Structural changes in a ferritic heat-resistant steel after long-term service

Materials Science and Engineering, 62 (1984) 129-136

Structural

129

Changes in a Ferritic Heat-resistant

Steel After Long-term Service

R. A. VARIN* and J. HAFTEK¢

Institute of Materials Science and Engineering, The Warsaw Technical University, Warsaw (Poland) (Received April 29, 1983; in revised form June 8, 1983)

SUMMARY

The microstructures o f A S T M A335-P11 1Cr-O.5Mo steel specimens removed from tubes which operated in a crude oil reformer for 1.05 ×105 and 1.22X 105 h at about 520 °C were examined by transmission electron microscopy. The main structural elements, ferrite grains, post-pearlite grains and ferrite grain boundaries, were investigated. It is s h o w n that ferrite grains after 1.05 X 105 h in service contain mainly needle-like Mo2C precipitates and a very small a m o u n t o f Mz3C6 carbides. The a m o u n t of M23C6 carbides in ferrite grains is significantly higher after 1.22 × 105 h in service. A t ferrite grain boundaries after both service periods a large a m o u n t o f M7C3 and M23C6 carbides was found. A f t e r 1.22 X 105 h in service, M6C carbides were also observed at some ferrite grain boundaries. S o m e specimens were also normalized at 930 °C after 1.05 × 105 h service and then tempered at 725 °C. Transmission electron microscopy o f heat-treated steel revealed small spherical precipitates in ferrite grains instead o f needle-like m o l y b d e n u m carbides and the majority o f ferrite grain boundaries were free o f precipitates. The results o f microstructural investigations are discussed in terms o f the eventual changes in creep properties. 1. INTRODUCTION As pointed out by Bolton et al. [1], a question of growing concern is how long a high *Present address: Department of Mechanical Engineering, University of Waterloo, Waterloo, Ontario N2L 3G1, Canada. ¢Present address: Institute of Precision Mechanics, Duchnicka 3, 00-967 Warsaw, Poland. 0025-5416/84/$3.00

temperature component can be kept in service. Can a component be run safely beyond its design,life? Because most steels used for long-service components gradually change their microstructure throughout their service life [2-5] and as a consequence their creep properties are also changed [1, 6, 7], metallographic methods for assessing residual life are particularly attractive because they require only small samples and measure the damage micromechanisms directly [1]. In order, however, to derive such a method, a knowledge of structural changes taking place during longterm service must be obtained. Results showing the changes in the structure of components which have been in service for 105 h or longer are especially valuable. Unfortunately, information is scarce. Williams and Wilshire [6] investigated the changes occurring in the form, size and spacing of particles in 0.5Cr-0.5Mo-0.25V steel (where the composition of this and the other steels discussed in this paper are specified approximately in weight per cent) after 105 h at 565 °C. Coarse acicular Mo2C was observed in the matrix of the steel after service compared with an initial fine distribution of vanadium carbides. Very recently, Uehara and Takano [4] investigated by the replica m e t h o d and analytical electron microscopy the structural changes in a boiler tube 2.25Cr1Mo steel after 10 s h at 580 °C. They f o u n d M23C~ and Mo2C precipitates in the matrix. The M23C~ particles after long-term service had an increased chromium content compared with the chromium content in M23C6 in steel before use. In contrast, needle-like Mo2C precipitates in steel before use became granular and contained a greater amount of silicon after use. Abdel-Latif et al. [5] used accelerated aging to obtain a series of microstructures similar to those observed after long-term service. Specimens of 2.25Cr-lMo steel were Q Elsevier Sequoia/Printed in The Netherlands

130

aged at 630 °C for various periods to produce microstructures which simulated those obtained after 25 000, 50 000, 75 000 and 100 000 h at 540 °C. Abdel-Latif et al. concluded that the majority of carbides analysed by transmission electron microscopy and energy-dispersive X-ray techniques were either M23C6 or MeC. They found no trace of Mo2C. The present work is part of a study of the assessment of the remaining life of tubes after long-term service in the petrochemical industry. The main purpose was to obtain microstructural information which could be helpful in the assessment of remaining life. A detailed microstructural analysis by transmission electron microscopy of a tube of 1Cr-0.5Mo steel, whose composition and properties conformed to ASTM alloy A335P l l , after exposure for 1.05 X 105 and 1.22 X 105 h in crude oil reformers was carried out. Mechanical testing of the tubes as received, after exposure in service and after heat treatment was also performed.

2. EXPERIMENTAL DETAILS

Specimens were cut from tubes of crude oil reformers which had operated for 1.05 X 105 and 1.22 X 105 h. Under normal operating conditions the pressure inside the tubes is 4 MPa. The calculated hoop stress is a b o u t 50 MN m -2 in straight tubes and about 80 MN m -2 in bent portions. The average service temperature was a b o u t 520 °C (+20 °C). Some specimens were normalized at 930 °C (0.5 h) after service and then tempered at 725 °C for 0.5 h. The aim of this heat treatment was to determine whether the structure acquired during long-term service could be changed in such a manner as to obtain mechanical properties at least comparable with those of as-received material. Accelerated stress rupture tests were carried o u t to measure the differences between the mechanical behaviours of specimens with different structures. Thin foils prepared by double-jet polishing in a Struers-Tenupol apparatus were examined in an EM300 Philips electron microscope operating at 100 kV.

3. RESULTS

The structure of the steel was divided into three elements: post-pearlite grains, ferrite grains and ferrite grain boundaries. These structural elements were examined in each specimen. The structure of as-received tubes was ferrite-pearlite, characteristic of low carbon, low alloy steel. 3.1. Ferrite grains 3.1.1. Steel after 1.05 X 105 h service Figure 1 is an electron micrograph of the typical structure of a ferrite grain. Fine needle-like precipitates (about 0.5 #m in length) are clearly visible. The average distance between them is about 0.4 #m. The precipitates were identified by electron diffraction as Mo2C [8]. Those precipitates are often observed in similar types of ferritic heatresistant steels [4, 7]. Other precipitates, denoted by the arrows in Fig. 1, were also observed in some ferrite grains. Their density, however, was significantly lower than that of the M o 2 C precipitates. Because of their relatively small size, their crystallographic structure could not be identified by electron diffraction. The same t y p e of precipitate but with a larger size and a greater number density was observed after 1.22 X 105 h service; this precipitate was identified as mainly M23Ce carbides (see below).

Fig. 1. Electron micrograph of a typical microstructure of a ferrite grain in 1Cr-0.5Mo steel after 1.05 × 105 h at about 520 °C. The needle-like precipitates are Mo2C. The arrows s h o w other precipitates observed in s o m e grains.

131

Fig. 2. Electron micrograph of a typical microstructure of a ferrite grain in 1Cr-0.5Mo steel after 1.22 x 105 h. Needle-like Mo2C co-exists with the coarser M23C6 (denoted by the arrows).

Fig. 5. Post-pearlite grain in steel after 1.05 X105h service. There is a substructure with embedded Fe3C.

3.1.2. Steel after 1.22X 105 h service A typical s t r u c t u r e o f ferrite grains a f t e r this e x p o s u r e is s h o w n in Fig. 2. T h e coarser s p h e r o i d i z e d precipitates visible were identified as M2~C6 carbides (arrows}. B o t h the n u m b e r d e n s i t y and the size increased comp a r e d with the s t r u c t u r e after 1.05 X 105 h service. It m u s t be p o i n t e d out, h o w e v e r , t h a t in some grains o n l y M23C6 was observed and n o Mo2C particles were f o u n d (Fig. 3).

Fig. 3. Ferrite grain in steel after 1.22 x 105h service without M02C. Only M23C6 is visible. The b~ight triangle shows the position of the diffraction aperture.

3.1.3. Specimens normalized and tempered after 1.05 X 10 5 h service Needle-like m o l y b d e n u m carbides (Mo2C) were n o t seen in ferrite grains b u t instead v e r y small u n i d e n t i f i e d carbides were f o u n d (den o t e d b y the arrows in Fig. 4), v e r y o f t e n linked to dislocations. 3.2. Post-pearlite grains 3.2.1. Steel after 1.05 X 105 h service T h e s t r u c t u r e o f post-pearlite grains consisted o f s p h e r o i d i z e d precipitates e m b e d d e d in a well-developed s u b s t r u c t u r e (Fig. 5). T h e precipitates were identified b y selected area e l e c t r o n d i f f r a c t i o n as Fe3C. Selected area d i f f r a c t i o n p a t t e r n s t a k e n f r o m a few tens o f individual carbides did n o t reveal the presence o f a n y o t h e r carbides.

Fig. 4. Structure of ferrite grain in steel normalized and tempered after 1.05 x 105 h service. The small precipitates are linked to dislocations (arrowed).

3.2.2. Steel after 1.22 X 105 h service A s u b s t r u c t u r e with well-developed subgrains and a relatively low d e n s i t y o f precipitates was observed in post-pearlite areas (Fig. 6). T h e p r e c i p i t a t e m a r k e d b y the

132

Fig. 6. Post-pearlite grain in steel after 1.22 × 105 h service. The precipitate denoted by the arrow is M7C3.

length) precipitates. A large number of electron diffraction patterns were recorded and indexed to characterize the crystal structure of those precipitates. Three types of carbides with different shapes, examples of which are shown in Fig. 8, were found: M7C3 with a lamellar shape (Figs. 8(a) and 8(b)), allotriomorphic M23C6 (Figs. 8(c) and 8(d)) and triangular Fe3C (Fig. 8(e)). Some fine contrast effects resembling microtwins or fine stacking faults were frequently observed in the structure of carbides (denoted by arrows in Figs. 8(c) and 8(d)). An example of a carbide with a high density of very fine contrast is shown in Fig. 8(b). A selected area diffraction pattern taken from the carbide clearly shows elongated streaks perpendicular to the lines of fine contrast. The presence of streaks on the diffraction pattern confirms the existence of a plate-like structure in the carbide.

3.3.2. Steel after 1.22 X 105 h service As after 1.05 X 105h exposure, the majority of ferrite boundaries contained coarse precipitates. Some of the precipitates were identified as M6C carbides (Fig. 9). This type of carbide was not found in steel after 1.05 X 105 h service. Fig. 7. Post-pearlite grain in steel normalized and tempered after 1.05 × 105h service. There is a substructure with Fe3C.

arrow in Fig. 6 was found to be M~C3. This t y p e of carbide was not found after 1.05 × 105 h service.

3.2.3. Steel normalized and tempered after 1.05 X 105 h service There was no significant difference between the appearances of post-pearlite grains after heat treatment or after service. The structure consisted of a well-developed substructure with embedded spheroidized particles (Fig. 7), identified by electron diffraction as Fe3C. 3.3. Ferrite grain boundaries 3. 3.1. Steel after 1.05 X 105 h service The majority of ferrite grain boundaries contained relatively large (about 1 #m in

3.3.3. Steel normalized and tempered after 1.05 X 105 h service After normalizing and tempering, the great majority of ferrite grain boundaries were free of precipitates. Only occasional medium-sized precipitates (about 0.5/~m in length) were observed (Fig. 10). They were identified as Fe3C.

4. DISCUSSION

4.1. Structural changes As seen from the electron micrographs, pronounced changes in the microstructure t o o k place during service. The pearlite colonies observed in the asreceived tubes changed during long-term service into post-pearlite areas with spheroidized Fe3C particles embedded in fairly welldeveloped subgrains. The only difference between the structures of post-pearlite areas after 1.05 X 105 and after 1.22 X 105 h service is the existence of single MTC 3 carbides after

Fig. 8. E l e c t r o n m i c r o g r a p h s o f carbides observed at ferrite grain boundaries in steel after 1.05 X 1 0 5 h service: (a), (b) M7C3; (c), (d) M23C6; (e) Fe3C. Contrasts resembling m i c r o t w i n s or very fine stacking faults are d e n o t e d by the arrows in (c) and (d). An e x a m p l e of a carbide w i t h a very high density of fine contrast is ~hown in (b). The inset in (b) s h o w s the selected area diffraction pattern w i t h characteristic "streaks" perpendicular to the lines o f fine contrast.

==t

134

Fig. 9. M6C at a ferrite grain boundary after 1.22 × 105 h service.

Fig. 10. Fe3C at a ferrite grain boundary in steel normalized and tempered after 1.05 × 105 h service.

1.22 × 105 h. This observation may suggest that the transformation of Fe3C to M7C3 starts in former pearlite colonies after about 1.22 × 105 h service. The most important microstructural changes took place in ferrite grains and at ferrite grain boundaries. A high density of M o 2 C w a s present in ferrite grains after both 1.05 × 105 and 1.22 × 105 h exposure. These carbide particles are believed to be the major factor responsible for most of the creep strength of ferritic creepresistant steels [9]. A significant density of M23C 6 carbides co-existing with Mo2C was also found in ferrite grains after 1.22 × 105 h at temperature. This observation is in evident disagreement with the results reported by Abdel-Latif e t al. [5] who found that the majority of carbides in simulated structures were either M23C 6 or M6C. No evidence of Mo2C was provided.

According to Baker and Nutting [10] the precipitation sequence in ferrite during tempering of 2 . 2 5 C r - l M o steel is as follows: M2C -~ M6C. As suggested by Abdel-Latif et al. [5], M6C is an equilibrium carbide and as such it is expected that, after long-term exposure, equilibrium would be achieved with the formation of M6C carbides in ferrite grains. Hence, it could be concluded that in our steel the equilibrium conditions in ferrite grains have not been achieved even after 1.22 × 105 h in service because M6C carbides were not observed in ferrite. The question, however, arises why there is a discrepancy between the behaviour of the steel investigated by AbdelLatif e t al. [5] and that of our steel. The obvious explanation is the lower chromium and m o l y b d e n u m content of our steel. Also we did not observe any significant spheroidization of M o 2 C in steel after 1.22 × 105 h, contrary to the results of Uehara and Takano [4] for 2 . 2 5 C r - l M o steel after use for about 105 h. The sequence of carbide precipitation at ferrite grain boundaries during the long-term service of A335-P11 steel can be summarized as Fe3C ~ MTC 3 + M23C 6 (1.05 X 105 h) M7C3 + M23C6 + M6C (1.22 × 105 h). In this case the equilibrium M6C carbides nucleated at ferrite grain boundaries between 1.05 × 105 and 1.22 × 105 h service. Most probably, during the extension of service time significantly b e y o n d 1.22 × 105 h, M23C6 (M7C3) precipitates could be completely replaced by M6C carbides. The structural changes taking place in ferrite grains and at grain boundaries must influence to a certain extent the creep properties of the steel. It is generally assumed that interaction solid solution strengthening is responsible for most of the initial creep strength of normalized Cr-Mo creep-resistant steels [9]. During creep the contribution from interaction solid solution strengthening decays because of a decrease in molybdenum, chromium and carbon contents in solution because of the precipitation of Mo2C. Precipitation strengthening increases in magnitude as the density of the Mo2C precipitates increases [9]. Hence, it may be concluded that in the steel used up to 1.05 × 105 h the m o l y b d e n u m content in the ferrite matrix must be very small and therefore the principal strengthening mechanism up to 1.05 × 105 h service

135 results from the existence of Mo2C particles. Precipitation of M23C 6 carbides between 1.05 × 105 and 1.22 × 10 ~ h could slightly m o d i f y the creep strength because of the slightly different shape and size of M23C 6 particles co mp ar ed with the needle-like Mo2C. However, changes in creep strength due to M23C 6 precipitation may not be significant because M23C 6 particles are still rather small and the total precipitate density (M23C6 plus Mo2C) is still high. It may, therefore, be concluded that the time in service after which M23C 6 particles appear in ferrite grains is not a critical point indicating that steel cannot be used b e y o n d that limit. We believe that a more i m p o r t a n t factor for the estimation of residual life of A 3 3 5 - P l l steel is the appearance of M6C particles at ferrite grain boundaries, after 1.22 X 105 h at 520 °C. M6C is basically a molybdenum-rich carbide [11] and as such a significant n u m b e r o f m o l y b d e n u m atoms are necessary for its nucleation and growth. These may be provided by a continuous dissolution of M23C 6 particles already formed at ferrite grain boundaries, but only to a limited e x t e n t because the solubility of m o l y b d e n u m in M23C6 carbide is small [5, 11]. It seems, therefore, that the major part of the m o l y b d e n u m must be provided by the ferrite matrix as a result of dissolution o f Mo2C particles. Such a process would lead to a significant diminution in the creep strength. F r o m this point of view, the first appearance of M6C carbides at ferrite grain boundaries may be useful in the estimation o f the residual creep life of A 3 3 5 - P l l steel. We suggest that, after M6C appears at ferrite grain boundaries, the density and interparticle spacing of Mo2C in ferrite should

be m o n i t o r e d to determine whether this carbide is dissolving. The observation that some ferrite grains after 1.22 × 105 h service do not contain Mo2C particles m ay suggest that solution of Mo2C has already taken place there.

4.2.

Creep properties

As seen from Table 1, preliminary results of increased-stress tests clearly indicate that the time to rupture for heat-treated steel is significantly higher than that for steel after use for 1.05 × 105 h and for as-received material, for almost all test conditions. Only for T = 500 °C and o = 153 MN m -2 was the time to rupture of as-received material larger than that for heat-treated steel. After about 1000 h of creep testing, the specimens of heat-treated steel were still in the secondary creep stage. It may, therefore, be concluded that proper heat t r e a t m e n t of c o m p o n e n t s after long-term service can lead to significant i m provem ent in their creep properties and that " r e c o v e r e d " c o m p o n e n t s could be reused. The residual life of steel after 1.05 X 105 h service can be roughly estimated from the simple life fraction relationship [12] t~ --

tfs

t. +

--=

1

(1)

tfp

where ts is the time in service, t~s is the rupture time under service conditions, tp is the time to rupture in the accelerated test of steel removed from service and t~p is the rupt ure time of new material under accelerated test conditions. Using data from Table 1 we obtain for a test t e m p e r a t u r e of 538 °C the following residual lives: 15 366 h (for a = 173

TABLE 1 Test

conditions and results of increased-stress tests

o a (MN m-2)

173 133 194 153

T a (oC)

538 538 500 500

tf (h), as received

235 810 620 ~ 1500

tf (h), after 1.05 × 105 h at about 520 °C 30 210 274 800

t (h), normalized and tempered after 1.05 × 105 h service

868 966 819 966

tf, time to rupture; t, time after which the creep test was interrupted (at this time the heat-treated specimens were still in the secondary creep stage). aincreased_stress test conditions.

136 MN m -2) and 36 750 h (for o = 133 MN m-2). For a test temperature of 500 °C the residual life is estimated as 83 150 h (for o = 194 MN m-2). When the fact that the average service hoop stress is only 50 MN m -2 (80 MN m -2 in bent parts of the tube) is taken into account, it may safely be assumed that the residual life of the tube is equal to about 40 000 h.

5. CONCLUSIONS (1) Needle-like Mo2C particles and a very small a m o u n t of M23C6 carbides were observed in ferrite grains of A 3 3 5 - P l l steel after 1.05 × 105 h in service at about 520 °C. (2) Ferrite grains after 1.22 × 105 h service exhibited a large a m o u n t of M23C6 carbides co-existing with needle-like Mo2C particles. In some grains, only M23C6 precipitates were observed. (3) Mo2C was not observed in ferrite grains in steel heat treated (normalized at 930 °C; iempered at 725 °C) after 1.05 × 105 h in service. Instead, some small unidentified precipitates were found. (4) The sequence of carbide precipitation at ferrite grain boundaries of A335-P11 steel during long-term service can be summarized as follows: Fe.~C~ MvC3 + M23C6(1.05 × 105h) ~M7C3 + M23C6 + M6C (1.22 × 105 h). (5) The appearance of M6C at ferrite grain boundaries after 1.22 X 105 h service may be an important factor in residual life predictions because this carbide needs a large number of m o l y b d e n u m atoms and as such its nucleation and growth may lead to the dissolution of Mo2C in the ferrite matrix. (6) The time to rupture obtained in an increased-stress test of steel heat treated after 1.05 X 105 h service is significantly higher than that for as-received material, which indicates that by proper heat treatment the tubes after

long-term service may be "recovered" and reused. (7) It is roughly estimated from the life fraction relationship that the residual life of A 3 3 5 - P l l steel after 1.05 × 105 h at about 520 °C may be no less than 40 000 h.

ACKNOWLEDGMENTS The authors thank the Institute of Energetics, Pruszkow, and the Technical Supervision Office, Poznafi, for performing the increased-stress tests. The authors are also grateful to Mr. W. Zielinski for very valuable technical assistance. This work was supported by a grant from the Ministry of Science, Higher Education and Technology under Contract MR.I.21. REFERENCES 1 C.J. Bolton, B. F. Dyson and K. R. Williams, Mater. Sci. Eng., 46 (1980) 231. 2 D. Lonsdale and P. E. J. Flewitt, Mater. Sci. Eng., 39 (1979) 217.

3 J. K. L. Lai, D. J. Chastell and P. E. J. Flewitt, Mater. Sei. Eng., 49 (1981) 19. 4 K. Uehara and Y. Takano, JEOL News, 20E (1982) 2. 5 A.M. Abdel-Latif, J. M. Corbett and D. M. R. Taplin, Met. Sci., 16 (1982) 90. K. R. Williamsand B. Wilshire,Mater. Sci. Eng., 28 (1977) 289. 7 K. R. Williamsand B. J. Cane, Mater. Sci. Eng., 38 (1979) 199. 8 W. Zielinski, K. J. Kurzydlowski and M. W. Grabski, Proc. 6th Pol. Conf. on Electron Microscopy and the Solid State, Polish Academy of Sciences, Krynica, 1981, p. 195. 9 J. D. Baird, A. Jamieson, R. R. Preston and R. C. Cochrane, Proc. Conf. on Creep Strength in Steel and High Temperature Alloys, Sheffield, 1972,

Metals Society, London, 1973, p. 207. 10 R. G. Baker and J. Nutting, J. Iron Steel Inst., London, 192 (1959} 257. 11 K. Kuo, Acta Metall., 1 (1953) 301. 12 R. V. Hart, Met. Technol., 3 (1976) 1.