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Microstructural degeneration of simulated heat-affected zone in 2.25Cr–1Mo steel during high-temperature exposure
Materials Science and Engineering A340 (2003) 15 /32 www.elsevier.com/locate/msea
Microstructural degeneration of simulated heat-affected zone in 2.25Cr 1Mo steel during high-temperature exposure /
M.C. Tsai 1, J.R. Yang * Department of Materials Science and Engineering, Institute of Materials Science, College of Engineering, National Taiwan University, 1 Roosevelt Rd., Sec. 4, Taipei, Taiwan, ROC Received 28 August 2001; received in revised form 25 January 2002
Abstract The microstructural of simulated heat-affected zone (HAZ) in 2.25Cr /1Mo steels has been investigated. The experiments with different heat inputs were carried out in a high speed dilatometer; the thermal cycles used corresponded to the actual thermal cycles that occurred in the coarse-grained region of the real HAZ. It was found that with a heat input of 20 kJ cm 1, the simulated HAZ microstructure gave larger amounts of lower bainite with significant amounts of martensite. The 20 kJ cm 1 heat input specimens were tempered at 700 8C for different time intervals ranging from 1 to 50 h. A sequence for corresponding microstructural degradation has been proposed. # 2002 Elsevier Science B.V. All rights reserved. Keywords: Cr /Mo steel; Coarse-grained heat-affected zone; Thermal simulation; High-temperature exposure; Microstructural change
1. Introduction The 2.25Cr/1Mo steel is the most frequently used in power generation industry for containment vessels, superheater tubes and steam pipes. These components are often used in the temperature range 480/565 8C with the stresses of about 15 /30 MPa over time periods of some 30 years [1,2]. The main factor responsible for the good creep resistance of this low alloy steel is the formation of fine and highly stable dispersions of alloy carbides, although a significant contribution also comes from solid solution strengthening. During fabrication of the vessels or pipes, the submerged-arc welding process (using different heat inputs in the range of 10 /60 kJ cm 1) is a necessary process [3]. As a consequence of
* Corresponding author. Tel.: /886-2-23632756; fax: /886-223634562 E-mail addresses: [email protected] (M.C. Tsai), [email protected] (J.R. Yang). 1 Tel.: /886-2-23620601; fax: /886-2-23634562.
this severe thermal cycle, the heat-affected zone (HAZ) occurs near the fusion boundary. That the existence of the HAZ causes a serious deterioration in toughness has been received much attention, especially the coarsegrained HAZ associated with the brittle microstructure [4]. Considerable effort has been devoted to assessment of creep life by extrapolating the results of laboratory short-term creep tests to the long-term service lives. Such parametric techniques, however, rarely consider microstructural change accompanying exposure to high temperatures. In fact, the microstructure is not stable and changes gradually with time [5 /8], especially in the HAZ. It is evident that microstructural degeneration during service depends on the initial microstructure. Although a lot of research work has been done on the 2.25Cr/1Mo steel, there is no systematic study on phase transformations in HAZ in this steel during hightemperature exposures. In the present work, the simulated HAZ (with the heat input of 20 kJ cm 1) have been employed to investigate the details of microstructural changes and carbide transformations during tempering at 700 8C for the time up to 50 h.
0921-5093/02/$ - see front matter # 2002 Elsevier Science B.V. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 0 2 ) 0 0 0 8 1 - 3
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Table 1 Chemical composition of the 2.25Cr /1Mo steel (wt.%) Fe
C
Cr
Mo
Mn
P
S
Si
Ni
Bal.
0.15
2.20
1.05
0.56
0.010
B 0.005
0.054
0.053
2. Experimental procedure Specimens of 2.25Cr /1Mo steel were received as hotrolled plates. The composition of the steel is listed in Table 1. In this work, the simulated HAZ experiments were performed on a Dilatronic III RDP deformation dilatometer of Theta Industries, Inc. Before thermal simulation, the representative weld thermal cycles of actual welding had been obtained as described below. Test beads were deposited on steel plates 20 mm thick using the bead on plate technique by submerged-arc welding with three different heat inputs, 20, 50 and 80 kJ cm 1. The thermal cycles were measured at the location of coarse-grained region in the HAZ close to the fusion boundary. The thermal cycles are shown in Fig. 1; the peak temperatures were approaching 1350 8C for the three cycles and cooling time between 800 and 500 8C, Dt8/5, for the heat inputs of 20, 50 and 80 kJ cm 1 were 16, 102 and 220 s, respectively. Before preparation of dilatometer specimens, the pieces of 2.25Cr/1Mo steel were homogenized at 1200 8C for 3 days while sealed in a quartz tube containing a partial pressure of pure argon. The specimens were then machined to 3 mm diameter cylindrical rods of 6 mm length. The dilatometer was connected to a computer workstation with a PDP 11/55 central processor to analyze the resulting data. A software package (also provided by theta
Fig. 1. Thermal cycles of HAZ simulations corresponding to real thermal cycles for heat inputs with 20, 50 and 80 kJ cm 1.
Industries) can give a flexible environment to execute identical thermal cycles for HAZ simulation. The length, time and temperature information can be recorded at microsecond intervals and the level of vacuum can reach 10 5 torr to protect specimens from oxidation. The
Fig. 2. Optical metallography showing microstructure obtained after HAZ simulations corresponding to real thermal cycles for heat input with: (a) 20; (b) 50; (c) 80 kJ cm 1.
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Fig. 3. Calculated TTT diagram for steel studied: after Bhadeshia [9].
Fig. 4. Dilatometric curves for the simulated coarse-grained HAZ specimen with the heat inputs of 20, 50 and 80 kJ cm 1.
Table 2 Phase transformation temperatures measured from dilatometric curves Heat input (kJ cm 1)
Transformation
Start temperature (8C) Finish temperature (8C)
20
50
80
485 335
535 405
540 380
Fig. 5. Transmission electron micrographs obtained from: (a) 20; (b) 50; (c) 80 kJ cm 1 heat input specimen.
high-temperature exposures for simulated HAZ specimens were carried out at 700 8C for 1, 2, 5, 10, 20 and 50 h in an air furnace, while the specimens were being
Table 3 Vickers hardness measurement for the specimens treated by three different heat inputs Region
Gray region White region Possible microstructures
Heat input (kJ cm 1) 20
50
80
35294 39997 Martensitelower bainite
34993 32195 Lower bainiteupper bainite
33394 31892 Upper bainiteallotriomorphic ferrite
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Fig. 6. Optical micrographs of tempering at 700 8C for: (a) 1; (b) 2; (c) 5; (d) 10; (e) 20; (f) 50 h.
sealed in quartz capsules under a partial pressure of argon. The specimens for optical metallography were prepared from dilatometer specimens and high-temperature exposured specimens. They were mechanically polished and then etched in 2% nital solution. Hardness measurement was made on optical specimens, using a Vickers hardness tester. A load of 300 g was used in order to make the indentation on the individual phase. Transmission electron microscopy (TEM) specimens were prepared from 0.25 mm thick discs. The discs were thinned to 0.07 mm by abrasion on SiC papers and then
twin-jet electropolished using a mixture of 5% perchloric acid, 25% glycerol, and 70% ethanol at /5 8C, using a 60 V polishing potential. They were examined using a JEM-100CX II TEM operated at 100 kV. Carbon extraction replicas were prepared from specimens that had been polished for optical microscopy. The specimens were heavily etched in 2% nital solution, and then a carbon film (about 15/30 nm in thickness) was deposited by evaporation. The carbon film of the specimen was scored with a sharp knife into a 1.5 mm square grid pattern, etched again with 5% nital solution to dissolve the matrix, and then the replicas were
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Fig. 7. Transmission electron micrographs showing the microstructural degeneration in lower bainitic region after tempering at 700 8C for: (a) 1; (b) 20; (c) 50 h.
‘floated off’ on the surface of ethanol. Finally, the replicas were picked up on copper grids and dried on filter paper. To investigate the morphology, size distribution, chemical microanalysis, and crystal structure of carbides, the replicas were examined on JOEL JEM2000EX STEM and Philips CM 200 FEG-TEM with energy dispersive X-ray (EDX) spectrometer (STEMEDS) operated at 200 kV.
3. Results and discussion 3.1. Simulated HAZ microstructures To investigate the effect of heat input on the microstructural constituents in the 2.25Cr /1Mo steel, three different thermal simulations (Fig. 1) were carried out on a dilatometer for optical metallography specimens
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Fig. 8. Transmission electron micrographs showing the microstructural evolution in martensitic region after tempering at 700 8C for: (a) 1; (b) 20; (c) 50 h.
and transmission electron micrograph specimens. The representative microstructures of coarse-grained HAZ under different three heat input conditions are shown in Fig. 2; when the heat input is larger, the microstructure becomes much coarser. However, in the optical metallographs, it is very difficult to reveal the microstructural details for the bainite and martensite.
For the purpose of elucidating the simulated HAZ microstructures, further investigations on the time/ temperature-transformation (TTT) diagram, dilatometry and transmission electron micrography were carried out. Fig. 3 shows the calculated TTT diagram, based on the model developed by Bhadeshia [9]. In such a diagram, the upper C curve represents the time taken
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Fig. 9. (a) Transmission electron micrograph of a replica showing the microstructure of bainitic region obtained after tempering at 700 8C for 1 h; (b) diffraction pattern of o-carbide; (c) diffraction pattern of M3C; (d) diffraction pattern of M7C3.
for the initiation of diffusion transformations such as allotriomorphic ferrite and pearlite, whereas the lower C curve represents the time taken for the initiation of displacive transformation such as bainite. The TTT diagram shows that the martensite start temperature (Ms) is about 421 8C and the bainite start temperature Bs about 547 8C. Fig. 4 shows the dilatometric curves obtained from the simulated HAZ specimens with the heat inputs of 20, 50 and 80 kJ cm 1, respectively. As the heat input is larger, the cooling time from 800 to 500 8C is longer and the transformation start temperature is higher. The start and finish temperatures of transformation were measured and listed in Table 2. Fig. 5 show the TEM micrographs obtained from the 20, 50 and 80 kJ cm 1 heat input specimens. In the case of the specimen with heat input of 20 kJ cm 1, it illustrates that the microstructure is composed chiefly of lower bainite (with a single variant of sheetlike intralath
carbide) with some amounts of martensite. The evidence suggests that the gray and white etching structures in Fig. 2(a) should be separately lower bainite and martensite. The result is consistent with that reported by Miranda et al. who investigated the microstructure of HAZ with 20 kJ cm 1 heat input in a 2.25Cr/1Mo steel [10]. The transmission electron micrographs obtained from the specimens with 50 and 80 kJ cm 1 show the existence of sub-unit of upper bainite with martensite/ austenite (M/A) constituents. The Vickers hardness value of the gray etching structure is higher than that of the white etching structure in the specimens with heat input 50 and 80 kJ cm 1. The possible microstructures in specimens with three kinds of heat inputs are listed in Table 3. The result shows that higher volume fractions of upper bainite were obtained in the cases of heat input with 50 and 80 kJ cm 1, and that a mixture of lower bainite and martensite were obtained in the case of heat
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Fig. 10. (a) Transmission electron micrograph of a replica showing the microstructure of bainitic region obtained after tempering at 700 8C for 10 h; (b) diffraction pattern of M7C3.
input with 20 kJ cm 1. As a part of the program for the research on the simulated HAZ in a 2.25Cr /1Mo steel, the work reported in Sections 3.2 and 3.3 deal only with microstructural degeneration in 20 kJ cm 1 heat input specimens during exposure at 700 8C. 3.2. Microstructural evolution during tempering The optical metallographs obtained from the 20 kJ cm 1 heat input specimens after tempering at 700 8C for different intervals (1, 2, 5, 10, 20 and 50 h) are presented in Fig. 6. The corresponding metallographs do not reveal any clear changes in the microstructure during the early stages of tempering (Fig. 10(a /d)), although fine dots of carbide precipitate can be detected after tempering at 700 8C for 20 h (Fig. 10(e /f)). Attention is therefore focused on TEM. Fig. 7 shows the detailed microstructural change in lower bainite region during exposure at 700 8C for 1, 20 and 50 h, respectively. At the initial stage of tempering (Fig. 7(a)), prior cementite carbide in the lower bainite
Fig. 11. (a) Transmission electron micrograph of a replica showing the microstructure of bainitic region obtained after tempering at 700 8C for 20 h; (b) diffraction pattern of M7C3; (c) diffraction pattern of M2C.
sub-unit became coarse and still kept a strong alignment in a specific direction. The recovery phenomenon did not happen until 20 kJ cm 1 heat input specimen was tempered at 700 8C for 20 h (Fig. 7(b)); the characteristic plate morphology of bainitic ferrite sub-units remained although a slight recovery occurred. Prolonged tempering at 700 8C for 50 h (Fig. 7(c)) made carbides spherodized and coalesced, and made the phenomenon of recovery much more significant. However, the broad faces of the platelets of bainitic ferrite were less affected, in spite of the fact that the tips of the platelets of bainitic ferrite became round. The microstructural evolution in martensitic region during tempering was much different to that in bainitic region. Fig. 8 shows the transmission electron micrographs of the martensitic region during tempering at 700 8C for different time intervals. In the case of tempering at 700 8C for 1 h (Fig. 8(a)), the recovery occurred obviously. After tempering at 700 8C for 20 h (Fig. 8(b)), recrystallized ferrite subgrains formed within martensite laths. Prolonged tempering at 700 8C for 50 h (Fig. 8(c)) led to the complete recrystallization in prior martensitic matrix.
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Fig. 12. (a) Transmission electron micrograph of a replica showing the microstructure of bainitic region obtained after tempering at 700 8C for 50 h; (b) diffraction pattern of M7C3; (c) diffraction pattern of M2C; (d) diffraction pattern of M23C6.
As compared the tempered structures of lower bainite and martensite in 20 kJ cm 1 heat input specimens, bainite seems to be more stable than martensite structure. In fact, martensite is the more deformed structure than bainite. The sub-units of bainitic ferrite possess a plate shape which is a consequence of displacive transformation with an invariant plane strain shape deformation (similar to that of martensite). However, the stored energy in martensite is much larger than that in bainite. The plate morphologies of bainite and martensite are metastable and could easily change if the specimens were to be reheated to a high temperature. When compared with martensite, bainite grows at relatively high temperature at which the microstructure
undergoes some recovery during transformation; the extent of recovery with bainite is much larger than that associated with martensite. As a result, that the tempering response is more sensitive in martensite than in bainite is expected.
3.3. Carbide evolution during tempering It has been suggested that carbide may be replaced by another carbides in two basic modes [11]: 1) ‘In situ transformation’: The new alloy carbide nucleates at the interface between the original
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Table 4 The precipitation sequence of carbides in bainitic region Region
Aging time (h) 1
2
5
10
20
50
Bainite region Boundary In sub-unit
o-Carbide, M3C o-Carbide, M3C, M7C3
M3C o-Carbide, M3C, M7C3
M3C, M7C3 M3C, M7C3
M7C3 M7C3
M7C3 M7C3, M2C
M7C3, M23C6 M2C, M23C6
Aging time (h)
Carbides
Fe (wt.%)
Cr (wt.%)
Mo (wt.%)
Bainite lath boundary 1 M3C o-Carbide 2 M3C 5 M7C3 M3C 10 M7C3 20 M7C3 50 M7C3 M23C6
66.8 49.1 66.8 35.9 51.8 34.3 30.8 43.2 35.2
24.5 36.3 22.2 52.9 33.9 51.4 53.4 50.2 54.8
8.7 14.6 11.0 11.1 14.4 14.3 15.8 6.7 10.0
Aging time (h)
Fe (wt.%)
Cr (wt.%)
Mo (wt.%)
60.2 38.3 66.8 37.8 49.6 2.0 36.9 34.7 36.1 1.6 37.8 2.3
28.4 54.0 22.2 52.1 36.0 12.6 53.9 54.4 53.6 13.2 56.0 11.7
11.4 7.8 11.0 10.1 14.4 85.3 9.2 10.8 10.3 85.2 6.1 86.0
Carbides
Within bainite lath 1 M3C M7C3 2 M3C M7C3 o-Carbide M2C 5 M7C3 10 M7C3 20 M7C3 M2C 50 M23C6 M2C
carbide and matrix, and the new carbide grows until the original carbide disappears. 2) ‘Separate nucleation’: The new alloy carbides nucleates separately at new sites primarily on dislocations, and the original carbide particles dissolve in the ferrite matrix.
3.3.1. Bainitic region The identification of carbide structures in the 20 kJ cm 1 heat input specimens during exposure at 700 8C have been carried out by analyzing the diffraction patterns and the chemical microanalyses from STEMEDX data. Fig. 9 shows the transmission electron micrograph of replica obtained from the specimen tempered at 700 8C for 1 h. It indicates that o-carbide and M3C were located on grain boundaries of bainitic ferrite. Besides, M7C3 particles were found within bainitic sub-units. Baker and Nutting [12] had proposed the sequences of carbide transformation in the tempered-bainite microstructure in the normalized 2.25Cr/ 1Mo steel as follows:
8 M C 0 M7 C 3 0 M 6 C > -carbide < 3 M3 C gM3 C 0 gM23 C6 0 M6 C > : M3 C M2 C Although Baker and Nutting did not detect M7C3 carbides until tempering at 700 8C for 5 h, M7C3 carbide particles have been observed in the simulated HAZ sample after tempering at 700 8C for 1 h. After tempering at 700 8C for 10 h, almost all carbides have transformed to M7C3 either on the bainitic sub-unit boundaries or within the sub-units. TEM micrograph of Fig. 10 shows that lump-shaped M7C3 carbide particles were located within bainitic subunits, and rod-shaped M7C3 carbide particles were discovered along bainitic sub-unit boundaries. The shape of M7C3 carbide appears to be rod-like and is very different from that of the M3C carbides (the circular-shape particles in sub-units and the film along the boundaries). Fig. 11 presents the TEM micrograph of replica after tempering at 700 8C for 20 h; it shows that tiny M2C particles were scattered within the bainitic sub-units. Since not any carbide can be found to be
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Fig. 13. (a) Transmission electron micrograph of a replica showing the microstructure of martensitic region obtained after tempering at 700 8C for 2 h; (b) diffraction pattern of o-carbide; (c) diffraction pattern of M3C.
adjacent to the M2C carbide, the mechanism of the M2C creation is suggested to be the type of ‘separate nucleation’ instead of the type of ‘in situ transformation’. The evolution of carbides in bainitic matrix obtained from the specimen tempered at 700 8C for 50 h is illustrated in Fig. 12; M23C6 particles were detected on the sub-unit boundaries and within the bainitic sub-units. The size of M23C6 is slightly smaller than that of M7C3 and the corresponding chemical compositions of M23C6 particles detected are listed with Fig. 12. The EDX data show that the Cr content in M23C6 carbide is lower than that in M7C3, and that Mo content in M23C6 carbide is higher vice versa. The sequences of carbide formations in bainitic region and
the chemical microanalyses for corresponding carbides are listed in Table 4. It clearly indicates that in the bainitic region in HAZ specimens o-carbide formed at initial stage and M23C6 carbide existed at the final stage. 3.3.2. Martensitic region Fig. 13 presents the replica micrograph obtained from the specimen tempered at 700 8C for 2 h; the M3C carbide particles were located on martensite lath boundaries and within martensite laths. o-carbide particles can also be detected when the specimen was tempered at 700 8C for 2 h. Fig. 14 illustrates that o-carbide transformed to M3C by a manner of in situ nucleation in the martensite region in the simulated HAZ specimen
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¯ reflection; (c) Fig. 14. Transmission electron micrograph of replicas showing: (a) bright field image; (b) dark field image using o-carbide (11¯ 1) ¯ reflection; (g) corresponding corresponding diffraction pattern; (d) interpretation of (c); (e) bright field image; (f) dark field image using M3C (220) diffraction pattern; (h) interpretation of (g).
after tempering at 700 8C for 2 h. Transmission electron micrographs of the replica show the bright and dark field images with the corresponding diffraction patterns for o-carbide (as indicated by position A in Fig. 14(a)) and M3C (as indicated by position B in Fig. 14(e)). The measured composition of o-carbide was 73.7Fe /
19.3Cr/7.0Mo (wt.%), and that of M3C was 63.7Fe / 28.0Cr/8.3Mo (wt.%). These two carbides contacted each other and their chemical compositions were only slightly different. It is suggested that as the M3C carbide grows, alloying elements are provided by the adjacent ocarbide, which gradually disappears.
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Fig. 15. (a) Transmission electron micrograph of a replica showing the microstructure of martensitic region obtained after tempering at 700 8C for 10 h; (b) diffraction pattern of M7C3.
After tempering for 10 h, the carbides almost transformed to M7C3 carbides (Fig. 15). The M2C and M23C6 carbides were discovered in the specimen tempered at 700 8C for 50 h (Fig. 16). It is suggested that the mechanisms of M2C and M23C6 nucleation are a type of ‘separate nucleation’. The sequences of carbide transformations in martensite region and the chemical microanalyses for the corresponding carbides are given in Table 5. Fig. 17 also present the details for M3C 0/M7C3 transformation in the way of ‘in situ transformation’ in the martensite region in the simulated HAZ specimen
after tempering at 700 8C for 10 h (M3C and M7C3 are indicated by positions C and D in Fig. 17, respectively); the measured composition of M3C was 53.6Fe /31.5Cr / 14.9Mo (wt.%), and that of M7C3 was 32.5Fe /48.8Cr / 18.7Mo (wt.%). These two carbides contacted each other and the concentrations of Cr and Mo were slightly higher in M7C3 than in M3C. It is suggested that as the M7C3 carbide grows, alloy elements are provided by the adjacent M3C carbide, which gradually disappears. From the above observation, it shows two kinds of mechanisms: in situ transformation and separate nucleation for carbide transformations in simulated HAZ
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Fig. 16. (a) Transmission electron micrograph of a replica showing the microstructure of martensitic region obtained after tempering at 700 8C for 50 h; (b) diffraction pattern of M7C3; (c) diffraction pattern of M23C6; (d) diffraction pattern of M2C.
samples in 2.25Cr /1Mo steel. The carbide transformation sequences in martensitic and bainitic regions are similar; they are proposed as follows:
3.4. The characteristic X-ray spectra of EDX Morphological differences can sometimes act as a guide in the identification of carbides in the steel studied. The carbide with a distinctive needle shaped
M23 C6 f-carbideg 0fM3 Cg 0fM7 C3 g 0 M2 C in situ
in situ
separate nucleation
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Table 5 The precipitation sequence of carbides in martensitc region Region
Aging time (h) 1
2
5
10
20
50
Martensite region Boundary In lath
M3C, o-carbide M3C, o-carbide
M3C, o-carbide M3C
M3C, M7C3 M7C3
M7C3 M7C3
M7C3 M7C3
M23C6 M7C3, M23C6, M2C
Aging time (h)
Carbides
Fe (wt.%)
Cr (wt.%)
Mo (wt.%)
M3C M3C M2C M7C3 M3C M7C3 M7C3 M2C M7C3 M23C6
64.7 67.5 14.0 37.5 54.1 33.5 37.9 27.8 35.1 35.7
26.7 24.7 20.8 50.7 31.5 53.0 53.0 9.7 55.8 51.5
8.6 7.9 65.2 11.8 14.3 13.4 9.1 62.5 9.0 12.8
Carbides
Fe (wt.%)
Cr (wt.%)
Mo (wt.%)
62.4 70.8 33.0 74.8 36.0 37.2 32.5 32.7 35.2 36.8 3.5
29.0 19.7 49.9 18.8 52.9 53.4 49.2 51.3 52.6 50.1 16.0
8.6 9.5 17.1 6.4 11.0 9.4 18.3 15.9 12.2 13.1 80.6
At boundary 1 2 5 10 20 50
Aging time (h)
Within martensite lath 1 M3C o-Carbide M7C3 2 M3C M7C3 5 M7C3 10 M7C3 20 M7C3 50 M7C3 M23C6 M2C
morphology is presumed to be M2C, but both M7C3 and M23C6 have the similar morphologies with the spherical form. Thus, it cannot be concluded that the shape of the particle alone is a guide to identification. Most of the carbide phases can certainly be identified using electron diffraction, but this is very difficult to examine large numbers of small particles. The carbides are often too thick to provide resolvable diffraction patterns. Another alternative method to identify the type of carbides is based on EDX microanalysis. This technique for carbide identification has been usually carried out in the TEM and permits chemical analyses to be taken from individual carbides. Although electron diffraction, microanalysis and morphological observation have been carried out to identify the carbides in this work, an attempt to use the standard EDS spectra for identifying the carbide has been done. The result is consistent with that reported by Hippsley [8] and Pilling and Ridley [13]. The typical EDS spectra taken from a range of carbides investigated are shown in Fig. 18. The significant peaks in Fig. 18 are for the peak energies of molybdenum (L), chromium
(Ka , Kb ) and iron (Ka , Kb ). The spectra were obtained from the various carbides: o-carbide, M3C, M7C3, M23C6 and M2C, respectively; the corresponding crystal structures were identified by analyses of electron diffraction patterns. These spectra can provide an access to identify the type of carbide in 2.25Cr /1Mo steel in this work.
4. Conclusions (1) The initial microstructure obtained from simulated coarse-grained HAZ with heat 20 kJ cm 1 input was lower bainite with martensite. (2) Bainite structure is more stable than martensite structure during tempering at elevated temperatures for long term exposures. (3) The sequences of carbide transformation were different at the regions on the bainite boundary, within bainitic sub-unit, on martensite lath boundary and within martensite lath. Their differences are suggested as follows:
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¯ reflection; (c) Fig. 17. Transmission electron micrograph of replicas showing: (a) bright field image; (b) dark field image using M3C (021) ¯ reflection; (g) corresponding corresponding diffraction pattern; (d) interpretation of (c); (e) bright field image; (f) dark field image using M7C3 (122) diffraction pattern; (h) interpretation of (g).
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Fig. 18. Typical energy-dispersive X-ray spectra taken from five different types of carbide detected.
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(a) Bainitic sub-unit boundary in situ
in situ
-carbide 0 M3 C 0
Acknowledgements
separate nucleation Fe3 C 0M7 C3 0 M23 C6 M7 C3
(b) Bainitic sub-unit
8 ( ( separate nucleation <-carbide M3 C in situ 0 M C3 0 M23 C6 M3 C 0 0 M7 C3 separate7nucleation M7 C3 : 0 M2 C M7 C3
This work was carried out with financial support from the National Science Council of the Republic of China, Taiwan, under Contract NSC 87-TPC-E-002-032. The authors are grateful for the support of this research.
(c) Martensite lath boundary
separate nucleation -carbidein situ M3 C 0M7 C3 0 M23 C6 0 M3 C M7 C3
(d) Martensite lath 8 > <
-carbidein situ
> :M 3 C
in situ
0 M3 C 0 M7 C3
8separate nucleation > 0 M23 C6 < > :
separate nucleation
0 0 M7 C3
M2 C
(4) The final equilibrium precipitate for 2.25Cr /1Mo steel studied is M23C6 carbide. M6C carbide was not encountered in the microstructural investigation. (5) The characteristic EDX spectra for five various carbides have been used to identify type of carbides in 2.25Cr /1Mo steel.
References [1] H.K.D.H. Bhadeshia, Bainite in steels, The Institute of Materials, London, 1992, pp. 329 /346. [2] J.R. Davis, et al., Metals Handbook, vol. 1, 10th ed., ASM International, USA, 1990, pp. 617 /652. [3] R.M. Miranda, M.A. Fortes, Mater. Sci. Eng. A108 (1989) 1 /8. [4] W.H. Rearns (Ed.), Welding Handbook Metals and Their Weldability, vol. 4, American Welding Society, Miami, FL, 1980. [5] J. Yu, C.J. Mcmahon, Jr., Metall. Trans. A 11A (1980) 277 /289. [6] J. Yu, C.J. Mcmahon, Jr., Metall. Trans. A 11A (1980) 291 /300. [7] J. Pilling, N. Ridley, Metall. Trans. A 13A (1982) 557 /563. [8] C.A. Hippsley, Metal Sci. (1981) 137 /147. [9] H.K.D.H. Bhadeshia, Metal Sci. 16 (1982) 159 /165. [10] R.M. Miranda, M.A. Fortes, Mater. Sci. Eng. A108 (1989) 1 /8. [11] R.W.K. Honeycombe, H.K.D.H. Bhadeshia, Steels Microstructure and Properties, Edward Arnold, London, 1995, p. 187. [12] R.G. Baker, J. Nutting, J. Iron Steel Inst. 192 (1959) 257 /268. [13] J. Pilling, N. Ridley, Metall. Trans. A 13A (1982) 557 /563.
Microstructural degeneration of simulated heat-affected zone in 2.25Cr 1Mo steel during high-temperature exposure /
M.C. Tsai 1, J.R. Yang * Department of Materials Science and Engineering, Institute of Materials Science, College of Engineering, National Taiwan University, 1 Roosevelt Rd., Sec. 4, Taipei, Taiwan, ROC Received 28 August 2001; received in revised form 25 January 2002
Abstract The microstructural of simulated heat-affected zone (HAZ) in 2.25Cr /1Mo steels has been investigated. The experiments with different heat inputs were carried out in a high speed dilatometer; the thermal cycles used corresponded to the actual thermal cycles that occurred in the coarse-grained region of the real HAZ. It was found that with a heat input of 20 kJ cm 1, the simulated HAZ microstructure gave larger amounts of lower bainite with significant amounts of martensite. The 20 kJ cm 1 heat input specimens were tempered at 700 8C for different time intervals ranging from 1 to 50 h. A sequence for corresponding microstructural degradation has been proposed. # 2002 Elsevier Science B.V. All rights reserved. Keywords: Cr /Mo steel; Coarse-grained heat-affected zone; Thermal simulation; High-temperature exposure; Microstructural change
1. Introduction The 2.25Cr/1Mo steel is the most frequently used in power generation industry for containment vessels, superheater tubes and steam pipes. These components are often used in the temperature range 480/565 8C with the stresses of about 15 /30 MPa over time periods of some 30 years [1,2]. The main factor responsible for the good creep resistance of this low alloy steel is the formation of fine and highly stable dispersions of alloy carbides, although a significant contribution also comes from solid solution strengthening. During fabrication of the vessels or pipes, the submerged-arc welding process (using different heat inputs in the range of 10 /60 kJ cm 1) is a necessary process [3]. As a consequence of
* Corresponding author. Tel.: /886-2-23632756; fax: /886-223634562 E-mail addresses: [email protected] (M.C. Tsai), [email protected] (J.R. Yang). 1 Tel.: /886-2-23620601; fax: /886-2-23634562.
this severe thermal cycle, the heat-affected zone (HAZ) occurs near the fusion boundary. That the existence of the HAZ causes a serious deterioration in toughness has been received much attention, especially the coarsegrained HAZ associated with the brittle microstructure [4]. Considerable effort has been devoted to assessment of creep life by extrapolating the results of laboratory short-term creep tests to the long-term service lives. Such parametric techniques, however, rarely consider microstructural change accompanying exposure to high temperatures. In fact, the microstructure is not stable and changes gradually with time [5 /8], especially in the HAZ. It is evident that microstructural degeneration during service depends on the initial microstructure. Although a lot of research work has been done on the 2.25Cr/1Mo steel, there is no systematic study on phase transformations in HAZ in this steel during hightemperature exposures. In the present work, the simulated HAZ (with the heat input of 20 kJ cm 1) have been employed to investigate the details of microstructural changes and carbide transformations during tempering at 700 8C for the time up to 50 h.
0921-5093/02/$ - see front matter # 2002 Elsevier Science B.V. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 0 2 ) 0 0 0 8 1 - 3
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Table 1 Chemical composition of the 2.25Cr /1Mo steel (wt.%) Fe
C
Cr
Mo
Mn
P
S
Si
Ni
Bal.
0.15
2.20
1.05
0.56
0.010
B 0.005
0.054
0.053
2. Experimental procedure Specimens of 2.25Cr /1Mo steel were received as hotrolled plates. The composition of the steel is listed in Table 1. In this work, the simulated HAZ experiments were performed on a Dilatronic III RDP deformation dilatometer of Theta Industries, Inc. Before thermal simulation, the representative weld thermal cycles of actual welding had been obtained as described below. Test beads were deposited on steel plates 20 mm thick using the bead on plate technique by submerged-arc welding with three different heat inputs, 20, 50 and 80 kJ cm 1. The thermal cycles were measured at the location of coarse-grained region in the HAZ close to the fusion boundary. The thermal cycles are shown in Fig. 1; the peak temperatures were approaching 1350 8C for the three cycles and cooling time between 800 and 500 8C, Dt8/5, for the heat inputs of 20, 50 and 80 kJ cm 1 were 16, 102 and 220 s, respectively. Before preparation of dilatometer specimens, the pieces of 2.25Cr/1Mo steel were homogenized at 1200 8C for 3 days while sealed in a quartz tube containing a partial pressure of pure argon. The specimens were then machined to 3 mm diameter cylindrical rods of 6 mm length. The dilatometer was connected to a computer workstation with a PDP 11/55 central processor to analyze the resulting data. A software package (also provided by theta
Fig. 1. Thermal cycles of HAZ simulations corresponding to real thermal cycles for heat inputs with 20, 50 and 80 kJ cm 1.
Industries) can give a flexible environment to execute identical thermal cycles for HAZ simulation. The length, time and temperature information can be recorded at microsecond intervals and the level of vacuum can reach 10 5 torr to protect specimens from oxidation. The
Fig. 2. Optical metallography showing microstructure obtained after HAZ simulations corresponding to real thermal cycles for heat input with: (a) 20; (b) 50; (c) 80 kJ cm 1.
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Fig. 3. Calculated TTT diagram for steel studied: after Bhadeshia [9].
Fig. 4. Dilatometric curves for the simulated coarse-grained HAZ specimen with the heat inputs of 20, 50 and 80 kJ cm 1.
Table 2 Phase transformation temperatures measured from dilatometric curves Heat input (kJ cm 1)
Transformation
Start temperature (8C) Finish temperature (8C)
20
50
80
485 335
535 405
540 380
Fig. 5. Transmission electron micrographs obtained from: (a) 20; (b) 50; (c) 80 kJ cm 1 heat input specimen.
high-temperature exposures for simulated HAZ specimens were carried out at 700 8C for 1, 2, 5, 10, 20 and 50 h in an air furnace, while the specimens were being
Table 3 Vickers hardness measurement for the specimens treated by three different heat inputs Region
Gray region White region Possible microstructures
Heat input (kJ cm 1) 20
50
80
35294 39997 Martensitelower bainite
34993 32195 Lower bainiteupper bainite
33394 31892 Upper bainiteallotriomorphic ferrite
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Fig. 6. Optical micrographs of tempering at 700 8C for: (a) 1; (b) 2; (c) 5; (d) 10; (e) 20; (f) 50 h.
sealed in quartz capsules under a partial pressure of argon. The specimens for optical metallography were prepared from dilatometer specimens and high-temperature exposured specimens. They were mechanically polished and then etched in 2% nital solution. Hardness measurement was made on optical specimens, using a Vickers hardness tester. A load of 300 g was used in order to make the indentation on the individual phase. Transmission electron microscopy (TEM) specimens were prepared from 0.25 mm thick discs. The discs were thinned to 0.07 mm by abrasion on SiC papers and then
twin-jet electropolished using a mixture of 5% perchloric acid, 25% glycerol, and 70% ethanol at /5 8C, using a 60 V polishing potential. They were examined using a JEM-100CX II TEM operated at 100 kV. Carbon extraction replicas were prepared from specimens that had been polished for optical microscopy. The specimens were heavily etched in 2% nital solution, and then a carbon film (about 15/30 nm in thickness) was deposited by evaporation. The carbon film of the specimen was scored with a sharp knife into a 1.5 mm square grid pattern, etched again with 5% nital solution to dissolve the matrix, and then the replicas were
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Fig. 7. Transmission electron micrographs showing the microstructural degeneration in lower bainitic region after tempering at 700 8C for: (a) 1; (b) 20; (c) 50 h.
‘floated off’ on the surface of ethanol. Finally, the replicas were picked up on copper grids and dried on filter paper. To investigate the morphology, size distribution, chemical microanalysis, and crystal structure of carbides, the replicas were examined on JOEL JEM2000EX STEM and Philips CM 200 FEG-TEM with energy dispersive X-ray (EDX) spectrometer (STEMEDS) operated at 200 kV.
3. Results and discussion 3.1. Simulated HAZ microstructures To investigate the effect of heat input on the microstructural constituents in the 2.25Cr /1Mo steel, three different thermal simulations (Fig. 1) were carried out on a dilatometer for optical metallography specimens
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Fig. 8. Transmission electron micrographs showing the microstructural evolution in martensitic region after tempering at 700 8C for: (a) 1; (b) 20; (c) 50 h.
and transmission electron micrograph specimens. The representative microstructures of coarse-grained HAZ under different three heat input conditions are shown in Fig. 2; when the heat input is larger, the microstructure becomes much coarser. However, in the optical metallographs, it is very difficult to reveal the microstructural details for the bainite and martensite.
For the purpose of elucidating the simulated HAZ microstructures, further investigations on the time/ temperature-transformation (TTT) diagram, dilatometry and transmission electron micrography were carried out. Fig. 3 shows the calculated TTT diagram, based on the model developed by Bhadeshia [9]. In such a diagram, the upper C curve represents the time taken
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Fig. 9. (a) Transmission electron micrograph of a replica showing the microstructure of bainitic region obtained after tempering at 700 8C for 1 h; (b) diffraction pattern of o-carbide; (c) diffraction pattern of M3C; (d) diffraction pattern of M7C3.
for the initiation of diffusion transformations such as allotriomorphic ferrite and pearlite, whereas the lower C curve represents the time taken for the initiation of displacive transformation such as bainite. The TTT diagram shows that the martensite start temperature (Ms) is about 421 8C and the bainite start temperature Bs about 547 8C. Fig. 4 shows the dilatometric curves obtained from the simulated HAZ specimens with the heat inputs of 20, 50 and 80 kJ cm 1, respectively. As the heat input is larger, the cooling time from 800 to 500 8C is longer and the transformation start temperature is higher. The start and finish temperatures of transformation were measured and listed in Table 2. Fig. 5 show the TEM micrographs obtained from the 20, 50 and 80 kJ cm 1 heat input specimens. In the case of the specimen with heat input of 20 kJ cm 1, it illustrates that the microstructure is composed chiefly of lower bainite (with a single variant of sheetlike intralath
carbide) with some amounts of martensite. The evidence suggests that the gray and white etching structures in Fig. 2(a) should be separately lower bainite and martensite. The result is consistent with that reported by Miranda et al. who investigated the microstructure of HAZ with 20 kJ cm 1 heat input in a 2.25Cr/1Mo steel [10]. The transmission electron micrographs obtained from the specimens with 50 and 80 kJ cm 1 show the existence of sub-unit of upper bainite with martensite/ austenite (M/A) constituents. The Vickers hardness value of the gray etching structure is higher than that of the white etching structure in the specimens with heat input 50 and 80 kJ cm 1. The possible microstructures in specimens with three kinds of heat inputs are listed in Table 3. The result shows that higher volume fractions of upper bainite were obtained in the cases of heat input with 50 and 80 kJ cm 1, and that a mixture of lower bainite and martensite were obtained in the case of heat
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Fig. 10. (a) Transmission electron micrograph of a replica showing the microstructure of bainitic region obtained after tempering at 700 8C for 10 h; (b) diffraction pattern of M7C3.
input with 20 kJ cm 1. As a part of the program for the research on the simulated HAZ in a 2.25Cr /1Mo steel, the work reported in Sections 3.2 and 3.3 deal only with microstructural degeneration in 20 kJ cm 1 heat input specimens during exposure at 700 8C. 3.2. Microstructural evolution during tempering The optical metallographs obtained from the 20 kJ cm 1 heat input specimens after tempering at 700 8C for different intervals (1, 2, 5, 10, 20 and 50 h) are presented in Fig. 6. The corresponding metallographs do not reveal any clear changes in the microstructure during the early stages of tempering (Fig. 10(a /d)), although fine dots of carbide precipitate can be detected after tempering at 700 8C for 20 h (Fig. 10(e /f)). Attention is therefore focused on TEM. Fig. 7 shows the detailed microstructural change in lower bainite region during exposure at 700 8C for 1, 20 and 50 h, respectively. At the initial stage of tempering (Fig. 7(a)), prior cementite carbide in the lower bainite
Fig. 11. (a) Transmission electron micrograph of a replica showing the microstructure of bainitic region obtained after tempering at 700 8C for 20 h; (b) diffraction pattern of M7C3; (c) diffraction pattern of M2C.
sub-unit became coarse and still kept a strong alignment in a specific direction. The recovery phenomenon did not happen until 20 kJ cm 1 heat input specimen was tempered at 700 8C for 20 h (Fig. 7(b)); the characteristic plate morphology of bainitic ferrite sub-units remained although a slight recovery occurred. Prolonged tempering at 700 8C for 50 h (Fig. 7(c)) made carbides spherodized and coalesced, and made the phenomenon of recovery much more significant. However, the broad faces of the platelets of bainitic ferrite were less affected, in spite of the fact that the tips of the platelets of bainitic ferrite became round. The microstructural evolution in martensitic region during tempering was much different to that in bainitic region. Fig. 8 shows the transmission electron micrographs of the martensitic region during tempering at 700 8C for different time intervals. In the case of tempering at 700 8C for 1 h (Fig. 8(a)), the recovery occurred obviously. After tempering at 700 8C for 20 h (Fig. 8(b)), recrystallized ferrite subgrains formed within martensite laths. Prolonged tempering at 700 8C for 50 h (Fig. 8(c)) led to the complete recrystallization in prior martensitic matrix.
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Fig. 12. (a) Transmission electron micrograph of a replica showing the microstructure of bainitic region obtained after tempering at 700 8C for 50 h; (b) diffraction pattern of M7C3; (c) diffraction pattern of M2C; (d) diffraction pattern of M23C6.
As compared the tempered structures of lower bainite and martensite in 20 kJ cm 1 heat input specimens, bainite seems to be more stable than martensite structure. In fact, martensite is the more deformed structure than bainite. The sub-units of bainitic ferrite possess a plate shape which is a consequence of displacive transformation with an invariant plane strain shape deformation (similar to that of martensite). However, the stored energy in martensite is much larger than that in bainite. The plate morphologies of bainite and martensite are metastable and could easily change if the specimens were to be reheated to a high temperature. When compared with martensite, bainite grows at relatively high temperature at which the microstructure
undergoes some recovery during transformation; the extent of recovery with bainite is much larger than that associated with martensite. As a result, that the tempering response is more sensitive in martensite than in bainite is expected.
3.3. Carbide evolution during tempering It has been suggested that carbide may be replaced by another carbides in two basic modes [11]: 1) ‘In situ transformation’: The new alloy carbide nucleates at the interface between the original
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Table 4 The precipitation sequence of carbides in bainitic region Region
Aging time (h) 1
2
5
10
20
50
Bainite region Boundary In sub-unit
o-Carbide, M3C o-Carbide, M3C, M7C3
M3C o-Carbide, M3C, M7C3
M3C, M7C3 M3C, M7C3
M7C3 M7C3
M7C3 M7C3, M2C
M7C3, M23C6 M2C, M23C6
Aging time (h)
Carbides
Fe (wt.%)
Cr (wt.%)
Mo (wt.%)
Bainite lath boundary 1 M3C o-Carbide 2 M3C 5 M7C3 M3C 10 M7C3 20 M7C3 50 M7C3 M23C6
66.8 49.1 66.8 35.9 51.8 34.3 30.8 43.2 35.2
24.5 36.3 22.2 52.9 33.9 51.4 53.4 50.2 54.8
8.7 14.6 11.0 11.1 14.4 14.3 15.8 6.7 10.0
Aging time (h)
Fe (wt.%)
Cr (wt.%)
Mo (wt.%)
60.2 38.3 66.8 37.8 49.6 2.0 36.9 34.7 36.1 1.6 37.8 2.3
28.4 54.0 22.2 52.1 36.0 12.6 53.9 54.4 53.6 13.2 56.0 11.7
11.4 7.8 11.0 10.1 14.4 85.3 9.2 10.8 10.3 85.2 6.1 86.0
Carbides
Within bainite lath 1 M3C M7C3 2 M3C M7C3 o-Carbide M2C 5 M7C3 10 M7C3 20 M7C3 M2C 50 M23C6 M2C
carbide and matrix, and the new carbide grows until the original carbide disappears. 2) ‘Separate nucleation’: The new alloy carbides nucleates separately at new sites primarily on dislocations, and the original carbide particles dissolve in the ferrite matrix.
3.3.1. Bainitic region The identification of carbide structures in the 20 kJ cm 1 heat input specimens during exposure at 700 8C have been carried out by analyzing the diffraction patterns and the chemical microanalyses from STEMEDX data. Fig. 9 shows the transmission electron micrograph of replica obtained from the specimen tempered at 700 8C for 1 h. It indicates that o-carbide and M3C were located on grain boundaries of bainitic ferrite. Besides, M7C3 particles were found within bainitic sub-units. Baker and Nutting [12] had proposed the sequences of carbide transformation in the tempered-bainite microstructure in the normalized 2.25Cr/ 1Mo steel as follows:
8 M C 0 M7 C 3 0 M 6 C > -carbide < 3 M3 C gM3 C 0 gM23 C6 0 M6 C > : M3 C M2 C Although Baker and Nutting did not detect M7C3 carbides until tempering at 700 8C for 5 h, M7C3 carbide particles have been observed in the simulated HAZ sample after tempering at 700 8C for 1 h. After tempering at 700 8C for 10 h, almost all carbides have transformed to M7C3 either on the bainitic sub-unit boundaries or within the sub-units. TEM micrograph of Fig. 10 shows that lump-shaped M7C3 carbide particles were located within bainitic subunits, and rod-shaped M7C3 carbide particles were discovered along bainitic sub-unit boundaries. The shape of M7C3 carbide appears to be rod-like and is very different from that of the M3C carbides (the circular-shape particles in sub-units and the film along the boundaries). Fig. 11 presents the TEM micrograph of replica after tempering at 700 8C for 20 h; it shows that tiny M2C particles were scattered within the bainitic sub-units. Since not any carbide can be found to be
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Fig. 13. (a) Transmission electron micrograph of a replica showing the microstructure of martensitic region obtained after tempering at 700 8C for 2 h; (b) diffraction pattern of o-carbide; (c) diffraction pattern of M3C.
adjacent to the M2C carbide, the mechanism of the M2C creation is suggested to be the type of ‘separate nucleation’ instead of the type of ‘in situ transformation’. The evolution of carbides in bainitic matrix obtained from the specimen tempered at 700 8C for 50 h is illustrated in Fig. 12; M23C6 particles were detected on the sub-unit boundaries and within the bainitic sub-units. The size of M23C6 is slightly smaller than that of M7C3 and the corresponding chemical compositions of M23C6 particles detected are listed with Fig. 12. The EDX data show that the Cr content in M23C6 carbide is lower than that in M7C3, and that Mo content in M23C6 carbide is higher vice versa. The sequences of carbide formations in bainitic region and
the chemical microanalyses for corresponding carbides are listed in Table 4. It clearly indicates that in the bainitic region in HAZ specimens o-carbide formed at initial stage and M23C6 carbide existed at the final stage. 3.3.2. Martensitic region Fig. 13 presents the replica micrograph obtained from the specimen tempered at 700 8C for 2 h; the M3C carbide particles were located on martensite lath boundaries and within martensite laths. o-carbide particles can also be detected when the specimen was tempered at 700 8C for 2 h. Fig. 14 illustrates that o-carbide transformed to M3C by a manner of in situ nucleation in the martensite region in the simulated HAZ specimen
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¯ reflection; (c) Fig. 14. Transmission electron micrograph of replicas showing: (a) bright field image; (b) dark field image using o-carbide (11¯ 1) ¯ reflection; (g) corresponding corresponding diffraction pattern; (d) interpretation of (c); (e) bright field image; (f) dark field image using M3C (220) diffraction pattern; (h) interpretation of (g).
after tempering at 700 8C for 2 h. Transmission electron micrographs of the replica show the bright and dark field images with the corresponding diffraction patterns for o-carbide (as indicated by position A in Fig. 14(a)) and M3C (as indicated by position B in Fig. 14(e)). The measured composition of o-carbide was 73.7Fe /
19.3Cr/7.0Mo (wt.%), and that of M3C was 63.7Fe / 28.0Cr/8.3Mo (wt.%). These two carbides contacted each other and their chemical compositions were only slightly different. It is suggested that as the M3C carbide grows, alloying elements are provided by the adjacent ocarbide, which gradually disappears.
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Fig. 15. (a) Transmission electron micrograph of a replica showing the microstructure of martensitic region obtained after tempering at 700 8C for 10 h; (b) diffraction pattern of M7C3.
After tempering for 10 h, the carbides almost transformed to M7C3 carbides (Fig. 15). The M2C and M23C6 carbides were discovered in the specimen tempered at 700 8C for 50 h (Fig. 16). It is suggested that the mechanisms of M2C and M23C6 nucleation are a type of ‘separate nucleation’. The sequences of carbide transformations in martensite region and the chemical microanalyses for the corresponding carbides are given in Table 5. Fig. 17 also present the details for M3C 0/M7C3 transformation in the way of ‘in situ transformation’ in the martensite region in the simulated HAZ specimen
after tempering at 700 8C for 10 h (M3C and M7C3 are indicated by positions C and D in Fig. 17, respectively); the measured composition of M3C was 53.6Fe /31.5Cr / 14.9Mo (wt.%), and that of M7C3 was 32.5Fe /48.8Cr / 18.7Mo (wt.%). These two carbides contacted each other and the concentrations of Cr and Mo were slightly higher in M7C3 than in M3C. It is suggested that as the M7C3 carbide grows, alloy elements are provided by the adjacent M3C carbide, which gradually disappears. From the above observation, it shows two kinds of mechanisms: in situ transformation and separate nucleation for carbide transformations in simulated HAZ
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Fig. 16. (a) Transmission electron micrograph of a replica showing the microstructure of martensitic region obtained after tempering at 700 8C for 50 h; (b) diffraction pattern of M7C3; (c) diffraction pattern of M23C6; (d) diffraction pattern of M2C.
samples in 2.25Cr /1Mo steel. The carbide transformation sequences in martensitic and bainitic regions are similar; they are proposed as follows:
3.4. The characteristic X-ray spectra of EDX Morphological differences can sometimes act as a guide in the identification of carbides in the steel studied. The carbide with a distinctive needle shaped
M23 C6 f-carbideg 0fM3 Cg 0fM7 C3 g 0 M2 C in situ
in situ
separate nucleation
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Table 5 The precipitation sequence of carbides in martensitc region Region
Aging time (h) 1
2
5
10
20
50
Martensite region Boundary In lath
M3C, o-carbide M3C, o-carbide
M3C, o-carbide M3C
M3C, M7C3 M7C3
M7C3 M7C3
M7C3 M7C3
M23C6 M7C3, M23C6, M2C
Aging time (h)
Carbides
Fe (wt.%)
Cr (wt.%)
Mo (wt.%)
M3C M3C M2C M7C3 M3C M7C3 M7C3 M2C M7C3 M23C6
64.7 67.5 14.0 37.5 54.1 33.5 37.9 27.8 35.1 35.7
26.7 24.7 20.8 50.7 31.5 53.0 53.0 9.7 55.8 51.5
8.6 7.9 65.2 11.8 14.3 13.4 9.1 62.5 9.0 12.8
Carbides
Fe (wt.%)
Cr (wt.%)
Mo (wt.%)
62.4 70.8 33.0 74.8 36.0 37.2 32.5 32.7 35.2 36.8 3.5
29.0 19.7 49.9 18.8 52.9 53.4 49.2 51.3 52.6 50.1 16.0
8.6 9.5 17.1 6.4 11.0 9.4 18.3 15.9 12.2 13.1 80.6
At boundary 1 2 5 10 20 50
Aging time (h)
Within martensite lath 1 M3C o-Carbide M7C3 2 M3C M7C3 5 M7C3 10 M7C3 20 M7C3 50 M7C3 M23C6 M2C
morphology is presumed to be M2C, but both M7C3 and M23C6 have the similar morphologies with the spherical form. Thus, it cannot be concluded that the shape of the particle alone is a guide to identification. Most of the carbide phases can certainly be identified using electron diffraction, but this is very difficult to examine large numbers of small particles. The carbides are often too thick to provide resolvable diffraction patterns. Another alternative method to identify the type of carbides is based on EDX microanalysis. This technique for carbide identification has been usually carried out in the TEM and permits chemical analyses to be taken from individual carbides. Although electron diffraction, microanalysis and morphological observation have been carried out to identify the carbides in this work, an attempt to use the standard EDS spectra for identifying the carbide has been done. The result is consistent with that reported by Hippsley [8] and Pilling and Ridley [13]. The typical EDS spectra taken from a range of carbides investigated are shown in Fig. 18. The significant peaks in Fig. 18 are for the peak energies of molybdenum (L), chromium
(Ka , Kb ) and iron (Ka , Kb ). The spectra were obtained from the various carbides: o-carbide, M3C, M7C3, M23C6 and M2C, respectively; the corresponding crystal structures were identified by analyses of electron diffraction patterns. These spectra can provide an access to identify the type of carbide in 2.25Cr /1Mo steel in this work.
4. Conclusions (1) The initial microstructure obtained from simulated coarse-grained HAZ with heat 20 kJ cm 1 input was lower bainite with martensite. (2) Bainite structure is more stable than martensite structure during tempering at elevated temperatures for long term exposures. (3) The sequences of carbide transformation were different at the regions on the bainite boundary, within bainitic sub-unit, on martensite lath boundary and within martensite lath. Their differences are suggested as follows:
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¯ reflection; (c) Fig. 17. Transmission electron micrograph of replicas showing: (a) bright field image; (b) dark field image using M3C (021) ¯ reflection; (g) corresponding corresponding diffraction pattern; (d) interpretation of (c); (e) bright field image; (f) dark field image using M7C3 (122) diffraction pattern; (h) interpretation of (g).
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Fig. 18. Typical energy-dispersive X-ray spectra taken from five different types of carbide detected.
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(a) Bainitic sub-unit boundary in situ
in situ
-carbide 0 M3 C 0
Acknowledgements
separate nucleation Fe3 C 0M7 C3 0 M23 C6 M7 C3
(b) Bainitic sub-unit
8 ( ( separate nucleation <-carbide M3 C in situ 0 M C3 0 M23 C6 M3 C 0 0 M7 C3 separate7nucleation M7 C3 : 0 M2 C M7 C3
This work was carried out with financial support from the National Science Council of the Republic of China, Taiwan, under Contract NSC 87-TPC-E-002-032. The authors are grateful for the support of this research.
(c) Martensite lath boundary
separate nucleation -carbidein situ M3 C 0M7 C3 0 M23 C6 0 M3 C M7 C3
(d) Martensite lath 8 > <
-carbidein situ
> :M 3 C
in situ
0 M3 C 0 M7 C3
8separate nucleation > 0 M23 C6 < > :
separate nucleation
0 0 M7 C3
M2 C
(4) The final equilibrium precipitate for 2.25Cr /1Mo steel studied is M23C6 carbide. M6C carbide was not encountered in the microstructural investigation. (5) The characteristic EDX spectra for five various carbides have been used to identify type of carbides in 2.25Cr /1Mo steel.
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