Precipitation Characterization and Creep Strength at 600°C for Creep Resistant Cr–Mo Steel

Abstract

The creep strength of 5Cr-0.5Mo steel was determined at 600°C and 78–170 MPa, as well as its relation to the microstructural changes during the creep tests. The microstructural characterization showed that the creep tests were conducted under the presence of a mixture of both intergranular and intragranular M₇C₃ and M₂₃C₆ carbides dispersed in the ferrite matrix. The n exponent of Norton-Bailey law suggested that the creep deformation process occurred through the ferrite grains, which conducted to a transgranular ductile- fracture mode after creep testing. The creep strength of this steel is directly related to the average radius and number density of carbides present during the test. The ferrite grain size of 5 µm seemed to cause an enhancement of the creep strength for this steel in comparison to that of other similar steels reported in the literature.

1. Introduction

Heat and creep resistant Cr–Mo steels have been used since a long time ago with great success for applications in the power generation and petrochemical industries. Typical components of these industries are boilers, heaters, heat exchangers and hydrocrackers,^1,2,3) which are usually fabricated as heavy wall pressure vessels. The basic Cr–Mo steels are 0.5Cr-0.5Mo, 1Cr-0.5Mo, 2.25Cr-1Mo, 5Cr-0.5Mo, 9Cr-1Mo and 12Cr-1Mo alloyed steels.⁴⁾ The use of alloying elements such as, V, W, Ni, Ti, Nb, B and/or N allows to reach the current grades such as the T/P22, T/P23, T/P24, T/P91, T/P92 and VM 12-SHC steels.^1,2) Many of these new grade steels have been applied successfully in industry, but they are also still under development. This type of steels has been designed^1,2,3,4) to work at temperatures between 350 and 600°C and pressures of about 15–30 MPa for life time of about 250000 h. Creep resistant steels can support certain stress at specific service temperatures without passing a specified amount of elongation. The maximum stress to rupture at a specific temperature after a specific time, e.g. 600°C and 100000 h, is referred as creep rupture stress. For instance, an engineering design criterion for a power plant could require a minimal stress of 100 MPa for 100000 h at service temperature.^1,2,3,4) The basic idea is that the vessel keeps original its size and shape during service for up to 20 to 30 years.⁴⁾ The creep resistance of a Cr–Mo steel is based on the formation of stable precipitates such as, alloyed carbides in a ferritic, bainitic and/or martensitic matrix phase in the normalized steel condition. Because of the subsequent tempering treatment, a stable microstructure with precipitates is formed remaining stable at the service temperature for which the steel has been developed.^4,5,6,7) The precipitates formed will lock the grain-boundaries and avoid sliding of the slip-planes to give the desired creep resistance properties.^1,2) These precipitates should present the correct morphology, volume fraction and dispersion in order to obtain a homogeneous structure with homogeneous properties. These characteristics depend on the alloying element level and the heat treatments, which permit to form different volume fractions and types of precipitates. Fe₃C, M₂X, M₅C₂ M₇C₃, M₂₃C₆, M₂C, M₆C, Chi and Laves are some of the precipitates present in the Cr–Mo steels. As previously mentioned, their mechanical properties such as, excellent mechanical strength, toughness and creep strength, are mainly based on^2,3,4,5,6,7) the dispersion of different precipitates in the ferritic matrix. Nevertheless, the operation conditions of industrial components usually involve the exposure at high temperature for prolonged times under stress, which promote different microstructure changes such as, grain boundary segregation (temper embrittlement), brittle-phases formation, as well as the coarsening of precipitates.^2,3,4,5,6,7) These changes may alter the mechanical properties of the steel.⁶⁾ Therefore, the evaluation of deterioration in mechanical properties such as, toughness and creep strength is an important issue to be considered for the safety operation of the equipment. Thus, the precipitate of this type of steel plays an important role to obtain good creep strength. For instance, the ferritic 5Cr-0.5Mo steel^{8,9,10,11,12,13,14,15)} is considered as a member of the Cr–Mo steels and it has the presence of an abundant precipitation from the as-received condition since this is usually normalized and tempered, which promotes the formation of different carbides, Fe₃C, M₇C₃ and M₂₃C₆.^1,9,11) Moreover, this steel is usually operated at temperatures higher than 400°C for prolonged times, which also contributes to the precipitate formation during the early stages of isothermal heating. Furthermore, the carbide precipitation has been reported^11,12,13) to be almost composed of M₂₃C₆ carbides after aging of the steel at 600°C for 400 h. Additionally, its prolonged exposures at high temperatures also originate the coarsening of carbides, which might have detrimental effects on the creep strength because of the lower fraction and increase in size of precipitates facilitating the dislocation movement and grain sliding.^1,2) Besides, there are few works related to the precipitation and creep strength of this steel in comparison to the other Cr–Mo steels. Most of these works were pursued at test stresses lower than 80 MPa during the creep test.

Thus, the purpose of the present work is to evaluate the creep strength at 600°C and test stresses of 78–170 MPa for a 5Cr-0.5Mo steel, as well as the role of the precipitation evolution on the creep strength.

2. Experimental Procedure

Table 1 shows the chemical composition of present work steel. This composition corresponds to a 5Cr-0.5Mo steel ASTM A387 Grade 5 Class 2.⁸⁾ The steel plates have a size of 200×200×12 mm. These were normalized from 950°C and subsequently tempered at 715°C for 15 min. Tension test specimens of 86 mm length, 30 mm gauge-length and 6 mm diameter were prepared from samples of 12×12×100 mm, cut paralell to the rolling direction. Tension tests of steel were conducted at 20°C and 600°C according to JIS Z 2241 and G 0567 (ISO 6892-2) standards,^16,17) respectively.

Table 1. Chemical composition (mass%) of 5Cr-0.5Mo steel.

Fe	C	Cr	Mo	Mn	P	S	Si	Cu	Ni	Ti
Bal.	0.096	4.531	0.454	0.356	0.009	0.0005	0.344	0.047	0.117	0.002

Uniaxial creep test specimens of 90 mm length, diameter gauge-length of 30 mm and 6 mm diameter were also prepared by machining. The conventional creep tests were carried out at 600±3°C under different stresses in the range of 78–170 MPa according to JIS Z 2271(ISO 204) standard.¹⁸⁾ The testing temperature was stabilized in the furnace for 4 h and an atmosphere of high-purity Ar gas was used during testing. The strain was recorded using an extensometer located in the gauge area. Transversal and longitudinal sections of tested specimens were metallographically prepared, etched with Nital reagent and subsequently observed with High- Resolution Scaning Electron Microscope (HR-SEM) and Conventional (SEM) at 20 kV. In order to compare the precipitation evolution of the tested creep specimens, as-received steel specimens of 12×12×12 mm were aged at 600°C for times between 0 and 6000 h. These were also metallographically prepared, etched and observed with light microscope, Conventional-SEM and HR-SEM. Besides, the precipitates were extracted by electrochemical dissolution of the ferrite matrix from both the as-received and aged specimens. The electrolyte was composed of 2–4 vol.% nitric acid in ethanol with a graphite cathode at room temperature and 2–5 V (d.c.). The extracted residues were analyzed with an X-ray diffractometer with monochromated Cu Kα radiation. The present phases and their amount were determined with commercial software. Quantitative metallography of heat treated specimens was pursued to determine number density and size of precipitates using HR-SEM images and commercial software. The Vickers hardness, HV 0.1/12 s, was determined in all the aged specimens according to ASTM E-92 standard.¹⁹⁾

3. Results and Discussion

3.1. Tensile Properties

Table 2 summarizes the tensile properties of steel at 20 and 600°C. The tensile properties at 20°C are in very good agreement with those expected from a 5Cr-0.5Mo steel ASTM A387 Grade 5 Class 2.⁸⁾ The tensile properties at 600°C are also consistent with those at 600°C.²⁰⁾

Table 2. Tensile properties at 20 and 600°C of 5Cr-0.5Mo steel.

Temperature (°C)	Yield strength (MPa)	Tensile strength (MPa)	Elongation (%)	Young modulus (GPa)
20	441.0	572.0	34.0	188.0
600	268.0	304.0	31.1	142.6

3.2. Creep Properties

The creep curves, strain (%) versus time (h), at 600°C are shown in Figs. 1(a) and 1(b) for the test stresses 120–170 and 78–95 MPa, respectively. The second and third stages of creep are clearly observed in the creep curves. As expected, the time-to-rupture, duration of creep-secondary stage, and total strain increase with the decrease in test stress.

Fig. 1.

Creep curves at 600°C for stresses between (a) 120–170 and (b) 78–95 MPa.

Table 3 summarizes the total elongation and area reduction for each test. The elongation increases with the decrease in test stress and the reduction in area shows an increase with the test stress.

Table 3. Creep parameters of tested specimens.

Test stress (MPa)	Area reduction (%)	Total elongation (%)
78	83.7	48.7
88	85.8	42.1
95	87.6	36.6
120	89.1	37.4
130	89.9	37.1
170	92.7	34.0

Figures 2(a) and 2(b) show the plot of strain rate, $\dot{\epsilon }$, versus time for the test stresses 120–170 and 78–95 MPa, respectively, where the three creep stages are clearly identified. That is, the strain rate, $\dot{\epsilon }$, decreases with time, as shown for short times in Fig. 2(a). The decreasing strain rate in the creep-primary stage has been related to the strain hardening or to a decrease in free or mobile dislocations.^1,2) In the creep-secondary stage, the strain rate remains almost constant. This strain rate is designated as a steady-state creep rate, ${\dot{\epsilon }}_s$, which is commonly associated with a balance between the rate of generation of dislocations contributing to the hardening and the rate of recovery contributing to the softening.^3,4) In the creep-tertiary stage, the strain rate increases with time until rupture at rupture time, t_r, and rupture strain, ε_r.¹⁾ This figure also shows that the duration of creep-secondary and creep-tertiary stages is shorter as the test stress increases. Likewise, the strain rate at the creep-tertiary stage increases with the test stress, Fig. 2(a).

Fig. 2.

Plot of strain rate, $\dot{\epsilon }$, versus time for stresses between (a) 120–170 and (b) 78–95 MPa.

Since the minimum creep rate, ${\dot{\epsilon }}_{min}$, is preferred by engineers and researches to characterize the secondary-creep stages, this was determined from Fig. 2. The stress dependence of minimum creep rate, ${\dot{\epsilon }}_{min}$, is usually expressed by a power law, known as Norton-Bailey Law:^1,2,3,4)   

$${\dot{\epsilon }}_{min} =A{\sigma }^n$$(1)

where A and n are exponents stress-independent. Figure 3 shows the stress dependence of minimum creep rate, ${\dot{\epsilon }}_{min}$. This log-log plot shows a clear linear relationship between the minimum creep rate, ${\dot{\epsilon }}_{min}$, and the test stress σ with an R² value of approximately 0.9961. This figure also indicates that the stress exponent n is equal to 6.29. The n values have been observed to be from 2.5 to 4 in the case of low stresses, 60–100 MPa for low-alloy ferritic1.25Cr-0.5Mo and 2.25Cr-1Mo steels. In contrast, n values of 6–12 have been reported^1,2,3) for high stresses at temperatures of 510–620°C. Besides, the fracture mode changed from intergranular to transgranular at high stresses. Furthermore, a range of n values from 1 at low stresses to 14 at high stresses has been observed to occur in both ferritic and austenitic steels creep-tested at temperatures between 510 and 710°C. Most of these works^1,2,3,4) reported that the deformation process is controlled by the grain boundary sliding at low stresses and by the matrix deformation at high stresses. Thus, an n value of approximately 6.29 suggests that the creep deformation mechanism took place by the deformation of ferrite matrix grains, promoting transgranular ductile-fracture mode for this steel. In the case of low-stress values, 15–50 MPa, the expected creep mechanism for the present work steel may cause the grain boundary sliding, which could originate the presence of intergranular fracture.

Fig. 3.

Plot of minimum creep rate, ${\dot{\epsilon }}_{min}$, versus test stress σ.

Additionally, the steady-state minimum creep rate ${\dot{\epsilon }}_{min}$ is also inversely to the time-to-rupture t_r:^1,2)   

$${\dot{\epsilon }}_{min}=\frac{C}{t_r^m}$$(2)

where C is a constant depending on the total elongation during creep, m is a constant often nearly close to one and t_r the time-to-rupture. This equation is known as Monkman-Grant relationship. Figure 4 shows the log-log plot of the minimum creep rate, ${\dot{\epsilon }}_{min}$, against time-to-rupture, t_r. The m constant was determined to be about 1.0165, which agrees with the expected value for this equation.²⁾

Fig. 4.

Monkman-Grant plot of log ${\dot{\epsilon }}_{min}$ versus log t_r.

Figure 5 presents the plot of the test stress versus the Larson-Miller Parameter (LPM). LMP is defined as T(C+log t_r) where T is the temperature expressed in K and t_r the time-to-rupture in h, and C a material constant equal to 20 for this type of steels.²¹⁾ A clear linear regression can be noted for the present work results with an R² value of approximately 0.9878. Thus, this relation can be used to estimate the time-to-rupture t_r for a different test temperature.

Fig. 5.

Plot of test stress σ versus Larson-Miller parameter.

3.3. SEM Micrographs of Tested Creep Specimens

Figures 6(a)–6(d) show the SEM fractograph of tested specimens for 78, 95, 120 and 170 MPa, respectively. A transgranular ductile-fracture mode can be observed for all the tested specimens. The fracture surface shows typical dimpled characteristics of the ductile fracture by microvoid coalescence.^1,2) The ductile-mode fracture shows an increase with the test stress. Besides, the presence of larger isolated microvoids can also be noted. These were formed after the microvoid coalescence process of ductile fracture. The size and number of these isolated microvoids decreased with the test stress. This behavior seems to be related with the precipitation size, as will be shown later. These fracture characteristics are consistent with the deformation mechanism related to the exponent n, as described in the previous section.

Fig. 6.

SEM fractographs of the tested specimens with (a) 78, (b) 95, (c) 120 and (d) 170 MPa.

SEM micrographs for a longitudinal region of the necked specimen are shown in Figs. 7(a)–7(d), for instance, after testing at 78, 95, 120 and 170 MPa, respectively. These micrographs show both intragranular and intergranular carbides dispersed in the ferrite matrix. Furthermore, the lowest number density and the largest size of precipitates are observed to occur in the necked region of the 78 MPa-tested specimen, Fig. 7(a). This testing condition originated the longest time-to-rupture t_r, which promoted the formation of coarsened precipitates. In the case of 95 MPa-tested specimen, the number density and size of precipitates are slightly higher and smaller, respectively, than the previously described specimen. The 120 MPa-tested specimen shows smaller size and higher density number of precipitates. Likewise, the 170 MPa-tested specimen has similar size distribution with the previous one, but it presents the smallest size because the time-to-rupture t_r is too short to permit either the growth or coarsening of precipitates.

Fig. 7.

SEM micrographs for the longitudinal section of tested specimens with (a) 78, (b) 95, (c) 120 and (d) 170 MPa.

The creep deformation occurred through the ferrite grains, which is also in agreement with the mechanism corresponding to n exponent value. The increase in deformation of ferrite grains is more notorious as the test stress increases, which seems to be related to a higher area reduction noted as the test stress increases, see Table 3.

3.4. Microstructure Evolution during Aging at 600°C

SEM micrographs of the steel aged at 600°C for 0 (as-received), 300, 1000 and 2000 h are shown in Figs. 8(a)–8(d), respectively. The microstructure of the as-received steel shows the presence of both intergranular and intragranular precipitates dispersed in the ferrite matrix, Fig. 8(a). Bainite constituent is expected to be present according to the Time-Temperature-Transformation (TTT) diagram for this type of steel;²²⁾ however, it is not observable in this SEM micrograph because of the high magnification. Therefore, Fig. 9 shows the optical microscope OM micrograph for the as-received steel and it can be noted the presence of ferrite grains and bainite, which is circled in red. In general, the size for intergranular precipitates is larger than that corresponding to the intragranular ones. It is important to notice that the precipitate size for the as-received steel is larger than that corresponding to that of the 300 h-aged steel because the former precipitates were formed during the heat treatment of normalizing and tempering at temperatures higher than 600°C. However, most of these carbides correspond to the metastable M₇C₃, as will be shown later, and thus, they are dissolved during the aging process to form the stable M₂₃C₆ carbide. The effect of aging process at 600°C can be noted in the microstructure evolution, as shown in Figs. 8(b)–8(d). That is, there is an increase in the size, from 55 to 65 nm and a decrease in the number density of precipitates, from 1.9 × 10¹² to 1.8 × 10¹² particles m⁻², in the steel after aging for 1000 h in comparison to 300 h aging, Fig. 8(b). This fact suggests that some precipitates were dissolved, while others were being formed during the aging process. As the aging process progresses, the precipitate number density decreases, while the precipitate size continues increasing, as shown in Fig. 10. This characteristic clearly indicated the presence of precipitate coarsening.²³⁾

Fig. 8.

SEM micrographs of the steel aged at 600°C for (a) 0 h, (b) 300 h, (c) 1000 h and (d) 2000 h.

Fig. 9.

Optical microscope micrograph of the as-received steel. (Online version in color.)

Fig. 10.

Plot of average radius and number density against aging time.

To identify the precipitate type, the XRD patterns of extracted residues from the specimens corresponding to Fig. 8 are shown in Fig. 11. The identified precipitates were Cr₇C₃, PDF 00-35-0783 ICDD, and Cr₂₃C₆, PDF 00-36-1482 ICDD, carbides²⁴⁾ for the as-received steel, Fig. 11. These carbides were formed after the normalizing and tempering treatments. Several works^4,5,6,7) have reported that the precipitation sequence that occurs in the ferrite matrix is as follows: M₃C → M₇C₃ → M₂₃C₆. The M₃C carbide is similar to the Fe₃C cementite phase with an orthorhombic structure (space group Pnma),²⁵⁾ M₇C₃ is a metastable Cr-rich carbide with an orthorhombic structure (space group Pnma)²⁵⁾ and M₂₃C₆ is the equilibrium stable Cr-rich carbide, with cubic structure (space group Fm3m),²⁶⁾ at 600°C for the chemical composition of this steel.²⁷⁾ It is important to note that M in both carbides means that the main element is Cr, but also it contains Fe and Mo.²⁷⁾ This precipitation reaction is expected to occur not only by isothermal heating, but also by continuous cooling from normalizing temperature.⁷⁾ That is, the presence of M₇C₃ and M₂₃C₆ carbides for the as-received steel can be attributable to the cooling process of normalizing and the isothermal heating of tempering at 715°C for 15 min. The absence of M₃C cementite for the as-received steel seems to be related to the fact that this phase can be formed at the early stages of precipitation in the air-cooling of the normalizing treatment. This cementite M₃C is subsequently dissolved to permit the nucleation of M₇C₃ during the tempering treatment. M₇C₃ and M₂₃C₆ carbides can be present simultaneously since the transformation of M₇C₃ carbide into M₂₃C₆ carbide is a concomitant process.²³⁾ As the aging process progresses, see Fig. 11, the X-ray diffraction intensity from peaks corresponding to the M₇C₃ carbide decreases, while those from M₂₃C₆ show an increase in intensity. Furthermore, the XRD of steel aged for 2000 h mainly shows the presence of peaks corresponding to M₂₃C₆ carbide. Table 4 indicates the weight percent of M₇C₃ and M₂₃C₆ carbides determined from the XRD pattern of Fig. 11. This table makes more evident that the transformation of M₇C₃ carbide into M₂₃C₆ carbide takes place during aging at 600°C.

Fig. 11.

XRD patterns of the extracted residues from the steel aged at 600°C for 0, 300, 1000 and 2000 h.

Table 4. Mass percent of carbides from XRD.

Aging time (h)	% Cr7C3	% Cr23C6
0	52	48
300	56	44
1000	58	42
2000	20	80
6000	0	100

On the other hand, the SEM micrographs of the extracted carbides, corresponding to Figs. 8(a)–8(d) are shown in Figs. 12(a)–12(d), respectively. Figure 12(a) indicates that there is a mixture of particles with polygonal and cuboid shape for the as-received steel. The EDS spectrum of these particles is also shown, for example, in Fig. 12(a). It is confirmed that they are Cr-rich carbides, as previously described. The Cr-rich carbide also shows the presence of Fe and Mo, which in agreement with the chemical analysis reported in the literature.^5,15) The volume fraction of the polygonal particles decreases with aging time, in comparison to the other ones, Figs. 12(b) and 12(c). Besides, an increase in particle size of both particles can be noticed with the increase in aging time. The morphology of particles becomes cuboids and square rods for the specimen aged for 2000 h. The polygonal shape corresponds to the M₇C₃ carbides, which is in agreement with the crystal shape expected for the point group $D_{2h}^{16}$ corresponding to the space group Pnma.²⁸⁾ In contrast, the shape of cuboids or like square rods corresponds to the M₂₃C₆ carbides, which is also consistent with the cubic shape of crystal expected for the point group $O_h^5$ corresponding to the space group Fm3m.²⁸⁾ It is important to mention that the elastic-strain energy between precipitate and matrix also plays an important role on the precipitate morphology.²³⁾ Both polygonal and cuboid shapes present planar faces, which may form coherent interfaces with the ferrite matrix. This fact may favor the creep strength for this type of steel.

Fig. 12.

The extracted residues from the steel aged at 600°C for (a) 0 and its corresponding EDS analysis, (b) 300, (c) 1000 and (d) 2000 h.

The plot of Vickers hardness as a function of the aging time is shown in Fig. 13. The decrease in hardness can be noted as the time increases. The initial decrease in hardness can be related to the disappearance of bainite for the as-received steel. The subsequent decrease in hardness with time is a result of the transformation of M₇C₃ carbide into M₂₃C₆ carbide, since the hardness of Cr₇C₃ carbide is higher than that of Cr₂₃C₆ carbide,²⁵⁾ which suggests that the dislocations move easier under the presence of the latter carbide, promoting more ductility. The coarsening of M₂₃C₆ carbide can be present for aging times longer than 2000 h, which may also cause a decrease in hardness.

Fig. 13.

Plot of Vickers hardness vs. time for the steel aged at 600°C.

3.5. Effect of Microstructure Evolution on Creep Properties

The previous results show the decrease in time-to-rupture t_r and the increase in minimum creep rate ${\dot{\epsilon }}_{min}$ with the test stress and Table 5 summarizes the relation of theses creep parameters to the microstructure evolution during the creep test. According to the XRD results, the steel microconstituents during creep tests were mainly composed of the ferrite matrix and a dispersion of both intergranular and intragranular M₇C₃ and M₂₃C₆ carbides, as shown in Table 5. This fact suggests that the different conditions of creep tests at 600°C are comparable each other. Nevertheless, the 170 MPa-test could have had the presence of certain amount of bainite during the creep test in comparison to the other test stresses, where the test time could have been enough to promote its transformation into allotromorphous or ideomorphous ferrite and cementite. The presence of bainite could have also contributed to the short duration of secondary-stage creep at high stresses, since the strengthening mechanisms were superior to those of the softening.^1,2,3) In general, a clear relationship can be noted between the time-to-rupture t_r and minimum creep rate ${\dot{\epsilon }}_{min}$ with the number density and average radius of precipitates. That is, the creep strength decreases with the increase in average radius and the decrease in number density of precipitates. This behavior is important because it provokes that the dislocation movement is favored by the presence of a lower number of precipitates, which facilitates the deformation process during the creep-secondary stage.

Table 5. Creep parameters of tested specimens related to microstructure.

Test stress (MPa)	Time-to-rupture (h)	${\dot{\epsilon }}_{min}$ (h−1)	Precipitates	Average radius (nm)	Number Density (m−2)	Time-to-rupture (h) [10]
78	1301	1.78 × 10−4	M7C3 + M23C6	67	1.8 × 1012	450–1080
88	640	4.22 × 10−4	M7C3 + M23C6	55	1.8 × 1012	220–550
95	510	6.20 × 10−4	M7C3 + M23C6	52	1.9 × 1012	90–400
120	95	23.00 × 10−4	M7C3 + M23C6	45	2.0 × 1012	–
130	60	53.30 × 10−4	M7C3 + M23C6	40	2.1 × 1012	–
170	11	242.00 × 10−4	M7C3 + M23C6	31	2.1 × 1012	–

On the other hand, the highest number density and lowest size of precipitates are responsible for the highest amount of ductile-mode fracture for specimens tested at high stresses, since the formation and coalescence of microvoids are facilitated for this precipitation.

Finally, Table 5 also shows that the time-to-rupture t_r is longer for test stresses of 78, 88 and 95 MPa than that reported in the literature¹⁰⁾ for similar steel compositions. This fact suggests that the creep resistance of present work is higher, which may be attributed to the fine grain size of about 5 μm for ferrite, in comparison to the grain size of approximately 25 to 50 μm for the steel reported in the literature.¹⁰⁾ This fact may also promote a more abundant intergranular precipitation of carbides. Besides, the as-received steel has the presence of small amount of bainite. In contrast, the micrographs of some steels, reported in reference,¹⁰⁾ indicate only the occurrence of ferrite grains. The presence of bainite may also cause higher strengthening for present work steel, but its effect would be no considerable for low-stress tests, since the long time-to-rupture t_r could cause its transformation to ferrite.

4. Conclusions

The creep properties for creep resistant 5Cr-0.5Mo steel were determined at 600°C, as well as its relation to the precipitation evolution during test, and the conclusions are:

(1) The n exponent of 6.29 determined from the Norton-Bailey law, indicated that the deformation process took place through the ferrite grains causing transgranular ductile-fracture mode after creep tests, which was observed in the creep-tested specimens.

(2) The XRD results indicate that the creep test at 600°C and 78–170 MPa are conducted under the presence of a mixture of M₇C₃ and M₂₃C₆ carbides.

(3) The duration of secondary-creep stage and time-to-rupture t_r are closely related to the radius and number density of carbides.

(4) The creep strength, expressed by time-to-rupture t_r, of this work steel was determined to be higher than those reported in the literature for this type of steel tested at the same stress level.

Acknowledgements

The authors wish to acknowledge the financial support from SIP-IPN, Conacyt and Kagoshima University.

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