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A quantitative X-ray diffraction and analytical electron microscopy study of service-exposed 2.25Cr–1Mo steels
Materials Characterization 47 (2001) 17 – 26
A quantitative X-ray diffraction and analytical electron microscopy study of service-exposed 2.25Cr–1Mo steels D.R.G. Mitchell*, C.J. Ball Materials Division, ANSTO, Lucas Heights Science and Technology Centre, New Illawarra Road, Lucas Heights, PMB 1, Menai, NSW 2234, Australia Received 6 February 2001; accepted 11 April 2001
Abstract 2.25Cr – 1Mo power plant headers, with exposures spanning the as-installed condition through to 190,000 h, have been characterised with analytical transmission electron microscopy and quantitative X-ray diffraction. Microscopy revealed a number of service-induced changes including M2C development within pro-eutectoid ferrite, and its dissolution along prior austenite grain boundaries. Carbide coarsening and development of Mo- and Cr-rich carbide phases were also apparent in materials aged for long periods. Microscopy data showed carbide populations that evolved in accordance with the Baker and Nutting sequence, with M6C being the thermodynamically favoured end-point carbide in the bainitic regions, and M2C in the pro-eutectoid ferrite. X-ray diffraction showed that microscopy-based identification of carbides on the basis of composition alone was a valid approach, and good agreement between assessment of relative carbide populations was achieved with the two techniques. Trends in carbide populations as a function of service exposure derived using X-ray diffraction showed some scatter, and this probably stems from microstructural and compositional variations between the steels studied. D 2001 Elsevier Science Inc. All rights reserved.
1. Introduction Low alloy steels, such as 2.25Cr – 1Mo, find widespread use in the power generation industry [1,2]. The good scaling and creep resistance imparted by Cr and Mo additions, respectively [2], are essential for high temperature applications such as steam headers. These are thick-walled ( 0.1 m), large diameter ( 1 m) manifolds that collect steam from banks of boiler tubes. These structures are usually forged and welded, and postweld stress relief heat treatment prior
* Corresponding author. Tel.: +61-2-9717-3456; fax: +61-2-9543-7179. E-mail address: [email protected] (D.R.G. Mitchell).
to service is required. Such steels typically have a mixed microstructure made up of pro-eutectoid ferrite and bainite [2]. Typically the carbide makeup comprises acicular M2C in the pro-eutectoid ferrite and various phases in the bainite, which can include M7C3 and M23C6. M3C is lost very rapidly at normal postweld stress relief heat treatment temperatures [2 – 5]. This mixed bainitic/pro-eutectoid ferritic microstructure imparts good creep resistance in the short-term [2]. However, it is metastable, and changes occur with exposure to service temperature and the effects of creep [1,6]. Changes in microstructure are typically quite slow over service lives that may extend beyond 20 years at operating temperatures of 520 – 560C [7 – 11]. These microstructural changes modify the mechanical properties of the
1044-5803/01/$ – see front matter D 2001 Elsevier Science Inc. All rights reserved. PII: S 1 0 4 4 - 5 8 0 3 ( 0 1 ) 0 0 1 4 7 - 4
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ination during routine shutdowns. The compositions of the materials are shown in Table 1. Due to the large wall thickness ( 100 mm) it was possible to chain drill prismatic sections ( 40 10 10 mm) from the outer surface across longitudinal welds. TEM extraction replicas were produced by a single stage technique from specimens prepared for optical microscopy. Analytical transmission electron microscopy was performed on a JEOL 2000FXII transmission electron microscope fitted with an Oxford Instruments ISIS X-ray analytical system. This was calibrated to allow quantitative analysis of the elements of interest by using experimentally derived k factors from appropriate titanate reference materials, rather than theoretical values supplied with the system [13]. Carbides were extracted from these steels using the electrolytic dissolution technique described by Stevens and Lonsdale [17]. Areas were selected for dissolution by protecting surrounding areas with a nonconductive lacquer. Only the parent plate and weld could be investigated in this way, as the heataffected zone (HAZ) was too narrow to permit useful quantities of carbides to be extracted. After dissolution of 0.5 g of steel, the specimen was rinsed, then ultrasonically cleaned to liberate any adhered carbides. All washings were added to the extract. Centrifuging and washing of this carbide extract was repeated several times to permit separation of carbides from supernatant liquor. The extracted carbides were dispersed on a nonreflective substrate and X-ray diffraction patterns were obtained with a Scintag diffractometer, using Cu Ka radiation, over the range 10 < 2q < 80. The patterns were analysed using the program SIROQUANT to determine the proportions of the phases present.
steel [10,11]. One of these, creep resistance, is of particular concern in headers, since it is often the life-determining factor [1]. There have been many detailed studies of ageing in 2.25Cr – 1Mo [3,5,12 – 15]. However, these have invariably been carried out on material that has undergone accelerated ageing by exposure to temperatures much higher than those encountered typically in service. This is a necessary expedient, particularly in creep testing, in order to ensure realistic experimental time frames. Even so, such tests may run for many thousands of hours. One disadvantage of accelerated testing is that at higher temperature both the kinetics and thermodynamics of carbide evolution change [3]. The resulting microstructures may therefore not be representative of those encountered in service-exposed material. In this program of work we have examined seven in-service superheater outlet headers from various power stations. Detailed analytical TEM studies [16] have enabled the microstructure and carbide compositional changes due to service to be mapped. We have extended this work to include a quantitative X-ray diffraction study on carbide phases extracted from the headers. The aim has been to understand the evolution of carbide populations during ageing, and to identify any carbide parameters that may be useful indicators of service exposure, and may thus find application to remnant life assessment. Also of interest was to assess the accuracy of our carbide identification based on analytical TEM data.
2. Experimental Seven superheater outlet headers from five different power stations were made available for examTable 1 Composition (wt.%) and exposure of headers (h) Header A B C D E F G
Service exposure/h P W P W P W P W P W P W P W
0 70,200 75,800 114,000 119,000 119,000 190,000
P = parent plate, W = weld, N/A = not available.
Cr
Mo
C
P
Mn
Si
Ni
2.13 2.28 2.35 2.24 2.01 2.21 1.71 1.97 1.78 1.95 2.40 2.20 2.14 N/A
0.96 1.02 0.92 0.98 0.97 1.00 1.04 1.01 1.04 1.04 1.01 1.03 0.92
0.08 0.10 0.19 0.23 0.18 0.10 0.16 0.06 0.13 0.07 0.17 0.12 0.07
0.008 0.011 0.008 0.010 0.007 0.027 0.007 0.006 0.018 0.008 0.002 0.013 0.017
0.53 0.87 0.57 0.77 0.47 1.14 0.47 0.80 0.47 0.79 0.50 0.63 0.45
0.27 0.24 0.25 0.24 0.04 0.33 0.26 0.40 0.29 0.38 0.24 0.36 0.21
0.20 0.10 0.11 0.05 0.15 0.05 0.12 0.25 0.14 0.25 0.17 0.20 0.09
D.R.G. Mitchell, C.J. Ball / Materials Characterization 47 (2001) 17–26
3. Results 3.1. Parent plate Detailed analytical electron microscopy of these materials has been reported previously [16], and showed that the microstructure was unstable as a function of service exposure. Fig. 1 shows the parent plate microstructure in four headers in the virgin (asinstalled) condition, and those with service exposures of 75,800, 119,000 and 190,000 h (21.7 years). Carbides were selected at random for quantitative analysis, which permitted identification on the basis of composition [4,5,12,13,15,16]. Pro-eutectoid ferritic regions were characterised by a dispersion of acicular M2C, which coarsened with increasing service exposure. Prior austenite grain boundary carbides also developed, and alongside this in the pro-eutectoid ferrite a zone denuded in
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M2C formed. The width of this zone as a function of service exposure (Fig. 2) showed a dramatic increase after about 120,000 h exposure. In the bainite, the carbides were either globular or lath-like. Analytical electron microscopy showed these to be M23C6 and M7C3 at short exposures. However, at longer exposures some spheroidisation of the bainitic carbides occurred (cf. Fig. 1a,b and c,d), and at the longest exposures large faceted M6C carbides were found to dominate the microstructure (Fig. 1d). Quantitative EDS analysis of carbides permitted ternary plots to be constructed (Fig. 3a – f ). In the early stages of service, the M2C carbides were too small to analyse individually, and so clusters of 100 or so carbides were analysed and the average compositions are shown. The frequency of appearance of M2C carbides in the ternary plots, therefore, bears no relation to the frequency of occurrence in the steel. In the bainitic regions the major carbides present
Fig. 1. Parent plate microstructures of (a) virgin material — Header A — 0 h exposure, (b) Header C — 75,800 h, (c) Header F — 119,000 h and (d) Header G — 190,000 h. The microstructure is unstable at normal service exposure temperatures. Acicular M2C develops within the pro-eutectoid ferrite, while coarsening and transformation of bainitic carbides occurs, culminating in large populations of large faceted M6C. A zone denuded in carbides (DZ) develops along the prior austenite grain boundaries.
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Fig. 2. Width of M2C denuded zone in pro-eutectoid ferrite regions adjoining prior austenite grain boundaries in the parent plate as a function of service exposure.
included M7C3, M23C6 and M6C. These were all of similar dimensions, being much larger than the M2C needles in the pro-eutectoid ferrite (a few M2C carbides were also present in the bainite of short-term exposure materials, but disappeared with ageing). With the limited sampling feasible with quantitative analytical electron microscopy, there is probably a loose correlation between the frequency of occurrence of the larger bainitic carbides in the alloy and in the ternary plots. A qualitative assessment of the frequency of the various bainitic carbide types (ignoring M2C) was made from the ternary diagrams and the results are shown in (Table 2). This suggested that M7C3 and M23C6 dominated the early parent plate bainitic carbide population. With increasing exposure M23C6 diminished and M6C became the dominant carbide at very long exposures. One header (Header F, Fig. 3e) showed unusual behaviour both in terms of carbide compositions and carbide makeup. M23C6 was the almost exclusive bainitic carbide detected in this header despite its considerable service exposure (119,000 h). Comparable exposure headers [D (Fig. 3c) and E (Fig. 3d)], which were of very similar vintage and design and came from the same power station, showed M7C3 and M6C as the dominant phases (Tables 2 and 3). The results of the X-ray diffraction-based carbide measurements are shown in Table 3. Since the X-ray diffraction technique measured carbides from both bainite and pro-eutectoid ferrite, in many cases a significant fraction of M2C was detected. This was the exclusive intraferritic carbide phase (Fig. 1a – d). In order to compare the X-ray diffraction data with those derived from analytical electron microscopy carbide assessments [which were based only on larger ( 0.1 – 1 mm) carbides from bainitic regions — i.e., the fine M2C in the pro-eutectoid ferrite was ignored],
the M7C3, M6C and M23C6 carbide fractions determined with X-ray diffraction were ranked with M2C omitted. Comparison of the data from the two methods (Table 2) showed that composition-based classification of carbides using analytical electron microscopy produced good agreement with the more accurate X-ray diffraction method. In nearly all cases analytical electron microscopy confirmed the presence of carbides identified using X-ray diffraction. The only exception was Header E (Fig. 3d), where no M23C6 was located with analytical electron microscopy, but where X-ray diffraction showed that nearly a third of the carbide was this phase (Table 2). The analytical electron microscopy technique is highly sensitive to any microstructural inhomogeneity in the steel, owing to the very tiny volumes sampled. The sampling statistics for the X-ray diffraction technique are significantly better. In general, the consistency between the ranking derived from the two data sets improved for the more aged materials. Agreement as to the most dominant carbide phase determined by X-ray diffraction and analytical electron microscopy was achieved for Headers B – G, while for A the two most dominant carbide phases were correctly identified, albeit with a different ranking of frequency of occurrence. In some instances X-ray diffraction did not detect minority phases that were located by analytical electron microscopy (e.g., M6C in Headers C and F). This was simply due to the low sensitivity of the X-ray diffraction technique. X-ray diffraction (Table 3) confirmed the analytical electron microscopy result, which showed that Header F was unusual compared with its sister headers (D and E) in that it contained only M23C6 despite being quite heavily aged (Fig. 3e). The X-ray diffraction data for the various carbides in the parent plate is plotted as a function of service exposure in Fig. 4a,b. All the data showed considerable scatter, though trends did emerge. The percentage of M2C generally diminished with service exposure (Fig. 4a), falling from a peak of 50% down to 15% in the most heavily aged material. Header F showed the lowest M2C levels ( 5%). The percentage of M6C showed a general increase from 0% in the virgin material to 56% in the most heavily aged steel. However, three aged headers contained very little M6C; Header C (75,800 h) contained no M6C, while a comparable exposure header (B — 70,200 h) contained 15%. This erratic behaviour was also shown by Headers D (114,000 h — 21% M6C) and E (119,000 h — 2.4% M6C). The X-ray diffraction results for M7C3 and M23C6 are shown in Fig. 4b. Data for Header F was again a long way off the trends for M7C3 and M23C6, with the latter carbide being almost the exclusive parent plate carbide
D.R.G. Mitchell, C.J. Ball / Materials Characterization 47 (2001) 17–26
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Fig. 3. Ternary composition plots for carbides found in the parent plates of the headers. Typical carbide compositions from the data of Moss and Kelly [20] ( ) and Titchmarsh [4] (~) are also shown: (a) A = 0 h, (b) C = 75,800 h, (c) D = 114,000 h, (d) E = 119,000 h, (e) F = 119,000 h, and (f) G = 190,000 h.
(95%). The M7C3 fraction changed little with service exposure, while after 75,000 h service, M23C6
declined from up to 42% to zero in the most heavily aged material.
Table 2 Carbide frequency in bainitic regions based on random AEM and XRD analysis — M2C omitted Header
Parent (AEM)
Parent (XRD)
Weld (AEM)
Weld (XRD)
A B C D E F G
M7C3 > M23C6 M7C3>M6C>M23C6 M7C3 M23C6>M6C M7C3 >> M6C>M23C6 M7C3 >> M6C M23C6 >> M6C M6C>M7C3
M23C6>M7C3 M7C3>M6C M23C6>M7C3 M7C3 >> M6C>M23C6 M7C3 >> M23C6 >> (M6C) M23C6 M6C >> M7C3
M7C3 M6C>M7C3 M6C M6C >> M7C3 M6C >> M7C3 M6C >> M7C3 N/A
M7C3 >> (M23C6) M7C3>M6C M6C >> (M7C3) M6C >> (M7C3) M6C >> (M7C3) N/A N/A
Carbides shown in parentheses represent < 15% of the total, N/A = not available.
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Table 3 Quantitative XRD results for carbides extracted from service-exposed 2.25Cr – 1Mo (P = parent, W = weld) Header A B C D E F G
P W P W P W P W P W P W P
M23C6
M6C
M2C
M7C3
119,000
31.2 12.2 0 21.8 42.2 0 17.2 0 20.1 0 94.4
0 0 15.2 42 0 92.9 21.4 92.5 2.4 89.7 0
44.0 18.8 49.6 14.7 32.2 0 15.7 0 39.0 0 5.6
24.9 69.0 35.2 21.5 25.6 7.1 45.7 7.5 38.5 10.3 0
190,000
0
56.0
14.6
29.3
Exposure/h 0 70,200 75,800 114,000 119,000
with quantitative X-ray diffraction due to the volumes of available material, HAZ data is not presented here
3.2. Welds Analytical electron microscopy characterisation was also performed on the HAZs and welds. Since only the parent and weld regions could be analysed
Fig. 4. Percentiles of (a) M2C and M6C and (b) M7C3 and M23C6 in the parent plate as a function of service exposure, determined using X-ray diffraction. In (a) the M6C increases while the M2C decreases with increasing exposure, albeit with considerable scatter, while in (b) M23C6 diminishes with service exposure while M7C3 changes little.
Fig. 5. TEM micrographs of the weld microstructures of headers after service exposures of (a) 0, (b) 75,800 and (c) 119,000 h (Header F). Spherical Mn/Si-rich weld inclusions (WI) 1 mm in diameter are evident. With increasing exposure fine scale M2C disappears, while spheroidisation and coalescence of larger M6C occurs.
D.R.G. Mitchell, C.J. Ball / Materials Characterization 47 (2001) 17–26
(see Ref. [16] for analytical electron microscopy details). Fig. 5a – c shows a selection of weld microstructures from headers exposed for 0 (Fig. 5a), 75,800 (Fig. 5b), and 119,000 h (Fig. 5c). All welds were quite similar, comprising a coarse-grained fully bainitic microstructure. As well as carbides, rounded weld inclusions of 1 mm diameter were also present. These were typically rich in manganese and silicon. Analytical electron microscopy showed that the weld carbides were M6C and/or M7C3 with the former dominating in the longer aged materials (Table 2). No M23C6 was found, however very small rod-like M2C carbides were present in the virgin header and those exposed for 75,800 h (indicated in Fig. 5a). In materials aged for 114,000 h or more there were no very fine M 2C particles present (Fig. 5c). Ageing during service caused some coarsening of the larger weld carbides, with submicron particles being consumed by the larger carbides (cf. Fig. 5a,b and c). Qualitative ranking of the preponderance of various carbides in the welds (Table 2) showed that service exposure induced the formation of M6C quite early in the service life, but thereafter little further change occurred. Comparison of the carbide type and frequency of occurrence (M2C omitted) in the welds as determined by X-ray diffraction and analytical electron microscopy (Table 2) showed quite good agreement. In four out of five cases, analytical electron microscopy correctly predicted the carbide types and their frequency of occurrence when compared with the more accurate X-ray diffraction results. In Header A, a small amount of M23C6 (17% of the total with M2C excluded) was detected by X-ray diffraction but not by analytical electron microscopy. However, M7C3 was shown to be the dominant phase by both techniques in this material. X-ray diffraction analysis of the weld of Header F was not possible due to insufficient material being available. The plot of the percentage of M2C as a function of service exposure determined with X-ray diffraction (Fig. 6a) showed that up to 70,000 h M2C comprised up to 20% of the carbide, but thereafter it declined to zero. A concomitant and dramatic increase in M6C was found (Fig. 6a) with all welds with over 75,800 h exposure containing 90% or more of this carbide. The corresponding plots for M23C6 and M7C3 (Fig. 6b) showed that low levels of M23C6 were present in the virgin header weld, but diminished to zero after 75,800 h of service exposure. A significant decrease in M7C3 levels occurred, which mirrored the increase of M6C. Initial M7C3 levels of 70% fell to 10% after 75,800 h exposure. These trends indicated that in the welds M6C grew at the expense of M2C and M7C3, and that M23C6 was rapidly lost upon service exposure.
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Fig. 6. Percentiles of (a) M2C and M6C and (b) M7C3 and M23C6 in the welds as a function of service exposure, determined using X-ray diffraction. Upon ageing, M6C rapidly dominates the carbide population by completely replacing M23C6 and M2C. Only low levels ( 5 wt.%) of M7C3 persist after 75,000 h of service exposure.
3.3. Supernatant liquor analysis and lattice parameter measurements The supernatant liquor from the carbide extractions was analysed for the principal alloying elements to determine the proportion of various solutes present in the matrix. The carbide compositional data from analytical electron microscopy was combined with the carbide fraction data from X-ray diffraction to determine the proportion of various solutes present in the carbide. Combining these two proportions and dividing by the known solute concentration in the alloy should yield a value of unity for each solute providing a check for experimental error. In most cases the values for Mo and Cr were between 0.9 and 1.1, which was quite reasonable considering the number of steps required to obtain the final result. However, in a few cases, errors of up to 50% were found. Plots of various solute fractions in either the matrix or carbide as a function of service exposure (not shown) yielded no simple trends. No trends in the carbide lattice parameters as a function of service exposure were found.
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4. Discussion The microstructural instability of 2.25Cr – 1Mo upon ageing can lead to a general deterioration of mechanical properties [1,6,11]. The usual method for creep testing involves use of temperatures higher than those used in service, to accelerate the creep rate. Microstructural changes during this accelerated testing have been widely reported [13,18]. The drawback with such data is that the kinetics and thermodynamics of carbide change at a testing temperature of say 700C, may be quite different to those at the service temperature of 540C. One of the major advantages of the present study is that it has been carried out on service-exposed material, with exposures of up to 22 years. The data is therefore directly applicable to other service-exposed materials. In general there is a paucity of data derived from actual service-exposed material owing to its commercial sensitivity. One disadvantage of a study such as this is that it compares steels with subtly differing compositions, different thermal processing histories, as well as operating regimes that may differ. Meaningful test certificate and operational data for these headers was completely absent. This is not at all uncommon, and has been identified as a major limiting factor in the assessment of operational plant [19]. The carbide compositional parameters derived for these steels previously [16] indicated some quite strong compositional trends as a function of service exposure. This suggested that a comparative study of the various header steels was worthwhile, even with the various experimental uncertainties. In the analytical electron microscopy work, carbides were identified on the basis of composition alone. Several authors [4,13] have performed both electron diffraction and analytical electron microscopy studies on carbides, and have shown that the various carbide phases in low alloy steels have discrete compositional ranges, which permits ready identification. Ternary diagrams (e.g., Fig. 3a – f ) are a convenient way of showing the clustering of compositions for different carbide types. They are also useful for highlighting any particles with anomalous compositions. These most likely stem from superposition of different carbides as a result of nucleation and growth of new carbides on preexisting carbides. Ternary plots also show any compositional clustering or spreading upon ageing, as well as any general shifts in the carbide population. All the analytical electron microscopy work performed here was done quantitatively, ostensibly to determine mean carbide compositions. Typically, 20 – 30 carbides were analysed. This is a small sample for assessing the carbide population, and is subject to considerable scatter. This scatter is also exacerbated
by the fact that although there does not appear to be a dependence of carbide composition on size (within the range of carbide sizes studied here), relative abundance of carbides do vary in relation to the prior austenite grain boundaries [16]. Despite these limitations, the analytical electron microscopy and X-ray diffraction data showed excellent agreement with regard to the types and relative amounts of carbide present (Table 2). The X-ray diffraction data also confirmed the unusual behaviour of Header F. There appears to have been few quantitative X-ray diffraction studies of either accelerated aged [17] or ex-service 2.25Cr – 1Mo [10]. The precipitation sequence in 2.25Cr – 1Mo was first determined by Baker and Nutting [3], and has since been confirmed by several others [5,12 – 15]. Upon cooling from the austenitising temperature, Fe-rich M3C is the kinetically favoured carbide in the bainite of normalised steels [2]. M2C nucleates independently within the pro-eutectoid ferrite. Subsequent tempering causes the bainitic M3C to transform to more thermodynamically favoured carbides, rich in Cr and Mo. These include M7C3, M23C6 and M6C, which nucleate and grow on preexisting bainitic carbides. M2C in the pro-eutectoid ferrite has been reported to transform to M6C [3,21]. However, it is clearly very resistant to service-induced ageing, as it has been shown to persist in steels with exposures of 190,000 h in the present work (Fig. 1d) as well as in other studies of long-term aged steels [7,9]. Titchmarsh [4] has shown that tempering 2.25Cr – 1Mo for 10 h at 700C, conditions typically used to postweld heat-treat 2.25Cr – 1Mo, resulted in all the M3C being converted to alloy carbides. All the present materials had been postweld stress relief heat-treated, and therefore M3C had transformed to M23C6 and M7C3 before entering service (Fig. 3a), Table 3. Subsequent service-induced ageing results in the nucleation of M6C on existing carbides that are then rapidly consumed [3]. This sequence of events was reflected in the parent plate X-ray diffraction data (Fig. 4a,b), where the principal changes were loss of M23C6 and M2C while M7C3 levels changed little, but M6C increased considerably. The M6C data were interesting in that not all heavily aged materials developed M6C to the same degree (Table 3, Fig. 4a). For instance, in Header B (70,200 h) 15 wt.% of the carbides were M6C, while in Header C (75,800 h) there was none (Fig. 4a). A similar pattern emerged for Headers D and E/F. This may simply be a reflection of different nucleation kinetics for the M6C phase, although there are no obvious, microstructural or compositional factors that might account for this. No simple trends were present in the parent plate X-ray diffraction data as a function of service exposure. The reasons are most likely related to
D.R.G. Mitchell, C.J. Ball / Materials Characterization 47 (2001) 17–26
differences in parent plate composition, combined with differing fractions of bainite, and differing relative fractions of prior austenite grain boundary within the material (the ASTM grain sizes for the ferrite and bainite were quite similar for the respective materials, although these are of course average values). Since the M 2 C is primarily restricted to the pro-eutectoid ferrite, it will be very sensitive to the proportion of this phase in the steel, which in turn will depend on the alloy carbon content Attempts to derive correlations between the fraction of various carbides, such as M2C, and compositional parameters, such as the carbon levels, were not successful. X-ray diffraction data showed that M2C diminished with increasing service exposure (Fig. 4a), even though TEM showed this acicular carbide elongated with time at operating temperature (cf. Fig. 1a and d). This apparent contradiction can be explained by the fact that M2C was found to disappear rapidly from within bainitic regions upon ageing, due to reaction with other bainitic carbides. More importantly, considerable M2C dissolution from the near-prior austenite grain boundary zones occurred due to growth of alloy carbides (especially Mo-rich M6C) along the boundaries (Figs. 1a – d and 2). TEM of the welds showed a fully bainitic microstructure (Fig. 5a – c). The carbide makeup appeared more advanced in terms of thermal ageing than in the respective parent plate, which was to be expected owing to the additional tempering effects of multipass welding, prior to postweld heat treatment. Analytical electron microscopy data (Table 2) suggested an initial population of only M7C3, but this gave way to a population of >90% M6C (the most thermodynamically stable carbide) after the first 75,800 h of service. X-ray diffraction data was in very good agreement with analytical electron microscopy data (Table 2). The development of M6C upon tempering resulted in a loss of the fine scale (sub-50 nm) M2C carbide (cf. Fig. 5a and c). This dissolution of M2C as a result of the development of M6C was also observed in the parent plate near prior austenite grain boundary regions (Fig. 1d). M6C is more thermodynamically stable than M2C and has been shown to develop at its expense [3,21]. The growth of M6C and the loss of M2C in the weld were also observed with X-ray diffraction (Fig. 6a). Also noted were the complete loss of M23C6 after 75,800 h of service exposure (Fig. 6b), and a dramatic decrease in M7C3. The weld X-ray diffraction data showed much higher M6C levels at a given service exposure compare with parent plate (cf. Figs. 4a and 6a), again reflecting the greater effective ageing within the weld, due to the heat treatment effects of multipass welding.
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In general, the X-ray diffraction population data in the welds showed smoother trends compared with those of the parent plate (cf. Figs. 4 and 6). A lower level of scatter was also apparent in the mean compositional data of the weld carbides compared with the parent plate [16]. The reason for this is probably due to the dominance of a single carbide species in the virgin weld material ( 70% M7C3 — Table 3). The corresponding parent plate contained 31% M23C6, 44% M2C and 25% M7C3. This is a more diverse carbide environment within which a wider range of carbide precipitation, growth and dissolution reactions can occur.
5. Conclusions Analytical electron microscopy shows that 2.25Cr – 1Mo service-exposed for up to 190,000 h is microstructurally unstable. Within the parent plate, extensive development of acicular M2C occurs within the pro-eutectoid ferrite, while along prior austenite grain boundaries, growth of alloy carbides leads to the development of zones denuded in M2C. Within bainitic regions, some spheroidisation and carbide transformation reactions occur. Early carbide populations dominated by M7C3 and M23C6, give way to the more thermodynamically stable M6C upon ageing. Carbide identification and evolution has been studied by quantitative analytical electron microscopy, with ternary compositional diagrams offering a convenient way of sorting and classifying carbides. Within the fully bainitic welds, microstructural changes with service exposure were more subtle, predominantly involving loss of fine scale M2C. Early carbide populations were dominated by M7C3, which transformed to M6C upon ageing. Comparative quantitative X-ray diffraction analysis of extracted carbides showed that analytical electron microscopy correctly identified carbides on the basis of composition alone. Despite the poor counting statistics of the analytical electron microscopy method, qualitative assessment of carbide populations was generally in good agreement with that determined by the X-ray diffraction technique. Carbide evolution was in accordance with the Baker and Nutting sequence. Attempts to derive correlations between the weight fraction of individual carbides and service exposure showed general trends, but no close correlation was found, precluding any potential application for remnant life assessment. This lack of correlation was probably due to the combined effects of variables such as carbon content/fraction of proeutectoid ferrite/bainite, grain size/percentage of material in or near grain boundaries, composition and heat treatment.
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Acknowledgments The authors are indebted to Pacific Power International for the provision of materials and for funding part of this work, and to G. Smith and A. Croker of ANSTO for technical assistance and advice.
References [1] Bhadeshia HKDH, Strang A, Gooch DJ. Ferritic power plant steels: remanent life assessment and approach to equilibrium. Int Mater Rev 1998;43:45 – 69. [2] Lundlin CD, Kruse BJ, Pendley MR. High temperature properties of 2.25Cr – 1Mo weld metal. Weld Res Counc BullMay 1982;277:1 – 27. [3] Baker RG, Nutting J. The tempering of 2.25Cr – 1Mo steel after quenching and normalising. J Iron Steel Inst 1959;192:257 – 68. [4] Titchmarsh JM. The identification of second-phase particles in steels using an analytical transmission electron microscope. UKAE Harwell Report AERER-9661, December 1979, pp. 1 – 12. [5] Pilling J, Ridley N. Tempering of 2.25Pct Cr – 1 Pct Mo low carbon steels. Metall Trans 1982;13A:557 – 63. [6] Masuyama F, Nishimura N, Takeda Y. Carbide changes and sulphur segregation in CrMo steel during creep. High Temp Technol 1990;8:257 – 60 (November). [7] Yang JR, Huang CY, Yang CN, Horng JL. Microstructural examination of 2.25Cr – 1Mo steel steam pipes after extended service. Mater Charact 1993;30:75 – 88. [8] Benvenuti A, Ricci N, Fedeli G. Evaluation of microstructural parameters for the characterization of 2 1/4Cr 1Mo steel operating at elevated temperature. . Proc. Fourth Int. Conf. on Creep and Fracture of Engineering Materials and Structures, Swansea, UK (April 1 – 6). [9] Pigrova GD, Sedov VM, Archakov YuI. Carbide transformations in Cr – Mo steels in the process of long-term aging and operation. Metal Sci Heat Treat 1997;39: 371 – 5.
[10] Wada T, Biss VA. Restoration of elevated temperature tensile strength in 2.25Cr1Mo steel. Metall Trans 1983;14A:845 – 55. [11] Lin C-A, Horng J-L. Degradation phenomena of Cr – Mo steels after long-term high temperature service. MRI Bull Res Dev 1990;4:59 – 65. [12] Hippsley CA. Precipitation sequences in heat-affected zone of 2.25Cr – 1Mo steel during stress-relief heat treatment. Met Sci 1981;15:137 – 47. [13] Abdel-Latif AM, Corbett JM, Taplin DMR. Analysis of carbides formed during accelerated aging of 2.25Cr – 1Mo steel. Met Sci 1982;16:90 – 6. [14] Thomson RC, Bhadeshia HKDH. Changes in chemical composition of carbides in 2.25Cr – 1Mo power plant steel: Part 1. Bainitic microstructure. Mater Sci Technol 1994;10:193 – 203. [15] Shaw BJ. A study of carbides formed in low alloy Cr – Mo steels. In: Swift RA, editor. Research on chrome – moly steels — MPC-21. New York: ASME, 1984. pp. 117 – 28. [16] Mitchell DRG, Small R. Microstructural evolution in seven 2.25Cr – 1Mo superheater outlet headers due to service exposure. Sci Technol Weld Joining 2001;6: 1 – 9 (March). [17] Stevens RA, Lonsdale D. Isolation, identification and quantification by X-ray diffraction of carbide phases in 2.25Cr – 1Mo steel. J Mater Sci 1985;20:3631 – 8. [18] Lonsdale D, Flewitt PEJ. Damage accumulation and microstructural changes occurring during the creep of a 2.25Cr – 1Mo steel. Mater Sci Eng 1979;39: 217 – 29. [19] Evans RW, Williams KR, Wilshire B. A microstructural approach to remanent life assessment. Proc. Fourth Int. Conf. on Creep and Fracture of Engineering Materials and Structures, Swansea, UK (April 1 – 6). [20] Moss CJ, Kelly PM. The mechanisms and detection of embrittlement in Cr – Mo pressure vessel steels. Fatigue Fract Eng Mater Struct 1994;17:369 – 80. [21] Kurzydlowski KJ, Zielinski W. Mo2C ! M6C carbide transformation in low alloy Cr – Mo ferritic steels. Met Sci 1984;18:223 – 4.
A quantitative X-ray diffraction and analytical electron microscopy study of service-exposed 2.25Cr–1Mo steels D.R.G. Mitchell*, C.J. Ball Materials Division, ANSTO, Lucas Heights Science and Technology Centre, New Illawarra Road, Lucas Heights, PMB 1, Menai, NSW 2234, Australia Received 6 February 2001; accepted 11 April 2001
Abstract 2.25Cr – 1Mo power plant headers, with exposures spanning the as-installed condition through to 190,000 h, have been characterised with analytical transmission electron microscopy and quantitative X-ray diffraction. Microscopy revealed a number of service-induced changes including M2C development within pro-eutectoid ferrite, and its dissolution along prior austenite grain boundaries. Carbide coarsening and development of Mo- and Cr-rich carbide phases were also apparent in materials aged for long periods. Microscopy data showed carbide populations that evolved in accordance with the Baker and Nutting sequence, with M6C being the thermodynamically favoured end-point carbide in the bainitic regions, and M2C in the pro-eutectoid ferrite. X-ray diffraction showed that microscopy-based identification of carbides on the basis of composition alone was a valid approach, and good agreement between assessment of relative carbide populations was achieved with the two techniques. Trends in carbide populations as a function of service exposure derived using X-ray diffraction showed some scatter, and this probably stems from microstructural and compositional variations between the steels studied. D 2001 Elsevier Science Inc. All rights reserved.
1. Introduction Low alloy steels, such as 2.25Cr – 1Mo, find widespread use in the power generation industry [1,2]. The good scaling and creep resistance imparted by Cr and Mo additions, respectively [2], are essential for high temperature applications such as steam headers. These are thick-walled ( 0.1 m), large diameter ( 1 m) manifolds that collect steam from banks of boiler tubes. These structures are usually forged and welded, and postweld stress relief heat treatment prior
* Corresponding author. Tel.: +61-2-9717-3456; fax: +61-2-9543-7179. E-mail address: [email protected] (D.R.G. Mitchell).
to service is required. Such steels typically have a mixed microstructure made up of pro-eutectoid ferrite and bainite [2]. Typically the carbide makeup comprises acicular M2C in the pro-eutectoid ferrite and various phases in the bainite, which can include M7C3 and M23C6. M3C is lost very rapidly at normal postweld stress relief heat treatment temperatures [2 – 5]. This mixed bainitic/pro-eutectoid ferritic microstructure imparts good creep resistance in the short-term [2]. However, it is metastable, and changes occur with exposure to service temperature and the effects of creep [1,6]. Changes in microstructure are typically quite slow over service lives that may extend beyond 20 years at operating temperatures of 520 – 560C [7 – 11]. These microstructural changes modify the mechanical properties of the
1044-5803/01/$ – see front matter D 2001 Elsevier Science Inc. All rights reserved. PII: S 1 0 4 4 - 5 8 0 3 ( 0 1 ) 0 0 1 4 7 - 4
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ination during routine shutdowns. The compositions of the materials are shown in Table 1. Due to the large wall thickness ( 100 mm) it was possible to chain drill prismatic sections ( 40 10 10 mm) from the outer surface across longitudinal welds. TEM extraction replicas were produced by a single stage technique from specimens prepared for optical microscopy. Analytical transmission electron microscopy was performed on a JEOL 2000FXII transmission electron microscope fitted with an Oxford Instruments ISIS X-ray analytical system. This was calibrated to allow quantitative analysis of the elements of interest by using experimentally derived k factors from appropriate titanate reference materials, rather than theoretical values supplied with the system [13]. Carbides were extracted from these steels using the electrolytic dissolution technique described by Stevens and Lonsdale [17]. Areas were selected for dissolution by protecting surrounding areas with a nonconductive lacquer. Only the parent plate and weld could be investigated in this way, as the heataffected zone (HAZ) was too narrow to permit useful quantities of carbides to be extracted. After dissolution of 0.5 g of steel, the specimen was rinsed, then ultrasonically cleaned to liberate any adhered carbides. All washings were added to the extract. Centrifuging and washing of this carbide extract was repeated several times to permit separation of carbides from supernatant liquor. The extracted carbides were dispersed on a nonreflective substrate and X-ray diffraction patterns were obtained with a Scintag diffractometer, using Cu Ka radiation, over the range 10 < 2q < 80. The patterns were analysed using the program SIROQUANT to determine the proportions of the phases present.
steel [10,11]. One of these, creep resistance, is of particular concern in headers, since it is often the life-determining factor [1]. There have been many detailed studies of ageing in 2.25Cr – 1Mo [3,5,12 – 15]. However, these have invariably been carried out on material that has undergone accelerated ageing by exposure to temperatures much higher than those encountered typically in service. This is a necessary expedient, particularly in creep testing, in order to ensure realistic experimental time frames. Even so, such tests may run for many thousands of hours. One disadvantage of accelerated testing is that at higher temperature both the kinetics and thermodynamics of carbide evolution change [3]. The resulting microstructures may therefore not be representative of those encountered in service-exposed material. In this program of work we have examined seven in-service superheater outlet headers from various power stations. Detailed analytical TEM studies [16] have enabled the microstructure and carbide compositional changes due to service to be mapped. We have extended this work to include a quantitative X-ray diffraction study on carbide phases extracted from the headers. The aim has been to understand the evolution of carbide populations during ageing, and to identify any carbide parameters that may be useful indicators of service exposure, and may thus find application to remnant life assessment. Also of interest was to assess the accuracy of our carbide identification based on analytical TEM data.
2. Experimental Seven superheater outlet headers from five different power stations were made available for examTable 1 Composition (wt.%) and exposure of headers (h) Header A B C D E F G
Service exposure/h P W P W P W P W P W P W P W
0 70,200 75,800 114,000 119,000 119,000 190,000
P = parent plate, W = weld, N/A = not available.
Cr
Mo
C
P
Mn
Si
Ni
2.13 2.28 2.35 2.24 2.01 2.21 1.71 1.97 1.78 1.95 2.40 2.20 2.14 N/A
0.96 1.02 0.92 0.98 0.97 1.00 1.04 1.01 1.04 1.04 1.01 1.03 0.92
0.08 0.10 0.19 0.23 0.18 0.10 0.16 0.06 0.13 0.07 0.17 0.12 0.07
0.008 0.011 0.008 0.010 0.007 0.027 0.007 0.006 0.018 0.008 0.002 0.013 0.017
0.53 0.87 0.57 0.77 0.47 1.14 0.47 0.80 0.47 0.79 0.50 0.63 0.45
0.27 0.24 0.25 0.24 0.04 0.33 0.26 0.40 0.29 0.38 0.24 0.36 0.21
0.20 0.10 0.11 0.05 0.15 0.05 0.12 0.25 0.14 0.25 0.17 0.20 0.09
D.R.G. Mitchell, C.J. Ball / Materials Characterization 47 (2001) 17–26
3. Results 3.1. Parent plate Detailed analytical electron microscopy of these materials has been reported previously [16], and showed that the microstructure was unstable as a function of service exposure. Fig. 1 shows the parent plate microstructure in four headers in the virgin (asinstalled) condition, and those with service exposures of 75,800, 119,000 and 190,000 h (21.7 years). Carbides were selected at random for quantitative analysis, which permitted identification on the basis of composition [4,5,12,13,15,16]. Pro-eutectoid ferritic regions were characterised by a dispersion of acicular M2C, which coarsened with increasing service exposure. Prior austenite grain boundary carbides also developed, and alongside this in the pro-eutectoid ferrite a zone denuded in
19
M2C formed. The width of this zone as a function of service exposure (Fig. 2) showed a dramatic increase after about 120,000 h exposure. In the bainite, the carbides were either globular or lath-like. Analytical electron microscopy showed these to be M23C6 and M7C3 at short exposures. However, at longer exposures some spheroidisation of the bainitic carbides occurred (cf. Fig. 1a,b and c,d), and at the longest exposures large faceted M6C carbides were found to dominate the microstructure (Fig. 1d). Quantitative EDS analysis of carbides permitted ternary plots to be constructed (Fig. 3a – f ). In the early stages of service, the M2C carbides were too small to analyse individually, and so clusters of 100 or so carbides were analysed and the average compositions are shown. The frequency of appearance of M2C carbides in the ternary plots, therefore, bears no relation to the frequency of occurrence in the steel. In the bainitic regions the major carbides present
Fig. 1. Parent plate microstructures of (a) virgin material — Header A — 0 h exposure, (b) Header C — 75,800 h, (c) Header F — 119,000 h and (d) Header G — 190,000 h. The microstructure is unstable at normal service exposure temperatures. Acicular M2C develops within the pro-eutectoid ferrite, while coarsening and transformation of bainitic carbides occurs, culminating in large populations of large faceted M6C. A zone denuded in carbides (DZ) develops along the prior austenite grain boundaries.
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Fig. 2. Width of M2C denuded zone in pro-eutectoid ferrite regions adjoining prior austenite grain boundaries in the parent plate as a function of service exposure.
included M7C3, M23C6 and M6C. These were all of similar dimensions, being much larger than the M2C needles in the pro-eutectoid ferrite (a few M2C carbides were also present in the bainite of short-term exposure materials, but disappeared with ageing). With the limited sampling feasible with quantitative analytical electron microscopy, there is probably a loose correlation between the frequency of occurrence of the larger bainitic carbides in the alloy and in the ternary plots. A qualitative assessment of the frequency of the various bainitic carbide types (ignoring M2C) was made from the ternary diagrams and the results are shown in (Table 2). This suggested that M7C3 and M23C6 dominated the early parent plate bainitic carbide population. With increasing exposure M23C6 diminished and M6C became the dominant carbide at very long exposures. One header (Header F, Fig. 3e) showed unusual behaviour both in terms of carbide compositions and carbide makeup. M23C6 was the almost exclusive bainitic carbide detected in this header despite its considerable service exposure (119,000 h). Comparable exposure headers [D (Fig. 3c) and E (Fig. 3d)], which were of very similar vintage and design and came from the same power station, showed M7C3 and M6C as the dominant phases (Tables 2 and 3). The results of the X-ray diffraction-based carbide measurements are shown in Table 3. Since the X-ray diffraction technique measured carbides from both bainite and pro-eutectoid ferrite, in many cases a significant fraction of M2C was detected. This was the exclusive intraferritic carbide phase (Fig. 1a – d). In order to compare the X-ray diffraction data with those derived from analytical electron microscopy carbide assessments [which were based only on larger ( 0.1 – 1 mm) carbides from bainitic regions — i.e., the fine M2C in the pro-eutectoid ferrite was ignored],
the M7C3, M6C and M23C6 carbide fractions determined with X-ray diffraction were ranked with M2C omitted. Comparison of the data from the two methods (Table 2) showed that composition-based classification of carbides using analytical electron microscopy produced good agreement with the more accurate X-ray diffraction method. In nearly all cases analytical electron microscopy confirmed the presence of carbides identified using X-ray diffraction. The only exception was Header E (Fig. 3d), where no M23C6 was located with analytical electron microscopy, but where X-ray diffraction showed that nearly a third of the carbide was this phase (Table 2). The analytical electron microscopy technique is highly sensitive to any microstructural inhomogeneity in the steel, owing to the very tiny volumes sampled. The sampling statistics for the X-ray diffraction technique are significantly better. In general, the consistency between the ranking derived from the two data sets improved for the more aged materials. Agreement as to the most dominant carbide phase determined by X-ray diffraction and analytical electron microscopy was achieved for Headers B – G, while for A the two most dominant carbide phases were correctly identified, albeit with a different ranking of frequency of occurrence. In some instances X-ray diffraction did not detect minority phases that were located by analytical electron microscopy (e.g., M6C in Headers C and F). This was simply due to the low sensitivity of the X-ray diffraction technique. X-ray diffraction (Table 3) confirmed the analytical electron microscopy result, which showed that Header F was unusual compared with its sister headers (D and E) in that it contained only M23C6 despite being quite heavily aged (Fig. 3e). The X-ray diffraction data for the various carbides in the parent plate is plotted as a function of service exposure in Fig. 4a,b. All the data showed considerable scatter, though trends did emerge. The percentage of M2C generally diminished with service exposure (Fig. 4a), falling from a peak of 50% down to 15% in the most heavily aged material. Header F showed the lowest M2C levels ( 5%). The percentage of M6C showed a general increase from 0% in the virgin material to 56% in the most heavily aged steel. However, three aged headers contained very little M6C; Header C (75,800 h) contained no M6C, while a comparable exposure header (B — 70,200 h) contained 15%. This erratic behaviour was also shown by Headers D (114,000 h — 21% M6C) and E (119,000 h — 2.4% M6C). The X-ray diffraction results for M7C3 and M23C6 are shown in Fig. 4b. Data for Header F was again a long way off the trends for M7C3 and M23C6, with the latter carbide being almost the exclusive parent plate carbide
D.R.G. Mitchell, C.J. Ball / Materials Characterization 47 (2001) 17–26
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Fig. 3. Ternary composition plots for carbides found in the parent plates of the headers. Typical carbide compositions from the data of Moss and Kelly [20] ( ) and Titchmarsh [4] (~) are also shown: (a) A = 0 h, (b) C = 75,800 h, (c) D = 114,000 h, (d) E = 119,000 h, (e) F = 119,000 h, and (f) G = 190,000 h.
(95%). The M7C3 fraction changed little with service exposure, while after 75,000 h service, M23C6
declined from up to 42% to zero in the most heavily aged material.
Table 2 Carbide frequency in bainitic regions based on random AEM and XRD analysis — M2C omitted Header
Parent (AEM)
Parent (XRD)
Weld (AEM)
Weld (XRD)
A B C D E F G
M7C3 > M23C6 M7C3>M6C>M23C6 M7C3 M23C6>M6C M7C3 >> M6C>M23C6 M7C3 >> M6C M23C6 >> M6C M6C>M7C3
M23C6>M7C3 M7C3>M6C M23C6>M7C3 M7C3 >> M6C>M23C6 M7C3 >> M23C6 >> (M6C) M23C6 M6C >> M7C3
M7C3 M6C>M7C3 M6C M6C >> M7C3 M6C >> M7C3 M6C >> M7C3 N/A
M7C3 >> (M23C6) M7C3>M6C M6C >> (M7C3) M6C >> (M7C3) M6C >> (M7C3) N/A N/A
Carbides shown in parentheses represent < 15% of the total, N/A = not available.
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D.R.G. Mitchell, C.J. Ball / Materials Characterization 47 (2001) 17–26
Table 3 Quantitative XRD results for carbides extracted from service-exposed 2.25Cr – 1Mo (P = parent, W = weld) Header A B C D E F G
P W P W P W P W P W P W P
M23C6
M6C
M2C
M7C3
119,000
31.2 12.2 0 21.8 42.2 0 17.2 0 20.1 0 94.4
0 0 15.2 42 0 92.9 21.4 92.5 2.4 89.7 0
44.0 18.8 49.6 14.7 32.2 0 15.7 0 39.0 0 5.6
24.9 69.0 35.2 21.5 25.6 7.1 45.7 7.5 38.5 10.3 0
190,000
0
56.0
14.6
29.3
Exposure/h 0 70,200 75,800 114,000 119,000
with quantitative X-ray diffraction due to the volumes of available material, HAZ data is not presented here
3.2. Welds Analytical electron microscopy characterisation was also performed on the HAZs and welds. Since only the parent and weld regions could be analysed
Fig. 4. Percentiles of (a) M2C and M6C and (b) M7C3 and M23C6 in the parent plate as a function of service exposure, determined using X-ray diffraction. In (a) the M6C increases while the M2C decreases with increasing exposure, albeit with considerable scatter, while in (b) M23C6 diminishes with service exposure while M7C3 changes little.
Fig. 5. TEM micrographs of the weld microstructures of headers after service exposures of (a) 0, (b) 75,800 and (c) 119,000 h (Header F). Spherical Mn/Si-rich weld inclusions (WI) 1 mm in diameter are evident. With increasing exposure fine scale M2C disappears, while spheroidisation and coalescence of larger M6C occurs.
D.R.G. Mitchell, C.J. Ball / Materials Characterization 47 (2001) 17–26
(see Ref. [16] for analytical electron microscopy details). Fig. 5a – c shows a selection of weld microstructures from headers exposed for 0 (Fig. 5a), 75,800 (Fig. 5b), and 119,000 h (Fig. 5c). All welds were quite similar, comprising a coarse-grained fully bainitic microstructure. As well as carbides, rounded weld inclusions of 1 mm diameter were also present. These were typically rich in manganese and silicon. Analytical electron microscopy showed that the weld carbides were M6C and/or M7C3 with the former dominating in the longer aged materials (Table 2). No M23C6 was found, however very small rod-like M2C carbides were present in the virgin header and those exposed for 75,800 h (indicated in Fig. 5a). In materials aged for 114,000 h or more there were no very fine M 2C particles present (Fig. 5c). Ageing during service caused some coarsening of the larger weld carbides, with submicron particles being consumed by the larger carbides (cf. Fig. 5a,b and c). Qualitative ranking of the preponderance of various carbides in the welds (Table 2) showed that service exposure induced the formation of M6C quite early in the service life, but thereafter little further change occurred. Comparison of the carbide type and frequency of occurrence (M2C omitted) in the welds as determined by X-ray diffraction and analytical electron microscopy (Table 2) showed quite good agreement. In four out of five cases, analytical electron microscopy correctly predicted the carbide types and their frequency of occurrence when compared with the more accurate X-ray diffraction results. In Header A, a small amount of M23C6 (17% of the total with M2C excluded) was detected by X-ray diffraction but not by analytical electron microscopy. However, M7C3 was shown to be the dominant phase by both techniques in this material. X-ray diffraction analysis of the weld of Header F was not possible due to insufficient material being available. The plot of the percentage of M2C as a function of service exposure determined with X-ray diffraction (Fig. 6a) showed that up to 70,000 h M2C comprised up to 20% of the carbide, but thereafter it declined to zero. A concomitant and dramatic increase in M6C was found (Fig. 6a) with all welds with over 75,800 h exposure containing 90% or more of this carbide. The corresponding plots for M23C6 and M7C3 (Fig. 6b) showed that low levels of M23C6 were present in the virgin header weld, but diminished to zero after 75,800 h of service exposure. A significant decrease in M7C3 levels occurred, which mirrored the increase of M6C. Initial M7C3 levels of 70% fell to 10% after 75,800 h exposure. These trends indicated that in the welds M6C grew at the expense of M2C and M7C3, and that M23C6 was rapidly lost upon service exposure.
23
Fig. 6. Percentiles of (a) M2C and M6C and (b) M7C3 and M23C6 in the welds as a function of service exposure, determined using X-ray diffraction. Upon ageing, M6C rapidly dominates the carbide population by completely replacing M23C6 and M2C. Only low levels ( 5 wt.%) of M7C3 persist after 75,000 h of service exposure.
3.3. Supernatant liquor analysis and lattice parameter measurements The supernatant liquor from the carbide extractions was analysed for the principal alloying elements to determine the proportion of various solutes present in the matrix. The carbide compositional data from analytical electron microscopy was combined with the carbide fraction data from X-ray diffraction to determine the proportion of various solutes present in the carbide. Combining these two proportions and dividing by the known solute concentration in the alloy should yield a value of unity for each solute providing a check for experimental error. In most cases the values for Mo and Cr were between 0.9 and 1.1, which was quite reasonable considering the number of steps required to obtain the final result. However, in a few cases, errors of up to 50% were found. Plots of various solute fractions in either the matrix or carbide as a function of service exposure (not shown) yielded no simple trends. No trends in the carbide lattice parameters as a function of service exposure were found.
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D.R.G. Mitchell, C.J. Ball / Materials Characterization 47 (2001) 17–26
4. Discussion The microstructural instability of 2.25Cr – 1Mo upon ageing can lead to a general deterioration of mechanical properties [1,6,11]. The usual method for creep testing involves use of temperatures higher than those used in service, to accelerate the creep rate. Microstructural changes during this accelerated testing have been widely reported [13,18]. The drawback with such data is that the kinetics and thermodynamics of carbide change at a testing temperature of say 700C, may be quite different to those at the service temperature of 540C. One of the major advantages of the present study is that it has been carried out on service-exposed material, with exposures of up to 22 years. The data is therefore directly applicable to other service-exposed materials. In general there is a paucity of data derived from actual service-exposed material owing to its commercial sensitivity. One disadvantage of a study such as this is that it compares steels with subtly differing compositions, different thermal processing histories, as well as operating regimes that may differ. Meaningful test certificate and operational data for these headers was completely absent. This is not at all uncommon, and has been identified as a major limiting factor in the assessment of operational plant [19]. The carbide compositional parameters derived for these steels previously [16] indicated some quite strong compositional trends as a function of service exposure. This suggested that a comparative study of the various header steels was worthwhile, even with the various experimental uncertainties. In the analytical electron microscopy work, carbides were identified on the basis of composition alone. Several authors [4,13] have performed both electron diffraction and analytical electron microscopy studies on carbides, and have shown that the various carbide phases in low alloy steels have discrete compositional ranges, which permits ready identification. Ternary diagrams (e.g., Fig. 3a – f ) are a convenient way of showing the clustering of compositions for different carbide types. They are also useful for highlighting any particles with anomalous compositions. These most likely stem from superposition of different carbides as a result of nucleation and growth of new carbides on preexisting carbides. Ternary plots also show any compositional clustering or spreading upon ageing, as well as any general shifts in the carbide population. All the analytical electron microscopy work performed here was done quantitatively, ostensibly to determine mean carbide compositions. Typically, 20 – 30 carbides were analysed. This is a small sample for assessing the carbide population, and is subject to considerable scatter. This scatter is also exacerbated
by the fact that although there does not appear to be a dependence of carbide composition on size (within the range of carbide sizes studied here), relative abundance of carbides do vary in relation to the prior austenite grain boundaries [16]. Despite these limitations, the analytical electron microscopy and X-ray diffraction data showed excellent agreement with regard to the types and relative amounts of carbide present (Table 2). The X-ray diffraction data also confirmed the unusual behaviour of Header F. There appears to have been few quantitative X-ray diffraction studies of either accelerated aged [17] or ex-service 2.25Cr – 1Mo [10]. The precipitation sequence in 2.25Cr – 1Mo was first determined by Baker and Nutting [3], and has since been confirmed by several others [5,12 – 15]. Upon cooling from the austenitising temperature, Fe-rich M3C is the kinetically favoured carbide in the bainite of normalised steels [2]. M2C nucleates independently within the pro-eutectoid ferrite. Subsequent tempering causes the bainitic M3C to transform to more thermodynamically favoured carbides, rich in Cr and Mo. These include M7C3, M23C6 and M6C, which nucleate and grow on preexisting bainitic carbides. M2C in the pro-eutectoid ferrite has been reported to transform to M6C [3,21]. However, it is clearly very resistant to service-induced ageing, as it has been shown to persist in steels with exposures of 190,000 h in the present work (Fig. 1d) as well as in other studies of long-term aged steels [7,9]. Titchmarsh [4] has shown that tempering 2.25Cr – 1Mo for 10 h at 700C, conditions typically used to postweld heat-treat 2.25Cr – 1Mo, resulted in all the M3C being converted to alloy carbides. All the present materials had been postweld stress relief heat-treated, and therefore M3C had transformed to M23C6 and M7C3 before entering service (Fig. 3a), Table 3. Subsequent service-induced ageing results in the nucleation of M6C on existing carbides that are then rapidly consumed [3]. This sequence of events was reflected in the parent plate X-ray diffraction data (Fig. 4a,b), where the principal changes were loss of M23C6 and M2C while M7C3 levels changed little, but M6C increased considerably. The M6C data were interesting in that not all heavily aged materials developed M6C to the same degree (Table 3, Fig. 4a). For instance, in Header B (70,200 h) 15 wt.% of the carbides were M6C, while in Header C (75,800 h) there was none (Fig. 4a). A similar pattern emerged for Headers D and E/F. This may simply be a reflection of different nucleation kinetics for the M6C phase, although there are no obvious, microstructural or compositional factors that might account for this. No simple trends were present in the parent plate X-ray diffraction data as a function of service exposure. The reasons are most likely related to
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differences in parent plate composition, combined with differing fractions of bainite, and differing relative fractions of prior austenite grain boundary within the material (the ASTM grain sizes for the ferrite and bainite were quite similar for the respective materials, although these are of course average values). Since the M 2 C is primarily restricted to the pro-eutectoid ferrite, it will be very sensitive to the proportion of this phase in the steel, which in turn will depend on the alloy carbon content Attempts to derive correlations between the fraction of various carbides, such as M2C, and compositional parameters, such as the carbon levels, were not successful. X-ray diffraction data showed that M2C diminished with increasing service exposure (Fig. 4a), even though TEM showed this acicular carbide elongated with time at operating temperature (cf. Fig. 1a and d). This apparent contradiction can be explained by the fact that M2C was found to disappear rapidly from within bainitic regions upon ageing, due to reaction with other bainitic carbides. More importantly, considerable M2C dissolution from the near-prior austenite grain boundary zones occurred due to growth of alloy carbides (especially Mo-rich M6C) along the boundaries (Figs. 1a – d and 2). TEM of the welds showed a fully bainitic microstructure (Fig. 5a – c). The carbide makeup appeared more advanced in terms of thermal ageing than in the respective parent plate, which was to be expected owing to the additional tempering effects of multipass welding, prior to postweld heat treatment. Analytical electron microscopy data (Table 2) suggested an initial population of only M7C3, but this gave way to a population of >90% M6C (the most thermodynamically stable carbide) after the first 75,800 h of service. X-ray diffraction data was in very good agreement with analytical electron microscopy data (Table 2). The development of M6C upon tempering resulted in a loss of the fine scale (sub-50 nm) M2C carbide (cf. Fig. 5a and c). This dissolution of M2C as a result of the development of M6C was also observed in the parent plate near prior austenite grain boundary regions (Fig. 1d). M6C is more thermodynamically stable than M2C and has been shown to develop at its expense [3,21]. The growth of M6C and the loss of M2C in the weld were also observed with X-ray diffraction (Fig. 6a). Also noted were the complete loss of M23C6 after 75,800 h of service exposure (Fig. 6b), and a dramatic decrease in M7C3. The weld X-ray diffraction data showed much higher M6C levels at a given service exposure compare with parent plate (cf. Figs. 4a and 6a), again reflecting the greater effective ageing within the weld, due to the heat treatment effects of multipass welding.
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In general, the X-ray diffraction population data in the welds showed smoother trends compared with those of the parent plate (cf. Figs. 4 and 6). A lower level of scatter was also apparent in the mean compositional data of the weld carbides compared with the parent plate [16]. The reason for this is probably due to the dominance of a single carbide species in the virgin weld material ( 70% M7C3 — Table 3). The corresponding parent plate contained 31% M23C6, 44% M2C and 25% M7C3. This is a more diverse carbide environment within which a wider range of carbide precipitation, growth and dissolution reactions can occur.
5. Conclusions Analytical electron microscopy shows that 2.25Cr – 1Mo service-exposed for up to 190,000 h is microstructurally unstable. Within the parent plate, extensive development of acicular M2C occurs within the pro-eutectoid ferrite, while along prior austenite grain boundaries, growth of alloy carbides leads to the development of zones denuded in M2C. Within bainitic regions, some spheroidisation and carbide transformation reactions occur. Early carbide populations dominated by M7C3 and M23C6, give way to the more thermodynamically stable M6C upon ageing. Carbide identification and evolution has been studied by quantitative analytical electron microscopy, with ternary compositional diagrams offering a convenient way of sorting and classifying carbides. Within the fully bainitic welds, microstructural changes with service exposure were more subtle, predominantly involving loss of fine scale M2C. Early carbide populations were dominated by M7C3, which transformed to M6C upon ageing. Comparative quantitative X-ray diffraction analysis of extracted carbides showed that analytical electron microscopy correctly identified carbides on the basis of composition alone. Despite the poor counting statistics of the analytical electron microscopy method, qualitative assessment of carbide populations was generally in good agreement with that determined by the X-ray diffraction technique. Carbide evolution was in accordance with the Baker and Nutting sequence. Attempts to derive correlations between the weight fraction of individual carbides and service exposure showed general trends, but no close correlation was found, precluding any potential application for remnant life assessment. This lack of correlation was probably due to the combined effects of variables such as carbon content/fraction of proeutectoid ferrite/bainite, grain size/percentage of material in or near grain boundaries, composition and heat treatment.
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Acknowledgments The authors are indebted to Pacific Power International for the provision of materials and for funding part of this work, and to G. Smith and A. Croker of ANSTO for technical assistance and advice.
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