2018 Volume 58 Issue 6 Pages 1146-1154
New ferritic heat-resistant steels with high nitrogen content were prototyped and their microstructures and mechanical properties at high temperature were evaluated. The addition of 0.3 mass% N into ferritic steels was achieved without the formation of blowholes by applying pressurized melting methods under an atmosphere of up to 4.0 MPa. The high-nitrogen ferritic heat-resistant steels contained several kinds of nitrides within the lath martensitic structure. V-rich coarse particles were identified as crystallized MN. Fine VN or Cr2N particles were precipitated on the martensitic grain boundaries such as prior-austenite grain boundary, packet boundary, block boundary and lath boundary depending on the V content. The martensitic structure of the high-nitrogen steels contained a hierarchical microstructure including martensitic laths, blocks, packets, and prior-austenitic grains. These martensitic structures satisfied the Kurdjumov–Sachs relationship as with conventional carbon steel. The creep strengths of the prototyped steels were comparable with those of Gr. 91 steel, albeit lower than those of Gr. 92. Additional precipitates other than nitrides are required for further strengthening of the developed steels.
High-chromium ferritic heat-resistant steels are used as high-temperature structural materials for thermal power plants. The body-centered cubic ferrite phase has a low thermal expansion coefficient and excellent thermal conductivity, making it suitable for high-temperature components of thermal power plants that must be resistant to repeated cycles of heating and cooling. The efficiency of such plants is improved by increasing the steam temperature. In an effort to meet the demand for superior high-temperature strength and oxidation resistance, research into improving the chemical composition and microstructure of ferritic heat-resistant steels has continued.1) ASME Grade P/T 92 (Gr. 92) steel is the strongest of the commercially available ferritic heat-resistant steels, with a maximum operating temperature of approximately 620°C.2) A key characteristic on the chemical component of Gr. 92 steel is that it contains approximately 9% Cr (hereinafter, all of the added element amounts are expressed in units of mass%) to improve the oxidation resistance. Furthermore, Cr is a major component of M23C6 carbide, which plays an important role in stabilizing the microstructures in the steel at high temperatures.3,4) MX-type carbonitrides are also important precipitates that are responsible for the high-temperature strength of Gr. 92 steel. MX-type carbonitrides are more stable than M23C6 carbides at high temperature, and are particularly important for maintaining long-term creep strength.3) Tempered lath martensite is a characterized microstructure present in high-Cr ferritic heat-resistant steel. Martensite lath, which is the smallest structural unit of the martensite microstructure, is the main cause of creep strengthening of the steel,3,5) and precipitates directly contribute to its strength as obstacles to dislocation motion as well as by maintaining a fine lath structure during creep deformation. Based on the basic theory of creep strength, the martensite microstructure, with its many high-angle grain boundaries such as prior-austenite grain boundaries, packet boundaries and block boundaries, and low-angle grain boundaries such as lath boundaries and subgrain boundaries (Here, in this paper “grain boundary” means all these boundaries unless otherwise noted.), is disadvantageous in terms of creep strength. However, material properties such as toughness are also important for the construction and operation of plants, so practical steels have a martensite microstructure.6)
Currently, the development of an advanced ultra super critical pressure (A-USC) plant that raises the steam temperature up to 700°C is being promoted.2) Since there are no ferritic heat-resisting steels capable of withstanding long-term use at 700°C, the adoption of Fe–Ni based alloys and nickel based heat resisting alloys is considered for the boiler material of the A-USC plant.7) Meanwhile, as mentioned above, since ferritic heat-resisting steels excellent in basic physical properties are suitable as materials for thermal power boilers, efforts are being made to improve the high-temperature strength and oxidation resistance of ferritic heat-resistant steels. Ferritic heat resistant steels having ferrite microstructure instead of tempered lath martensite microstructure has been developed as a material assumed to be used at 700°C or higher corresponding to A-USC condition.8,9) Novel steels excellent in creep strength which are strengthen by precipitation of intermetallic compounds such as the Laves phase and the χ phase8) or NiAl and Ni2TiAl9) has been developed. However, significant reduction in toughness is a problem in the steels with ferritic microstructure.10) To the best of the authors’ knowledge, as a ferritic heat resistant steel whose matrix phase is a tempered lath martensite structure, only a steel developed in National Institute for Materials Science (NIMS) achieves sufficient creep strength in the temperature range up to 650°C.2,6,11)
The upper temperature limit of ferritic heat-resistant steels is defined by not only the creep strength but also the oxidation resistance. The oxidation resistance of conventional ferritic heat-resistant steels is generally improved by the addition of Cr. However, to simultaneously achieve the formation of the martensite microstructures in the presence of the large added amounts of Cr, it is also necessary to add austenite-stabilizing elements to counteract the ferrite-stabilizing effect of Cr. Masuyama et al. investigated the influence of N addition on the oxidation resistance of high-Cr ferritic heat-resistant steel at 600–650°C, and they reported that the addition of about 0.15% N was extremely effective for improving the oxidation resistance of the steel.12) In other words, N addition represents a promising method in alloy design to achieve both improvement of the oxidation resistance and formation of the martensite microstructure. In addition, MX particles, which are the main strengthening phase in conventional steels, precipitate as MN-type nitrides in nitrogen-added steels. Consequently, if high concentrations of N could be added to ferritic heat-resistant steels, there exists the possibility that a large amount of MN particles could be precipitated as the strengthening phase.
Although the addition of N to ferritic heat-resistant steels has the potential to improve the material properties as described above, research into nitrogen-added steels has so far been limited by the difficulty of adding N into the ferrite phase. Considering the Fe–N binary system, the amount of N that can be dissolved in the ferrite phase is about 0.1% at most. Therefore, when the melting is performed at atmospheric pressure, even if a large amount of N is added to the molten steel, it is gasified during solidification. To suppress this gasification and achieve the addition of large amounts of N into the ferrite phase, it is necessary to forcibly dissolve N in the solid phase using the pressurized melting method.13) In this study, we report on the microstructural characterization of nitrogen-added steel and the evaluation of its high-temperature strength to provide fundamental information regarding its potential application to heat-resistant steels.
Thermal equilibrium calculations were performed using the Thermo-Calc software to determine the appropriate alloy composition for forming the dispersion-strengthened lath martensite microstructure. SSOL 5 was used as the thermodynamic database. LIQUID (liquid phase), FCC_A1 (austenite phase or MX-type carbonitride), BCC_A2 (ferrite phase), HCP_A3 (Cr2N nitride), M23C6 (M23C6 carbide), and LAVES_C14 (Fe2W Laves phase) were input as the possible equilibrium phases. From the results of the thermal equilibrium calculations, three types of steels with different chemical compositions were selected and prototyped.
The pressurized electro-slag remelting (P-ESR)13) and pressurized induction melting14) methods were used to produce high-nitrogen ferritic heat-resistant steel. In the P-ESR method, since N is added as a solid nitride, it is necessary to account for the dissolution of N from the pressurized gas into the molten steel in order to accurately adjust the N addition amount to the target value. Therefore, the dissolution amount of N with respect to the gas pressure was measured in a preliminary test prior to the actual steel prototype. For the preliminary test, the high-Cr ferritic stainless steel SUS410 (12Cr-0.13C) was used. Two P-ESR electrodes were prepared: an electrode consisting of only SUS410 and an electrode filled with FeCrN powder in an amount calculated to correspond to 0.3% N added to SUS410. The P-ESR melting was performed at gas pressures of 0.5 MPa, 1.0 MPa, and 2.0 MPa. The composition of the pressurized gas used was 90% He/10% N2.
Based on the results of the preliminary tests, in the production of the nitrogen-added ferritic heat-resistant steel by the P-ESR method, an N-free ESR electrode and a steel pipe filled with FeCrN powder corresponding to the target N amount were melted together under a total pressure of 4.0 MPa and a nitrogen partial pressure of 0.4 MPa. The slag was 4N purity CaF2. To melt the slag, the current and voltage were set to 2.8 kA and 24 V, respectively. The molten steel was poured into the water-cooled mold through the melted slag. The P-ESR ingots had a cylindrical shape with a diameter of approximately 100 mm, a length of 350 mm, and a weight of about 15 kg. After dye penetrant testing, the P-ESR ingot was cut into upper, middle, and lower portions, and quantitative analysis of the nitrogen amount in each portion was performed. For one ingot in which cracking had confirmed by the dye penetrant testing, the fractured portion was cut and subjected to hot forging at a temperature of 900–1200°C. Finally, it was formed into a bar of 12–15 mm square by hot rolling at 1200°C.
To determine whether N addition could be performed regardless of the manufacturing method, nitrogen-added ferritic steel with the same composition as the steel produced by P-ESR was prepared by the pressurized induction melting method. The gas pressure was 2.0 MPa and the composition of the pressurized gas was 100% N2. In the pressurized induction melting method, N was introduced into the N-free molten steel from a pressurized N2 gas atmosphere. Molten steel from the pressurized melt was sampled and the amount of N was determined. When the amount of N in the molten steel had reached a predetermined level, the steel was cast into a mold while maintaining a N2 gas pressure of 2.0 MPa. The weight of an ingot prepared by pressure induction melting was approximately 500 kg. After cutting off the lower portion of the ingot, annealing treatment was performed at 700°C for 5 h. Subsequently, hot forging was performed up to φ 20 mm or φ 80 mm. Both the P-ESR material and the pressurized induction melting material were subjected to homogenization at 1200°C for 30 min and tempering at 780°C for 1 h.
Microstructural observation using a scanning electron microscope (Carl Zeiss, Ultra 55) was performed for each steel sample after the heat treatment. The specimens for the microstructural observation were cut from the steel rod and mirror finished by wet polishing. Then, colloidal silica (50 nm) polishing was performed. The microstructural evaluation was performed by observation of the backscattered electron image, elemental analysis of the precipitated phase by energy-dispersive X-ray spectroscopy (EDS), and analysis of the lath martensite structure by the electron backscattered diffraction (EBSD) method. The SEM observation was carried out at an acceleration voltage of 15 kV. For the EDS measurement, the plane analysis was performed in the region including the second phase particle using N Kα, V Kα, Cr Kα, Fe Kα, and Nb Lα radiation. The step interval for the EBSD measurement was 0.5 μm, and four fields of view of 200 μm × 200 μm were measured for each steel sample. The crystal orientation maps were drawn with a regular hexagonal lattice pattern.
Tensile tests and creep tests were performed on the prototype steel. The shapes of the test pieces used for the tensile and creep tests were uniaxial with a gauge portion diameter and length of 6 mm and 30 mm, respectively. The strain rate used in the tensile test was 0.3%/min up to a strain of 1.0%, then 7.5%/min until breaking, and the temperature range of the test was from room temperature to 750°C. The creep test was carried out in the stress ranges of 80–140 MPa at 650°C and 40–100 MPa at 700°C.
The detailed chemical compositions of the prototype nitrogen-added ferritic heat-resistant steels were determined as follows. First, it was decided to add 9% Cr to impart an oxidation resistance equivalent to those of conventional high-Cr ferritic heat-resistant steels. Among the group IV and group V elements that tend to form compounds with N as MN-type nitrides, V was selected with the aim of strengthening the steels by precipitating vanadium nitride (VN). From the thermal equilibrium state calculations, it was expected that the VN would crystallize in the molten steel when excessive amounts of V and N were added. Therefore, to suppress VN crystallization, the added amounts of V and N were limited to a maximum of 1.3% and 0.3%, respectively. To suppress the precipitation of carbides, the added amount of C was set to be below 0.01%. Based on the above considerations, the target chemical compositions of the prototype steels were established as summarized in Table 1. HN-A contains 1.3% V and 0.3% N. HN-B has a similar composition to HN-A but also contains 1% W for solid solution strengthening and 2% Co to stabilize the austenite phase. HN-C was prepared with a reduced amount of V compared with that of HN-B to allow investigation of the microstructure and strength when Cr nitride is the major precipitate. These HN-X series steels were produced using the P-ESR method. In addition, to allow evaluation of the importance of the production method, the pressure induction melting method was used to prepare the PN-B sample with the same chemical composition as HN-B. Figure 1 shows the relationships between the phase fraction and temperature for each phase of the prototype steels, as calculated using the Thermo-Calc software. Based on these results, the major precipitates for HN-A and HN-B were expected to be MX nitrides (VN), whereas that for HN-C to be Cr2N nitrides. For HN-B and HN-C, precipitation of the Laves phase and small amounts of M23C6 were also expected for lower temperatures below 750°C. It is known that it precipitates as M23(C,N)6 in nitrogen-containing steels,15) but in the result of Thermo-Calc calculation, N contained in M23C6 was zero. This is considered to be a database problem used in this research. However, since the carbon content of the prototype steels is as low as 0.01%, almost no M23(C,N)6 was observed in the microstructure. Therefore, it is considered that the inclusion of N in M23(C,N)6 has almost no influence on the formation of other nitrides.
Relationships between the phase fractions and temperature for the prototype steels, as calculated using the Thermo-Calc software.
Figure 2 shows the amount of N in the materials prepared at different gas pressures in the preliminary tests. Even when the total pressure was increased to 2.0 MPa and the nitrogen partial pressure was increased to 0.2 MPa, no increase in the N amount was observed for the steel sample not containing the FeCrN powder. These results demonstrate that dissolution of N from the pressurized gas atmosphere into the molten steel did not occur. During P-ESR melting, the molten steel is shielded from the pressurized gas atmosphere by the slag and the gas pressure is added to the molten steel through the molten slag. Therefore, it was considered that the dissolution of N from the pressurized gas to the molten steel was suppressed by the shielding effect of the slag. In contrast, in the steel containing FeCrN powder, approximately 0.21% N was added to the material at a total pressure of 0.5 MPa or higher. However, even when the total pressure was increased to 2.0 MPa, the amount of added N did not increase or reach the target value of 0.3%. These results indicated that the supersaturated N was gasified and desorbed when passing through the δ-ferrite region with low nitrogen solubility during the solidification process, even at a pressure of 2.0 MPa. Therefore, for the production of the prototype steels by the P-ESR method, the gas pressure was increased to 4.0 MPa. In addition, in anticipation of the partial desorption of N, the charged N amount was set to 0.33% with respect to the target N addition amount of 0.3%.
Relationship between the gas pressure and nitrogen addition amount.
Table 2 summarizes the production conditions of the prototype steels and the actual N content at the upper, middle and lower parts of the ingots. N2 blowholes were not detected in any of the samples, and the analytical results confirmed that all the prototype steels contained approximately 0.3% N. For the steel sample produced by the P-ESR method, the desorption of N during solidification was suppressed by setting the gas pressure to 4.0 MPa. The N was added homogeneously to the steel without segregation between the upper, middle, and lower parts of the P-ESR ingot. Among the three P-ESR samples, HN-C was found to contain a lower N amount, which was ascribed to the lower V amount in this sample; in other words, it was considered that V, which was not present in the preliminary test material SUS410, contributes to the introduction of N into the steel. Furthermore, PN-B, for which the N was introduced from a pressurized gas atmosphere, was also found to contain the target amount of N and no blowholes were observed in the ingot. Therefore, nitrogen-added steel was also successfully produced by pressurized induction melting at 2.0 MPa.
Dye penetrant testing was performed for the HN-C ingot, for which large cracks were observed after quenching. The appearance of the dyed HN-C ingot is presented in Fig. 3. As shown in Fig. 3(a), cracks could be clearly observed from the bottom to the top of the ingot in the longitudinal direction. On the upper surface of the ingot shown in Fig. 3(b), the crack had reached the center of the ingot and branched out in two directions. The cross sections of the ingot are shown in Fig. 3(c). Although the crack had reached the center of the ingot in the upper portion, it became shallower toward the lower portion. Consequently, hot forging was performed for HN-C after cutting at the positions indicated by the red solid line and blue broken lines in Fig. 3(c). Since the cracking had occurred in the quenched specimen, the cracks generated in the ingot of HN-C are considered to be due to quench cracking owing to martensitic transformation during cooling. In contrast, no cracks were observed in the ingots of HN-A, HN-B, or PN-B. The amount of solid solution N in HN-C after quenching should be larger than that in the other steels because the addition amount of V was relatively small. Therefore, the transformation strain accompanying the martensitic transformation became greater, and this is assumed to be the cause of the quench cracking in HN-C. To prevent the quench cracking of HN-C, the lower temperature limit of the hot forging was set to be above its A3 point and tempering was promptly performed after quenching. As a result, quench cracking could be suppressed with the final bar of HN-C.
Photographs of the dyed HN-C ingot: (a) crack in the longitudinal direction, (b) cracks on the upper surface, and (c) cracks on the cross sections.
Figure 4 shows the backscattered electron (BSE) images of samples of each of the steels after tempering. It is understood that the lath martensite structure is formed by the addition of N in all of the four steels, and most of the matrix exhibited a lath martensite structure. Coarse rectangular particles with a particle size of several microns were observed in all of the steels. These particles are considered to have crystallized at the time of melting, since they were also observed in the ingots after melting. In addition, in the case of HN-C, many spherical particles with diameters of several hundred nanometers were also observed within the lath grain. So, these particles probably precipitated during normalization or crystallized at the time of solidification. Bright particles are also observed in the BSE images. Figure 5 is a higher magnification BSE image. Many fine dark particles were observed, mainly on the grain boundaries of the martensitic structures. The particles on the high-angle grain boundaries tended to be coarsened compared to those on the lath boundaries. The EDS elemental analysis results for the dispersed particles observed in the BSE images are also shown in Fig. 5. Figure 5(a) presents the EDS maps for the coarse particles observed in HN-A. These coarse rectangular particles are composed of Cr, V, Nb and N. A EBSD measurement revealed the crystal structure of the particle is FCC. Therefore, these particles were confirmed to be crystallized MN (M: Cr, V, Nb). Constituent elements of the bright particles are V and N. So, these bright particles considered to be VN, and its brightness is caused not by backscattered electrons but by secondary electrons emitted from protruded particle from the sample surface because of its coarseness. Figure 5(b) shows the EDS analysis results for the fine particles observed on the grain boundaries in tempered HN-A. These fine particles were also composed of V and N. They were identified as precipitated VN from the tempering, since they were found on the grain boundaries of the martensite structures. Figure 5(c) presents the EDS maps of the fine particles observed on the grain boundaries in HN-C. These particles contained Cr as a main constituent element, so they were identified as Cr2N based on the Thermo-Calc calculation results shown in Fig. 1(c). In addition, as shown in Fig. 5(c), it was confirmed that the spherical particles with diameters of several hundred nanometers were nitrides containing V, Cr and Nb, although it is unclear whether these particles were crystallized or precipitated. The above observations demonstrate that a structure in which nitride particles were dispersed in the martensite structure was obtained for all of the prototype steels.
SEM backscattered electron images of the prototype steels. Martensitic lath structure and rectangular coarse particles in (a) HN-A, (b) HN-B, (c) HN-C, and (d) PN-B.
SEM-EDS element maps of the second phase particles in the HN-A and HN-C steels. (a) V-rich coarse particles in HN-A, (b) V-rich fine particles in HN-A, and (c) Cr-rich fine particles in HN-C.
Figure 6 shows the high-angle grain boundary maps obtained from the EBSD measurements of samples of each of the steels. The martensite structures of the prototype steels were distinguished and are shown on the block boundaries, the packet boundaries, and the prior-austenite grain boundaries, assuming that the martensite structures in these steels satisfied the Kurdjumov–Sachs (K–S) relationship.16) From Fig. 6, the martensite structures of the prototype steels with added N also satisfied the K–S relationship as well as the conventional steel such as Gr. 91 and Gr. 92. The prior-austenite grain sizes of HN-A, HN-B, and PN-B were approximately 50 μm, whereas that of HN-C was slightly finer at about 10 μm. This fine prior-austenite grain sizes for HN-C was ascribed to the suppression of grain boundary migration by the many spherical particles which did not solute during hot working and normalization. In addition, comparing HN-B and PN-B, the crystal grain size of PN-B was found to be slightly coarser. This finding was attributed to the fact that the PN-B was produced from a large ingot that was held at a high temperature for a long time during the hot-working process.
High-angle grain boundary maps obtained by SEM-EBSD analysis. (a) HN-A, (b) HN-B, (c) HN-C, and (d) PN-B. The red, blue, and black lines represent martensitic block boundaries, martensitic packet boundaries, and prior-austenitic grain boundaries, respectively.
For lath martensitic steel, it is known that the creep strength increases as the width of the martensite lath becomes narrower.3) According to the theory of quantitative microscopy, the lath width is inversely proportional to the lath boundary length per unit area on the observation surface.17) Figure 7 shows the boundary length per unit area for the lath boundary, block boundary, packet boundary, and prior-austenite grain boundary as determined from the EBSD data. For comparison, the results obtained for Gr. 92 steel18) are also shown. However, the Gr. 92 steel used here for comparison was tempered at 760°C, which is lower than the prototype steels. It should be noted here that, in this study, the low-angle boundaries with a misorientation of 1° to 5° were defined as the lath boundary. As for the high-angle boundary lengths per unit area presented in Fig. 7(a), HN-A, HN-B, and Gr. 92 steel had a similar scale of the martensite structure. On the other hand, HN-C has a remarkably higher packet boundary density than other steels. As shown in Fig. 6(c), for HN-C, the fine prior-austenite grains are divided into several packets, which are composed of single or a few blocks. Therefore, packet boundary density became very high despite the block boundary density is not so high. According to Fig. 7(b), the lath boundary lengths per unit area for each of HN-A and HN-B were not significantly different from that of the Gr. 92 steel. Considering that the tempering temperature of HN-A and HN-B is 20°C higher than that of Gr. 92, recovery of lath during tempering is sufficiently suppressed in these prototype steels. In contrast, the lath boundary lengths per unit area for HN-C and PN-B were obviously low. It is considered because Cr2N precipitated on lath boundaries in HN-C has a less suppressive effect on recovery of lath compared to VN in HN-A and HN-B. For PN-B, there is a possibility that precipitation of VN during annealing and hot forging performed for the large ingot influenced the dispersion state of VN after tempering. Although the quantitative evaluation is not carried out, if the dispersion state of VN is sparse, it can be understood that the coarsening of lath in PN-B. Considering the coarse lath structure of the prototype steels, high creep strength exceeding Gr. 92 steel cannot be expected for these prototyped steels, even in the as-tempered state.
Comparison of the grain boundary densities in the developed steels and Gr. 92 steel. (a) Densities of block, packet, and prior-austenitic grain boundaries, and (b) density of martensitic lath boundaries.
Figures 8(a) and 8(b) show the temperature dependence of the 0.2% proof strength (0.2% PS) and rupture elongation, respectively, obtained by tensile tests for HN-A, HN-B, and HN-C. Among these three steels, HN-B and HN-C exhibited the highest and lowest values of 0.2% PS, respectively, at all of the temperatures tested. The higher 0.2% PS of HN-B relative to HN-A was considered to be due to solid solution strengthening by W. However, although it increased by 40 MPa or more at room temperature, the increment became smaller at temperatures above 600°C. Therefore, the solid solution strengthening effect of W is not large at high temperatures. HN-C also contains the same amount of W as HN-B, but the 0.2% PS of HN-C was similar to that of HN-A at room temperature and lower at high temperature. This suggests that the strengthening ability of Cr2N, which is the major strengthening phase in HN-C, is lower than that of VN, which is the major strengthening phase of HN-A and HN-B. The prototype steels exhibited sufficient ductility, with rupture elongation values of approximately 15% or higher. In other words, there was no decrease in ductility due to the coarse crystallized MN present in the prototyped steel. Furthermore, the greater rupture elongation of HN-C relative to HN-A and HN-B was considered to be due to its fine prior-austenite grains and packets.
Tensile test results for the developed steels from room temperature to 750°C: (a) 0.2% proof stress and (b) rupture elongation.
Figure 9 shows the relationship between the creep rupture time and the stress obtained from the creep test at 650°C and 700°C. The creep rupture strengths of Gr. 91 steel19) and Gr. 92 steel20) are also shown for comparison. The creep strength of HN-C was the lowest among the prototype steels, and those of HN-A and HN-B were comparable. In comparison with the conventional steel, the creep strengths of HN-A and HN-B were similar to that of Gr. 91 steel but lower than that of Gr. 92 steel. Similar to the results of the tensile tests, the creep strength results also indicated that the contribution of Cr2N to creep strengthening is smaller than that of VN. In addition, the shorter lath boundary length per unit area for HN-C compared with the other steels was also considered to be partially responsible for the decrease in creep strength. The creep rupture times of HN-A and HN-B was shorter than that of Gr. 92 steel, although the lath widths of these prototype steels prior to deformation were comparable with that of Gr. 92 steel. This suggests that the suppressing effect for the recovery of lath structure during creep deformation by the precipitated VN particles in the prototype steels is smaller than that of the M23C6 particles in Gr. 92 steel. By adding 1.5% W, the M23C6 particles in Gr. 92 steel are stabilized and the creep strength of the steel is improved.4) In contrast, in the comparison between HN-A and HN-B, no increase in creep strength due to the addition of W was observed. This was considered to be because W does not dissolve in the nitride, so the addition of W does not contribute to the stabilization of the nitride. It is also apparent that solid solution strengthening of W do not contribute to creep strength at 650–700°C. The creep strength would be improved if the amount of precipitation of VN could be increased with respect to HN-A and HN-B. However, according to the Thermo-Calc calculation results, it is difficult to further increase the creep strength using VN precipitates, because the addition of 1.3% or more of V causes an increase in coarse crystallized MN. Consequently, it can be concluded that to improve the creep strength of nitrogen-added steels, it is necessary to perform the strengthening using precipitates other than VN and Cr2N.
Relationship between creep stress and time to rupture at 650°C and 700°C.
To determine the influence of the addition of N to ferritic heat-resistant steel, nitrogen-added ferritic heat-resistant steels were prepared using the pressurized melting method and their microstructures and high-temperature strengths were evaluated. Consequently, the following conclusions were obtained:
(1) By applying the 4.0 MPa P-ESR method and the 2.0 MPa pressurized induction melting method, 0.3% N was added to 9% Cr ferritic steel and ingots were successfully obtained without the formation of blowholes.
(2) Coarse crystallized MN was present in the ferritic steel containing 0.3% N. Through tempering, the VN was precipitated on the martensite grain boundary in the steels to which 1.3% V had been added, whereas Cr2N was precipitated in the steel to which 0.6% V had been added.
(3) The matrices of the nitrogen-added steels after normalizing and tempering had a martensite structure, which satisfied the K–S relationship and contained microstructures including prior-austenite grains, packets, blocks, and laths, similar to those of conventional steels.
(4) The creep strengths of the nitrogen-added steels were lower than that of Gr. 92 steel, and HN-A and HN-B strengthened with VN had almost the same strength as Gr. 91 steel. The addition of 1% W to nitrogen-added steels contributed slightly to solid solution strengthening, but a stabilizing effect of precipitated nitride was not observed and is not effective for creep strengthening.
This work was carried out as part of the research activities of the Advanced Low Carbon Technology Research and Development Program (ALCA). Financial support received from the Japan Science and Technology Agency (JST) is gratefully acknowledged. In addition, we used the equipment of the Institute for Materials Science and Materials to produce the steel by the P-ESR method. We are grateful to Mr. Satoshi Iwasaki, Mr. Takaaki Hibaru, and Mr. Syuji Kuroda for operational assistance.