature of 540°C, which was successfully increased in the late 1970s up to 595°C with the introduction of vanadium and niobium microalloying in the composition of the steel. This steel was used as a benchmark for the development of steels with upper-use temperatures of 600–620°C.

As it has been introduced in previous sections, the pioneer commercial 9-12Cr

However, it is difficult keep pushing the higher operating temperature too much. Therefore, to continue to exploit the advantages of ferritic steels, oxide dispersionstrengthened (ODS) steels [43–46] were introduced. The first successful alloy was presented in the 1960s, and, since then, it has been an active research field. ODS steels are strengthened by small oxide particles, but the complicated and expensive manufacturing route avoided the full implantation as structural material in the

Despite being around for about 40 years, the ODS steels are still in the development stage because of having mechanical property anisotropy [43, 45, 47–49]. Therefore, an alternative strategy to achieve a high number density of precipitates is needed. In this section, we present preliminary results that allow us to conclude that conventional thermomechanical control processing strategy is adequate to achieve

Lath martensite is a particular microstructure that ensures microstructural stability. Furuhara and Miyamoto [50] described the variety of crystalline size in lath martensite structures. A hierarchy of lath martensite structure is clearly identified particularly in low-carbon steels. A prior austenite (γ) grain is divided into "packets," each of which consists of a group of martensite laths with the same

the punch ball. All these components are made of Al2O3 ceramics. The clamping device is surrounded by an electrical heater and a thermal insulation. The upper plate carrying the additional dead weight is guided by two pillars with ball bearings. The temperature is measured in the lower die directly under the sample. The displacement is measured by a capacitive sensor between the upper plate and the thermal insulation with an accuracy of ±1 μm. A load cell is placed between the upper plate and the plunger rod. The disk deflection vs. time resulting from the SPCTs might be divided into three

*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*

the punch ball. All these components are made of Al2O3 ceramics. The clamping device is surrounded by an electrical heater and a thermal insulation. The upper plate carrying the additional dead weight is guided by two pillars with ball bearings. The temperature is measured in the lower die directly under the sample. The displacement is measured by a capacitive sensor between the upper plate and the thermal insulation with an accuracy of ±1 μm. A load cell is placed between the upper plate and the plunger rod.

The disk deflection vs. time resulting from the SPCTs might be divided into three different regions similarly to conventional strain vs. time creep curves obtained from uniaxial testing. However, the failure in SPCTs occurs away from the load line with cracks propagating in a circumferential direction due to membrane stretching. Therefore, the first part of the disk deflection vs. time curve corresponds to the loading region where the spherical indenter loads on a very small contact area of the sample are. Since the stresses will be higher than the yield stress of the material, local plasticity and an initial large deformation are produced. This large deformation is accumulated in a short period of time. The second stage corresponds to the steadystate region, which coincides with most of the sample life, where the disk deflection rate reached almost a minimum. Finally, the third stage consists of an acceleration of disk deflection and fracture region. The interpretation of this behavior is that 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. Another explanation might be due to the localized necking without crack presence. The deformation mechanism in the tertiary region is a mixture among accumulation of creep damage, geometric softening, and crack growth effect.
