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Enhanced mechanical properties induced by refined heat treatment for 9Cr–0.5Mo–1.8W martensitic heat resistant steel
Materials and Design 50 (2013) 174–180
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Enhanced mechanical properties induced by refined heat treatment for 9Cr–0.5Mo–1.8W martensitic heat resistant steel S.S. Wang, D.L. Peng, L. Chang, X.D. Hui ⇑ State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing 100083, PR China
a r t i c l e
i n f o
Article history: Received 25 October 2012 Accepted 31 January 2013 Available online 27 February 2013 Keywords: Heat resistant steels Heat treatment Orthogonal experimental design Mechanical property Strengthening mechanism
a b s t r a c t In this work, the phase transformation in 9Cr–1.8W–0.5Mo steel was investigated to provide the theoretical basis for the design of the temperature of heat treatment process. Refined heat treatment technologies including normalizing and tempering processes were designed by using orthogonal experiment method. By evaluating the room temperature and high temperature mechanical properties, optimum heat treatment parameters were determined. It is shown that under the optimum heat treatment condition, the room temperature elongation can reach 25% or above, while the tensile strength and yield strength at 650 °C were increased by 48.3% and 50%, respectively, compared with those of supplied steel. The strengthening mechanism of this steel was also investigated by microstructure observation. This work is hopeful to benefit the service safety and the life time extension of power plants. Ó 2013 Elsevier Ltd. All rights reserved.
1. Introduction Nowadays, the fossil-fired power is still occupying about 75% of the total energy product in the world, and it will be the main resource of the energy supply to mankind for a long time. To make the current power plants run with high efficiency, it is imperious to design and produce new heat resistant steels with higher performance and more safety allowance under more severe condition [1]. Ferritic/martensitic steels exhibit high strengths, high heat conductivity and low expansion coefficient, which make them desirable heat resistant steels used for supercritical coal-fired power plants, and furthermore, they have balanced mechanical properties and processing abilities such as welding performance, suitable toughness and hardness [3,4]. T/P92 steel [2], a typical Cr–W–Mo–V–Nb heat resistant ferritic/ martensitic steel with 8.5–9.5%Cr, 1.5–2.0%W, 0.30–0.60%Mo, has been used in the components of fossil-fired power plant more and more, such as main steam pipes or high temperature boilers within the temperature range from 580 °C to 625 °C. This alloy is strengthened by the composite alloying of W–Mo and V–Nb combined with a small quantity of N and B. Within a certain service temperature range, the creep strength of T/P92 steel is about 20% higher than that of traditional austenite steels which can partly be replaced [3]. Moreover, the welding performance of T/P92 steel with austenitic steels is also excellent [5,6].
⇑ Corresponding author. Tel.: +86 10 62332169; fax: +86 10 62333447. E-mail address: [email protected] (X.D. Hui). 0261-3069/$ - see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.matdes.2013.01.072
In the last few decades, great attention to 9–12%Cr steels was focused on the composition design by alloying, microstructural characterization and the performance during the service life [4]. By appropriate composition designing [7–9] and processing [10–12], the density, morphology, distribution and size of M23C6 (M = Fe, Cr, W, Mo) and MX (M = V, Nb; X = C, N) type carbides, nitrides and carbonitrides, which are very important to the ferritic/martensitic steels, have been controlled [13–15], so the room and high temperature mechanical strengths could be adjusted. The coarsening behavior of martensite laths for 9–12%Cr heat-resistant steels has also been investigated during high temperature creep processes [16–20]. It is found that the coarsening rate of laths decreases with tungsten element content increasing, and is correlated with the Ostwald ripening rate of M23C6 type carbides [20,21]. In addition, Laves phase [22,23] and Z-phase [24–26] were observed in ferritic/martensitic steels after creep deformation for dozens of hours. Although a lot of work has been done on T/P92 steel, the relationship among the microstructure, properties and processing is still incompletely clarified. Considering the matrix microstructure and the amounts, morphology and distribution of various second phases can be adjusted through heat treatment technology, so it is possible to improve the steel properties if we are able to select the heat treatment parameters reasonably. In this work, the evolution of microstructure of T/P92 steel during heating process was revealed by using differential scanning calorimeter (DSC), thermal dilatometer and high temperature X-ray diffraction (XRD). Based on the information of phase transformation, the heat treatment temperature and time for normalizing and tempering were deter-
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mined, and refined heat treatment mechanical tests were designed and performed by using orthogonal experimental method [27,28]. Through the measurement of room and high temperature mechanical properties, the influence of various parameters on the properties of this steel was evaluated, and optimized heat treat technology was obtained. By using optical microscope (OM) and transmission electronic microscope (TEM), the structure and strengthening mechanism of T/P92 steel was investigated.
Table 2 Factors and levels of the heat treatment experiments for T92 steel. Level
1 2 3
Factor Normalizing temperature (°C)
Normalizing time (min)
Tempering temperature (°C)
Tempering time (min)
1030 1060 1090
10 20 30
750 770 790
60 90 120
2. Experimental procedure 3. Results and discussion The material investigated in this work is 9Cr–1.5W–1Mo–VNb steel (ASME SA-213 T92). The chemical compositions of this steel are listed in Table 1. The material is supplied in tube with the diameter of 36 mm and the thickness of 8 mm. The original microstructure of T92 steel is martensite matrix processed by normalizing and tempering. The thermal stability measurement of the steel was performed by using STA 449C Jupiter differential scanning calorimeter according to Chinese Standard Method GB/T 13464-2008 [29]. The sample with masses of about 50 mg was first put into an Al2O3 crucible, and then heated to 1200 °C at the rate of 10 °C/min under the atmosphere of argon. The thermal expansion curve was tested using DIL805A mode of thermal dilatometer produced by Baehr Co. Ltd according to the standard ASTM: E831-12. The sample size is U5 10 mm, heating rate and the highest temperature is 10 °C/ min and 1200 °C, respectively. In this work, the high temperature XRD analysis was conducted by using TTR III mode of X-ray diffractometer. The polished specimen with a size of 20 10 2 mm was firstly heated to predefined temperature, held for 2 min, and at last scanned within the range of 30–90°. The phases were identified by comparing the peaks with those specified in standard PDF card. The heat treatment was designed by using orthogonal experiment method. We chose L9(34) [27] type orthogonal form, as shown in Table 2. The samples with a diameter of 5 mm were designed according to Chinese Standard Method GB/T 228.1-2010 [30]. The tensile tests were conducted by using CMT4105 type test machine. The samples for high temperature tensile experiments were prepared according to Chinese Standard Method GB/T 4338-2006 [31], and performed by using AG-50KNE mode of tensile machine at 550 °C, 600 °C and 650 °C, respectively. The samples were tested three times at each temperature. The Vickers hardness (HV) which is one of the parameters for evaluating the mechanical properties was also tested by using Leica VMHT 30M model microhardness instrument according to Chinese Standard Method GB/T 4340-2009 [32]. The specimens were grinded to 2000 mesh sandpapers followed by machinery polishing. The HV values at ten different positions for each sample were measured by imposing a load of 100 g for 10 s. Finally, the HV of each sample was evaluated by taking the average value of these ten values. The microstructure was analyzed by using OM and TEM. The metallographic analysis was examined by using Leica OM. The samples were firstly polished and corroded with Fe3Cl ethanol solution. The samples for TEM observation were prepared by linear cutting and machinery polishing. These samples were then reduced by using ion etching device. The TEM observation was conducted by using JEOL JEM-2100 mode of TEM, which is also equipped with Oxford INCA type spectrometer and GATAN 832 CCD image recorder.
3.1. The evolution of the microstructure and properties during heating process The DSC and thermal expansion curves of supplied steel during the heating process are shown in Fig. 1. From the DSC curve, it is seen that endothermic peaks appear at 742.8 °C and 864.3 °C, respectively, indicating that obvious phase transformation takes place. As known from the previous work [33], a slope change was found on the dilation–temperature curve during heating process, enhanced by slow heating rate, before the transformation from martensite to austenite, which is caused by the dissolution of precipitates and diffusion of alloying elements within a specific temperature range. This phenomenon is also found in this work. As shown in Fig. 1a, the slope of the thermal expansion curve changes at 744.7 °C, when the sample was heated at the rate of 10 °C/min, which is corresponding to the first peak of DSC curve. For another two slope changes of the thermal expansion curve at 844.0 °C and 879.3 °C, the first one is the starting temperature of austenitic transformation (Ac1) and the second is exactly its end (Ac3) [34], which must be corresponding to the second peak of DSC curve. It has been known that, among austenite, ferrite and martensite, the thermal expansion coefficient of austenite is the highest and that of martensite is the lowest, so the volume will be shrunk as the matrix is transformed into austenite. The difference between the specific volumes of the two phases might be reflected in the thermal expansion curve. Therefore, the evolution of the microstructure can be briefly deduced by DSC and thermal expansion curves as follows: most of carbides, nitrides and carbonitrides dissolve at 744.7 °C, and the matrix is transformed into austenite at 844.0 °C. The schematic for the austenitic transformation process is shown in Fig. 1b. The austenitic transformation can be further verified by the in situ XRD experiments performed at 25 °C (RT), 800 °C, 950 °C, and 1100 °C, respectively. As shown in Fig. 2, the crystal structure of supplied steel at RT is body-centered cubic (BCC). Considering that the final process is high temperature tempering treatment, the matrix should be tempered martensite. At 800 °C, the XRD pattern indicates that the matrix is mainly composed of BCC type structure, and it is also noticed that the peaks are shifted to lower angle compared with those at RT, which is caused by the heating and diffusion of alloying elements into matrix. The XRD pattern at 950 °C shows that the phase with face-centered cubic (FCC) type structure appears, indicating the matrix has been transformed into austenite and alloying elements dissolved into FCC matrix. The XRD pattern at 1100 °C shows that there is no BCC type peaks anymore, indicating that the matrix has been transformed into austenite completely.
Table 1 Chemical composition of T92 heat-resistant steel (mass fraction/%). Element Content
C 0.10
Cr 8.81
W 1.59
Mo 0.35
V 0.20
Nb 0.05
Mn 0.43
Si 0.38
B 0.0030
N 0.052
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and 1200 °C [9,35], respectively. Moreover, for 9–12%Cr steels, the upper limit of normalizing temperature should be around 1100 °C to avoid the coarsening of microstructure [11,36]. Next, let’s discuss the determination of tempering temperature. As the steel cooled after normalized, the microstructure is transformed from austenite to martensite with homologous grain sizes. To relieve the normalizing stress, improve the plasticity and the microstructure stability, high temperature tempering with the lower temperature than Ac1 should be employed. In addition, the tempering temperature must be no more than 800 °C, otherwise Zphase (CrNbN) [25] which is highly injurious to creep performance will come out inevitably. It is well known that low temperature tempering is able to improve room temperature strength and hardness, whereas high temperature tempering can obtain suitable plasticity and stable dislocation networks, which is helpful to the microstructural stability during long-time aging process. In addition, the higher the tempering temperature, the more the precipitated phases, which have excellent pinning effect on the martensite lath boundaries and dislocation networks and then benefit to high temperature behavior [37]. Therefore, after comprehensively considering the strength, plasticity and high temperature performance, we choose the tempering temperature from 744.7 °C, the maximum dissolution temperature, to 800 °C. It should be noticed that the quality of heat treatment is also dependent on the time duration of various process because the phase transformations are kinetic processes. The time durations for various processes were determined according to the temperature and feature of the process, which are shown in Table 2. 3.2. The heat treatment technology and mechanical properties Fig. 1. (a) DSC and thermal expansion curve and (b) magnified drawing of austenitic transformation of T92 steel.
Fig. 2. In-situ XRD patterns of T92 steel heated at different temperatures.
The above evolution of microstructure may provide the theoretical basis for the design of the heat treatment temperature. It is known that the ideal microstructure is lath-shaped martensite matrix strengthened by precipitated phases. To attain this aim, two temperature ranges should be determined for normalizing and tempering processes. As the starting point of heat treatment, homogeneous microstructure composed of single austenite phase with appropriate grain size should be obtained firstly and the most of carbides, nitrides and carbonitrides should also be dissolved into the matrix. Thus the normalizing temperature should be higher than its Ac3. On the other hand, from the XRD pattern at 1100 °C, it is seen that the single austenite phase has completely formed, and it is reported that dissolution of M23C6 and MX type precipitates begin at about 780 °C and 900 °C and finish at about 940 °C
The above analysis on T92 steel microstructural evolution has basically provided the temperature ranges for the heat treatment technology. To obtain stable microstructure and excellent mechanical properties, optimal heat treatment parameters are in great request. In this work, we designed the heat treatment experiments by using the orthogonal experimental method. L9(34) type orthogonal form was chosen. In the orthogonal form, the four factors are normalizing temperature and the soaking time; tempering temperature and the soaking time, respectively. For each factor, three levels were designed, such as the tempering temperatures are 750 °C, 770 °C, and 790 °C. Eighty-one experiments would be needed if the tests were designed just by round permutation and combination. But by using orthogonal design method, nine experiments are enough in the first inning. In the next inning, we designed two supplemental heat treatment experiments, which used the parameters optimized from the first inning, to inspect the validity of the orthogonal design method. To compare the mechanical properties of the samples processed by the present heat treatment parameters with those of supplied material, the tests of supplied steel were also performed. The room temperature tensile properties and HV values of T92 steel in present work and previously reported under different heat treatment conditions are summarized in Table 3. The corresponding heat treatment parameters are also listed in the table. Among the mechanical indexes (the tensile strength (Rm), yield strength (Rp0.2) and elongation (A) after fractured and microhardness) for T92 steel, all the Rm and Rp0.2 are much higher than those specified in standard ASME SA-213M. In this case, the elongation is the key index that limits the comprehensive mechanical properties. It is seen that the variation of the HV values with the condition of heat treatment is consistent with those of Rm and Rp0.2. Considering that the plasticity is critical to the safety, processability and long service of the whole setup, the elongation was chosen as the primary index for the evaluation of heat treatment quality in this work. And Rp0.2 is usually used as the permissible stress in the
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S.S. Wang et al. / Materials and Design 50 (2013) 174–180 Table 3 Parameters of heat treatment orthogonal experiments and mechanical properties of T92steel in present work and previously reported at room temperature. Sample number
Normalizing temperature (°C)
Normalizing time (min)
Tempering temperature (°C)
Tempering time (min)
Tensile strength (Rm) (MPa)
Yield strength (Rp0.2) (MPa)
Elongation (A) (%)
HV values
92-0 92-1 92-2 92-3 92-4 92-5 92-6 92-7 92-8 92-9 92-10 92-11 T/P92[38] T/P92[39]
– 1030 1030 1030 1060 1060 1060 1090 1090 1090 1060 1060 1070 –
– 10 20 30 10 20 30 10 20 30 10 10 – –
– 750 770 790 770 790 750 790 750 770 790 790 760 –
– 60 90 120 120 60 90 90 120 60 120 90 – –
685.00 816.67 733.33 708.33 728.33 731.67 803.33 716.67 793.33 771.67 742.52 749.56 692.3 726.0
521.67 681.67 563.33 525.00 550.00 550.00 656.67 531.67 645.00 618.33 571.86 590.07 512.9 566.7
22.67 19.33 23.83 25.50 24.83 25.00 20.67 25.50 19.83 22.00 23.07 21.89 20.1 26.0
243.0 286.3 263.1 243.3 246.2 247.5 272.1 244.5 269.0 254.3 273.1 269.5 – –
P620
P440
P20
–
ASME SA-213/SA-213 M (at room temperature)
Table 4 Range analysis of the specific elongation of T92 steel after heat treatment. Analysis
Normalizing temperature (°C)
Normalizing time (min)
Tempering temperature (°C)
Tempering time (min)
A-K1 A-K2 A-K3 Range
22.887 23.500 22.443 1.057
23.220 22.887 22.723 0.497
19.943 23.553 25.333 5.390
22.110 23.333 23.387 1.277
Fig. 3. Temperature dependence of the tensile strength Rm and yield strength Rp0.2 of T92 steel.
design of machinery parts, so we choose Rp0.2 as the secondary evaluation index. At last, Rm and HV values will be used to evaluate the performance of T92 steel. As shown in Table 3, all the tensile strengths processed in this work are higher but some elongations are lower than those of No. 92-0 condition. According to the sequence of the importance for the evaluation of heat treatment quality specified in the above section, we performed the direct-vision and range analysis on the elongation from No. 92-1 to 92-9, which are shown in Table 4. Here the A-K1, A-K2 and A-K3 represent the average elongations corresponding to each factor from level 1 to 3, respectively. Range is the difference between the maximum and minimum values of the elongations for a specified factor. It is easily seen that the sequence of the factor affecting the elongation is tempering temperature, tempering time, normalizing temperature and normalizing time. On the other hand, for each factor of normalizing temperature, normalizing time and tempering temperature, the average
Fig. 4. OM images of the microstructure of supplied steel (a) and (b) under heat treatment No. 92-5.
elongation is the best at level 2, 1 and 3 correspondingly. Therefore, we can acquire a group of heat treatment parameters as shown followed: normalizing temperature 1060 °C, normalizing time 10 min and tempering temperature 790 °C. As for the factor of tempering time, the elongations at level 2 and 3 are very close, so we designed two supplemental experiments Nos. 92-10 and 9211 in Table 3. However, the direct-vision of the data indicates that the elongations of the two supplemental experiments are not the best among all performed experiments. From Table 3, it is seen that the tensile properties of T92 steel in this work are higher than those of T/P 92 steel treated under the condition of normalizing at 1070 °C and tempering at 760 °C [38]. It is especially noted that the elongation of No. 92-6 is almost the same as value of steel trea-
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Fig. 5. Microstructure of T92 steel after heat treatment analyzed by TEM: (a) microstructure of martensite matrix and the precipitates; (b) TEM image of M23C6 displayed along the prior austenite grain boundaries; (c and d) TEM images of M23C6 and MX, respectively, distributed in the martensite laths; (e and f) SADPs of M23C6 and MX type precipitates; (g and h) the SADP analysis of M23C6, MX and the matrix with zone axis of [1 1 1], [0 1 2] and [1 2 5], respectively.
ted in [38], while the tensile strengths of present steel are about 100 MPa higher than those. It is also seen that under Nos. 92-3, 92-5 and 92-7 conditions, the tensile properties of the present steel are near close to those given by Baosteel Ltd. [39]. The elongation of our steel can reach 25% or above while Rm and Rp0.2 is far higher than those specified in standard ASME SA-213/SA-213M. There-
fore, we choose the heat treatment parameters of Nos. 92-3, 92-5 and 92-7 for high temperature tensile experiments. For high temperature application, the tensile and yield strengths represent the high temperature performance to a certain extent. Fig. 3 shows the temperature dependence of Rm and Rp0.2 for the three heat treatment conditions. Within the temperature
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range from 550 °C to 650 °C, both Rm and Rp0.2 treated under the condition of No. 92-5 are the best among the three experiments. It has been reported that Rm and Rp0.2 of the steel fabricated by Baosteel Ltd. are 280 MPa and 253 MPa, respectively, at 650 °C [39]. Compared with the values of Baosteel Ltd. at same temperature, the Rm and Rp0.2 of the present steel under the condition of No. 92-5 are increased by 48.3% and 50%, respectively. The room and high temperature mechanical properties of T92 steel obtained in this work shows that refined heat treatment process is highly necessary for the production and application. Based on the above experimental results, we can determine the optimized heat treatment parameters as follows: normalizing temperature 1060 °C and the soaking time 20 min; tempering temperature 790 °C and the soaking time 60 min. 3.3. The microstructure and strengthening mechanism Fig 4 shows the OM images of supplied steel and treated under the condition of No. 92-5 experiment. It is found that the sample treated under the condition of No. 92-5 has the grain size of 20– 50 lm (Fig. 4b), which is finer than that of supplied steel (Fig. 4a). The martensite laths were fragmented and the subgrains formed. Therefore, it is seen that the strengthening effect after heat treatment partly results from the refinement of martensite matrix. The refined matrix was obtained by appropriate control of austenization and the followed tempering process. The high content of Cr in T92 steel results in high hardenability. Therefore, martensite can form easily when the sample is cooled in air from normalizing temperature. From this point of view, the appropriate control of austenization is important to the strengthening by grain refinement. Besides controlling the grain size by austenization, the tempering process also affect the grain size. Of course, the lower the tempering temperature and the shorter the time, the smaller the grain size of martensite matrix. However, the selection of the temperature and time of normalizing and tempering is also affected by the comprehensive consideration of strength and plasticity. That means the grain size cannot be very fine. Moreover, in order to serve stably at high temperature, the appropriate grain size must be obtained. The TEM images and selected area electron diffraction pattern (SADP) of T92 steel processed under the condition of No. 92-5 are shown in Fig. 5. We can see that the structure is composed of martensite matrix, M23C6 and MX type precipitates. Fig. 5a shows the prior austenite grain boundaries and the M23C6 type carbides between them. In Fig. 5b and c, the M23C6 type carbides, which have particulate or short rod-like shape with a size of 100– 200 nm, are distributed along the prior austenite grain boundaries and in the martensite laths. A high density of dislocations with a straight or curved shape is accumulated near the precipitates. It is also found that the lath-shaped martensite was fragmented after tempered. From Fig. 5c and d, it is seen that the width of martensite laths is within a range of 200–400 nm. In these laths, there exist secondary phases and dislocation networks. MX type precipitates with very fine elliptical shape and a size of about 50 nm are distributed in the martensite laths. According to the energy dispersive Xray (EDX) spectrum and SADPs (Fig. 5e–h), both the two kinds of second phases are FCC type structure. M23C6 is composed of Fe, Cr, W, Mo and C, and MX is composed of V, Nb, C and N, which is coincident with K. Yamada’s results [40]. As previously reported, the secondary phases are beneficial to the mechanical properties when their size is in a specific range. As aging process goes on, the M23C6 type carbides are easy to grow up to worsen the performance of the steel [41]. However, during the heating process, M23C6 begins to dissolve at about 780 °C, while MX begins at about 900 °C. On the other hand, a certain ratio of W can hinder the MX grow up [21], so this type of precipitate is more stable. Thus the fine MX type precipitates are the chief strengthen secondary
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phases in this kind of steel, which is consistent with Kimmura Kazuhiro’s research during creep process [42]. Based on the above property analysis and structural observation, we can see that the strengthening effect is attributed to the following aspects. Firstly, the refined heat treatment ensures the accomplishment of phase transformation so as to obtain the martensite matrix with appropriate shape and size. The next is the strengthening effect by the subgrains and dislocations as shown in Fig. 5. It is easily understood that the subgrains are formed during the tempering process as internal stress releasing and martensite matrix being fragmented. Near the dense interfaces of subgrains, high density of dislocations, which can accumulate and block the deformation of material, are observed. Thirdly, the precipitation of the secondary phases plays a critical role in the strengthening effect. These phases are formed during the heat treatment process, especially the tempering process. It is supersaturation of C and N in the matrix that result in the formation of martensite after cooling from normalizing temperature. Since M23C6 and MX type carbides, nitrides and carbonitrides are fine enough and distributed randomly but not aggregately. These precipitates can strengthen the martensite matrix, for one thing, by blocking the movement of dislocations; for another, the precipitation of secondary phases at the interfaces of subgrains may also have strengthening effect. Carbides, nitrides and carbonitrides can be easily produced at the interfaces due to the feasibility of diffusivity of elements. During tempering or deformation process, the dislocations easily move toward the interfaces, and the precipitates at the interfaces can effectively block the dislocations as shown in Fig. 5b. It should be especially noticed that the fine MX type secondary phases are also precipitated in the martensite laths, which not only block the movement of dislocations but also result in the distortion of the matrix lattice. In the final analysis, the strengthening effect is resulted from the optimized heat treatment process.
4. Conclusions In this work, we investigated the evolution of microstructure of T/P92 steel during heating process. The orthogonal experimental method was applied to optimize the heat treatment process including the heat treatment temperature and time for normalizing and tempering. The influence of heat treatment parameters on the properties of this steel was evaluated by measuring the room and high temperature mechanical properties. The microstructure and strengthening mechanism of T/P92 steel was also discussed. Conclusive remarks can be deduced as follows: (1) The DSC, thermal expansion, in situ XRD experiments reveal the evolution of matrix transformation: most of precipitates dissolve at about 744.7 °C, then the matrix begins to transform to austenite at 844.0 °C, the austenization and the dissolution of carbides, nitride and carbonitrides are finished completely before 1100 °C. These results may provide the theoretical basis for the design of the temperature of heat treatment process. (2) The heat treatment experiments were designed by using L9(34) type orthogonal form. Based on the mechanical properties at room and high temperature, Rm, Rp0.2, elongation and HV values, the optimized heat treatment technology was determined: normalizing temperature 1060 °C and soaking time 20 min; tempering temperature 790 °C and soaking time 60 min. (3) Under the optimized heat treatment technology, the elongation at room temperature can reach 25% or above, and Rm and Rp0.2 are both far higher than those specified in standard ASME SA-213. The Rm and Rp0.2 at 650 °C are
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increased by 48.3% and 50%, respectively, compared with those of supplied steel. (4) The examination of OM and TEM shows that the strengthening effect is attributed to the formation of martensite matrix with appropriate shape and size, the subgrains and dislocations, and the precipitation of the secondary phases including M23C6 and MX type carbides, nitrides and carbonitrides. All the strengthening effects actually result from the optimized heat treatment technology.
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[20] Abe F. Coarsening behavior of lath and its effect on creep rates in tempered martensitic 9Cr–W steels. Mater Sci Eng, A 2004;387–389:565–9. [21] Hong SG, Lee WB, Park CG. The effects of tungsten addition on the microstructural stability of 9Cr–Mo steels. J Nucl Mater 2001;288:202–7. [22] Dimmler G, Weinert P, Kozeschnik E, Cerjak H. Quantification of the Laves phase in advanced 9–12%Cr steels using a standard SEM. Mater Charact 2003;51:341–52. [23] Korcakova Leona, Hald John, Somers Marcel AJ. Quantification of Laves phase particle size in 9Cr–W steel. Mater Charact 2001;47:111–7. [24] Ardeshir Golpayegani, Hans-Olof Andren, Hilmar Danielsen, John Hald. A study on Z-phase nucleation in martensitic chromium steels. Mater Sci Eng, A 2008;489:310–8. [25] Sawada K, Suzuki K, Kushima H, Tabuchi M, Kimeta K. Effect of tempering temperature on Z-phase formation and creep strength in 9Cr–1Mo–V–Nb–N steel. Mater Sci Eng, A 2008;480:558–63. [26] Hilmar Kjartansson Danielsen, John Hald. A thermodynamic model of the Zphase Cr (V, Nb) N. Comput Coupling Phase Diagrams Thermochem 2007;31:505–14. [27] Montgomery Douglas C. Design and analysis of experiments. New York of America: John Wiley & Sons Inc.; 2011. [28] Ruijiang Liu, Yewang Zhang, Chongwei Wen, Jian Tang. Study on the design and analysis methods of orthogonal experiment. Exp Technol Manage 2012;27:52–5. [29] Yingchun Chen, Zhenyu Deng, Ping Zhuo, Chengxian Gong. GB/T 13464-2008: thermal analysis test methods for thermal stability of materials. General Administration of Quality Supervision, Inspection and Quarantine of the People’s Republic of China & Standardization Administration of The People’s Republic of China; June 2008. [30] Yiwen Gao, Xinbang Liang, Li Dong, et al. GB/T 228.1-2010: metallic materials – tensile testing – Part1: method of test at room temperature. General Administration of Quality Supervision, Inspection and Quarantine of the People’s Republic of China & Standardization Administration of The People’s Republic of China; December 2010. [31] Junping Zhao, Jiulin Li, Tiezhu Zhu, Li Dong, Fusheng Wang. GB/T 4338-2006: metallic materials – tensile testing at elevated temperature. General Administration of Quality Supervision, Inspection and Quarantine of the People’s Republic of China & Standardization Administration of The People’s Republic of China; August 2006. [32] Ying Li, Jingang Shi, Yiwen Gao, et al. GB/T 4340-2009: metallic materials – Vickers hardness test. General Administration of Quality Supervision, Inspection and Quarantine of the People’s Republic of China & Standardization Administration of The People’s Republic of China; June 2009. [33] Kapoor R, Kumar Lalit, Batra IS. A dilatometric study of the continuous heating transformations in 18 wt.% Ni maraging steel of grade 350. Mater Sci Eng, A 2003;352:318–24. [34] Kumar Anish, Laha K, Jayakumar T, Bhanu Sankaro Rao K, Raj Baldev. Comprehensive microstructural characterization in modified 9Cr–1Mo ferritic steel by ultrasonic measurements. Metall Mater Trans A 2002;33:1617–26. [35] Jinrong Cao, Zhengdong Liu, Shichang Cheng, et al. Thermodynamic calculation and analysis on equilibrium phases transition in T122 heat-resistant steel. Spec Steel 2005;26:16–9. [36] Lei Mao, Rudong Li, Daogang Zhang, Feng LiJin. Effects of quenching temperature on microstructures and properties of 2Cr12NiMoWV. Hot Work Technol 1997;5:32–4. [37] Kimura K, Sawada K, Kushima H, Toda Y. Degradation behavior and long-term creep strength of 12Cr ferritic creep resistant steels. In: Proceedings of the 8th liege conference on materials for advanced power engineering. Belgium: Liege; 2006. p. 120–8. [38] Guo Xiaofeng, Gong Jianming, Jiang Yong, Rong Dongsong. The influence of long-term aging on microstructures and static mechanical properties of P92 steel at room temperature. Mater Sci Eng: A 2013;564:199–205. [39] Qijiang Wang, Jie Hong, Jiaqiang Gao, et al. R&D of T92 high-pressure boiler tube for ultra supercritical pressure boilers. Power Equip 2009;6:451–5. [40] Yamada Katsumi, Igarashi Masaak, Muneki Seiichi, Abe Fujio. Creep properties affected by morphology of MX in high-Cr ferritic steels. ISIJ Int 2001;41:116–20. [41] Masuyama Fujimitsu, Komai Nobuyoshi, et al. Three-Year Experience with 2.25Cr–1.6W (HCM2S) and 12Cr–0.4Mo–2W (HCM12A) steel tubes in a power boiler. JSME Int J Ser B, Fluids Thermal Eng 1998;41:1098–104. [42] Kazuhiro Kimmura, Fujio Abe. Degradation of Mod.9Cr–1Mo steel during longterm creep deformation. Tetsu to Hagane 1999;85:841–7.
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Materials and Design journal homepage: www.elsevier.com/locate/matdes
Enhanced mechanical properties induced by refined heat treatment for 9Cr–0.5Mo–1.8W martensitic heat resistant steel S.S. Wang, D.L. Peng, L. Chang, X.D. Hui ⇑ State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing 100083, PR China
a r t i c l e
i n f o
Article history: Received 25 October 2012 Accepted 31 January 2013 Available online 27 February 2013 Keywords: Heat resistant steels Heat treatment Orthogonal experimental design Mechanical property Strengthening mechanism
a b s t r a c t In this work, the phase transformation in 9Cr–1.8W–0.5Mo steel was investigated to provide the theoretical basis for the design of the temperature of heat treatment process. Refined heat treatment technologies including normalizing and tempering processes were designed by using orthogonal experiment method. By evaluating the room temperature and high temperature mechanical properties, optimum heat treatment parameters were determined. It is shown that under the optimum heat treatment condition, the room temperature elongation can reach 25% or above, while the tensile strength and yield strength at 650 °C were increased by 48.3% and 50%, respectively, compared with those of supplied steel. The strengthening mechanism of this steel was also investigated by microstructure observation. This work is hopeful to benefit the service safety and the life time extension of power plants. Ó 2013 Elsevier Ltd. All rights reserved.
1. Introduction Nowadays, the fossil-fired power is still occupying about 75% of the total energy product in the world, and it will be the main resource of the energy supply to mankind for a long time. To make the current power plants run with high efficiency, it is imperious to design and produce new heat resistant steels with higher performance and more safety allowance under more severe condition [1]. Ferritic/martensitic steels exhibit high strengths, high heat conductivity and low expansion coefficient, which make them desirable heat resistant steels used for supercritical coal-fired power plants, and furthermore, they have balanced mechanical properties and processing abilities such as welding performance, suitable toughness and hardness [3,4]. T/P92 steel [2], a typical Cr–W–Mo–V–Nb heat resistant ferritic/ martensitic steel with 8.5–9.5%Cr, 1.5–2.0%W, 0.30–0.60%Mo, has been used in the components of fossil-fired power plant more and more, such as main steam pipes or high temperature boilers within the temperature range from 580 °C to 625 °C. This alloy is strengthened by the composite alloying of W–Mo and V–Nb combined with a small quantity of N and B. Within a certain service temperature range, the creep strength of T/P92 steel is about 20% higher than that of traditional austenite steels which can partly be replaced [3]. Moreover, the welding performance of T/P92 steel with austenitic steels is also excellent [5,6].
⇑ Corresponding author. Tel.: +86 10 62332169; fax: +86 10 62333447. E-mail address: [email protected] (X.D. Hui). 0261-3069/$ - see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.matdes.2013.01.072
In the last few decades, great attention to 9–12%Cr steels was focused on the composition design by alloying, microstructural characterization and the performance during the service life [4]. By appropriate composition designing [7–9] and processing [10–12], the density, morphology, distribution and size of M23C6 (M = Fe, Cr, W, Mo) and MX (M = V, Nb; X = C, N) type carbides, nitrides and carbonitrides, which are very important to the ferritic/martensitic steels, have been controlled [13–15], so the room and high temperature mechanical strengths could be adjusted. The coarsening behavior of martensite laths for 9–12%Cr heat-resistant steels has also been investigated during high temperature creep processes [16–20]. It is found that the coarsening rate of laths decreases with tungsten element content increasing, and is correlated with the Ostwald ripening rate of M23C6 type carbides [20,21]. In addition, Laves phase [22,23] and Z-phase [24–26] were observed in ferritic/martensitic steels after creep deformation for dozens of hours. Although a lot of work has been done on T/P92 steel, the relationship among the microstructure, properties and processing is still incompletely clarified. Considering the matrix microstructure and the amounts, morphology and distribution of various second phases can be adjusted through heat treatment technology, so it is possible to improve the steel properties if we are able to select the heat treatment parameters reasonably. In this work, the evolution of microstructure of T/P92 steel during heating process was revealed by using differential scanning calorimeter (DSC), thermal dilatometer and high temperature X-ray diffraction (XRD). Based on the information of phase transformation, the heat treatment temperature and time for normalizing and tempering were deter-
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mined, and refined heat treatment mechanical tests were designed and performed by using orthogonal experimental method [27,28]. Through the measurement of room and high temperature mechanical properties, the influence of various parameters on the properties of this steel was evaluated, and optimized heat treat technology was obtained. By using optical microscope (OM) and transmission electronic microscope (TEM), the structure and strengthening mechanism of T/P92 steel was investigated.
Table 2 Factors and levels of the heat treatment experiments for T92 steel. Level
1 2 3
Factor Normalizing temperature (°C)
Normalizing time (min)
Tempering temperature (°C)
Tempering time (min)
1030 1060 1090
10 20 30
750 770 790
60 90 120
2. Experimental procedure 3. Results and discussion The material investigated in this work is 9Cr–1.5W–1Mo–VNb steel (ASME SA-213 T92). The chemical compositions of this steel are listed in Table 1. The material is supplied in tube with the diameter of 36 mm and the thickness of 8 mm. The original microstructure of T92 steel is martensite matrix processed by normalizing and tempering. The thermal stability measurement of the steel was performed by using STA 449C Jupiter differential scanning calorimeter according to Chinese Standard Method GB/T 13464-2008 [29]. The sample with masses of about 50 mg was first put into an Al2O3 crucible, and then heated to 1200 °C at the rate of 10 °C/min under the atmosphere of argon. The thermal expansion curve was tested using DIL805A mode of thermal dilatometer produced by Baehr Co. Ltd according to the standard ASTM: E831-12. The sample size is U5 10 mm, heating rate and the highest temperature is 10 °C/ min and 1200 °C, respectively. In this work, the high temperature XRD analysis was conducted by using TTR III mode of X-ray diffractometer. The polished specimen with a size of 20 10 2 mm was firstly heated to predefined temperature, held for 2 min, and at last scanned within the range of 30–90°. The phases were identified by comparing the peaks with those specified in standard PDF card. The heat treatment was designed by using orthogonal experiment method. We chose L9(34) [27] type orthogonal form, as shown in Table 2. The samples with a diameter of 5 mm were designed according to Chinese Standard Method GB/T 228.1-2010 [30]. The tensile tests were conducted by using CMT4105 type test machine. The samples for high temperature tensile experiments were prepared according to Chinese Standard Method GB/T 4338-2006 [31], and performed by using AG-50KNE mode of tensile machine at 550 °C, 600 °C and 650 °C, respectively. The samples were tested three times at each temperature. The Vickers hardness (HV) which is one of the parameters for evaluating the mechanical properties was also tested by using Leica VMHT 30M model microhardness instrument according to Chinese Standard Method GB/T 4340-2009 [32]. The specimens were grinded to 2000 mesh sandpapers followed by machinery polishing. The HV values at ten different positions for each sample were measured by imposing a load of 100 g for 10 s. Finally, the HV of each sample was evaluated by taking the average value of these ten values. The microstructure was analyzed by using OM and TEM. The metallographic analysis was examined by using Leica OM. The samples were firstly polished and corroded with Fe3Cl ethanol solution. The samples for TEM observation were prepared by linear cutting and machinery polishing. These samples were then reduced by using ion etching device. The TEM observation was conducted by using JEOL JEM-2100 mode of TEM, which is also equipped with Oxford INCA type spectrometer and GATAN 832 CCD image recorder.
3.1. The evolution of the microstructure and properties during heating process The DSC and thermal expansion curves of supplied steel during the heating process are shown in Fig. 1. From the DSC curve, it is seen that endothermic peaks appear at 742.8 °C and 864.3 °C, respectively, indicating that obvious phase transformation takes place. As known from the previous work [33], a slope change was found on the dilation–temperature curve during heating process, enhanced by slow heating rate, before the transformation from martensite to austenite, which is caused by the dissolution of precipitates and diffusion of alloying elements within a specific temperature range. This phenomenon is also found in this work. As shown in Fig. 1a, the slope of the thermal expansion curve changes at 744.7 °C, when the sample was heated at the rate of 10 °C/min, which is corresponding to the first peak of DSC curve. For another two slope changes of the thermal expansion curve at 844.0 °C and 879.3 °C, the first one is the starting temperature of austenitic transformation (Ac1) and the second is exactly its end (Ac3) [34], which must be corresponding to the second peak of DSC curve. It has been known that, among austenite, ferrite and martensite, the thermal expansion coefficient of austenite is the highest and that of martensite is the lowest, so the volume will be shrunk as the matrix is transformed into austenite. The difference between the specific volumes of the two phases might be reflected in the thermal expansion curve. Therefore, the evolution of the microstructure can be briefly deduced by DSC and thermal expansion curves as follows: most of carbides, nitrides and carbonitrides dissolve at 744.7 °C, and the matrix is transformed into austenite at 844.0 °C. The schematic for the austenitic transformation process is shown in Fig. 1b. The austenitic transformation can be further verified by the in situ XRD experiments performed at 25 °C (RT), 800 °C, 950 °C, and 1100 °C, respectively. As shown in Fig. 2, the crystal structure of supplied steel at RT is body-centered cubic (BCC). Considering that the final process is high temperature tempering treatment, the matrix should be tempered martensite. At 800 °C, the XRD pattern indicates that the matrix is mainly composed of BCC type structure, and it is also noticed that the peaks are shifted to lower angle compared with those at RT, which is caused by the heating and diffusion of alloying elements into matrix. The XRD pattern at 950 °C shows that the phase with face-centered cubic (FCC) type structure appears, indicating the matrix has been transformed into austenite and alloying elements dissolved into FCC matrix. The XRD pattern at 1100 °C shows that there is no BCC type peaks anymore, indicating that the matrix has been transformed into austenite completely.
Table 1 Chemical composition of T92 heat-resistant steel (mass fraction/%). Element Content
C 0.10
Cr 8.81
W 1.59
Mo 0.35
V 0.20
Nb 0.05
Mn 0.43
Si 0.38
B 0.0030
N 0.052
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and 1200 °C [9,35], respectively. Moreover, for 9–12%Cr steels, the upper limit of normalizing temperature should be around 1100 °C to avoid the coarsening of microstructure [11,36]. Next, let’s discuss the determination of tempering temperature. As the steel cooled after normalized, the microstructure is transformed from austenite to martensite with homologous grain sizes. To relieve the normalizing stress, improve the plasticity and the microstructure stability, high temperature tempering with the lower temperature than Ac1 should be employed. In addition, the tempering temperature must be no more than 800 °C, otherwise Zphase (CrNbN) [25] which is highly injurious to creep performance will come out inevitably. It is well known that low temperature tempering is able to improve room temperature strength and hardness, whereas high temperature tempering can obtain suitable plasticity and stable dislocation networks, which is helpful to the microstructural stability during long-time aging process. In addition, the higher the tempering temperature, the more the precipitated phases, which have excellent pinning effect on the martensite lath boundaries and dislocation networks and then benefit to high temperature behavior [37]. Therefore, after comprehensively considering the strength, plasticity and high temperature performance, we choose the tempering temperature from 744.7 °C, the maximum dissolution temperature, to 800 °C. It should be noticed that the quality of heat treatment is also dependent on the time duration of various process because the phase transformations are kinetic processes. The time durations for various processes were determined according to the temperature and feature of the process, which are shown in Table 2. 3.2. The heat treatment technology and mechanical properties Fig. 1. (a) DSC and thermal expansion curve and (b) magnified drawing of austenitic transformation of T92 steel.
Fig. 2. In-situ XRD patterns of T92 steel heated at different temperatures.
The above evolution of microstructure may provide the theoretical basis for the design of the heat treatment temperature. It is known that the ideal microstructure is lath-shaped martensite matrix strengthened by precipitated phases. To attain this aim, two temperature ranges should be determined for normalizing and tempering processes. As the starting point of heat treatment, homogeneous microstructure composed of single austenite phase with appropriate grain size should be obtained firstly and the most of carbides, nitrides and carbonitrides should also be dissolved into the matrix. Thus the normalizing temperature should be higher than its Ac3. On the other hand, from the XRD pattern at 1100 °C, it is seen that the single austenite phase has completely formed, and it is reported that dissolution of M23C6 and MX type precipitates begin at about 780 °C and 900 °C and finish at about 940 °C
The above analysis on T92 steel microstructural evolution has basically provided the temperature ranges for the heat treatment technology. To obtain stable microstructure and excellent mechanical properties, optimal heat treatment parameters are in great request. In this work, we designed the heat treatment experiments by using the orthogonal experimental method. L9(34) type orthogonal form was chosen. In the orthogonal form, the four factors are normalizing temperature and the soaking time; tempering temperature and the soaking time, respectively. For each factor, three levels were designed, such as the tempering temperatures are 750 °C, 770 °C, and 790 °C. Eighty-one experiments would be needed if the tests were designed just by round permutation and combination. But by using orthogonal design method, nine experiments are enough in the first inning. In the next inning, we designed two supplemental heat treatment experiments, which used the parameters optimized from the first inning, to inspect the validity of the orthogonal design method. To compare the mechanical properties of the samples processed by the present heat treatment parameters with those of supplied material, the tests of supplied steel were also performed. The room temperature tensile properties and HV values of T92 steel in present work and previously reported under different heat treatment conditions are summarized in Table 3. The corresponding heat treatment parameters are also listed in the table. Among the mechanical indexes (the tensile strength (Rm), yield strength (Rp0.2) and elongation (A) after fractured and microhardness) for T92 steel, all the Rm and Rp0.2 are much higher than those specified in standard ASME SA-213M. In this case, the elongation is the key index that limits the comprehensive mechanical properties. It is seen that the variation of the HV values with the condition of heat treatment is consistent with those of Rm and Rp0.2. Considering that the plasticity is critical to the safety, processability and long service of the whole setup, the elongation was chosen as the primary index for the evaluation of heat treatment quality in this work. And Rp0.2 is usually used as the permissible stress in the
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S.S. Wang et al. / Materials and Design 50 (2013) 174–180 Table 3 Parameters of heat treatment orthogonal experiments and mechanical properties of T92steel in present work and previously reported at room temperature. Sample number
Normalizing temperature (°C)
Normalizing time (min)
Tempering temperature (°C)
Tempering time (min)
Tensile strength (Rm) (MPa)
Yield strength (Rp0.2) (MPa)
Elongation (A) (%)
HV values
92-0 92-1 92-2 92-3 92-4 92-5 92-6 92-7 92-8 92-9 92-10 92-11 T/P92[38] T/P92[39]
– 1030 1030 1030 1060 1060 1060 1090 1090 1090 1060 1060 1070 –
– 10 20 30 10 20 30 10 20 30 10 10 – –
– 750 770 790 770 790 750 790 750 770 790 790 760 –
– 60 90 120 120 60 90 90 120 60 120 90 – –
685.00 816.67 733.33 708.33 728.33 731.67 803.33 716.67 793.33 771.67 742.52 749.56 692.3 726.0
521.67 681.67 563.33 525.00 550.00 550.00 656.67 531.67 645.00 618.33 571.86 590.07 512.9 566.7
22.67 19.33 23.83 25.50 24.83 25.00 20.67 25.50 19.83 22.00 23.07 21.89 20.1 26.0
243.0 286.3 263.1 243.3 246.2 247.5 272.1 244.5 269.0 254.3 273.1 269.5 – –
P620
P440
P20
–
ASME SA-213/SA-213 M (at room temperature)
Table 4 Range analysis of the specific elongation of T92 steel after heat treatment. Analysis
Normalizing temperature (°C)
Normalizing time (min)
Tempering temperature (°C)
Tempering time (min)
A-K1 A-K2 A-K3 Range
22.887 23.500 22.443 1.057
23.220 22.887 22.723 0.497
19.943 23.553 25.333 5.390
22.110 23.333 23.387 1.277
Fig. 3. Temperature dependence of the tensile strength Rm and yield strength Rp0.2 of T92 steel.
design of machinery parts, so we choose Rp0.2 as the secondary evaluation index. At last, Rm and HV values will be used to evaluate the performance of T92 steel. As shown in Table 3, all the tensile strengths processed in this work are higher but some elongations are lower than those of No. 92-0 condition. According to the sequence of the importance for the evaluation of heat treatment quality specified in the above section, we performed the direct-vision and range analysis on the elongation from No. 92-1 to 92-9, which are shown in Table 4. Here the A-K1, A-K2 and A-K3 represent the average elongations corresponding to each factor from level 1 to 3, respectively. Range is the difference between the maximum and minimum values of the elongations for a specified factor. It is easily seen that the sequence of the factor affecting the elongation is tempering temperature, tempering time, normalizing temperature and normalizing time. On the other hand, for each factor of normalizing temperature, normalizing time and tempering temperature, the average
Fig. 4. OM images of the microstructure of supplied steel (a) and (b) under heat treatment No. 92-5.
elongation is the best at level 2, 1 and 3 correspondingly. Therefore, we can acquire a group of heat treatment parameters as shown followed: normalizing temperature 1060 °C, normalizing time 10 min and tempering temperature 790 °C. As for the factor of tempering time, the elongations at level 2 and 3 are very close, so we designed two supplemental experiments Nos. 92-10 and 9211 in Table 3. However, the direct-vision of the data indicates that the elongations of the two supplemental experiments are not the best among all performed experiments. From Table 3, it is seen that the tensile properties of T92 steel in this work are higher than those of T/P 92 steel treated under the condition of normalizing at 1070 °C and tempering at 760 °C [38]. It is especially noted that the elongation of No. 92-6 is almost the same as value of steel trea-
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Fig. 5. Microstructure of T92 steel after heat treatment analyzed by TEM: (a) microstructure of martensite matrix and the precipitates; (b) TEM image of M23C6 displayed along the prior austenite grain boundaries; (c and d) TEM images of M23C6 and MX, respectively, distributed in the martensite laths; (e and f) SADPs of M23C6 and MX type precipitates; (g and h) the SADP analysis of M23C6, MX and the matrix with zone axis of [1 1 1], [0 1 2] and [1 2 5], respectively.
ted in [38], while the tensile strengths of present steel are about 100 MPa higher than those. It is also seen that under Nos. 92-3, 92-5 and 92-7 conditions, the tensile properties of the present steel are near close to those given by Baosteel Ltd. [39]. The elongation of our steel can reach 25% or above while Rm and Rp0.2 is far higher than those specified in standard ASME SA-213/SA-213M. There-
fore, we choose the heat treatment parameters of Nos. 92-3, 92-5 and 92-7 for high temperature tensile experiments. For high temperature application, the tensile and yield strengths represent the high temperature performance to a certain extent. Fig. 3 shows the temperature dependence of Rm and Rp0.2 for the three heat treatment conditions. Within the temperature
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range from 550 °C to 650 °C, both Rm and Rp0.2 treated under the condition of No. 92-5 are the best among the three experiments. It has been reported that Rm and Rp0.2 of the steel fabricated by Baosteel Ltd. are 280 MPa and 253 MPa, respectively, at 650 °C [39]. Compared with the values of Baosteel Ltd. at same temperature, the Rm and Rp0.2 of the present steel under the condition of No. 92-5 are increased by 48.3% and 50%, respectively. The room and high temperature mechanical properties of T92 steel obtained in this work shows that refined heat treatment process is highly necessary for the production and application. Based on the above experimental results, we can determine the optimized heat treatment parameters as follows: normalizing temperature 1060 °C and the soaking time 20 min; tempering temperature 790 °C and the soaking time 60 min. 3.3. The microstructure and strengthening mechanism Fig 4 shows the OM images of supplied steel and treated under the condition of No. 92-5 experiment. It is found that the sample treated under the condition of No. 92-5 has the grain size of 20– 50 lm (Fig. 4b), which is finer than that of supplied steel (Fig. 4a). The martensite laths were fragmented and the subgrains formed. Therefore, it is seen that the strengthening effect after heat treatment partly results from the refinement of martensite matrix. The refined matrix was obtained by appropriate control of austenization and the followed tempering process. The high content of Cr in T92 steel results in high hardenability. Therefore, martensite can form easily when the sample is cooled in air from normalizing temperature. From this point of view, the appropriate control of austenization is important to the strengthening by grain refinement. Besides controlling the grain size by austenization, the tempering process also affect the grain size. Of course, the lower the tempering temperature and the shorter the time, the smaller the grain size of martensite matrix. However, the selection of the temperature and time of normalizing and tempering is also affected by the comprehensive consideration of strength and plasticity. That means the grain size cannot be very fine. Moreover, in order to serve stably at high temperature, the appropriate grain size must be obtained. The TEM images and selected area electron diffraction pattern (SADP) of T92 steel processed under the condition of No. 92-5 are shown in Fig. 5. We can see that the structure is composed of martensite matrix, M23C6 and MX type precipitates. Fig. 5a shows the prior austenite grain boundaries and the M23C6 type carbides between them. In Fig. 5b and c, the M23C6 type carbides, which have particulate or short rod-like shape with a size of 100– 200 nm, are distributed along the prior austenite grain boundaries and in the martensite laths. A high density of dislocations with a straight or curved shape is accumulated near the precipitates. It is also found that the lath-shaped martensite was fragmented after tempered. From Fig. 5c and d, it is seen that the width of martensite laths is within a range of 200–400 nm. In these laths, there exist secondary phases and dislocation networks. MX type precipitates with very fine elliptical shape and a size of about 50 nm are distributed in the martensite laths. According to the energy dispersive Xray (EDX) spectrum and SADPs (Fig. 5e–h), both the two kinds of second phases are FCC type structure. M23C6 is composed of Fe, Cr, W, Mo and C, and MX is composed of V, Nb, C and N, which is coincident with K. Yamada’s results [40]. As previously reported, the secondary phases are beneficial to the mechanical properties when their size is in a specific range. As aging process goes on, the M23C6 type carbides are easy to grow up to worsen the performance of the steel [41]. However, during the heating process, M23C6 begins to dissolve at about 780 °C, while MX begins at about 900 °C. On the other hand, a certain ratio of W can hinder the MX grow up [21], so this type of precipitate is more stable. Thus the fine MX type precipitates are the chief strengthen secondary
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phases in this kind of steel, which is consistent with Kimmura Kazuhiro’s research during creep process [42]. Based on the above property analysis and structural observation, we can see that the strengthening effect is attributed to the following aspects. Firstly, the refined heat treatment ensures the accomplishment of phase transformation so as to obtain the martensite matrix with appropriate shape and size. The next is the strengthening effect by the subgrains and dislocations as shown in Fig. 5. It is easily understood that the subgrains are formed during the tempering process as internal stress releasing and martensite matrix being fragmented. Near the dense interfaces of subgrains, high density of dislocations, which can accumulate and block the deformation of material, are observed. Thirdly, the precipitation of the secondary phases plays a critical role in the strengthening effect. These phases are formed during the heat treatment process, especially the tempering process. It is supersaturation of C and N in the matrix that result in the formation of martensite after cooling from normalizing temperature. Since M23C6 and MX type carbides, nitrides and carbonitrides are fine enough and distributed randomly but not aggregately. These precipitates can strengthen the martensite matrix, for one thing, by blocking the movement of dislocations; for another, the precipitation of secondary phases at the interfaces of subgrains may also have strengthening effect. Carbides, nitrides and carbonitrides can be easily produced at the interfaces due to the feasibility of diffusivity of elements. During tempering or deformation process, the dislocations easily move toward the interfaces, and the precipitates at the interfaces can effectively block the dislocations as shown in Fig. 5b. It should be especially noticed that the fine MX type secondary phases are also precipitated in the martensite laths, which not only block the movement of dislocations but also result in the distortion of the matrix lattice. In the final analysis, the strengthening effect is resulted from the optimized heat treatment process.
4. Conclusions In this work, we investigated the evolution of microstructure of T/P92 steel during heating process. The orthogonal experimental method was applied to optimize the heat treatment process including the heat treatment temperature and time for normalizing and tempering. The influence of heat treatment parameters on the properties of this steel was evaluated by measuring the room and high temperature mechanical properties. The microstructure and strengthening mechanism of T/P92 steel was also discussed. Conclusive remarks can be deduced as follows: (1) The DSC, thermal expansion, in situ XRD experiments reveal the evolution of matrix transformation: most of precipitates dissolve at about 744.7 °C, then the matrix begins to transform to austenite at 844.0 °C, the austenization and the dissolution of carbides, nitride and carbonitrides are finished completely before 1100 °C. These results may provide the theoretical basis for the design of the temperature of heat treatment process. (2) The heat treatment experiments were designed by using L9(34) type orthogonal form. Based on the mechanical properties at room and high temperature, Rm, Rp0.2, elongation and HV values, the optimized heat treatment technology was determined: normalizing temperature 1060 °C and soaking time 20 min; tempering temperature 790 °C and soaking time 60 min. (3) Under the optimized heat treatment technology, the elongation at room temperature can reach 25% or above, and Rm and Rp0.2 are both far higher than those specified in standard ASME SA-213. The Rm and Rp0.2 at 650 °C are
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increased by 48.3% and 50%, respectively, compared with those of supplied steel. (4) The examination of OM and TEM shows that the strengthening effect is attributed to the formation of martensite matrix with appropriate shape and size, the subgrains and dislocations, and the precipitation of the secondary phases including M23C6 and MX type carbides, nitrides and carbonitrides. All the strengthening effects actually result from the optimized heat treatment technology.
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