Materials Science Forum Vols. 706-709 (2012) pp 841-846 © (2012) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/MSF.706-709.841
Effect of Tempering on Mechanical Properties and Microstructure of a 9% Cr Heat Resistant Steel Valeriy Dudkoa, Andrey Belyakovb and Rustam Kaibyshevc Laboratory of Mechanical Properties of Nanoscale Materials and Superalloys, Belgorod State University, Pobeda 85, Belgorod 308015, Russia Email:
[email protected],
[email protected],
[email protected] Keywords: high-chromium heat resistant steel, martensite, tempering, carbides, mechanical properties.
Abstract. Effect of tempering temperature ranged from 400 to 720°C on mechanical properties and microstructure of a P92-type creep resistant steel was investigated. The hardness value of 400 HB, which was obtained after solution treatment, increased to 430 HB with increasing the tempering temperature to 525°С. Further increase in the tempering temperature resulted in gradual decrease in hardness, which approached a level of about 220 HB after tempering at 720°С. The equiaxed particles of MX-type carbonitrides with a size of about 30 nm were precipitated randomly after tempering under all conditions. At temperatures below 525°C, the tempered martensite lath structure (TMLS) was characterized by a random distribution of fine M3C-type carbides and MXtype carbonitrides. The precipitation of M23C6 was observed after tempering at T ≥ 525°C. At 525°C, the M23C6 carbides appeared as thin films on high-angle boundaries (HAB), while M23C6 particles having almost equiaxed shape and located on various boundaries including low-angle lath boundaries precipitate at higher temperatures. Introduction High-chromium steels with TMLS became widely used materials for components of fossil power plants during the last decade [1]. The ability of these steels to offer high creep resistance depends critically on the distribution of dispersed precipitates within their microstructures that is mainly dictated by the tempering conditions [2, 3]. In general, the tempering treatment results in precipitation of various carbides and carbonitrides. The main parameters, i.e. nature, origin, size and distribution, of these secondary-phase particles control the creep behavior and, therefore, determine the exploitation temperature of martensitic steels [1-8]. These particles exert Zener drag force, which prevents migration of both HABs and boundaries of martensite laths as well [9]. However, there exist limited data [8] dealt with examination of precipitation sequences under tempering and its effect on microstructure and mechanical properties of martensitic steels. For instance, a P92 steel is most widely used creep resistant material [1]. However, there is no information in literature on microstructure evolution in this material under tempering. It is wellknown [1-7], that after standard tempering at 750 or 770°C the TMLS is formed in this steel and characterized by the presence of M23C6-type carbides located on boundaries of martensite laths and HABs of prior austenite grains (PAG), packet and blocks. In addition, MX-type carbonitrides are uniformly distributed within TMLS. Dispersion of these carbonitrides consists of V-rich M(C,N) and Nb-rich M(C,N) particles. This thermodynamically equilibrium two-phase separation of the carbonitrides is extremely important for creep resistance of 9%Cr martensitic steels because it provides a strong resistance of M(C,N) particles against coagulation under creep conditions. The aim of the present study is to consider the effect of tempering conditions on mechanical properties and TMLS of a P92-type creep resistant steel at T≤720oC. Specific attention is paid to examination of precipitation sequences of carbides and carbonitrides under tempering.
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Experimental A P92-type steel (Fe-0.1С-0.17Si-0.54Mn-8.75Cr-0.21Ni-0.51Mo-1.60W-0.23V-0.07Nb, all in mass %) was examined. The steel was subjected to solution treatment at 1050oC followed by air cooling. Next, the specimens were tempered for 3 hours at temperatures of 400, 450, 500, 525, 550, 600, 650, 720°С. Isothermal tensile tests were carried out right after the tempering at the same temperatures by using flat specimens with cross section area of 3×1.5 mm2 and a 12 mm gauge length. The structural characterizations were carried out using a Quanta 600 scanning electron microscope (SEM) and JEM-2100 transmission electron microscope (TEM) equipped with an INCA energy dispersive X-ray spectroscope (EDS). The TEM foils were prepared by double jet electro-polishing using a solution of 10% perchloric acid in glacial acetic acid. Extraction carbon replicas on copper grids were used for identification of precipitates to avoid matrix effects, when obtaining the EDS spectra. The specimens were mechanically polished by emery paper and 3 µm silica suspension, and then etched using a solution of 2.5% nitric acid and 1.5 % hydrofluoric acid in water. The lath/subgrain sizes were measured on TEM micrographs by the linear intercept method, including all clearly visible (sub)boundaries. The dislocation densities were evaluated by counting the individual dislocations within the grain/subgrain interiors on at least six arbitrarily selected typical TEM images for each data point. Differential scanning calorimetry (DSC) was performed on a ∼45 mg specimens during heating to 1050°C at a rate of 20°C min-1 and cooling in nitrogen atmosphere with the rate of 15°C min-1 by using an SDT Q600 (TA Instruments) calorimeter. Equilibrium mass fractions of phases and their chemical compositions were calculated with a version 5 of the Thermo-Calc software using the TCFE6 database.
Figure 1. Mechanical properties of a P92-type steel at different temperatures (temperature of tensile test and tempering is the same). Results and Discussion Mechanical Properties. Figure 1 illustrates the aging behavior of the P92-type steel. The hardness value of 400 HB after solution treatment increased to 430 HB with increasing the tempering temperature to 525°С. Further increase in the tempering temperature to 720°C resulted in gradual decrease of the hardness, which approached a level of about 220 HB after the tempering at 720°С (Fig. 1a). In the temperature range 400-525°C, the yield stress (YS) of about 960 MPa is almost constant, while the ultimate tensile strengths (UTS) decreased slowly with increasing temperature. An increase in the tempering temperature to 720°С resulted in fast degradation of strength properties (Fig. 1b).
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Phase Equilibrium. Thermo-Calc calculations were carried out to determine the influence of main alloying elements on the phase composition at different temperatures. The fractions of the main secondary phases, i.e. M23C6 carbide, Laves phase and MX carbonitride, are depicted in Fig. 2 for austenization temperature of 1050°C and different tempering temperatures. The calculations suggested complete dissolution of Cr23C6 and Laves phases at 1050°C, although a small amount of un-dissolved MX carbonitrides was predicted. In the temperature interval 350-720°C, the equilibrium fractions of Cr23C6 and MX were calculated about 1.8 and 0.25 mass%, respectively. The appearance of Laves phase is expected at temperatures below 720°C. DSC Analysis. Figure 3 shows the results of DSC for the both heating (Fig. 3a) and cooling (Fig. 3b) experiments. Under heating, weak exothermic reactions at temperatures of 340 to 420°С and 440 to 640°С are indicated as A and B, respectively, in Fig. 3a. Reactions of A and B suggest the precipitation of M3C and M23C6 carbides, respectively [8]. Then, two endothermic reactions were recorded at around 700 to 790°С and 840 to 900°C. The first endothermic reaction results from a magnetic transition from a ferromagnetic state to a paramagnetic state; the second one is attributed to austenite transformation (Ac1 and Ac3 points in Fig. 3a). The DSC curve obtained during cooling shows the sharp exothermic peak resulting from the martensitic transformation in temperature range of 389-324°C (Fig. 3b).
Figure 2. Equilibrium mass fractions of Cr23C6, MX and Laves phases at different temperatures calculated by Thermo-Calc.
Figure 3. DSC curves of a P92-type steel (a) on heating and (b) on cooling.
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TMLS. Representative micrographs of air cooled microstructure after solution treatment are shown in Fig. 4. The size of PAGs revealed by SEM is about 20 µm (Fig. 4a). PAGs are subdivided on blocks and packets. The fine substructure of martensite packets/blocks is composed of dislocation laths (Fig. 4b). The transverse lath size is about 250 nm. A rather high dislocation density of 7 × 1014 m-2 is observed within interiors of martensitic laths.
Figure 4. Martensitic structure after normalization. (a) SEM photograph, (b) TEM photograph The tempering at 400°С resulted in the appearance of two types of dispersoids. First, needle-like Fe-rich M3C carbides were found within martensitic laths. These dispersoids were identified as M3C carbides with orthorhombic structure (Fig. 5a). Their length and thickness are about 110 and 20 nm, respectively. Second type of dispersoids having equiaxed shape with an average size of about 30 nm was found within the martensite laths. These particles were identified as NbC carbides (Fig. 5b). The EDX data revealed that most of these particles are characterized by a large amount of Nb, which comprises about 90% of metallic atoms (Fig. 5b). No formation of V-rich carbonitrides was found at this temperature.
Figure 5. TEM micrographs of normalized and tempered at 400°C P92 steel, showing (a) M3C particles and (b) chemical composition (at.%) of MX particles on carbon replica. At 525°С, the M3C and NbC carbides also were found as fine particles randomly distributed within interiors of martensitic laths. Their sizes and shapes were almost the same as those evolved at the lower temperature. The formation of M23C6-type carbides, which appear as extended thin films
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along some HABs of PAGs and packets (Fig. 6a), was found at 525°C. Therefore, tempering behaviors of the P92-type steel and P911 steel with 3%Co additions [8] are almost the same. Increasing temperature to 625°С leads to the formation of round-shaped M23C6 carbides with an average size of about 100 nm (Fig. 6b). These carbides are located on HABs and boundaries of martensitic laths. It seems that these equiaxed carbides result from concurrent occurrence of two processes: a coagulation of film-like M23C6 carbide located on HABs and independent precipitation of equiaxed M23C6 carbides on block and lath boundaries. The second process is in dominant. It should be noted that the tempering in temperature range from 400 to 625°C does not lead to remarkable change in the dislocation density and the transverse lath size.
Figure 6. TEM micrographs of normalized and (a) tempered at 525°C, (b) tempered at 625°C P92type steel, showing the temperature effect on the distribution of M23C6-type carbides. At 720°C, the width of martensite lath increased to about 0.3 µm (Fig. 7a), while the dislocation density slightly decreased to 6.2 × 1014 m-1. Volume fraction of M23C6 carbides exhibiting essentially equiaxed shape attained about 2% that nearly matches Thermo-Calc calculations. Therefore, the thermodynamically equilibrium dispersion of M23C6 carbides evolved at 720oC. The mean size of M23C6 carbides is about 100 nm (Fig. 7b).
Figure 7. TEM micrographs of normalized and tempered at 720°C P92-type steel, showing (a) M23C6 on the lath boundaries, and (b) chemical composition (at.%) of M23C6 particles on carbon replica.
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Two different fractions of M23C6 carbides could be distinguished by their location. Coarse M23C6 carbides locate on HABs; and nano-scale M23C6 carbides locate on low-angle lath boundaries. This bimodal distribution of M23C6 carbides is unstable under creep conditions [9], extensive coagulation of coarse M23C6 carbides located on HABs of PAGs and packets and progressive dissolution of nano-scale M23C6 carbides located on lath boundaries is the main process deteriorating creep resistance of high-chromium steels. It is worth noting that no formation of Laves phase was detected that is in consistent with other works [1]. It is known [1, 9] Laves phase precipitates under creep conditions; their volume fraction attains thermodynamicall equilibrium value after 103 hours or more. Low-duration tempering is not enough for precipitations of Laves phase. The MX particles with a size of about 30 nm exhibiting a well-defined equiaxed shape were revealed. Therefore, no formation of V-rich M(C,N) carbonitrides having a well-known “wing” shape [1, 8] occurs at T
