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High temperature oxidation in the context of life assessment and microstructural degradation of weldments of 2.25Cr–1Mo steel
International Journal of Pressure Vessels and Piping 79 (2002) 585–590 www.elsevier.com/locate/ijpvp
High temperature oxidation in the context of life assessment and microstructural degradation of weldments of 2.25Cr– 1Mo steel R.K. Singh Raman*, B.C. Muddle School of Physics and Materials Engineering, Monash University, 3800 Clayton, Vic., Australia
Abstract The prevalence of in-service failures in the welds of chromium – molybdenum ferritic steels causes great concern in steam generating/handling systems of power plants, and components of petroleum/petrochemical industries. This paper is a review of the nonuniform scaling behaviour across microstructural gradients in weldments of pressure vessel steels in order to develop a global model for lifeassessment by relating oxide scale thickness with time – temperature history of in-service components. The paper also investigates gaseous corrosion-assisted deterioration of the weldment microstructure. q 2002 Elsevier Science Ltd. All rights reserved. Keywords: Life assessment; Oxide scales; Steam environment; 2.25%Cr–1%Mo steels
1. Introduction Life extension of ageing plants is largely concerned with materials degradation of high temperature components. Materials degradation and ageing is of particular concern [1,2] for the steam generating/handling systems in fossil fuel power plants, which are constructed commonly from ferritic/bainitic chromium– molybdenum steels (generally referred to as ‘Cr – Mo’ steels) [3,4]. As the economic value of life extension and the need to understand materials ageing/degradation of in-service components are becoming increasingly important, new approaches have been explored in recent years: (a) for life assessment and reliability analysis of aged components [1,2], and (b) for improved understanding of materials degradation. Low alloy ‘Cr – Mo’ steels, viz. 2.25Cr – lMo and lCr – 0.5Mo steels, are used extensively in the steam generating and handling systems of power plants (in the temperature range of 623 –873 K) because they satisfy the required mechanical properties, weldability, formability and corrosion resistance [3,4]. Common applications of ‘Cr – Mo’ steels also include reactors for refining and processing of petroleum, high-temperature, high-pressure vessels for thermal reforming, polymerisation, alkylation, hydrocracking and coal-gasification [5 – 7]. The strength of the weldments of these steels is reported to be poor [8,9], to the extent that about 80% of the failures are reported [10] to * Corresponding author.
take place in the weldments. Considerable research work of the past 3– 4 decades has often indicated the inferior creeprupture of the weldments to result primarily from the microstructural degradation of the heat affected zone (HAZ). However, there is little information available on the high temperature gaseous corrosion-assisted mirostructural degradation and their specific contribution in damaging creep life of the weldments of ‘Cr – Mo’ steels. Another important issue is the corrosion-assisted damage in the various zones in the repair welded plate which has been thermally modified by long term exposure at operating temperatures and pressures prior to welding. Significant structural changes occur during in-service ageing and the modified structure responds differently to the weld thermal cycle than virgin steels [11]. This paper reviews the issues in the development of a new life assessment method and microstructural degradation model in the context of inferior in-service performance of the welded components of ‘Cr – Mo’ steels, and identifies the areas where a clear lack of understanding exists.
2. Oxide scaling: a tool for component assessment As the life extension of in-service components and hence the issue of materials ageing/degradation are becoming increasingly important, new approaches for life assessment and reliability analysis of aged components have been explored in recent years [1,2].
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2.1. Life assessment by scale thickness measurements Life assessment of steam generating/handling systems by scale thickness measurement is an innovative use of existing information in the form of the oxide scale, which is generally present on the in-service high temperature components. This emerging tool has been effectively used for determining the temperature history of steam generator tubes in fossil-fuel power plants [1,12,13]. In a typical steam generator tube bank, while the bulk steam temperature may be within design limits, the local steam temperatures in some of the superheater and reheater tubing may be up to 100 8C higher than the bulk temperature. Exposure to this increased temperature can cause a rapid loss of creep life for these tubes. Non-destructive measurement of the oxide scale thickness across the tube bank will provide information on specific tubes that have developed an excessive scale. A greater-than-average scale thickness is taken as indicative of an excessively high operating temperature, which could deteriorate the remaining life of the particular tube and necessitate replacement of the tube, and in some cases the tube bank. 2.2. ‘Oxide dating’ for crack growth measurement and failure analysis ‘Oxide dating’ [14] utilises an examination of oxide scales which forms at the fracture surface of a failed high temperature component, to determine crack growth history. This technique has been applied in failure investigations of various ‘Cr –Mo’ steel components, such as boiler headers, steam pipes, steam chests/castings, weldments and turbine blades [14]. The technique involves measurement of the thickness of oxide scales at locations between which the crack velocity is to be determined [12]. Presuming the average temperature remains the same over the component life, the scale thickness at each location can be taken as a rough measure of the time duration from the moment when a bare surface was created at the given location by crack propagation. The distance between any two locations, and the time spent in traversing the distance (which can be determined from the growth kinetics of scale thickness, as shown in Fig. 1) will provide the crack velocity between the two points. Similar measurements over the entire fracture surface will provide the crack growth history.
Fig. 1. Diagram showing the crack growth measurement from oxide scale thickness over the fracture surface. (Crack velocity ðA ! BÞ ¼ d=ðtA 2 tB Þ; tA and tB, times for scale growth at the locations A and B, can be determined from the scale growth kinetics).
‘Cr–Mo’ steel weldments, and consequent formation of a less protective scale during subsequent oxidation [18,19] may result in development of a thicker oxide scale over the HAZ than the other regions of the weldment, viz. weld metal and base metal regions. Air-oxidation behaviour of the weldments of ‘Cr –Mo’ steels has been investigated extensively [18 – 22]. In one of the investigations [21], a surface profile (shown in Fig. 2) describing the difference in thickness of the oxide scales over the weld metal and HAZ of a 2.25Cr– 1Mo steel weldment (oxidised: air/773 K/500 h) has unambiguously established that a much thicker scale forms (viz. nearly 14 mm thicker) over the HAZ region. In relation to life assessment and ‘oxide dating’ by scale thickness measurement, it will be mandatory to have a complete understanding of the greater scale thickness developed exclusively over the HAZ, in steam and other environments. That a considerably thicker scale formed over the HAZ also calls for the development of an improved model for accurately predicting the temperature history of the weldment region of
2.3. Non-homogenous scaling across ‘Cr –Mo’ steel weldments As will be discussed in Section 3, microstructural changes due to welding of chromium –molybdenum steels include additional chromium rich precipitate formation and/ or enrichment of chromium in the existing secondary precipitates in the HAZ adjoining the weld metal [4,15 – 18]. Trapping of ‘free’ chromium (in the matrix), through chromium-rich carbide precipitation in the HAZ of
Fig. 2. A surface profile describing the difference in thickness of oxide scales over weld metal and HAZ of 2.25Cr–1Mo steel, oxidised in air at 773 K for 500 h.
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in-service components. Hence, it will be mandatory to examine a similar non-homogeneity in the scaling rate during steam oxidation. 2.4. Relevance to life assessment of welded components In relation to life assessment by scale thickness measurement [1,12,13], the possibility of a greater scale thickness developed over the HAZ of the weldment (as compared to the weld and base metal regions, as discussed in the preceding section), may misleadingly indicate that this region had experienced an excessively high temperature. Therefore, in using scale thickness as a tool for life assessment of welded components of ‘Cr – Mo’ steels, it may be necessary to take into account the scaling rate of the steel weldments [18,21]. It may be necessary to develop a simple model to relate the scale thickness to the time-temperature history, for life assessment of the different regions of weldments, and extend the model to account for temperature transients in typical steam generation, to more accurately assess the life of welded components. In the context of ‘oxide dating’ [14] through scale thickness measurement, the thicker scale developed over the HAZ of the weldment could overestimate the exposure time (and, hence, underestimate the crack velocity). Therefore, while determining the crack growth history of welded components, it will be necessary to take into account the different scaling rates of the steel weldments.
3. Corrosion and microstructural degradation of weldments 3.1. Microstructural degradation of ‘Cr – Mo’ steel weldments The microstructures of ‘Cr – Mo’ ferritic steels (viz. 2.25Cr–lMo and lCr – 0.5Mo steels) are highly susceptible to thermomechanical treatments, which are often exploited to develop carbide precipitates of the required morphology and distribution to effect precipitation hardening. However, due to their metastable structure and morphology, the strengthening precipitates are known to undergo undesirable transformations [15 –18] during service at elevated temperatures and/or thermomechanical treatments experienced during fabrication, viz. welding, forging, hot-rolling. The creep rupture life of the weldments of ‘Cr –Mo’ steels is reported to be poor [8 – 10], to the extent that the creeprupture of the welds is often the life-limiting factor. Since weldments are an indispensable part of most component fabrications, considerable effort has been directed in the past 3 –4 decades, to the correlation of the in-service failure of these steels with the microstructural degradation caused during welding. Microstructural changes due to welding include enrichment of Cr in the secondary precipitates and/
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or additional Cr-rich precipitate formation in the area adjoining the weld metal (i.e. the HAZ) [3,4,8,10,16,17]. Trapping of ‘free’ chromium (in the matrix) through Cr-rich precipitate may result in formation of a less protective scale and thicker oxide scale over the HAZ than the other regions of weldment of ‘Cr –Mo’ steels. However, variation in the scaling rate across the weldments may be of secondary importance to design engineers (primarily because of the less deleterious influence of surface scaling losses on the mechanical properties, and also because scaling rate generally follows parabolic kinetics and the rate decreases considerably after the initial period). 3.2. Influence of gaseous corrosion on mechanical properties It is well known that the gaseous environments encountered by the high temperature components can play a significant role in determining their mechanical properties [23,24]. The most common influence is the loss of component thickness due to scaling, which is generally taken into account when designing the high temperature components (particularly, thin walled components). The influence of general scaling on the mechanical properties also include: creep acceleration by vacancy injection into the matrix, jacking of cracks/notches facilitating crack propagation, and notch/crack formation in the underlying matrix resulting from cracking of the oxide scale. However, there can be a marked influence on the mechanical properties due to the localised corrosion and the associated high temperature phenomena. Localised corrosion is reported to influence the creep and fatigue properties, due to one or more of the following phenomena: (a) crack initiation due to grain boundary oxidation [23], (b) crack blunting and arrest due to excessive oxidation at the grain boundary crack tip [24 – 26], (c) grain boundary strengthening due to selective oxidation and precipitation [26 –28], (d) selective leaching of alloying elements from the areas adjacent to grain boundaries, and vacancy and void formation which can assist crack propagation [26 –28], and (e) grain boundary cavitation resulting from annihilation of vacancies created due to selective leaching [25,26,29]. For example, as shown in Fig. 3, localised effects of gaseous corrosion include, formation of corrosion notches and extensive grain boundary oxidation causing grain detachment during air-oxidation of a 2.25Cr– 1Mo steel [30]. In cross-sections taken through a circumferentially cracked 2.25Cr –1Mo steel waterwall boiler tube, Smith and Marder [31] found the presence of extraordinarily deep and circumferentially oriented corrosion notches and cracks.
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Fig. 3. SEM micrograph showing extensive grain boundary oxidation causing grain detachment during air-oxidation of a 2.25Cr–1Mo steel at 823 K.
3.3. Gaseous corrosion-assisted microstructural degradation in ‘Cr –Mo’ steel weldments Intergranular cavitation and rupture life of metallic materials are reported to be influenced by both the environment and secondary precipitation [19,32] as well as the internal precipitation caused by the residual oxygen in the alloy matrix [33,34]. This section discusses an investigation on the oxidation-assisted microstructural degradation in a 2.25Cr– 1Mo steel weldment. In a critical investigation [35], specimen coupons including the weld metal, HAZ and base metal regions were sectioned from the weldments of a 2.25Cr–1Mo steel. These coupons were then oxidised at 873 K in an environment containing a mixture of 35% steam and nitrogen, and the morphology and structure of the oxidised specimens were characterised. A cross-section through a typical oxidised specimen (base metal) is shown in Fig. 4(a), which suggests the presence of at least three distinct regions, viz. a thick compact outer layer, a sub-scale region consisting of extensive internal oxide precipitates, and a thin inner layer which formed between the external layer and the subscale. The subscale region was populated with internal oxide precipitates, with the internal oxide precipitates forming more extensively at the alloy grain boundaries. X-ray mapping of the cross-section confirmed the internal oxide precipitates to be chromium-rich. Crosssections of the oxidised weld metal had features similar to those in the oxidised base metal. However, the cross-section of the oxidised HAZ specimen (seen in Fig. 4(b)) could be distinguished from those of the weld metal and the base metal specimens by two features: a subscale region with extensive void formation and blocky internal precipitates, and an additional region in the alloy matrix adjacent to the subscale, with features of cavitation at the alloy grain boundaries. Higher magnification (see Fig. 5) showed these grain boundary features to be cavities. Subscale regions in the base and weld metal specimens (Fig. 4(a)) were much less densely populated with internal precipitates than the subscale in the HAZ. More importantly, no grain boundary
Fig. 4. Representative SEM micrographs showing cross-sections through oxide scales and subscale zones formed over: (a) the base metal/weld metal specimens, and (b) HAZ, during steam oxidation at 873 K.
cavitation was observed in the region adjacent to the subscale regions in the oxidised specimens of weld metal and base metal. Microchemical analysis across the scale and subscale regions in the cross-sections of the oxidised weld metal, HAZ and base metal specimens confirmed the formation of Cr-rich inner scales. However, a comparison of the Crprofiles suggested that the relative Cr content in the inner layer of the scale developed over the HAZ was considerably lower (, 4 at pct) than those detected in the similar layers of
Fig. 5. Magnified features in the sub-scale and adjacent zones in the HAZ shown in Fig. 4(b). The arrows indicate regions of extensive precipitation and void formation in the sub-scale (B) and grain boundary cavitation in the adjoining area (C).
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the scales developed over the base and the weld metal (, 12 at pct). This variation in the chromium content of the inner oxide layer is very important since the chromium content of the inner layer of Fe/Cr oxides in low Cr steels governs the effective protectiveness of the scale [18,36]. It takes a specific combination of oxygen partial pressure beneath the external scale, alloy microstructure, temperature and the nature of resulting external scale to establish and sustain internal oxidation. The more-protective inner scales (i.e. those scales with high Cr contents), which formed on both the weld metal and the base metal, resulted in a limited inward diffusion of oxygen ions, and hence a less extensive internal oxidation. A less-protective inner scale formed in the case of the HAZ, presumably permitted a greater inward diffusion of oxygen ions, thus facilitating a greater concentration of oxygen available for reaction with the chromium of the alloy matrix. More rapid diffusion of oxygen ions through the HAZ scale caused extensive internal oxidation and formation of a subscale zone densely populated with internal precipitates. Depletion of Cr, due to extensive internal precipitation in the subscale zone of oxidised HAZ, may necessitate diffusion of Cr from adjacent areas in the alloy matrix, which will potentially lead to the generation of excess vacancies. These vacancies, preferentially annihilating at grain boundaries, could lead to the grain boundary void formation, seen in Fig. 5. Grain boundary cavitation resulting from extensive internal oxidation can provide an easy path for crack propagation [29,32], and hence needs to be taken into account for high temperature component design because it has a direct bearing on the creep/fatigue life. In this context, the oxidation-assisted grain boundary formation in the alloy matrix neighbouring the subscale zone in the oxidised HAZ specimen is particularly important since inservice failures are commonly found to occur in the HAZ of the welded components of 2.25Cr–1Mo steel [8,10].
4. Conclusions For an accurate use of oxide scale thickness as a tool for life assessment of the welded components, this review recommends further investigations as listed below: † examining the applicability of the recent practice of life assessment by scale thickness to the welded components, by establishing kinetics of the scale thickness growth across the broadly different (microstructural) zones (viz. weld metal, HAZ and base metal) of the steel weldments, in the environments of steam and air, † developing a suitable model for life assessment by relating scale thickness with time – temperature history (and hence, creep damage) in the different zones of the steel weldments, † testing the validity of the model for steam generators in fossil-fuel power plants,
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For an improved understanding of the deterioration in the mechanical properties due to deleterious alloy microstructure, it will be necessary to carry out further investigations as listed below: † characterising corrosion-assisted microstructural degradation in the alloy matrix across the broadly different (microstructural) zones (viz. weld metal, HAZ and base metal) of the steel weldments, in the environments of air and steam, and thus investigating the role of the environment in facilitating microstructural degradation, † examining the role of corrosion-assisted microstructural degradation, by conducting creep tests on the simulated weld metal and HAZ, in steam and inert environments.
Acknowledgments The authors are grateful to Mr Andrew Croker (ANSTO, Sydney), for the references on the life evaluation models, and to Dr Paul James, Power Technology Centre, Nottingham (UK) and Dr Steve Pascoe, Yallourn Energy, Australia, for their help with the references on ‘oxide dating’.
References [1] EPRI. EPRI Report TR-102433-V2, Boiler Tube Failure Metallurgical Guide. Palo Alto, EPRI; 1993. [2] Proceedings of Operating Pressure Equipment V. Melbourne, IMEA and WTIA; April 1999. [3] Khare AK (editor). Ferritic steels for high temperature applications. ASM International Conference, Warrendale, PA; 1981. [4] Pugh SF, Little EA, editors. Proceedings of the BNES lnternational Conference on Ferritic Steels for Fast Reactor Steam Generators, London: BNES; 1978. [5] Bland J. Welding J—Res Suppl 1956;35:181s. [6] Mandich LS, Fogelmen EL, Gulya JA. Symposium Heat Treated Steel for Elevated Temperature Services. New Orleans, LA: ASME; 1956. [7] Colebeck EW, Rait JR. High temperature steel and alloys for gas turbines. ISt Special Rep 1952;43:107. [8] Laha K, Rao KBS, Mannan SL. Mater Sci Engng 1990;129:183. [9] ASME. ASME Boiler and Pressuer Vessel Code, Code Case N-47. New York: ASME; 1986. [10] Chew B, Harris P. Met Construction 1979;11(5):11. [11] Sterjovski Z, Dunne D, Finlay R, Crocker A. Proceedings of IMMA Conference on Operating Pressure Equipment, Brisbane; April 1997. [12] Pinder LW. Corr Sci 1981;21:749. [13] EPRI. Appendix E, EPRI Report TR-102433-V2, Boiler Tube Failure Metallurgical Guide. Palo Alto: EPRI; 1993. [14] Powergen, Brochure: oxide dating for failure analysis. Ratcliffe on soar. Nottingham: PowerGen; 1999. [15] Pilling J, Ridley N. Metall Trans 1982;13A:557. [16] Roy P, Lauritzen T. Welding J 1986;65:45s. [17] Mann SD, Muddle BC, Proceedings of Microstructure and Mechanical Properties of Ageing Material, Warrendale, PA: TMS; 1993. [18] Singh Raman RK. Metall Mater Trans 1995;26A:1847. [19] Singh Raman RK. Metall Mater Trans 1998;29A:577. [20] Singh Raman RK. Metall Mater Trans 2000;31A:3101.
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[21] Singh Raman RK, Tyagi AK. Mater Sci Engng A 1994;185:97. [22] Singh Raman RK, Gnanamoorthy JB. Welding J—Res Suppl 1995; April:133-s. [23] Phaniraj C, Valson M, Mannan SL, Rodriguez P. Proceedings of the Workshop on Oxidation of Metals and Alloys, Indian Institute of Metals, Kalpakkam Chapter; 1984. p. 171. [24] Rodriguez P. Trans Indian Inst Metals 1967;20:213. [25] Shahinian P. Trans ASM 1957;49:862. [26] Widmer R, Grant NJ. Trans ASME 1960;82:882. [27] McCoy HE, Martin WR, Weir JR. Proc Inst Environ Sci 1961;163.
[28] Kramer IR, Balasubramanian N. Metall Trans 1973;2:431. [29] Johnson ER, Barrett CR, Nix WD. Metall Trans 1972;3:695. [30] Singh Raman RK, Khanna AS, Tiwari RK, Gnanamoorthy JB. Oxid Met 1992;37:1. [31] Smith BJ, Marder AR. Corrosion 1982;48:29. [32] Scaife EC, James PL. Metal Sci J 1968;2:217. [33] Bleakney HH. Can J Technol 1952;30:340. [34] Price CE. Acta Metall 1966;14:1787. [35] Singh Raman RK, Gnanamoorthy JB. J Mater Sci 1992;27:3435. [36] Singh Raman RK. Metall Mater Trans 1999;30A:2103.
High temperature oxidation in the context of life assessment and microstructural degradation of weldments of 2.25Cr– 1Mo steel R.K. Singh Raman*, B.C. Muddle School of Physics and Materials Engineering, Monash University, 3800 Clayton, Vic., Australia
Abstract The prevalence of in-service failures in the welds of chromium – molybdenum ferritic steels causes great concern in steam generating/handling systems of power plants, and components of petroleum/petrochemical industries. This paper is a review of the nonuniform scaling behaviour across microstructural gradients in weldments of pressure vessel steels in order to develop a global model for lifeassessment by relating oxide scale thickness with time – temperature history of in-service components. The paper also investigates gaseous corrosion-assisted deterioration of the weldment microstructure. q 2002 Elsevier Science Ltd. All rights reserved. Keywords: Life assessment; Oxide scales; Steam environment; 2.25%Cr–1%Mo steels
1. Introduction Life extension of ageing plants is largely concerned with materials degradation of high temperature components. Materials degradation and ageing is of particular concern [1,2] for the steam generating/handling systems in fossil fuel power plants, which are constructed commonly from ferritic/bainitic chromium– molybdenum steels (generally referred to as ‘Cr – Mo’ steels) [3,4]. As the economic value of life extension and the need to understand materials ageing/degradation of in-service components are becoming increasingly important, new approaches have been explored in recent years: (a) for life assessment and reliability analysis of aged components [1,2], and (b) for improved understanding of materials degradation. Low alloy ‘Cr – Mo’ steels, viz. 2.25Cr – lMo and lCr – 0.5Mo steels, are used extensively in the steam generating and handling systems of power plants (in the temperature range of 623 –873 K) because they satisfy the required mechanical properties, weldability, formability and corrosion resistance [3,4]. Common applications of ‘Cr – Mo’ steels also include reactors for refining and processing of petroleum, high-temperature, high-pressure vessels for thermal reforming, polymerisation, alkylation, hydrocracking and coal-gasification [5 – 7]. The strength of the weldments of these steels is reported to be poor [8,9], to the extent that about 80% of the failures are reported [10] to * Corresponding author.
take place in the weldments. Considerable research work of the past 3– 4 decades has often indicated the inferior creeprupture of the weldments to result primarily from the microstructural degradation of the heat affected zone (HAZ). However, there is little information available on the high temperature gaseous corrosion-assisted mirostructural degradation and their specific contribution in damaging creep life of the weldments of ‘Cr – Mo’ steels. Another important issue is the corrosion-assisted damage in the various zones in the repair welded plate which has been thermally modified by long term exposure at operating temperatures and pressures prior to welding. Significant structural changes occur during in-service ageing and the modified structure responds differently to the weld thermal cycle than virgin steels [11]. This paper reviews the issues in the development of a new life assessment method and microstructural degradation model in the context of inferior in-service performance of the welded components of ‘Cr – Mo’ steels, and identifies the areas where a clear lack of understanding exists.
2. Oxide scaling: a tool for component assessment As the life extension of in-service components and hence the issue of materials ageing/degradation are becoming increasingly important, new approaches for life assessment and reliability analysis of aged components have been explored in recent years [1,2].
0308-0161/02/$ - see front matter q 2002 Elsevier Science Ltd. All rights reserved. PII: S 0 3 0 8 - 0 1 6 1 ( 0 2 ) 0 0 0 9 3 - 5
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2.1. Life assessment by scale thickness measurements Life assessment of steam generating/handling systems by scale thickness measurement is an innovative use of existing information in the form of the oxide scale, which is generally present on the in-service high temperature components. This emerging tool has been effectively used for determining the temperature history of steam generator tubes in fossil-fuel power plants [1,12,13]. In a typical steam generator tube bank, while the bulk steam temperature may be within design limits, the local steam temperatures in some of the superheater and reheater tubing may be up to 100 8C higher than the bulk temperature. Exposure to this increased temperature can cause a rapid loss of creep life for these tubes. Non-destructive measurement of the oxide scale thickness across the tube bank will provide information on specific tubes that have developed an excessive scale. A greater-than-average scale thickness is taken as indicative of an excessively high operating temperature, which could deteriorate the remaining life of the particular tube and necessitate replacement of the tube, and in some cases the tube bank. 2.2. ‘Oxide dating’ for crack growth measurement and failure analysis ‘Oxide dating’ [14] utilises an examination of oxide scales which forms at the fracture surface of a failed high temperature component, to determine crack growth history. This technique has been applied in failure investigations of various ‘Cr –Mo’ steel components, such as boiler headers, steam pipes, steam chests/castings, weldments and turbine blades [14]. The technique involves measurement of the thickness of oxide scales at locations between which the crack velocity is to be determined [12]. Presuming the average temperature remains the same over the component life, the scale thickness at each location can be taken as a rough measure of the time duration from the moment when a bare surface was created at the given location by crack propagation. The distance between any two locations, and the time spent in traversing the distance (which can be determined from the growth kinetics of scale thickness, as shown in Fig. 1) will provide the crack velocity between the two points. Similar measurements over the entire fracture surface will provide the crack growth history.
Fig. 1. Diagram showing the crack growth measurement from oxide scale thickness over the fracture surface. (Crack velocity ðA ! BÞ ¼ d=ðtA 2 tB Þ; tA and tB, times for scale growth at the locations A and B, can be determined from the scale growth kinetics).
‘Cr–Mo’ steel weldments, and consequent formation of a less protective scale during subsequent oxidation [18,19] may result in development of a thicker oxide scale over the HAZ than the other regions of the weldment, viz. weld metal and base metal regions. Air-oxidation behaviour of the weldments of ‘Cr –Mo’ steels has been investigated extensively [18 – 22]. In one of the investigations [21], a surface profile (shown in Fig. 2) describing the difference in thickness of the oxide scales over the weld metal and HAZ of a 2.25Cr– 1Mo steel weldment (oxidised: air/773 K/500 h) has unambiguously established that a much thicker scale forms (viz. nearly 14 mm thicker) over the HAZ region. In relation to life assessment and ‘oxide dating’ by scale thickness measurement, it will be mandatory to have a complete understanding of the greater scale thickness developed exclusively over the HAZ, in steam and other environments. That a considerably thicker scale formed over the HAZ also calls for the development of an improved model for accurately predicting the temperature history of the weldment region of
2.3. Non-homogenous scaling across ‘Cr –Mo’ steel weldments As will be discussed in Section 3, microstructural changes due to welding of chromium –molybdenum steels include additional chromium rich precipitate formation and/ or enrichment of chromium in the existing secondary precipitates in the HAZ adjoining the weld metal [4,15 – 18]. Trapping of ‘free’ chromium (in the matrix), through chromium-rich carbide precipitation in the HAZ of
Fig. 2. A surface profile describing the difference in thickness of oxide scales over weld metal and HAZ of 2.25Cr–1Mo steel, oxidised in air at 773 K for 500 h.
R.K. Singh Raman, B.C. Muddle / International Journal of Pressure Vessels and Piping 79 (2002) 585–590
in-service components. Hence, it will be mandatory to examine a similar non-homogeneity in the scaling rate during steam oxidation. 2.4. Relevance to life assessment of welded components In relation to life assessment by scale thickness measurement [1,12,13], the possibility of a greater scale thickness developed over the HAZ of the weldment (as compared to the weld and base metal regions, as discussed in the preceding section), may misleadingly indicate that this region had experienced an excessively high temperature. Therefore, in using scale thickness as a tool for life assessment of welded components of ‘Cr – Mo’ steels, it may be necessary to take into account the scaling rate of the steel weldments [18,21]. It may be necessary to develop a simple model to relate the scale thickness to the time-temperature history, for life assessment of the different regions of weldments, and extend the model to account for temperature transients in typical steam generation, to more accurately assess the life of welded components. In the context of ‘oxide dating’ [14] through scale thickness measurement, the thicker scale developed over the HAZ of the weldment could overestimate the exposure time (and, hence, underestimate the crack velocity). Therefore, while determining the crack growth history of welded components, it will be necessary to take into account the different scaling rates of the steel weldments.
3. Corrosion and microstructural degradation of weldments 3.1. Microstructural degradation of ‘Cr – Mo’ steel weldments The microstructures of ‘Cr – Mo’ ferritic steels (viz. 2.25Cr–lMo and lCr – 0.5Mo steels) are highly susceptible to thermomechanical treatments, which are often exploited to develop carbide precipitates of the required morphology and distribution to effect precipitation hardening. However, due to their metastable structure and morphology, the strengthening precipitates are known to undergo undesirable transformations [15 –18] during service at elevated temperatures and/or thermomechanical treatments experienced during fabrication, viz. welding, forging, hot-rolling. The creep rupture life of the weldments of ‘Cr –Mo’ steels is reported to be poor [8 – 10], to the extent that the creeprupture of the welds is often the life-limiting factor. Since weldments are an indispensable part of most component fabrications, considerable effort has been directed in the past 3 –4 decades, to the correlation of the in-service failure of these steels with the microstructural degradation caused during welding. Microstructural changes due to welding include enrichment of Cr in the secondary precipitates and/
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or additional Cr-rich precipitate formation in the area adjoining the weld metal (i.e. the HAZ) [3,4,8,10,16,17]. Trapping of ‘free’ chromium (in the matrix) through Cr-rich precipitate may result in formation of a less protective scale and thicker oxide scale over the HAZ than the other regions of weldment of ‘Cr –Mo’ steels. However, variation in the scaling rate across the weldments may be of secondary importance to design engineers (primarily because of the less deleterious influence of surface scaling losses on the mechanical properties, and also because scaling rate generally follows parabolic kinetics and the rate decreases considerably after the initial period). 3.2. Influence of gaseous corrosion on mechanical properties It is well known that the gaseous environments encountered by the high temperature components can play a significant role in determining their mechanical properties [23,24]. The most common influence is the loss of component thickness due to scaling, which is generally taken into account when designing the high temperature components (particularly, thin walled components). The influence of general scaling on the mechanical properties also include: creep acceleration by vacancy injection into the matrix, jacking of cracks/notches facilitating crack propagation, and notch/crack formation in the underlying matrix resulting from cracking of the oxide scale. However, there can be a marked influence on the mechanical properties due to the localised corrosion and the associated high temperature phenomena. Localised corrosion is reported to influence the creep and fatigue properties, due to one or more of the following phenomena: (a) crack initiation due to grain boundary oxidation [23], (b) crack blunting and arrest due to excessive oxidation at the grain boundary crack tip [24 – 26], (c) grain boundary strengthening due to selective oxidation and precipitation [26 –28], (d) selective leaching of alloying elements from the areas adjacent to grain boundaries, and vacancy and void formation which can assist crack propagation [26 –28], and (e) grain boundary cavitation resulting from annihilation of vacancies created due to selective leaching [25,26,29]. For example, as shown in Fig. 3, localised effects of gaseous corrosion include, formation of corrosion notches and extensive grain boundary oxidation causing grain detachment during air-oxidation of a 2.25Cr– 1Mo steel [30]. In cross-sections taken through a circumferentially cracked 2.25Cr –1Mo steel waterwall boiler tube, Smith and Marder [31] found the presence of extraordinarily deep and circumferentially oriented corrosion notches and cracks.
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Fig. 3. SEM micrograph showing extensive grain boundary oxidation causing grain detachment during air-oxidation of a 2.25Cr–1Mo steel at 823 K.
3.3. Gaseous corrosion-assisted microstructural degradation in ‘Cr –Mo’ steel weldments Intergranular cavitation and rupture life of metallic materials are reported to be influenced by both the environment and secondary precipitation [19,32] as well as the internal precipitation caused by the residual oxygen in the alloy matrix [33,34]. This section discusses an investigation on the oxidation-assisted microstructural degradation in a 2.25Cr– 1Mo steel weldment. In a critical investigation [35], specimen coupons including the weld metal, HAZ and base metal regions were sectioned from the weldments of a 2.25Cr–1Mo steel. These coupons were then oxidised at 873 K in an environment containing a mixture of 35% steam and nitrogen, and the morphology and structure of the oxidised specimens were characterised. A cross-section through a typical oxidised specimen (base metal) is shown in Fig. 4(a), which suggests the presence of at least three distinct regions, viz. a thick compact outer layer, a sub-scale region consisting of extensive internal oxide precipitates, and a thin inner layer which formed between the external layer and the subscale. The subscale region was populated with internal oxide precipitates, with the internal oxide precipitates forming more extensively at the alloy grain boundaries. X-ray mapping of the cross-section confirmed the internal oxide precipitates to be chromium-rich. Crosssections of the oxidised weld metal had features similar to those in the oxidised base metal. However, the cross-section of the oxidised HAZ specimen (seen in Fig. 4(b)) could be distinguished from those of the weld metal and the base metal specimens by two features: a subscale region with extensive void formation and blocky internal precipitates, and an additional region in the alloy matrix adjacent to the subscale, with features of cavitation at the alloy grain boundaries. Higher magnification (see Fig. 5) showed these grain boundary features to be cavities. Subscale regions in the base and weld metal specimens (Fig. 4(a)) were much less densely populated with internal precipitates than the subscale in the HAZ. More importantly, no grain boundary
Fig. 4. Representative SEM micrographs showing cross-sections through oxide scales and subscale zones formed over: (a) the base metal/weld metal specimens, and (b) HAZ, during steam oxidation at 873 K.
cavitation was observed in the region adjacent to the subscale regions in the oxidised specimens of weld metal and base metal. Microchemical analysis across the scale and subscale regions in the cross-sections of the oxidised weld metal, HAZ and base metal specimens confirmed the formation of Cr-rich inner scales. However, a comparison of the Crprofiles suggested that the relative Cr content in the inner layer of the scale developed over the HAZ was considerably lower (, 4 at pct) than those detected in the similar layers of
Fig. 5. Magnified features in the sub-scale and adjacent zones in the HAZ shown in Fig. 4(b). The arrows indicate regions of extensive precipitation and void formation in the sub-scale (B) and grain boundary cavitation in the adjoining area (C).
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the scales developed over the base and the weld metal (, 12 at pct). This variation in the chromium content of the inner oxide layer is very important since the chromium content of the inner layer of Fe/Cr oxides in low Cr steels governs the effective protectiveness of the scale [18,36]. It takes a specific combination of oxygen partial pressure beneath the external scale, alloy microstructure, temperature and the nature of resulting external scale to establish and sustain internal oxidation. The more-protective inner scales (i.e. those scales with high Cr contents), which formed on both the weld metal and the base metal, resulted in a limited inward diffusion of oxygen ions, and hence a less extensive internal oxidation. A less-protective inner scale formed in the case of the HAZ, presumably permitted a greater inward diffusion of oxygen ions, thus facilitating a greater concentration of oxygen available for reaction with the chromium of the alloy matrix. More rapid diffusion of oxygen ions through the HAZ scale caused extensive internal oxidation and formation of a subscale zone densely populated with internal precipitates. Depletion of Cr, due to extensive internal precipitation in the subscale zone of oxidised HAZ, may necessitate diffusion of Cr from adjacent areas in the alloy matrix, which will potentially lead to the generation of excess vacancies. These vacancies, preferentially annihilating at grain boundaries, could lead to the grain boundary void formation, seen in Fig. 5. Grain boundary cavitation resulting from extensive internal oxidation can provide an easy path for crack propagation [29,32], and hence needs to be taken into account for high temperature component design because it has a direct bearing on the creep/fatigue life. In this context, the oxidation-assisted grain boundary formation in the alloy matrix neighbouring the subscale zone in the oxidised HAZ specimen is particularly important since inservice failures are commonly found to occur in the HAZ of the welded components of 2.25Cr–1Mo steel [8,10].
4. Conclusions For an accurate use of oxide scale thickness as a tool for life assessment of the welded components, this review recommends further investigations as listed below: † examining the applicability of the recent practice of life assessment by scale thickness to the welded components, by establishing kinetics of the scale thickness growth across the broadly different (microstructural) zones (viz. weld metal, HAZ and base metal) of the steel weldments, in the environments of steam and air, † developing a suitable model for life assessment by relating scale thickness with time – temperature history (and hence, creep damage) in the different zones of the steel weldments, † testing the validity of the model for steam generators in fossil-fuel power plants,
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For an improved understanding of the deterioration in the mechanical properties due to deleterious alloy microstructure, it will be necessary to carry out further investigations as listed below: † characterising corrosion-assisted microstructural degradation in the alloy matrix across the broadly different (microstructural) zones (viz. weld metal, HAZ and base metal) of the steel weldments, in the environments of air and steam, and thus investigating the role of the environment in facilitating microstructural degradation, † examining the role of corrosion-assisted microstructural degradation, by conducting creep tests on the simulated weld metal and HAZ, in steam and inert environments.
Acknowledgments The authors are grateful to Mr Andrew Croker (ANSTO, Sydney), for the references on the life evaluation models, and to Dr Paul James, Power Technology Centre, Nottingham (UK) and Dr Steve Pascoe, Yallourn Energy, Australia, for their help with the references on ‘oxide dating’.
References [1] EPRI. EPRI Report TR-102433-V2, Boiler Tube Failure Metallurgical Guide. Palo Alto, EPRI; 1993. [2] Proceedings of Operating Pressure Equipment V. Melbourne, IMEA and WTIA; April 1999. [3] Khare AK (editor). Ferritic steels for high temperature applications. ASM International Conference, Warrendale, PA; 1981. [4] Pugh SF, Little EA, editors. Proceedings of the BNES lnternational Conference on Ferritic Steels for Fast Reactor Steam Generators, London: BNES; 1978. [5] Bland J. Welding J—Res Suppl 1956;35:181s. [6] Mandich LS, Fogelmen EL, Gulya JA. Symposium Heat Treated Steel for Elevated Temperature Services. New Orleans, LA: ASME; 1956. [7] Colebeck EW, Rait JR. High temperature steel and alloys for gas turbines. ISt Special Rep 1952;43:107. [8] Laha K, Rao KBS, Mannan SL. Mater Sci Engng 1990;129:183. [9] ASME. ASME Boiler and Pressuer Vessel Code, Code Case N-47. New York: ASME; 1986. [10] Chew B, Harris P. Met Construction 1979;11(5):11. [11] Sterjovski Z, Dunne D, Finlay R, Crocker A. Proceedings of IMMA Conference on Operating Pressure Equipment, Brisbane; April 1997. [12] Pinder LW. Corr Sci 1981;21:749. [13] EPRI. Appendix E, EPRI Report TR-102433-V2, Boiler Tube Failure Metallurgical Guide. Palo Alto: EPRI; 1993. [14] Powergen, Brochure: oxide dating for failure analysis. Ratcliffe on soar. Nottingham: PowerGen; 1999. [15] Pilling J, Ridley N. Metall Trans 1982;13A:557. [16] Roy P, Lauritzen T. Welding J 1986;65:45s. [17] Mann SD, Muddle BC, Proceedings of Microstructure and Mechanical Properties of Ageing Material, Warrendale, PA: TMS; 1993. [18] Singh Raman RK. Metall Mater Trans 1995;26A:1847. [19] Singh Raman RK. Metall Mater Trans 1998;29A:577. [20] Singh Raman RK. Metall Mater Trans 2000;31A:3101.
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[21] Singh Raman RK, Tyagi AK. Mater Sci Engng A 1994;185:97. [22] Singh Raman RK, Gnanamoorthy JB. Welding J—Res Suppl 1995; April:133-s. [23] Phaniraj C, Valson M, Mannan SL, Rodriguez P. Proceedings of the Workshop on Oxidation of Metals and Alloys, Indian Institute of Metals, Kalpakkam Chapter; 1984. p. 171. [24] Rodriguez P. Trans Indian Inst Metals 1967;20:213. [25] Shahinian P. Trans ASM 1957;49:862. [26] Widmer R, Grant NJ. Trans ASME 1960;82:882. [27] McCoy HE, Martin WR, Weir JR. Proc Inst Environ Sci 1961;163.
[28] Kramer IR, Balasubramanian N. Metall Trans 1973;2:431. [29] Johnson ER, Barrett CR, Nix WD. Metall Trans 1972;3:695. [30] Singh Raman RK, Khanna AS, Tiwari RK, Gnanamoorthy JB. Oxid Met 1992;37:1. [31] Smith BJ, Marder AR. Corrosion 1982;48:29. [32] Scaife EC, James PL. Metal Sci J 1968;2:217. [33] Bleakney HH. Can J Technol 1952;30:340. [34] Price CE. Acta Metall 1966;14:1787. [35] Singh Raman RK, Gnanamoorthy JB. J Mater Sci 1992;27:3435. [36] Singh Raman RK. Metall Mater Trans 1999;30A:2103.