2016 Volume 56 Issue 9 Pages 1662-1667
Mechanical properties of a 9%Cr-ferritic steel grade P92 experimental alloy are studied. The effect of cooling rate on the hardenability was determined by means of continuous cooling diagrams and data provided by hardness measurements and microstructure observations. A fully martensitic microstructure after the solubilization treatment over a wide range of cooling rates was revealed. As this grade of steels is mostly supplied in tempered condition, tensile tests to determine the variation of the strength and ductility at temperatures ranging from 20 to 650°C were carried out after a treatment of 3 h at 760°C. In addition, Charpy V-notch tests were conducted to characterize the impact toughness of the steel and the ductile-brittle transition temperature. Finally, the creep strength was determined from creep tests in the range 550 to 650°C.
During recent years, the energy sector has favored extensive work to protect the global environment by increasing the thermal efficiency1) and decreasing carbon dioxide emissions in fossil fuel power plants.2,3,4) Such demand can be achieved by raising the operating temperatures and pressures of the plant, but to withstand more severe service conditions, new materials with superior long term creep and thermal fatigue properties are needed for boiler and turbine components.5)
Several ferritic steels containing 9–12%Cr, including the P92 grade, have been developed and widely used for operating temperatures up to 620°C because their excellent combination of mechanical, thermal and oxidation-resistant properties as well as void swelling resistance,6,7,8) acceptable room temperature properties9,10) and also good toughness, weldability and hot workability.11,12) Microstructural investigations support that the high creep strength of P92 grade has a complex structure with a large amount of effective barriers to dislocation movement: a lath structure inside the prior austenite grains, fine dispersion of small second phase particles, high dislocation density, and the presence of elements in solid solution. Most of these barriers can be controlled by the proper election of alloying elements, especially carbon, and by solution treatment and subsequent tempering or aging.
The combined effects of these mechanisms are complex and their contribution to strengthening may change during long-term creep deformation at high temperatures. Microstructural degradation of the advanced high Cr ferritic steels occurring during creep involve subgrain growth, particle coarsening, dislocation rearrangement, reduction of elements in solution, and precipitation of complex nitride Z phase, Cr(V,Nb)N, and/or Laves phases which has been identified as the main cause of limiting the life of the components.
Since the optimal properties of P92 steels are obtained by a dispersion of fine and thermally stable precipitate particles,13) the design of these alloys is focused on the size and distribution of specific precipitates and phases. Furthermore, precipitation hardening, grain size hardening, and solid solution hardening influences these properties which in turn are controlled by the heat treatment. The heat treatment employed is relatively simple and economically attractive consisting in austenization and air cooling to produce a practically fully martensitic structure, with minimal amounts, if any, of delta ferrite, which is generally regarded as detrimental for high temperature strength properties. Tempering is subsequently conducted to cause the precipitation of carbon and nitrogen in the supersaturated solid solution as carbide and nitride particles. This treatment also improves the ductility, toughness and impact strength at low temperatures. Even though these steels are used at high temperature where toughness is not a matter of concern, it is important that the steels show a good toughness at room temperature for fabrication and construction of the plants and for their periodically start up and shut down. As a consequence, the best toughness/creep compromise has to be obtained to guarantee all the requirements. Therefore, the thermal treatment of conventional ferritic steel grade P92 consists of austenization at about 1100°C followed by tempering at about 750°C.
With the objective of further improving the mechanical properties of ferritic steel grade P92, in the present investigation dilatometric studies were conducted to determine the microstructure of the steel as a function of cooling rate from the austenitization temperature. In addition, mechanical and creep tests at various temperatures were conducted to evaluate the mechanical properties.
The composition of the experimental casting is shown in Table 1. This steel was manufactured in a 40 kg vacuum furnace VCM-030. This furnace allows for castings under very exact processing conditions and to adjust precisely the composition and its homogenization. The ingot was hot rolled at 1150°C to obtain a billet 17 mm thick and 60×60 mm2 in cross section.
| C | Mn | Cr | Ni | Mo | V | W | B | Nb | N |
|---|---|---|---|---|---|---|---|---|---|
| 0.09 | 0.41 | 8.62 | 0.23 | 0.35 | 0.10 | 1.11 | 0.0063 | 0.050 | 0.0345 |
Continuous cooling transformation diagrams were determined to avoid cooling rates leading to the formation of deleterious microstructures with a high volume fraction of delta ferrite. The slight expansion/contraction associated to phase transitions was measured with a high-resolution dilatometer and the nature of transformations was determined by optical metallography and hardness measurements. Dilatometric samples, 10 mm long and 3 mm thick, were obtained directly from the ingots. These samples were austenitized at 1050°C for 30 min, and then cooled at rates ranging from 10°C/s to 0.014°C/s. The microstructures produced at the various cooling rates were revealed by light optical microscopy in polished samples that were etched by Vilella’s reagent. Prior austenite grain size was also measured to determine the effect of cooling rate.
Tensile properties in cylindrical samples, 6.25 mm in diameter, were measured according to Standard EN 10002.5 in the temperature range 20° to 650°C. Previous to testing, the samples were austenitized at 1050°C/60 min and subsequent air-cooled and annealed for 3 h at 760°C and air-cooled at about 5°C/s. The toughness of the steels was investigated by Charpy impact testing according to Standard EN 10045.1. Charpy impact tests were conducted in samples obtained directly from the ingots in the temperature range from −60 to 20°C. The ductile to brittle transition temperature (DBTT) was determined for impact energy equal to one-half of the difference between the respective minimum- and maximum- impact energies.
Finally, creep rupture tests were performed at 550, 600 y 650°C according to ASTM A-370 using applied stresses ranging from 350 to 80 MPa. Cylindrical creep samples 4 mm in diameter and 25 mm gage length were obtained directly from the ingots and were also austenitized at 1050°C/60 min and air cooled and further annealed at 760°C/3 h.
Previous to the dilatometric and mechanical investigations the evolution of grain size with austenization temperature was determined since it can strongly affect the mechanical properties and the formation of non-equilibrium phases, like bainite, martensite and retained austenite. The austenite grain growth kinetics of the ferritic steel grade P92 used for this study is shown in Fig. 1. After 20 minutes treatment, a considerable grain growth for austenitizing temperatures between 1000 and 1100°C is observed, and only slight changes occur at higher temperatures. This grain growth kinetics was related to the inhibition of austenitic grain growth of the different second phase particles dispersed in the microstructure, indicating that above 1100°C most of MX type carbonitrides and M23C6 carbides are dissolved. In general, a fine-grain microstructure promotes better mechanical properties of materials at room temperature since fine grain sizes are usually associated to an increase of hardness, yield strength, tensile strength, fatigue strength, impact strength and toughness.14) However, a significant improvement of creep strength can be obtained by increasing the number of particles precipitated during tempering.15) Besides, coarse-grained steels have better creep properties. Grain boundaries are regions where the rate of diffusion is high, and consequently creep strength, which is a diffusion controlled process, decreases with increasing grain boundary area. Thus, a compromise in the selection of the austenitizing temperature has to be reached to optimize the creep properties and to guaranty enough toughness at room temperature.
Evolution of the prior austenite grain size with austenization temperature of P92 steel. Samples were kept 20 min at the given temperature.
Continuous cooling transformation (CCT) diagrams provide a significant contribution to the understanding of the transformation behavior of steels. Figure 2 shows the CCT diagram for the P92 steel. Previously, the steel was austenitized at 1050°C/30 min. Curves at cooling rates in the range from 10 to 0.014°C/s (50°C/h) are shown. Between 10 and about 0.1°C/s the austenite transforms completely to martensite. For lower cooling rates both ferrite and perlite are formed.
Curves CCT for the P92 steel.
The microstructures obtained at three cooling rates, 10°C/s, 0.1°C/s and 50°C/h are shown in Fig. 3. Figure 3(a) corresponds to a rapid cooling rate and a fully martensitic microstructure free of δ-ferrite is obtained. The microstructure of Fig. 3(b) consists mainly of martensite with some patches of ferrite-perlite located mainly at the grain boundaries which are strongly etched by Vilella’s etchant. Finally, the microstructure of Fig. 3(c) is mainly ferritic with some particles of M23C6 carbides. These optical micrographs are taken at low magnification to reveal the general aspect of the microstructure, which is similar to that present in most of the of the 9–12%Cr steels. Although the carbides are hardly revealed due to its fine size, grain and lath boundaries in those micrographs are decorated by tiny dark spots corresponding to M23C6 particles. Fine MX particles present can only be observed at higher magnifications in TEM micrographs.
Microstructure of the P92 steel at various cooling rates: (a) 10°C/s, (b) 0.1°C/s and (c) 50°C/h.
The hardness as a function of cooling rate is given in Fig. 4. The hardness increases rapidly up to a cooling rate of 0.1°C/s when the curves no longer pass through the ferrite+perlite field and is maintained about constant at higher cooling rates of about 410 HV. These hardness values obtained for low and high cooling rates agree with the microstructures described and are typical of self-quenching steels with good hardenability.6)
Effect of cooling rate on Vickers hardness of P92 steel.
Previous to tensile testing, the samples were austenitized at 1050°C/60 min and subsequent air-cooled at about 5°C/s and tempered and annealed for 3 h at 760°C and air-cooled at about 5°C/s. The initial microstructure is giving in Fig. 5 and consists of tempered martensite containing M23C6 carbide particles decorating the prior austenite grain boundaries and martensitic laths and a dispersion of fine MX particles within these subgrains, which cannot be recognized in these micrographs.
Microstructure previous to tensile testing. The steel was austenitized at 1050°C/60 min and subsequent air-cooled at about 5°C/s and tempered for 2 h at 760°C and air-cooled at about 5°C/s.
Tensile tests were carried out and the yield stress (σy), taken at 0.2% strain, ultimate tensile strength (UTS) and elongation to failure at room and elevated temperatures were determined. Figure 6 shows the yield stress and ultimate tensile stress as a function of temperature for short time tensile tests. It should be noted that the microstructure of all the samples during deformation should be the same since they were tempered at a temperature, 760°C, that is higher than any testing temperature and for a time, 2 h, that is longer than that needed to start the tensile test, about 20 min. It is commonly accepted that the strength at room temperature in the case of tempered martensite is the result of a typical microstructure with a fine lath substructure, M23C6 carbides along the lath and prior austenite grain boundaries, fine MX precipitates homogeneously distributed in the matrix that interact with dislocations due its small size, a high dislocation density and some W and Mo in solid solution.1,7) However, it is very difficult to evaluate the individual contribution of these strengthening mechanisms since it is very complex to characterize quantitatively this microstructure.
Yield stress and UTS as a function of temperature for the P92 steel.
Values close to 600 and 800 MPa for the σy and UTS, respectively, at room temperature are observed. These stress values are higher than those observed in the ferritic steel AISI 409, having even higher Cr concentration, 9 vs 11%Cr, and the X10CrWMoVNb9 steel (EN10216-2).16) This is attributed mainly to the structure developed through the austenitizing conditions used and the higher cooling rate reached during quenching to ambient temperature. Since the samples prepared were relatively small, the high heat extraction rate produced over the whole specimen guarantees that the precipitation during quenching is negligible, and a maximum volume fraction of particles can precipitate in the matrix during tempering. In addition, dislocation density will vary with tempering temperature as well as size and distribution of carbide particles.
Increasing the testing temperature reduces the strength, moderately up to 400°C and rapidly for higher temperatures. As mentioned before, the strengthening mechanisms determining the yield and tensile stress of the material involve mainly the tempered martensitic lath structure, a high dislocation density, and second phase particles on subboundaries and within the ferrite laths. Up to 400°C, moving dislocations pass athermally through these obstacles by the applied stress, and thus Fig. 6 exhibits a relatively slight linear decrease of stress of about 20%. In this temperature range, the Young’s modulus decreases approximately linearly about 14% with increasing temperature (see for instance Product data Bulletin for 11Cr-Cb Stainless Steel. AK Steel, West Chester, 2007). As the temperature increases, diffusion assists dislocations in passing through the obstacles by climbing and also causes a gradual rearrangement and annihilation of dislocations into low-energy configurations with the progress of deformation (recovery of dislocation structure). These processes are a characteristic of the so called dislocation creep that occurs by glide and climb aided by vacancy diffusion. Dislocation creep starts contributing at about 0.5 Tm (Tm being the absolute melting temperature), which corresponds to 500–600°C for 9%Cr ferritic steels. Thus, at temperatures above 400°C both yield stress and tensile strength decrease rapidly by increasing the temperature, as observed in Fig. 6.
The same contributions explain the effect of temperature on ductility. Elongations at maximum load and to failure are represented as a function of testing temperature in Fig. 7. A total elongation close to 20% at room temperature is relatively high considering the high strength of the steel. This value increases with increasing temperature, reaching 35% at 650°C. Between room temperature and this temperature a minimum is observed at about 300°C. This minimum in ductility has been observed in many iron and nickel alloys as well as in other metals and alloys.17,18,19,20,21) The depth and width of this minimum of ductility, as well as the temperature at which appears, has been related with many factors like alloy composition, impurities content, heat treatment, state of applied stress, strain rate and grain size. As reported by Zheng et al.,21) explanations for this effect are quite diverse and are based on intergranular precipitates, grain boundary shearing or sliding, gas phase embrittlement decohesion of glide planes, dynamic strain aging, and grain boundary segregation of impurities. However, at present a complete understanding of intermediate temperature embrittlement cannot be provided.
Strain to failure and strain to maximum as a function of temperature for the P92 steel.
It is worth noting the large difference of elongations, at maximum load and at fracture, observed in Fig. 7 at testing temperatures above the ductility minimum. It must be pointed out that the value of uniform elongation plotted in this figure corresponds to the elongation up to the maximum load determined according to the Considere criterion. However, at high temperatures, when the contribution of dislocation creep to the overall plastic deformation becomes important, strain rates show strong non-linear strain rate dependencies on flow stress, and then the initiation of necking takes place at strains significantly greater than the Considere strain. Necking is closely associated with strain hardening phenomenon, and a material becomes unstable when it starts to deform faster than it can strain harden. As the temperature is increased above 400°C, the recovering and softening processes start to be operative, with the result of a reduction of the strain hardening and the increase of the total elongation to failure.
The toughness may be important for the solicitations of the tubes in power stations operating during stops for various reasons, such as maintenance and breakdowns. Figure 8 shows the Charpy V-notch absorbed impact energy of the steel as a function of temperature. The impact energy is about 200 J/cm2 at room temperature, a value that is similar to other high Cr steels.22) The ductile-to brittle transition temperature occurs at about −20°C. These high values are attributed to the fine prior austenite grain size and the absence of a carbide network at grain boundaries. It is to note that the variation in impact energy between 20 and 0°C is small. This should be attributed to the small presence of delta ferrite if any since a strong drop in impact energy occurs for samples containing concentrations higher than 2–4% of delta ferrite.23,24)
Charpy V-notch impact energy as a function of temperature.
Creep tests were conducted to determine the long-term stability of the material. Figure 9 shows the effect of temperature on the creep rupture strength as a function of time. The stresses applied were much lower than the proportional limit determined by short-time tensile tests. In these long-term tests the creep resistance of the steel is determined by the microstructure of intermetallic compounds and MX and its resistance to degradation caused by nucleation and individual growth of precipitates that can easily occur assisted by diffusion. As mentioned before, under these conditions creep is controlled by climb of dislocations over the obstacles, a process that is dominated by diffusion and therefore is strongly influenced by temperature. This dependence is observed in Fig. 9 where rupture strength for 10000 h is 210, 130 and 80 MPa at 550, 600 and 650°C respectively. This last value compares favorably to other P92 steels25,26,27) where the rupture strength is lower than the 80 MPa value obtained in this work.
Time to rupture at various temperatures as a function of tensile stress.
(1) The continuous cooling transformation diagram provided the conditions to precisely obtain fully martensitic, ferrite-perlite and mixed microstructures.
(2) Values close to 600 and 800 MPa for the σy and UTS, respectively, at room temperature are observed. Up to 400°C, strengthening mechanisms involving solid solution, dislocation substructure, and second phase particles, determine the yield and tensile stresses of the material. At temperatures above 500°C yield and tensile stresses decrease due to substructure recovery, and plastic deformation is controlled by a slip creep mechanism.
(3) High values of impact energy are obtained which is attributed to the lack of delta ferrite. The ductile-brittle transition temperature occurs at about −20°C.
(4) High values of about 80 MPa for the creep rupture strength is obtained at 650°C attesting the stability of the microstructure at high temperature.
Financial support from CICYT Projects MAT2012-39124 and PET2007-0475 is gratefully acknowledged.
