9–12% Cr Heat-Resistant Martensitic Steels: History
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As a promising alloying approach, the modification of chemical composition by increasing the B content and decreasing the N content has been applied to improve the creep resistance of various 9–12% Cr heat-resistant martensitic steels. The 9–12% Cr steels have to exhibit high long-term creep strength, oxidation resistance in a high temperature steam, low cycle fatigue resistance, impact toughness, etc. The creep resistance is the main critical requirement: the minimum long-term creep rupture strength on the base of 100,000 h should be 100 MPa or higher at 650 °C. 


  • martensitic steels
  • chemical composition
  • boron
  • nitrogen

1. Introduction

In the fossil power plant industry, 9–12% Cr martensitic steels are widely used materials for such components as boilers, pipes, turbines, rotors, and blades, etc. [1,2].
An increase in the operation temperature to 650–720 °C and pressure of power units with ultrasupercritical (USC) and advanced ultrasupercritical (A-USC) conditions require the development of new generation steels with enhanced creep resistance. Heat-resistant materials, such as austenitic steels and nickel-based superalloys, are used only for some parts of fossil power units (approximately up to 15%) [3]. Therefore, the high-chromium heat-resistant martensitic/ferritic steels remain as the main materials for fossil power units due to their excellent combination of creep resistance and fatigue resistance and smaller thermal expansion and larger thermal conductivity compared to austenitic steels and nickel-based alloys, as well as their low cost [4].
The creep resistance of high-chromium martensitic steels is determined by the stability of the non-equilibrium hierarchical microstructure. Prior austenite grains (PAG), packets, and blocks with martensite laths are the structural elements of the typical martensite lath structure [7,8]. Two main dispersion strengthening phases are the nanoscale M23C6 (where M—Cr, Fe, W, etc.) carbides and the fine MX carbonitrides (where M—V, Nb and X—C, N). During tempering, the M23C6 carbides precipitated on boundaries, while the MX carbonitrides homogeneously precipitated in the lath interiors. Both M23C6 carbides and MX carbonitrides prevent the dislocation climb and slow down the migration of low-angle boundaries through suppressing the knitting reaction between the dislocations comprising lath boundaries and lattice dislocations [9,10]. The transformation of the lath structure into the subgrain structure results in degradation of creep resistance.
One of the most effective ways to enhance the creep resistance was suggested by researchers at the National Institute for Materials Science (NIMS) in Japan. The method consists of increasing the B content and decreasing the N content [11,12,13]. Microalloying by approximately 0.01 wt.% B can increase the long-term creep resistance of the steels [1,7,8,14,15,16,17]. Boron has a positive effect on the coarsening resistance of M23C6–type carbides [7,8,15,18]. A nitrogen removal to approximately 0.01 wt.% prevents the BN formation and can also positively affect the creep resistance [7,8,13,15].

2. Obtaining and Heat Treatment of Steels

2.1. Chemical Composition of Steels

Advanced steels contain 2–3% Co, which is known to have a positive effect on the microstructure and creep strength of high-Cr steels [19,20,21,22,23]. The main purpose of the Co addition is to suppress the formation of undesirable δ-ferrite during normalizing. Helis et al. studied the effect of 0–5% Co on the 9Cr-3W-0.2V-0.05Nb-0.08C-0.05N steel [19]. It was found that addition of 1% Co reduced the δ-ferrite fraction from 6% to 0.4%, while 3% Co completely eliminated the δ-ferrite. This fact is due to Co being an austenite-forming element, and it extends the austenite region on the phase diagram. The absence of δ-ferrite in high-Cr steels increases the stability of the tempered martensite lath structure [19].
On the other hand, although precipitates do not contain Co, the addition of 3% Co provides an increasing amount of MX carbonitrides and M23C6 particles. It was revealed that the number of precipitates around PAG boundaries significantly increased at 3% Co from 6 (at 0% Co) to 14 per μm2 [19]. Co also affects the chemical composition of precipitates, increasing the V content in MX particles, and the Fe, Cr, and W content in M23C6 particles. Therefore, Co indirectly affects the precipitation strengthening of 9–12% Cr steels.

2.1.1. The 9% Cr Steels

The 9% Cr steels are represented by the advanced CSEF steels, such as MARBN, G115, and SAVE12AD steels, and experimental 9Cr-1.5W-3Co steel.
The MARBN steel is the Japanese 9% Cr steel developed by National Institute of Materials Science (NIMS), in co-operation with private companies in Japan, for application to thick section boiler components in USC power plant [11,25,26]. MARBN is a MARtensitic 9Cr steel strengthened by B boron and N nitrides [11,25,26]. Various compositions of MARBN steel are presented in literature, differing in concentrations of B and N [11,25,26,27].
The G115 steel is the 9% Cr Chinese steel developed by the China Iron and Steel Research Institute (CISRI) and Bao Steel [28,29,30]. This 9Cr-3W-3Co-1CuVNbB steel is recommended for use in USC power plants at operation temperatures up to 650 °C in China owing to its excellent overall properties. It is reported that G115 with 2.8%W–3%Co is for piping and with 3% W–3% Co for tubing [30]. Composition of the G115 is similar to MARBN steel, while approximately 1% Cu is added in order for additional strengthening by fine Cu-rich particles, by analogy with P122 steel [6,31].

2.1.2. The 10% Cr Steels

The 10% Cr steels are represented by the experimental 10Cr [8,15] and 10Cr-0.2Re [39,40,41] steels designed on the base of TOS 110 steel. As known, TOS 110 steel was developed in Japan at Toshiba with the main composition of 10Cr-0.7Mo-1.8W-3Co-VNb-0.01B-0.02N for turbine rotor application at 630 °C [42,43,44]. The experimental 10Cr steel is a modification of TOS 110 steel by decreasing the N content to 0.003% and addition of Ti (0.002%) while maintaining the high B content of 0.008%. It was shown that this modification results in enhanced long-term creep rupture strength [8,15]. The NF12 [2] and HR1200 [43,45] steels were used for comparison of creep resistance of the 10Cr steel.

2.1.3. The 11–12% Cr Steels

The 11–12% Cr steels are represented by the TAF650, SuperVM12 steels, and experimental 12% Cr steels.
The TAF650 steel is the 12% Cr Japanese steel derived from the TAF steel [47]. The TAF steel was developed in 1956 by Toshio Fujita [48] and has the superior high temperature strength. The TAF650 steel was developed for improvement of poor weldability and hot workability. In the TAF650 steel, the part of Mo was replaced by W; Co and Ni were added; C and B contents were reduced from 0.21% to 0.1% and from 0.03% to 0.019%, respectively [49,50,51,52].

2.2. Vacuum Induction Melting of Steels

In contrast to high-chromium steels with conventional N content (~0.05%), the method of vacuum induction melting is used for producing steels with decreased nitrogen content (<0.01%). It is attributed to the fact that the nitrogen content in alloy is determined by the gas pressure on the molten alloy.
As is known, the solubility of the diatomic nitrogen gas (N2) in metal melt can be described by the reaction:
1/2 N2 = [N],
and obeys Sievert’s law [59], according to which the solubility of nitrogen gas in metal melts is proportional to the square root of the partial pressure of the gas (under constant temperature):
N = K N p N 2 ,
where [N] is the solubility of the nitrogen gas in metal melt at a given partial pressure of gas pN2; KN is the solubility constant (Sievert’s constant), which depends on temperature and the way concentration and pressure are expressed. 

2.3. Heat Treatment

Heat treatment of the steels usually consists in normalization and tempering in order to form a tempered martensite lath structure. Normalization of the considered steels is carried out in a wide temperature range of 1050 to 1170 °C (Figure 4). Heating is carried out in the austenite region for preventing the δ-ferrite formation after normalization. Duration of heating varies from 0.5 to 1 h. During cooling (by air or oil), the austenite transforms to martensite. The normalization leads to the formation of the lath martensite structure, in which the PAGs, packets, blocks, and laths with high density of lattice dislocations are distinguished. The normalization temperature affects the PAG size. The size of PAG in the steels significantly varies from 20 to 200 μm due to different normalization temperatures.
Figure 4. Scheme of heat treatment of considered high-chromium martensitic steels.
Typical tempering temperatures are 750–780 °C, although lower and higher temperatures in the range of 680–800 °C are encountered. Duration of tempering varies in the range 1–4 h. During tempering, the normalized martensite becomes tempered martensite and the tempered martensite lath structure is formed (Figure 1). The dislocation density decreases as compared to the normalized state but remains high, approximately 2–4 × 1014 m−2. The M23C6 carbides precipitate on boundaries of PAGs, packets, blocks, and laths, while the MX carbonitrides homogeneously precipitate in the lath interiors.

3. The Boron/Nitrogen Ratio in the Steels

Boron is known to improve the stability of tempered martensite lath structure through the stability of M23C6 carbides. Boron suppresses the coarsening of these carbides. As it is included in M23(B, C)6 carbides during tempering, it significantly reduces the diffusion processes and retains the fine carbides for a long time [7,14].
However, addition of B alone to the steels can have a negative effect. Thus, El-Kashif et al. [62] reported that, at conventional N content (0.05 wt.%), increasing B from conventional 0.001% to 0.006% in the 9Cr-3W-3Co steel reduced the time to rupture by two times at creep testing at 650 °C, 160 MPa. The change in B content was accompanied by a decrease in the PAG size from 160 to 120 μm. At the same time, when the N content was lower than conventional and comprised 0.02%, an increase in the B content to 0.01 wt. % increased time to rupture. It is interesting, that creep strength improved in this case despite the fact that the PAG size also decreased from 155 to 110 μm. Then, this change in the PAG size did not affect the creep resistance. In the high N–high B steel, the number density of precipitates containing B was the highest, whereas, in the low N–high B steel, more free B existed in the matrix [62]. Therefore, free B in solid solution is more effective on the stabilization of the microstructure and strengthening than B contained in precipitates [62].
Therefore, the B/N ratio should be controlled to increase the operating temperature. The following factors determine the contents of boron and nitrogen:
  • Preventing the formation of BN phase. On the one hand, an increased boron content enhances the coarsening resistance of M23C6-type carbides during creep. On the other hand, the presence and content of nitrogen strongly affects the efficiency of alloying with boron. Nitrogen affects the solubility of boron in the ferritic matrix. At excess nitrogen content, the undesirable BN phase is formed during normalization. The coarse BN precipitates can act as initiation sites for creep cavities that reduce the creep resistance and creep ductility. Formation of BN leads to a depletion of boron from the ferritic matrix.
The solubility of boron is determined by the solubility product for boron nitride in 9–12% Cr steels at a normalizing temperature of 1050–1150 °C, given by [63]:
log[%B] = −2.45 log [%N] − 6.81, 
where [%B] and [%N] are the concentration of soluble boron and soluble nitrogen in mass %, respectively.
The steels with conventional B and N contents, such as P92, P122, Save12, NF12, and others, are prone to BN formation. Therefore, in these steels, the soluble B content is low due to the fact that B is consumed to form the BN. Sakuraya et al. [63] showed that there are many BN type inclusions with sizes up to about 4 μm in the P122 and P92 steels. Moreover, in the P122 steel, the BN inclusions were revealed to agglomerate in large size colonies (20 μm).
BN particles reduce creep ductility that deteriorates creep strength at long times [64]. Abe et al. established that the BN particles are responsible for the degradation in reduction of area (RA) at low stresses and long times by accelerating the formation of creep voids at interfaces between the BN particles and alloy matrix [65].
Optimal fraction of MX phase. The nitrogen content determines the volume fraction of MX carbonitrides, namely N- and V-enriched precipitates, in steels. At optimal B content, the N content should not be too high or too low.
Thus, in the MARBN steel (9Cr-3W-3Co) with 140 ppm B, Abe et al. showed that 79 ppm N improved the creep rupture time as compared with conventional 650 ppm N, whereas the low N content of 15 and 34 ppm reduces creep rupture time  [11,26]. According to Equation (3), at 140 ppm B, 95 ppm N can dissolve in the matrix without any formation of BN during normalization. However, low N results in a low fraction of MX phase. Therefore, weak strengthening by MX carbonitrides results in a higher minimum creep rate and shorter time to rupture. It is interesting that, at 79 ppm N, most of the N dissolved in the matrix after tempering and then started to precipitate in very fine MX carbonitrides during 1000 h of creep, which contributed to the creep strength [11].
Preventing MX→Z-phase transformation. In the steels with 10–12% Cr, a low N concentration prevents the transformation of the MX phase into the undesirable coarse Z-phase particles (CrVN). Thus, in the experimental 12% Cr steels with 20 ppm N, the Z phase was not observed even after long-term creep for >20,000 h at 650 °C [17]. In the Super VM12 steel with 110 ppm N, the Z-phase was not revealed after long-term creep during 23,844 h at 650 °C [54].

4. Microstructure and Creep Properties of Advanced 9–12% Cr Steels

4.1. Creep Properties of the 9–12% Cr Steels

4.1. Creep Properties of the 9–12% Cr Steels

Figure 7 illustrates the creep data at 650 °C for the advanced 9–12% Cr steels with the increased B and decreased N contents in comparison with those for the conventional P92 and P122 steels. The data are from Refs. [5,6,8,11,16,18,20,22,25,27,32,39]. It is clearly seen that the time to rupture versus applied stress points for these steels are higher, mainly than those for the P92 and P122 steels, suggesting the higher creep resistance of new steels.
Figure 7. Time to rupture vs. stress curves of 9–12% Cr steels with the increased B and decreased N contents in comparison with the curves for the P92 and P122 steels. Data from [5,6,8,11,16,18,20,22,25,27,32,39].

4.2. The 9% Cr Steels

The advanced 9% Cr CSEF steels—the MARBN, G115, and SAVE12AD steels—demonstrate a high level of creep strength both at high stresses (in the short-term creep region) and at low stresses (in the long-term creep region) as compared with the conventional P92 steel (Figure 8). Regression lines of the experimental creep points predict the long-term creep rupture strength of these steels at 650 °C for 100,000 h in the range from 80 to 110 MPa [5]. These values are significantly higher than those for the P92 steel and 3%Co-modified P92 steels with 2% and 3%W (approximately 60–70 MPa) (Figure 8a).
Figure 8. Time to rupture vs. stress curves of 9% Cr steels with increased B and decreased N contents in comparison with curves for 9% Cr steels with conventional B, N contents (ad). Experimental data from [6,8,11,16,18,20,22,25,27,32].
The creep strength curves for the MARBN steel vary depending on the B and N content [11,25]. The MARBN steel with 140 ppm B/79 ppm N demonstrates the highest creep resistance, whereas a lower N content of 34 ppm reduces the creep strength (Figure 8b).
The G115 steels with slightly different B and N content (140–150 ppm B and 80–90 ppm N) studied by Liu et al. [16] and Xiao et al. [32,33] exhibit close stress/time to rupture points, which are approximated by a regression line predicting the long-term creep strength of 100 MPa (Figure 8c).
The SAVE12AD steel is predicted to have the long-term creep rupture strength of 80 MPa at 650 °C for 100,000 h (Figure 8d) [36].

4.3. The 10% Cr and 11–12% Cr Steels

The main problem of the 10–12% Cr steels is a drop in creep resistance in the long-term creep region. Creep degradation is associated with the instability of the tempered martensite lath structure under creep conditions, which can be caused by such microstructural changes as:
preferential recovery of martensitic microstructure near the PAG boundaries;
Z-phase formation and disappearance of MX strengthening precipitates;
Laves phase formation, etc.
Thus, the creep strength of the 11% Cr conventional P122 steel is the same as that of P92 steel at high stresses, whereas, at low stresses, creep resistance deteriorates. Therefore, for the developing 11–12% Cr steels with excellent oxidation resistance due to high Cr content, the main task is to reach the creep strength level at least of the conventional 9% Cr P92 steel.
The experimental 10Cr [8] and 10Cr-0.2Re steels [39,40,41] with 80 ppm B/20–30 ppm N, as compared with the 10% Cr NF12 steel (10Cr-2.5W-2Co) with conventional B and N contents, show advantages in the long-term creep region. These steels exhibit higher creep strength at low stresses (<140 MPa), whereas, at high stresses, new steels are significantly less resistant to creep than the NF12 steel and closer to the creep level of the P92/P122 steels.
The 10Cr steel does not demonstrate the creep strength breakdown up to approximately 40,000 h, which results in a high predicted long-term creep rupture strength of 110 MPa at 650 °C for 100,000 h [8]. This value is sufficiently higher than that for the TOS 110 steel, which was developed for application at 630 °C, and its long-term creep rupture strength for 100,000 h at 630 °C is 100 MPa and at 650 °C is nearly 80 MPa [70]. It can be concluded that modification of the TOS 110 steel by a decrease in the N content from conventional 200 ppm to 30 ppm resulted in a higher long-term creep strength of the 10Cr steel [8]. Unfortunately, there are no data on the creep strength curve of the TOS 110 steel in the literature. Therefore, the 10Cr steel is also compared with the HR1200 steel (11Cr-2.7W-2.7Co-200 ppm B/200 ppm N), which was developed at the next step for application at 650 °C by increasing the W content to 2.7% and B content to 200 ppm [43,45].
The creep strength of the experimental Re-modified 10Cr steel is significantly higher in comparison with the 10Cr steel at ≥140 MPa. Its time to rupture is about five to ten times longer than that of the 10Cr steel. However, a marked creep strength breakdown takes place at <140 MPa (~10,000 h), leading to a decrease in the long-term creep strength of the 10Cr-0.2Re steel [40,71]. It should be noted that, at low stresses ≤140 MPa, the creep strength curve for the 10Cr-0.2Re steel is almost the same as for the TOS 203 steel. In contrast, the short-term creep strength of the 10Cr-0.2 Re steel is significantly lower as compared with the TOS 203 steel.

5. Strengthening Factors in the 9–12% Cr Steels

5.1. Solid Solution Strengthening

The solid solution strengthening of considered steels is mainly caused by chromium, tungsten, molybdenum, and cobalt. A small amount of rhenium is also used as a solid solution strengthening element in the 10Cr-0.2Re and TOS 203 steels.
Cr is the most strengthening element in the high-chromium steels. An increase in the Cr content from 9 to 12% can markedly increase the creep strength in the short-term creep region, for example, for the 11–12% Cr TAF650, 12Cr, and 12Cr-Ta steels, but not for the SuperVM12 steel (Figure 15). The W content varies from 1.5 to 3% in the steels, the Mo content varies from 0 to 0.7%, and Co content varies from 2 to 3%. Most of the 9% Cr advanced steels contain 3%W-0%Mo. On the other hand, the 9Cr-1.5W-0.6Mo steel shows the same creep strength in the short-term region (Figure 7 and Figure 8). The 10Cr-0.2Re steel shows higher creep strength up to the appearance of the creep strength breakdown in comparison with the 10Cr steel, probably due to the higher W content (3% instead of 2%), although the Mo equivalent (Mo + 1/2 W) is the same in these steels (1.7%), as well as due to the presence of Re. In the 11–12% Cr steels, there is also no distinct correlation between the content of W, Mo, and Co and the creep strength.

5.2. Boundary and Sub-Boundary Strengthening

In the 9–12% Cr steels with tempered martensite lath structure, the sub-boundary hardening enhanced by fine distributions of precipitates along boundaries gives the most important strengthening mechanism for creep compared to the PAG boundaries [7]. The sub-boundary strengthening is inversely proportional to the width of lath or subgrains. The width of lath is nearly the same in most of the steels in the as-tempered state and comprises 300–400 nm (Figure 19). Evolution of lath width during creep depends on the stability of precipitates located on the lath boundaries, not only the M23C6 carbides but also the Laves phase. Figure 19 shows that the lath width slightly increases during creep. A sharp increase in the lath width or subgrain size above 1 μm usually corresponds to the transformation of the lath structure into the subgrain structure and the creep strength breakdown appearance as for the 9Cr-1.5W-3Co, 10Cr-0.2Re, and TAF650 steels.
Figure 19. Evolution of the lath width during creep in the 9–12% Cr steels. Data from [8,16,17,21,22,38,39,49,71].

5.3. Dislocation Strengthening

The dislocation density in the steels before creep depends on the tempering temperature. Figure 20 shows that the dislocation density comprised about 1–3 × 1014 m−2 in the as-tempered steels at 750–780 °C. Lower tempering temperatures increased the dislocation density, for example, up to 7.7 × 1014 m−2 in the TAF650 steel tempered at 680 °C [49].
Figure 20. Evolution of the dislocation density during creep in the 9–12% Cr steels. Data from [8,16,17,21,22,38,39,49,71].

5.4. Precipitation Strengthening

Threshold stress. The M23C6 carbides and Laves phase particles are the main phases precipitated on the boundaries of lath, blocks, packets and PAGs.

The stresses created by small precipitates cause the threshold stress for the onset of creep. Dudova et al. examined the creep behavior of the 10Cr steel in terms of threshold stress and found a high threshold stress of 111.5 MPa, which is about 30% larger than that in the P92-type steel [8]. The calculation of the stresses required for a dislocation to pass through particles at the minimum creep rate stage was carried out. The Orowan mechanism (to bow a dislocation between two particles), climb mechanism (to generate the additional length of dislocation to climb over an obstacle), and detachment mechanism (to detach the dislocation from an attractive particle after finishing the climb) were taken into account. It was revealed that the threshold stress is associated with the stress required for detachment of dislocations from M23C6 carbides, MX carbonitrides, and Laves phase particles after finishing the climb. Further, essentially stable M23C6 carbides exert the main part of threshold stress. The detachment stress can be calculated by equation [79]:

M23C6 carbides. Figure 23 shows that the mean sizes of M23C6 carbides in the as-tempered steels are in the range 50–100 nm. In most of the considered high B and low N steels, there is only a slight increase in the mean size at long-term creep over 10,000 h to about 120–150 nm, for example, in the MARBN, G115, 9Cr-1.5W-3Co, 10Cr, 10Cr-0.2Re, and SuperVM12 steels. Whereas, in the steels with conventional B content (the 9Cr-2W-3Co and 9Cr-3W-3Co steels), the pronounced coarsening of carbides to 200–300 nm occurs. Therefore, this confirms that the enrichment of steels by boron leads to a decrease in the size of carbides and an increase in their coarsening resistance under creep conditions.
Figure 23. Evolution of the mean size of M23C6 carbides during creep in the 9–12% Cr steels. Data from [7,8,16,17,21,22,27,38,39,49,54,71].
Laves phase. Laves phase particles grow during creep deformation, and their coarsening negatively affects the creep strength. Figure 24 presents the evolution of the mean size of Laves phase particles in some 9–12%Cr steels. Laves phase particles slightly grow up to approximately 1000 h of creep, and then the rapid coarsening occurs.
Figure 24. Evolution of the mean size of Laves phase particles during creep in the 9–12% Cr steels. Data from [8,16,17,21,22,27,38,39,49,71].
MX phase. A decrease in the N content can lead to a decrease in the fraction of the MX phase in the steels with a high B and a low N content as compared to steels with conventional B and N contents. Consequently, the precipitation strengthening is reduced due to a small fraction of fine, homogeneously distributed in the lath interiors, MX particles. It can mainly concern the steels with a N content of lower than 70 ppm.
Pinning pressure. Fine and stable boundary precipitates effectively prevent the migration of lath boundaries due to high pinning pressure [8,16,18].

6. Summary

The review of the 9–12% Cr heat-resistant martensitic steels with increased boron and decreased nitrogen contents in comparison with similar steels with conventional B/N contents shows the following:
The approach to alloying by the increased B (80–150 ppm) and decreased N (30–100 ppm) contents is successfully applied to advanced 9% Cr, as well as 10–12% Cr martensitic steels. The predicted long-term creep rupture strength at 650 °C for 100,000 h attained:
80–110 MPa for the 9% Cr steels, such as the MARBN (9Cr-3Co-3W-100–140 ppm B/30–80 ppm N), G115 (9Cr-3Co-3W-1Cu-150 ppm B/80 ppm N), and SAVE12AD (9Cr-3Co-3W-0.04Nd-100 ppm B/100 ppm N) steels, which sufficiently exceeds the creep strength for the 9% Cr Co-free P92 steel (~60 MPa) and 3% Co-modified P92 steels (~65–70 MPa) with conventional B (~50 ppm) and N (~500 ppm) contents;
110 MPa for the 10% Cr experimental steel (10Cr-3Co-2W-0.7Mo-80 ppm B/30 ppm N), which sufficiently exceeds the creep strength for the 11%Cr Co-free P122 steel (~45 MPa), 10% Cr Co-containing NF12 steel (10Cr-2Co-2.5W-50 ppm B/200 ppm N) (40 MPa), and advanced 10Cr-3Co-1.8W-100 ppm B/200 ppm N (TOS 110) steel (80 MPa);
Approximately 65 MPa for the 11–12% Cr steels, such as SuperVM12 steel (11Cr-1.8Co-2W-0.5Mo-140 ppm B/110 ppm N), which is sufficiently higher than that for the previous VM12-SHC steel (~50 MPa), 11%Cr Co-free P122 steel (~45 MPa), and slightly higher than that for 9% Cr P92 steel (~60 MPa) with conventional B (~50 ppm) and N (~500–600 ppm) contents;
An increase in the B content and a decrease in the N content enhance the creep resistance in the long-term region at a low stress, while the creep strength in the short-term region at the higher stresses corresponds to that for the steels with conventional B and N contents;
A high B content at a low N content effectively increases the coarsening resistance of M23C6 carbides during creep at 650 °C in all considered steels;
A positive effect of B on the creep strength is associated with enrichment of M23C6 carbides located near the PAG boundaries, which increases their coarsening resistance during creep. Stable M23C6 carbides are able to impede the recovery of the lath structure in the vicinity of PAG boundaries and, hence, retard the local deformation in the PAG boundary regions;
Even if the B content was already high in steel, then lowering the N content less than the solubility limit can increase the long-term creep rupture strength due to full utilization of soluble boron in the matrix and M23(B,C)6-type carbides;
Precipitation of small MX particles during the transient stage of long-term creep effectively reduces the creep rate and increases the time to rupture, as was shown in the MARBN steel with 140 ppm B and 79 ppm N and the 10Cr experimental steel with 80 ppm B and 30 ppm N;
An increase in the B content and a decrease in the N content is an effective way to enhance the dispersion strengthening by: fine and highly coarsening-resistant M23C6 carbides, which, in turn, can provide the slower coarsening of Laves phase particles; moreover, the MX phase, despite the low N content, can enhance the creep strength in the long-term region. This provides a stable tempered martensite lath structure over a long creep time and prevents the transformation of the lath structure into a subgrain structure.

This entry is adapted from the peer-reviewed paper 10.3390/met12071119

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