- Email: [email protected]
In-service behaviour of creep strength enhanced ferritic steels Grade 91 and Grade 92 – Part 2 weld issues
International Journal of Pressure Vessels and Piping xxx (2012) 1e12
Contents lists available at SciVerse ScienceDirect
International Journal of Pressure Vessels and Piping journal homepage: www.elsevier.com/locate/ijpvp
In-service behaviour of creep strength enhanced ferritic steels Grade 91 and Grade 92 e Part 2 weld issues Jonathan Parker EPRI, 1300 West WT Harris Blvd, Charlotte, NC 28262, USA
a r t i c l e i n f o
a b s t r a c t
Article history: Received 14 September 2012 Received in revised form 26 November 2012 Accepted 27 November 2012
In Creep Strength Enhanced Ferritic steels control of both composition and heat treatment of the parent steel is necessary to avoid producing components which have properties below the minimum expected by applicable codes. The degree of tempering involved in manufacture will modify the material hardness. While under most conditions hardness is reduced by tempering, exceeding the AC1 temperature can lead to an increase in hardness. In this heat treatment the properties will be relatively poor even though the measured hardness may be apparently acceptable. Thus, care should be exercised in imposing an acceptance test of components based on simple hardness alone. Differences in parent material heat treatment and composition apparently have remarkably little influence on the creep life of the heat affect zone (HAZ). Thus, Type IV cracking in the fine grained or intercritically heat treated regions of the HAZ does not appear to directly depend on the strength of the base steel. This form of in-service damage is relatively difficult to detect using traditional methods of non-destructive testing. Moreover, since repeated heat treatment leads to over tempering and a degradation of properties, specific procedures for making and then lifing repair welds are required. The present paper summarizes examples of damage and discusses best option repairs. Ó 2013 Elsevier Ltd. All rights reserved.
Keywords: Creep Damage Weld Service performance Steel Grade 91
1. Introduction Grade 91 steel is one member of the family of creep strength enhanced ferritic steels (CSEF) [1]. This steel is increasingly considered the material of choice for boiler, piping and header applications. Thus, this steel is being routinely installed in both fossil-fuel fired and combined cycle power generating units. The inservice performance of components manufactured from CSEF steels is reviewed; Paper I considered behaviour of parent with this paper focussing on welds. There are design and/or construction practises which result in local stress concentrations in creep-weak regions such as the heat affected zones (HAZ) in weldments. These issues can be a consequence of the approaches used for reinforcement of penetrations or because of the influences of other local stress concentrations. However, because welds will introduce significant complexity of microstructure and geometry it is apparent that welds are frequently the locations most susceptible to damage. Experience shows that the following locations are particularly susceptible to problems:
Non 90 branches (offset branches, welded “Y” pieces) End caps (fully or partially welded) Seam welds (whether re-normalized or not) Thermocouple pockets (specific arrangements) Dissimilar metal joints
A recent EPRI project has successfully established the background technology necessary to underpin significant improvements in approaches for ensuring that high quality components are supplied and installed [2]. In addition, this work delivered a life management strategy which is both practical and effective in preventing in-service failures. However, there are large numbers of components currently in operation which are susceptible to cracking. In the majority of cases the defects formed will need to be repaired using fusion welding techniques. The present paper briefly provides examples of highlights key factors which have been linked to in-service cracking and discusses challenges associated with repair and refurbishment. 2. Examples of service problems
Repair welds Terminal welds and branch welds E-mail address: [email protected].
A summary of typical damage found in welds manufactured in Grade 91 components has been presented in a recent EPRI report [3]. In the absence of gross problems associated with original weld
0308-0161/$ e see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.ijpvp.2012.11.004
Please cite this article in press as: Parker J, In-service behaviour of creep strength enhanced ferritic steels Grade 91 and Grade 92 e Part 2 weld issues, International Journal of Pressure Vessels and Piping (2012), http://dx.doi.org/10.1016/j.ijpvp.2012.11.004
2
J. Parker / International Journal of Pressure Vessels and Piping xxx (2012) 1e12
fabrication and Post Weld Heat Treatment, most reported weld related defects have been found in the Weld Heat Affected Zone (HAZ). P91 is known to be prone to in-service cracking known as Type IV. This damage has been found relatively early in life, particularly at locations where there is strain localization. The strain concentration may be a consequence of piping system stresses or local geometry or both. Type IV cracking thus typically describes creep-induced damage that has initiated in, and propagated through, the continuous fine-grained/intercritically heated zone of material. The creep strength of this zone of material is compromised because of changes that occur in response to the particular range of temperatures which occur during the welding cycle. Additional complications can be found when dissimilar metal welds are used. The transitions in properties associated with these welds are particularly complicated when the component dimensions, operational conditions and weld location serve to concentrate the stresses. Longitudinal seam welds in high temperature piping systems have also experienced damage in the weld metal and in the HAZ. The following summarizes in-service damage detected at different types of component. In general, the common factor associated with the presence of cracking is the design which is less than optimal for the prevailing operating conditions. Specific options for improved design are discussed. 2.1. Issues with nozzles Boiler and piping systems require that connections be manufactured between different pressurized components. These connections may be associated with a small bore pipe joined to a large bore, as in a drain line from a pipe, or may involve joining cylindrical components of similar dimensions, as with ‘Tees’ or Manifolds. In most cases, damage found at small bore connections is associated with thermal fatigue. The fatigue damage is often induced by pressure changes which cause relatively cold water to ‘quench’ a hot component. Many of these failures show a typical ‘star burst’ pattern, with transgranular cracks developed as a consequence of the repeated severe transients, Fig. 1. In general, inspection methods are available to detect these defects before component failure occurs. Since the damage is driven by repeated thermal cycling mitigation of this type of issue relates to removal of the condensate from the lines. Potential solutions involve improved gradients, increasing the size of the drain lines and introduction of thermocouples to detect the presence of water. Problems related directly to high temperature creep have been found in tube to header connections and in larger section nozzles. A
major investigation of damage in a replacement superheater outlet header has been reported [4,5]. In this case the damage found was at least in part explained by excessive strain in the header associated with the hoop stress. There was also at least a potential relationship between the in-service damage and the level of Aluminium present in the parent steel. Thus, because the measured Aluminium levels were high, and the associated Nitrogen to Aluminium ratios (N:Al) were relatively low, it was considered that the creep strength of some sections of the parent steel was at the bottom of the creep strength scatter band [5]. This relatively low parent strength was believed to be one factor in why so many of the connections developed cracking, Fig. 2a. The damage was initially detected by an inspection after around 50,000 h of operation. At that time surface defects were removed by grinding, Fig. 2b, but because of the levels of damage present the header was replaced [6]. Examination of a wider range of components has found that further creep damaged Grade 91 welds have occurred in-service in situations where the parent exhibited low N:Al ratios [5]. It is apparent that the potential problems caused by relatively high Al levels will be not be limited to influences on creep strength. It has been clearly demonstrated that high levels of fine inclusions in the steel can reduce ductility. Indeed, reductions in ductility with relatively low N:Al ratios have been reported [7]. This reduction occurs, particularly in steels with a fine grain size, because the large numbers of inclusions on grain boundaries act to initiate creep voids and cracks [8]. High levels of Al will therefore have the potential to promote Type IV cracking both by reducing the strength and by facilitating damage nucleation. In more general cases of cracking in the HAZ of nozzle welds, the problems can be linked at least in part to design approaches associated with reinforcement of the connection. It is normally the case that compensation must be used to reinforce an opening. However, even within a single design code there are several options for how the opening is reinforced. These include: 1. Adding thickness to the branch alone, i.e. the branch has a wall thickness above code minimum requirements, 2. Adding thickness to the main alone, i.e. the main has a wall thickness above code minimum requirements, 3. Or some combination where both the main and the branch are above minimum wall thickness requirements. The specific dimensions being based on detailed analysis to minimize stresses. A typical example of manufacture of a nozzle with a thickened branch is shown in Fig. 3. Clearly this type of Tee involving
Fig. 1. Magnetic particle inspection highlighting thermal fatigue cracking associated with a small bore attachment to a steam line (a) and optical micrograph showing the character of the defect revealed by laboratory metallographic examination (b).
Please cite this article in press as: Parker J, In-service behaviour of creep strength enhanced ferritic steels Grade 91 and Grade 92 e Part 2 weld issues, International Journal of Pressure Vessels and Piping (2012), http://dx.doi.org/10.1016/j.ijpvp.2012.11.004
J. Parker / International Journal of Pressure Vessels and Piping xxx (2012) 1e12
a
90
3
b
80 70
Numbers
60 50
Total Stubs
40
Cracked Stubs
30 20 10 0
1
2
3
4
5
6
Decreasing N:Al Ratio >>> Fig. 2. Histogram showing the relationship between incidence of in-service cracking and N:Al ratio for damage in a superheater outlet header (a) and a photograph showing the local grinding performed to remove the defects (b) [6].
Fig. 3. Typical connection on a main steam type component showing the reinforced branch before and after completion of welding.
significant weld deposition may not be optimal from a local stress situation. However, fabrication using welded construction often offers lower upfront cost than production of a forged component. There are many examples where connections of this
type have shown creep cracking, see for example Fig. 4. The detail in Fig. 4b shows that the local, predominantly intergranular local damage develops until micro and macro-cracks are formed.
Fig. 4. An example of in-service Type IV cracking at a reinforced connection, note that when this damage is related to creep strain in the main component the damage is typically primarily in the HAZ on that side of the weld (a), the cracking forms after the nucleation and development of individual creep voids (b) [6].
Please cite this article in press as: Parker J, In-service behaviour of creep strength enhanced ferritic steels Grade 91 and Grade 92 e Part 2 weld issues, International Journal of Pressure Vessels and Piping (2012), http://dx.doi.org/10.1016/j.ijpvp.2012.11.004
4
J. Parker / International Journal of Pressure Vessels and Piping xxx (2012) 1e12
2.2. Issues with wyes In an apparent attempt to avoid the extra fabrication cost associated with forged laterals a number of architect engineering firms specified the use of welded laterals for steam piping splits. These piping laterals were made using either craftspersonprepared sections of piping (to create the lateral connection and required reinforcement) or more commonly were made using preengineered/fabricated lateral fittings for the branch side of the connection. Fig. 5 provides an example showing the common crack prone regions of these saddle welds. It is typical to find cracks that are initiating from the weld root at the heel of the saddle weld and propagate along the side of the saddle welds in the base metal heat affected zone. Cracking in welded P91 steam piping laterals has occurred in many instances after approximately 35,000 h service, although as with all creep processes the specific timing of damage will depend on specific operating conditions. Detailed post service examination has revealed that the damage mechanism is predominantly Type IV cracking. Most, if not all of the problems would be overcome with a design involving balanced reinforcement. Thus, even if a welded construction is required, by using a main pipe that is thickened compared to code minimum wall, will greatly reduce the creep strain concentration. The thicker section could easily be forged or machined down near its ends, with a generous thickness taper, so that the ends at the girth welds would be matched with the adjoining pipe. This design concept would result in a welded P91 lateral producing equal or better performance than a forged F91 lateral. Of course care must always be used when increasing the wall thickness of components which will be thermally cycled. In general, cyclic loading will increase as the wall thickness increases. The message then is that welded steam piping connections need to be carefully designed to compensate for the risk of excessive creep in the base metal heat affected zone region.
reported in other end cap welds on a wider range of power plants. The high susceptibility to damage is a consequence of several factors; the flat end caps are relatively thick and thus constrain expansion of the header. This constraint can lead to high stress at the corner of the end fitting. Putting a relatively weak HAZ at a location of high stress then is less than optimal. These locations are increasingly being earmarked for onsite component inspection using advanced techniques such as Linear Array UT. 2.4. Issues with dissimilar steel welds Transition welds in boiler components are used in a very wide range of applications; these can be associated with changes in Geometry, changes from one ferritic type steel to another as well as joints which can be fabricated between ferritic and austenitic steels. There have been well documented studies which detail problems associated with these welds. In general problems are encountered where a comprehensive engineering design and procedure validation have not been performed. This type of evaluation should consider factors such as:
Joint design Type and thickness of weld metal Recommended weld process(es) and procedures Proper PWHT conditions for combinations of base metals/filler metals Metallurgical complexities related to the as fabricated microstructure and time/temperature dependent affects, e.g. Carbon Migration Initial quality assurance testing both for acceptance of a defect free structure and as a base line for follow up assessment of damage Future performance and damage evaluation Systems loading (Requirements for use of Cold Pull) Issues associated with Cyclic Operation
2.3. Issues with end caps and plates A series of plant failures have been reported associated with end caps and end plates in Grade 91 components operating at high temperature. One of the first which occurred, after about 20,000 h in-service, was attributed to incorrect heat treatment. Others in header end plate welds occurred after 36,500 h (with 469 hot and 72 cold starts). Damage in both cases was identified as Type IV cracking, Fig. 6 [4]. Nominal operating conditions for the headers were 568 C (1055 F), and 160 bar, and the failed weld had been inspected using manual UT and found clear of indications only 8663 h prior to failure. Following the failure, high sensitivity UT was carried out on three other Grade 91 headers on the affected unit and the equivalent headers on another unit. This led to the discovery of further damaged end plate welds. Failures have been
In general, the factors affecting the life of transition welds are understood. However, it is apparent from the examples given below that the understanding which exists is not always being applied to design and fabrication of components. Because this type of joint will continue to be necessary within pressure boundary components it is important to apply this understanding to new equipment and when remedial action is needed on existing plant. Potential opportunities for additional ways to solve problems continue to become available; these include the development and validation of new weld consumables as well as improvements to optimization of manufacturing process control. There are often requirements to join Grade 91 type steels to low alloy components. This type of joint creates questions regarding the consumables (normally a joint will use weld consumables which
Fig. 5. Cracking typically seen associated with saddle welded laterolets used in hot reheat systems. Significant Type IV cracking has been observed after around 35,000 h of service.
Please cite this article in press as: Parker J, In-service behaviour of creep strength enhanced ferritic steels Grade 91 and Grade 92 e Part 2 weld issues, International Journal of Pressure Vessels and Piping (2012), http://dx.doi.org/10.1016/j.ijpvp.2012.11.004
J. Parker / International Journal of Pressure Vessels and Piping xxx (2012) 1e12
5
Fig. 6. Schematic diagram showing the header end cap geometry and the crack location (a) and a metallographic section showing detail of the Type IV cracking (b) [4].
match the parent but when there are two different parents in a joint it is impossible to match both), the requirements for post weld heat treatment and the geometry. In many cases the best engineering solution is to fabricate a transition piece which will accommodate the changes in materials and geometry while minimizing fabrication problems. However, it is frequently the case in plant that these transitions have both changes in alloy and geometry at the same location. As shown in Fig. 7, when a weld between low alloy steel and P91 is fabricated with a Grade 91 type consumable care must be exercised over joint details. In the present example, the desire to ‘match’ thickness at the weld resulted in over stress conditions in the low alloy steel and relatively short term creep failure. Problems when the constituents are metallurgical very different are even more complicated. There have been a number of early service life failures of high energy steam piping dissimilar metal welds in combined cycle plants. Particular issues have been reported in joints which involve dissimilar metal girth welds made between P91 piping and austenitic stainless steel flow nozzles. These joints generally involve some combination of ENiCrFe-3 (Inconel 182) and ERNiCr-3 (Inconel 82) weld consumables [9]. Cracking has been found in relatively simple full thickness welds, and in joints made after application of Inconel 182 butter passes on the P91 side of the weld joint. The buttered layer is typically subjected to a shop post weld heat treatment before completion of the final weld. No additional PWHT is normally applied to these joints after completion of the weld fill passes. In other cases the entire final weld has been post weld heat treated. An example of an inservice crack is shown in Fig. 8.
A further complicating factor in assessing these components is the fact that almost no two flow elements have identical design and construction details. However, there are two broadly similar approaches used. Either the flow element is a short section of austenitic steel with an appropriate geometry (see Fig. 9) or the flow element is an austenitic steel insert within the P91 pipe spools. In the either case, an Inconel weld consumable is used to join the P91 to austenitic stainless steel flow element. Failures (leaks, severed pipe connections) of these welded connections have occurred after only 20,000 to 40,000 operating hours. 3. Factors affecting weld performance As far as the parent steel is concerned there is clearly a need to ensure that both the composition and all aspects of the heat treatment are performed in accordance with the latest standards. However, since in boilers there is a requirement to weld sections of parent, it is important that some consideration is given here to the effects of parent composition and heat treatment on the creep performance of welds. It is now well established that for subcritically heat treated welds the primary concern for long term creep life is Type IV cracking. The microstructural regions within a bead on plate weld were fully characterized by work performed at Marchwood Engineering laboratories by Alberry, Jones and co-workers [10e12]. The primary regions present have been represented based on the specific thermal cycles developed during welding, Fig. 10. The effect of the specific thermal cycles on final microstructure will depend on the
Fig. 7. In-service damage developed in a dissimilar weld between a low alloy steel casing and a Grade 91 pipe, a schematic diagram is presented in (a) with a photograph of the weld and the defect shown in (b). The cracking developed as a result of creep damage in the low alloy steel HAZ, (c).
Please cite this article in press as: Parker J, In-service behaviour of creep strength enhanced ferritic steels Grade 91 and Grade 92 e Part 2 weld issues, International Journal of Pressure Vessels and Piping (2012), http://dx.doi.org/10.1016/j.ijpvp.2012.11.004
6
J. Parker / International Journal of Pressure Vessels and Piping xxx (2012) 1e12
Fig. 8. Photograph showing a flow nozzle arrangement where the Austenitic steel flow element had been welded into the P91 piping system using Inconel filler. The inset picture shows damage development at the fusion boundary.
Fig. 9. Schematic diagram of an Austenitic Stainless Steel Flow nozzle welded into a Grade 91 steel pipe using Inconel 182 welds. Cracking was detected on the Grade 91 side of one of the welds [9].
details of the metallurgic system being considered. However, the general features in relation to the microstructures of martensitic steels are as follows: (i) Coarse grain region (CGHAZ): Material near the fusion boundary that reaches a temperature well above AC3 during welding. Precipitates which constitute the main obstacle to
Fig. 10. Schematic diagram illustrating how the thermal cycles during manufacture of a simple bead on plate type weld will influence the microstructure resulting in the formation of a Heat Affected Zone of complex gradients in metallurgical structure.
growth of the austenite grains must be dissolved, so that once austenite is formed growth of austenite grains is relatively rapid. In the 9e12 Cr steels, the region containing relatively coarse grained austenite transforms into martensite on cooling under typical conditions, Fig. 11a. (ii) Fine grain region (FGHAZ): As the distance from the fusion line increases the peak temperature achieved is above AC3 so that complete transformation to austenite occurs but this temperature is not sufficient to dissolve all of the precipitates present. Thus, austenite grain growth is limited by the incomplete dissolution of carbides, nitrides or carbo-nitrides. The fine grained austenite produced will mostly transform to martensite in the 9e12 Cr steels under typical cooling. However, in locations where the prior austenite grain size is very small the ability of these very fine grains to undergo a fully shear type transformation may be limited, Fig. 11b. (iii) Intercritical region (ICHAZ): There will be a region where the peak temperature exceeds the AC1 but not the AC3. Thus there will be some new austenite formed during the heating cycle. Moreover, some of the precipitates may dissolve but the combination of time and temperature is insufficient for complete dissolution of precipitates. There will thus be only a partial reversion to austenite on heating. This new austenite will predominantly transform to martensite on cooling, but it should be emphasized that the description ‘intercritical’ can be used to describe a wide range of conditions. Moreover, there will likely be some variation of specific microstructural features within an ‘intercritical’ region, Fig. 11c. (iv) Highly tempered region: There will be a region adjacent to the parent where the time; temperature combination will have
Please cite this article in press as: Parker J, In-service behaviour of creep strength enhanced ferritic steels Grade 91 and Grade 92 e Part 2 weld issues, International Journal of Pressure Vessels and Piping (2012), http://dx.doi.org/10.1016/j.ijpvp.2012.11.004
J. Parker / International Journal of Pressure Vessels and Piping xxx (2012) 1e12
7
Fig. 11. General appearance of the microstructure across the HAZ of a Grade 91 type steel weld, with optical micrographs detailing the structures in the relatively coarse grained region (a), the fine grained region (b) and part of the HAZ where incomplete transformation took place (c). This region is typically referred to as intercritical because the peak thermal cycles are between AC1 and AC3.
modified the local substructure. In CSEF steels this modification may relate to changes in precipitates or the locally high dislocation density or both. In either case these changes will likely not be easily resolved using optical metallography and will only be determined using advanced electron optics. Of course with multipass welds the relatively simple situation represented by bead on plate is significantly more complicated. Thus, depending on factors such as heat input, electrode size and angle, degree of overlap etc there can be regions of the parent which are exposed to multiple thermal cycles. Moreover, regions of deposited weld will also be thermally influenced and these locations greatly increase the possible variations. The final process associated with weld fabrication is normally PWHT. In 9e12% Cr steels this heat treatment tempers the new martensite introduced by the welding thermal cycles. However, gradients in the microstructure and mechanical properties, extending typically over a few millimetres from the fusion boundary are likely to be present even after this treatment. As the extent of the tempering during PWHT increases there is the potential that the transitions in properties will also be reduced. An example of the increase in tempering during PWHT on microhardness of a Grade 91 weld is illustrated in Fig. 12 [13]. As shown, the lowest degree of tempering, 2 h at 732 C, shows the highest hardness in the weld metal and the lowest hardness in the HAZ. As the level of tempering increases the hardness of the weld and parent both exhibit reductions. However, even at the highest degree of tempering, 40 h at 732 C the lowest hardness values continue to be measured in the HAZ. It is apparent from cross weld creep testing of CSEF steels that for conditions approaching in-service stresses the time to Type IV cracking falls below that of the parent [14]. This behaviour is illustrated for Grade 92 steel in Fig. 13. It is apparent that all CSEF steels are highly susceptible to Type IV cracking. This susceptibility comes about since while most of the component and weld have
excellent strength (conferred by the well controlled heat treatment which avoids deleterious effects) there are regions in the HAZ where the thermal cycles due to welding result in microstructures with low creep strength. Simply put, there are regions in the HAZ with a very fine grain size and no martensitic microstructure. Furthermore, it has also been has shown [14]that M23C6 precipitates and Laves phases form faster in the fine grain HAZ region in 9Cr martensitic type steels compared with the other regions of the weldment. This metallurgical effect further increases the vulnerability of the Type IV region. Since not only are matrix-strengthening elements such as Cr, Mo and W depleted but the Laves phase offers potential sites for the nucleation of creep voids. There are many other examples of Type IV cracking in other CSEF steels. The pictures in Fig. 14 are from a major European
Fig. 12. Microhardness traverse across a SA weldment produced at ORNL. Both the base plate and filler wire were of the modified 9Cr-lMo composition. Post weld heat treatments were for the times shown at 732 C [13].
Please cite this article in press as: Parker J, In-service behaviour of creep strength enhanced ferritic steels Grade 91 and Grade 92 e Part 2 weld issues, International Journal of Pressure Vessels and Piping (2012), http://dx.doi.org/10.1016/j.ijpvp.2012.11.004
8
J. Parker / International Journal of Pressure Vessels and Piping xxx (2012) 1e12
Fig. 13. Comparison of the Creep Rupture behaviour of Grade 92 steel parent and the results of Cross Weld Creep tests at 650 C [based on data from Ref. [14], micrograph from Ref. [4]].
Programme studying the behaviour of E911 [15]. These investigations again show that the most susceptible region of the HAZ is the fine grained location. The failure occurred where the HAZ microstructure consisted of very small equiaxed prior austenite grains where the grain boundaries were decorated with large precipitates. No martensitic lath structure was observed within this region. Creep voids mainly formed along prior austenite grain boundaries. Final fracture occurred by interconnection of single creep voids to form micro-cracks, which themselves finally coalesced to form macro-cracks. As shown in Fig. 14, high densities of creep voids are developed over the HAZ, with crack formation and final propagation occurring only very late in creep life.
There has been some disagreement over details of the formation of creep voids associated with Type IV cracking. In general, where loading is reasonably uniform across the section, i.e. in the absence of significant bending loads, it is now clear that creep voids nucleate relatively early in the creep life. These voids are initiated at locations below the component surface so that even a properly performed examination at the surface of a weld will often under estimate the damage present. These observations have been seen in both laboratory cross weld experiments and in post service examination of pressure boundary welds. The results shown in Figs. 14 and 15 are typical of laboratory generated information. When a relatively large specimen is tested under the appropriate conditions damage will typically be developed in both the HAZs within the gauge length. Thus, preparation of a sample allows assessment of both the fractured HAZ and the interface where damage levels are relatively close to fracture. This situation is demonstrated in Fig. 15. The HAZ to the right of the weld has failed in a classic Type IV manner and Microdamage is clearly present in the left hand HAZ. Detailed optical metallography, Fig. 15a, reveals that individual voids were associated with the damage at both HAZs. In addition to variations in metallographic preparation techniques a wide variety of methods have been used to examine and quantify the creep voids present. It is apparent that even with careful preparation there will be variation in the number of voids detected depending on the magnification used for examination. As a general rule the number density of voids will increase with the magnification used at least in the range 200 times to around 2000 times. There have been several comprehensive investigations of void development reported. Fig. 16 shows post test examination results from a set of long term Grade 91 cross weld tests which
Fig. 14. Illustration of the Type IV damage development in the CSEF steel Grade E911 in Cross Weld Creep tests performed at 600 C [15]. The examination using Electron Backscatter Diffraction allows high resolution, detailed mapping of features such as grains, subgrains and martensite laths.
Please cite this article in press as: Parker J, In-service behaviour of creep strength enhanced ferritic steels Grade 91 and Grade 92 e Part 2 weld issues, International Journal of Pressure Vessels and Piping (2012), http://dx.doi.org/10.1016/j.ijpvp.2012.11.004
J. Parker / International Journal of Pressure Vessels and Piping xxx (2012) 1e12
9
Fig. 15. Details of creep voids formed during a cross weld creep test which was interrupted prior to failure (a), and a montage showing details of Type IV cracking (b). The right hand HAZ shows through section cracking with the left hand HAZ showing microcracking.
Fig. 16. Post test examination showing the development of creep voids in long term cross weld creep tests with increasing life fraction [16].
were interrupted at known life fractions. It is clear there is a relatively high density of creep voids present in the Type IV region after a creep exposure of approximately 70%. Even at this level of creep life fraction void damage does not appear to extend to the sample surface and there is no clear evidence of cracking. This is only apparent after around 90% of life. The data presented in Fig. 17 compare the creep lives observed in relatively low strength parent, the data from bar 257, with results obtained from cross weld creep testing. It is apparent the cross weld lives observed in the bar 257 steel fall well below the expected scatter band for parent steel. Consideration of the behaviour of a second steel, identified as bar 817, with typical parent behaviour indicates that the cross weld life is similar to that of the welds in bar 257. It thus appears that variations in chemical composition and heat treatment have a greater effect on the creep strength of the parent than they do in the HAZ. The high aluminium, weak Grade 91 material bar 257, which substantially underperforms mean P91 parent creep data, showed a much lesser shortfall in its HAZ
Fig. 17. Comparison of the creep rupture lives of cross weld samples which failed by Type IV cracking in the HAZ with Parent behaviour, all data are for Grade 91 steels tested at 600 C. [Based on data from Ref. [4]].
Please cite this article in press as: Parker J, In-service behaviour of creep strength enhanced ferritic steels Grade 91 and Grade 92 e Part 2 weld issues, International Journal of Pressure Vessels and Piping (2012), http://dx.doi.org/10.1016/j.ijpvp.2012.11.004
10
J. Parker / International Journal of Pressure Vessels and Piping xxx (2012) 1e12
performance. It may be that high aluminium, which is known to combine with nitrogen and suppress the formation of key creep strengthening carbonitride precipitates, has a more harmful effect on parent properties than on those of the HAZ. It could therefore be that, whereas optimized precipitation strengthening is the key to optimized parent performance, such strengthening mechanisms become less effective in the non-optimal and lath-free fine grained HAZ microstructure. Other factors, for example simple solution strengthening, may perhaps become more important in determining HAZ performance. Moreover, as discussed later there could also be an influence on Type IV cracking susceptibility due to welding residual stresses. 4. Discussion It is apparent that welding needs to be carried out carefully using qualified staff and appropriate procedures, further background is provided in reference [17]. Assessment of the creep performance of welds is then frequently studied by programmes which include cross weld testing. In these programmes it is important to select test conditions and specimen geometries that produce failures where the damage mechanisms are relevant to long terms service. In general, it is now widely accepted that in creep tests at relatively high stress and temperature, the results are not typical of long term damage in component welds. Thus, results from tests under these conditions should not be used as a guide to in-service behaviour. The concept of a weld strength reduction factor (WSRF) has been introduced in design codes for high temperature nuclear power and for process piping. It is defined as a ratio of the stresses causing rupture at the same time in welded joint and base metal. Although no explicit employment has been made in design codes for non-nuclear plants so far, the introduction of a WSRF is regarded necessary to avoid failure at welded joints. In general, the procedure involves obtaining a welded joint “allowable stress” based on cross weld creep rupture data and then calculating WSRF as the ratio of this value to the allowable stress for parent [18]. Thus, the average stress value and 95% lower value of 105 h rupture stress were calculated at each temperature using the equations for lower stress region. Then, the welded joint “allowable stress” was given as minimum value of (the average 0.67) and (95% lower value 0.8). The calculated WSRF for Grade 91, Grade 92 and Grade 122 are listed in Table 1 [18]. Residual stresses are often introduced unintentionally during fabricationdfor example, during welding or heat treatment. A few elegant experiments illustrate how phase transformations interact with the buildup of residual stress with the results summarized in Fig. 18 [19]. Using bainitic, martensitic, and stable austenitic steels, Jones and Alberry [20,21] demonstrated that transformation plasticity during the cooling of a uniaxially constrained sample from the austenite phase field acts to relieve the buildup of thermal stress as
Fig. 18. Interpretation of experimental data showing how residual stresses develop on cooling for an austenitic steel (no transformation), a bainitic low alloy steel (relatively high temperature of transformation) and a martensitic steel (relatively low temperature of transformation). From Ref. [19] based on data from Refs. [20] and [21].
the sample cools. By contrast, the non transforming austenitic steel exhibited a continuous increase in residual stress with decreasing temperature, as might be expected from the thermal contraction of a constrained sample. When the steels were transformed to bainite or martensite, the transformation strain compensated for any thermal contraction strains that arose during cooling. Significant residual stresses were therefore found to build upbuildup only after transformation was completed and the specimens approached ambient temperature [22,23]. There have been considerable efforts to create welding consumables which on solid state phase transformation partly compensate for the stresses which develop when a constrained weld cools to ambient temperatures [24,25]. Many of these studies have considered on structural steels which are ferritic. There is however, some work reported which has used alloy design methods to create a stainless steel welding consumable which solidifies as ferrite, transforms almost entirely into austenite which then undergoes martensitic transformation at a low temperature of about 220 C. Based on consideration of the effect of martensitic transformation on welding residual stresses there now could be a further reason why 9 to 12%Cr CSEF steels are universally susceptible to Type IV cracking. As shown earlier, there will be a band within the HAZ where martensitic transformation does not take place. This could be in the fine grained location or where peak temperatures were in the intercritical region. In either case it appears that the lack of a martensitic transformation will lead to locally high residual stresses. These stresses will relax by strain, however, since martensitic steels are strain softening these relaxation processes could lead to further reductions in strength. In any event it is likely that the relaxation strain would be relatively high so that there would be effective reduction in subsequent creep life.
Table 1 Evaluation of weld strength reduction factors for Grade 91, Grade 92 and Grade 122 steels on the basis of a recent reevaluation of the results of long term creep testing [18]. Material
Gr.91
Gr.122
Gr.92
Weld strength reduction factor
KA-SCMV28 KA-STPA28 KA-SFVAF28 KA-SUS410J3 KA-SUS410J3TP KA-SUSF410J3 KA-STPA29 KA-SFVAF29
Note
525 C
550 C
575 C
600 C
625 C
650 C
1.00 1.00 1.00 1.00 1.00
0.90 0.90 0.90 0.84 0.84
0.82 0.74 0.74 0.68 0.60
0.79 0.67 0.68 0.57 0.50
0.79 0.65 0.65 0.50 0.50
0.79 0.65 0.65 0.50 0.50
1.00
1.00
0.74
0.62
0.53
0.53
<¼76 mm >76 mm
Please cite this article in press as: Parker J, In-service behaviour of creep strength enhanced ferritic steels Grade 91 and Grade 92 e Part 2 weld issues, International Journal of Pressure Vessels and Piping (2012), http://dx.doi.org/10.1016/j.ijpvp.2012.11.004
J. Parker / International Journal of Pressure Vessels and Piping xxx (2012) 1e12
11
Table 2 Summary of key variables being considered in the EPRI project assessing the factors affecting the performance of repair welds in Grade 91 steel. Weld
1A, 1B 2A, 2B 3A, 3B 4A, 4B 5A, 5B 6A, 6B 7A, 7B 8A, 8B 9A, 9B 10A, 10B
Base material
As-received Grade 91 (A) Re-normalized Grade 91 (B)
Weld metal AWS Desig.
Trade name
E9015-B9 H4 E9015-B9 H4 E9015-B9 H4 E9015-B9 H4 E8018-B8 E9015-G E9015-G E9018-B3 H4 EPRIP87 ENiCrFe-2
Thermanit Chromo Thermanit Chromo Thermanit Chromo Thermanit Chromo 9Cre1Mo Thermanit P23 Thermanit P23 Bohler E9018-B3 EPRIP87 INCO-WELD A
It is clear that Grade 91 type steels are widely used in boilers and piping. Thus, there will be an increasing need to perform remedial action of high temperature components manufactured from this steel. To optimize the processes associated with remediation, it is critical to establish key information on removal of damage and on approaches for repair welding. Selected variables and weld processes are under investigation in a current EPRI project. This work includes weldment testing linking high temperature performance to details of how the repair was carried out. This aspect is considered important since in contrast to traditional low alloy steels which are relatively ‘user friendly’ as far as repairs are concerned; Grade 91 steels introduce additional complications. Because the properties and performance of the base metal will be critically dependent on the original composition and the full heat treatment history, there have been concerns expressed that the reparability of Grade 91 steel will be limited by the condition of the base metal. Because there are a very large number of variables to be evaluated the project will be performed in two phases. Phase 1 will provide a Ranking of Repair performance. The following will be the primary factors for consideration: Discussion of Methods and Extent of Excavation, Weld Procedure Considerations (including consumables) and Heat Treatment Post Repair evaluation of microstructure, damage etc Specimen Geometry and Testing Conditions including development of test matrix Analysis to identify best option repairs e ‘best option’ to be based on factors such as speed of welding, initial quality, creep life The key considerations for the first Phase are summarized in Table 2. The general use of summary phrases such as “Normal” and “Temperbead” in Table 1 should not be taken to suggest that the details were the same for each weld fabricated. These specifics were customized as necessary for each weld produced. The parent sections where of the same composition. One set of welds, designated A, were manufactured in steel in the ex-service condition with the other set, designated B, in the same sections after full Renormalizing and tempering heat treatment. It is clearly the expectation that by understanding the interaction of thermal cycles, composition and microstructural development the repair methods developed will have the potential for welding without the need for post weld heat treatment. The second Phase of this project will involve analysis of the phase 1 results to identify best option repair methods. The preferred methods will then be applied to sections of an ex-service header which were removed after around 75,000 h of operation. The performance of this header was well documented in several publications (e.g. 6). The published work has been supplemented by extensive post service NDE and selective laboratory metallography. The main tasks
9V 9V 9V 9V
Mod. Mod. Mod. Mod.
Welding procedure
PWHT
Normal þ Rec’d. PWHT Normal þ Min. PWHT Temperbead Poor Practice Temperbead Temperbead Temperbead Normal þ Rec’d. PWHT Temperbead Temperbead Temperbead
1375 25 F/2 h 1260 10 F/2 h None None None None 1375 25 F/2 h None None None
involved in Phase 2 will be associated with agreeing methods and extent of excavation, specifics of the best option repair methods and performance assessment of the manufactured welds. 5. Concluding remarks Grade 91 steel is one of the family of creep strength enhanced ferritic steels (CSEF). This steel is now routinely installed in both fossil-fuel fired and combined cycle units. Recent in-service experience has demonstrated that cracking can occur in Grade 91 steel early in life [1]. The results also show that while differences in material heat treatment prior to welding clearly affect parent behaviour, they have remarkably little influence on HAZ performance. Thus, it is very likely that the primary damage location inservice will be Type IV cracking. Potential problems associated with this form of in-service damage include challenges over detection and appropriate methodologies for repair. Since repeated heat treatment leads to over tempering and a degradation of properties specific procedures for performing and then lifing repair welds are required. The following issues need to be considered when considering performance: 1. It is now widely accepted that in creep tests at relatively high stress and temperature the results of cross weld creep testing are not typical of long term damage in component welds. 2. Clearly then it is important to select test conditions and specimen geometries for laboratory test programs so as to produce failures where the damage mechanisms are relevant to long terms service. 3. Using these conditions it is apparent that failure occurs as a consequence of the nucleation, growth and link up of creep voids. It appears that the damage is significantly greater within the volume of the specimen where relatively high constraint conditions are developed. 4. The Type IV life is significantly below that of the parent under the same conditions. The effects of weld thermal cycle on microstructural are clearly important to the susceptibility for creep damage to occur. However, there is at least inferential evidence that transformation behaviour is effecting the distribution of weld residual stress as well as the local creep strength. The significance of residual stresses to the susceptibility for Type IV cracking in 9e12%Cr steels is one explanation for why the creep strength of the parent has only a second order effect on Type IV performance. Further work is in progress to examine Grade 91 welded samples which have been tested to different creep life fractions. Advanced characterization techniques will be used to establish further details of creep cavity nucleation and growth within the weld HAZ.
Please cite this article in press as: Parker J, In-service behaviour of creep strength enhanced ferritic steels Grade 91 and Grade 92 e Part 2 weld issues, International Journal of Pressure Vessels and Piping (2012), http://dx.doi.org/10.1016/j.ijpvp.2012.11.004
12
J. Parker / International Journal of Pressure Vessels and Piping xxx (2012) 1e12
Acknowledgements Thanks are due to the participants of the recent EPRI project “Life Management of CSEF steels” for much stimulating discussion on all aspects of the fabrication, installation and performance assessment of Grade 91 and Grade 92 steels. References [1] DiStefano JR, Sikka VK. Summary of modified 9Cr-1Mo steel development program, 1975e1985. ORNL-6303. Oak Ridge National Laboratory; October 1986. [2] Guidelines and specifications for high-reliability fossil power plants: best practise guideline for manufacturing and construction of Grade 91 steel components. Palo Alto, CA: EPRI; 2011. 1023199. [3] Service experience with Grade 91 components. Palo Alto, CA: EPRI; 2009. 1018151. [4] Brett SJ, Allen DJ, Buchanan LW. The Type IV creep strength of Grade 91 materials. In: Proceedings of the third international conference on integrity of high temperature welds. Institute of Materials; 2007. p. 409e20. [5] Brett SJ, Bates JS, Thomson RC. Aluminium nitride precipitation in low strength Grade 91 power plant steels. In: Proceedings of the fourth international conference on advances in materials technology for fossil power plants, October 25 to 28. Hilton Head Island, South Carolina: EPRI; 2005. p. 1183e97. [6] Brett SJ. Service experience with a retrofit modified 9Cr (Grade 91) steel header. In: Proceedings of EPRI conference on advanced materials, Paper 4B04; 2010. [7] Foldyna V, Kubon Z. Consideration of the role of Nb, Al and trace elements in creep resistance and embrittlement susceptibility of 9-12% Cr steel. In: Conference proceedings, York, UK; 1994. p. 175e87. [8] Parker JD, Parsons AWJ. High temperature deformation and fracture processes in 2 1/4Cr1Mo-1/2Cr 1/2Mo 1/4V weldments. International Journal of Pressure Vessel and Piping 1995;63(1):45e54. [9] Patterson S, Geanaldi R, Cronin M. Early service life cracking of steam piping welds in combined cycle power plants. In: Proceedings of the EPRI Conference on Boiler Assessment; 2010. [10] Alberry PJ, Jones WKC. Diagram for the prediction of weld heat affected zone microstructure. Materials Technology 1977;4:360e4. [11] Alberry PJ, Jones WKC. Comparison of mechanical properties of 2Cr-Mo and 0.5Cr-Mo-V simulated heat affected zones. Metals Technology 1977;4: 45e51.
[12] Alberry PJ, Brunnstrom RRL, Jones KE. Computer model for predicting heat affected zone structures in mechanized tungsten inert gas weld deposits. Metals Technology 1983;10:28e38. [13] King JF, Sikka VK, Santella ML, Turner JF, Pickering EW. Weldability of modified 9Cr-1Mo steel. ORNL-6299. Oak Ridge National Laboratory; September 1986. [14] Abe F. Critical issues for development of high-Cr ferritic steels for USC power plant at 650 C. In: Proceedings of CREEP8: eighth international conference on creep and fatigue at elevated temperatures July 22e26, 2007; 2007. San Antonio, Texas; ASME Paper CREEP 2007 e 26255. [15] Cerjak H, Holzer I, Mayr P, Pein C, Sonderegger B, Kozeschnik E. Application of a comprehensive R&D concept to improve long-term creep behaviour of martensitic 9-12% Cr steels. In: Proceedings of the Fifth International Conference on Advances in materials technology for fossil power plants, October 2007. Marco Island, Florida: Electric Power Research Institute; 2007. Paper 4B e 07. [16] Li Y, Hongo H, Tabuchi M, Takahashi Y, Monma Y. Evaluation of creep damage in heat affected zone of thick welded joint for Mod.9Cr-1Mo steel. International Journal of Pressure Vessel and Piping 2009;86:585e92. [17] Creep strength enhanced ferritic (CSEF) steel welding guide. Palo Alto, CA: EPRI; 2011. 1024713. [18] Yaguchi M, Matsumura T, Hoshino K. Evaluation Of Long-term Creep Strength of Welded Joints of ASME Grades 91, 92 and 122 type steels. In: Proceedings of the ASME 2012 Pressure Vessels & Piping Conference, Paper PVP; 2012. 78393. [19] Bhadeshia HKDH. Effect of materials and processing: material factors. Handbook of residual stress and deformation of steel. Ohio, USA: ASM International; 2002. [20] Jones WKC, Alberry PJ. Ferritic steels for fast reactor steam generators. British Nuclear Engineering Society; 1977. p. 1e4 and Metals Technology 11: 557e566. [21] Jones WKC, Alberry PJ. Residual stresses in welded constructions. The Welding Institute; 1977. Paper 2. [22] Deng D, Murakawa H. Prediction of welding residual stress in multi-pass buttwelded modified 9Cre1Mo steel pipe considering phase transformation effects. Computational Materials Science 2006;37:209e19. [23] Francis JA, Bhadeshia HKDH, Withers PJ. Welding residual stresses in ferritic power plant steels. Materials Science and Technology 2007;23(9):1009e20. [24] Shirzadi AA, Bhadeshia HKDH, Karlsson L, Withers PJ. Stainless steel weld metal designed to mitigate residual stresses. Science and Technology of Welding and Joining 2009;14(6). [25] Payares-Asprino MC, Katsumoto H, Liu S. Effect of martensite start and finish temperature on residual stress development in structural steel welds. Welding Journal 2008;87:279e89.
Please cite this article in press as: Parker J, In-service behaviour of creep strength enhanced ferritic steels Grade 91 and Grade 92 e Part 2 weld issues, International Journal of Pressure Vessels and Piping (2012), http://dx.doi.org/10.1016/j.ijpvp.2012.11.004
Contents lists available at SciVerse ScienceDirect
International Journal of Pressure Vessels and Piping journal homepage: www.elsevier.com/locate/ijpvp
In-service behaviour of creep strength enhanced ferritic steels Grade 91 and Grade 92 e Part 2 weld issues Jonathan Parker EPRI, 1300 West WT Harris Blvd, Charlotte, NC 28262, USA
a r t i c l e i n f o
a b s t r a c t
Article history: Received 14 September 2012 Received in revised form 26 November 2012 Accepted 27 November 2012
In Creep Strength Enhanced Ferritic steels control of both composition and heat treatment of the parent steel is necessary to avoid producing components which have properties below the minimum expected by applicable codes. The degree of tempering involved in manufacture will modify the material hardness. While under most conditions hardness is reduced by tempering, exceeding the AC1 temperature can lead to an increase in hardness. In this heat treatment the properties will be relatively poor even though the measured hardness may be apparently acceptable. Thus, care should be exercised in imposing an acceptance test of components based on simple hardness alone. Differences in parent material heat treatment and composition apparently have remarkably little influence on the creep life of the heat affect zone (HAZ). Thus, Type IV cracking in the fine grained or intercritically heat treated regions of the HAZ does not appear to directly depend on the strength of the base steel. This form of in-service damage is relatively difficult to detect using traditional methods of non-destructive testing. Moreover, since repeated heat treatment leads to over tempering and a degradation of properties, specific procedures for making and then lifing repair welds are required. The present paper summarizes examples of damage and discusses best option repairs. Ó 2013 Elsevier Ltd. All rights reserved.
Keywords: Creep Damage Weld Service performance Steel Grade 91
1. Introduction Grade 91 steel is one member of the family of creep strength enhanced ferritic steels (CSEF) [1]. This steel is increasingly considered the material of choice for boiler, piping and header applications. Thus, this steel is being routinely installed in both fossil-fuel fired and combined cycle power generating units. The inservice performance of components manufactured from CSEF steels is reviewed; Paper I considered behaviour of parent with this paper focussing on welds. There are design and/or construction practises which result in local stress concentrations in creep-weak regions such as the heat affected zones (HAZ) in weldments. These issues can be a consequence of the approaches used for reinforcement of penetrations or because of the influences of other local stress concentrations. However, because welds will introduce significant complexity of microstructure and geometry it is apparent that welds are frequently the locations most susceptible to damage. Experience shows that the following locations are particularly susceptible to problems:
Non 90 branches (offset branches, welded “Y” pieces) End caps (fully or partially welded) Seam welds (whether re-normalized or not) Thermocouple pockets (specific arrangements) Dissimilar metal joints
A recent EPRI project has successfully established the background technology necessary to underpin significant improvements in approaches for ensuring that high quality components are supplied and installed [2]. In addition, this work delivered a life management strategy which is both practical and effective in preventing in-service failures. However, there are large numbers of components currently in operation which are susceptible to cracking. In the majority of cases the defects formed will need to be repaired using fusion welding techniques. The present paper briefly provides examples of highlights key factors which have been linked to in-service cracking and discusses challenges associated with repair and refurbishment. 2. Examples of service problems
Repair welds Terminal welds and branch welds E-mail address: [email protected].
A summary of typical damage found in welds manufactured in Grade 91 components has been presented in a recent EPRI report [3]. In the absence of gross problems associated with original weld
0308-0161/$ e see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.ijpvp.2012.11.004
Please cite this article in press as: Parker J, In-service behaviour of creep strength enhanced ferritic steels Grade 91 and Grade 92 e Part 2 weld issues, International Journal of Pressure Vessels and Piping (2012), http://dx.doi.org/10.1016/j.ijpvp.2012.11.004
2
J. Parker / International Journal of Pressure Vessels and Piping xxx (2012) 1e12
fabrication and Post Weld Heat Treatment, most reported weld related defects have been found in the Weld Heat Affected Zone (HAZ). P91 is known to be prone to in-service cracking known as Type IV. This damage has been found relatively early in life, particularly at locations where there is strain localization. The strain concentration may be a consequence of piping system stresses or local geometry or both. Type IV cracking thus typically describes creep-induced damage that has initiated in, and propagated through, the continuous fine-grained/intercritically heated zone of material. The creep strength of this zone of material is compromised because of changes that occur in response to the particular range of temperatures which occur during the welding cycle. Additional complications can be found when dissimilar metal welds are used. The transitions in properties associated with these welds are particularly complicated when the component dimensions, operational conditions and weld location serve to concentrate the stresses. Longitudinal seam welds in high temperature piping systems have also experienced damage in the weld metal and in the HAZ. The following summarizes in-service damage detected at different types of component. In general, the common factor associated with the presence of cracking is the design which is less than optimal for the prevailing operating conditions. Specific options for improved design are discussed. 2.1. Issues with nozzles Boiler and piping systems require that connections be manufactured between different pressurized components. These connections may be associated with a small bore pipe joined to a large bore, as in a drain line from a pipe, or may involve joining cylindrical components of similar dimensions, as with ‘Tees’ or Manifolds. In most cases, damage found at small bore connections is associated with thermal fatigue. The fatigue damage is often induced by pressure changes which cause relatively cold water to ‘quench’ a hot component. Many of these failures show a typical ‘star burst’ pattern, with transgranular cracks developed as a consequence of the repeated severe transients, Fig. 1. In general, inspection methods are available to detect these defects before component failure occurs. Since the damage is driven by repeated thermal cycling mitigation of this type of issue relates to removal of the condensate from the lines. Potential solutions involve improved gradients, increasing the size of the drain lines and introduction of thermocouples to detect the presence of water. Problems related directly to high temperature creep have been found in tube to header connections and in larger section nozzles. A
major investigation of damage in a replacement superheater outlet header has been reported [4,5]. In this case the damage found was at least in part explained by excessive strain in the header associated with the hoop stress. There was also at least a potential relationship between the in-service damage and the level of Aluminium present in the parent steel. Thus, because the measured Aluminium levels were high, and the associated Nitrogen to Aluminium ratios (N:Al) were relatively low, it was considered that the creep strength of some sections of the parent steel was at the bottom of the creep strength scatter band [5]. This relatively low parent strength was believed to be one factor in why so many of the connections developed cracking, Fig. 2a. The damage was initially detected by an inspection after around 50,000 h of operation. At that time surface defects were removed by grinding, Fig. 2b, but because of the levels of damage present the header was replaced [6]. Examination of a wider range of components has found that further creep damaged Grade 91 welds have occurred in-service in situations where the parent exhibited low N:Al ratios [5]. It is apparent that the potential problems caused by relatively high Al levels will be not be limited to influences on creep strength. It has been clearly demonstrated that high levels of fine inclusions in the steel can reduce ductility. Indeed, reductions in ductility with relatively low N:Al ratios have been reported [7]. This reduction occurs, particularly in steels with a fine grain size, because the large numbers of inclusions on grain boundaries act to initiate creep voids and cracks [8]. High levels of Al will therefore have the potential to promote Type IV cracking both by reducing the strength and by facilitating damage nucleation. In more general cases of cracking in the HAZ of nozzle welds, the problems can be linked at least in part to design approaches associated with reinforcement of the connection. It is normally the case that compensation must be used to reinforce an opening. However, even within a single design code there are several options for how the opening is reinforced. These include: 1. Adding thickness to the branch alone, i.e. the branch has a wall thickness above code minimum requirements, 2. Adding thickness to the main alone, i.e. the main has a wall thickness above code minimum requirements, 3. Or some combination where both the main and the branch are above minimum wall thickness requirements. The specific dimensions being based on detailed analysis to minimize stresses. A typical example of manufacture of a nozzle with a thickened branch is shown in Fig. 3. Clearly this type of Tee involving
Fig. 1. Magnetic particle inspection highlighting thermal fatigue cracking associated with a small bore attachment to a steam line (a) and optical micrograph showing the character of the defect revealed by laboratory metallographic examination (b).
Please cite this article in press as: Parker J, In-service behaviour of creep strength enhanced ferritic steels Grade 91 and Grade 92 e Part 2 weld issues, International Journal of Pressure Vessels and Piping (2012), http://dx.doi.org/10.1016/j.ijpvp.2012.11.004
J. Parker / International Journal of Pressure Vessels and Piping xxx (2012) 1e12
a
90
3
b
80 70
Numbers
60 50
Total Stubs
40
Cracked Stubs
30 20 10 0
1
2
3
4
5
6
Decreasing N:Al Ratio >>> Fig. 2. Histogram showing the relationship between incidence of in-service cracking and N:Al ratio for damage in a superheater outlet header (a) and a photograph showing the local grinding performed to remove the defects (b) [6].
Fig. 3. Typical connection on a main steam type component showing the reinforced branch before and after completion of welding.
significant weld deposition may not be optimal from a local stress situation. However, fabrication using welded construction often offers lower upfront cost than production of a forged component. There are many examples where connections of this
type have shown creep cracking, see for example Fig. 4. The detail in Fig. 4b shows that the local, predominantly intergranular local damage develops until micro and macro-cracks are formed.
Fig. 4. An example of in-service Type IV cracking at a reinforced connection, note that when this damage is related to creep strain in the main component the damage is typically primarily in the HAZ on that side of the weld (a), the cracking forms after the nucleation and development of individual creep voids (b) [6].
Please cite this article in press as: Parker J, In-service behaviour of creep strength enhanced ferritic steels Grade 91 and Grade 92 e Part 2 weld issues, International Journal of Pressure Vessels and Piping (2012), http://dx.doi.org/10.1016/j.ijpvp.2012.11.004
4
J. Parker / International Journal of Pressure Vessels and Piping xxx (2012) 1e12
2.2. Issues with wyes In an apparent attempt to avoid the extra fabrication cost associated with forged laterals a number of architect engineering firms specified the use of welded laterals for steam piping splits. These piping laterals were made using either craftspersonprepared sections of piping (to create the lateral connection and required reinforcement) or more commonly were made using preengineered/fabricated lateral fittings for the branch side of the connection. Fig. 5 provides an example showing the common crack prone regions of these saddle welds. It is typical to find cracks that are initiating from the weld root at the heel of the saddle weld and propagate along the side of the saddle welds in the base metal heat affected zone. Cracking in welded P91 steam piping laterals has occurred in many instances after approximately 35,000 h service, although as with all creep processes the specific timing of damage will depend on specific operating conditions. Detailed post service examination has revealed that the damage mechanism is predominantly Type IV cracking. Most, if not all of the problems would be overcome with a design involving balanced reinforcement. Thus, even if a welded construction is required, by using a main pipe that is thickened compared to code minimum wall, will greatly reduce the creep strain concentration. The thicker section could easily be forged or machined down near its ends, with a generous thickness taper, so that the ends at the girth welds would be matched with the adjoining pipe. This design concept would result in a welded P91 lateral producing equal or better performance than a forged F91 lateral. Of course care must always be used when increasing the wall thickness of components which will be thermally cycled. In general, cyclic loading will increase as the wall thickness increases. The message then is that welded steam piping connections need to be carefully designed to compensate for the risk of excessive creep in the base metal heat affected zone region.
reported in other end cap welds on a wider range of power plants. The high susceptibility to damage is a consequence of several factors; the flat end caps are relatively thick and thus constrain expansion of the header. This constraint can lead to high stress at the corner of the end fitting. Putting a relatively weak HAZ at a location of high stress then is less than optimal. These locations are increasingly being earmarked for onsite component inspection using advanced techniques such as Linear Array UT. 2.4. Issues with dissimilar steel welds Transition welds in boiler components are used in a very wide range of applications; these can be associated with changes in Geometry, changes from one ferritic type steel to another as well as joints which can be fabricated between ferritic and austenitic steels. There have been well documented studies which detail problems associated with these welds. In general problems are encountered where a comprehensive engineering design and procedure validation have not been performed. This type of evaluation should consider factors such as:
Joint design Type and thickness of weld metal Recommended weld process(es) and procedures Proper PWHT conditions for combinations of base metals/filler metals Metallurgical complexities related to the as fabricated microstructure and time/temperature dependent affects, e.g. Carbon Migration Initial quality assurance testing both for acceptance of a defect free structure and as a base line for follow up assessment of damage Future performance and damage evaluation Systems loading (Requirements for use of Cold Pull) Issues associated with Cyclic Operation
2.3. Issues with end caps and plates A series of plant failures have been reported associated with end caps and end plates in Grade 91 components operating at high temperature. One of the first which occurred, after about 20,000 h in-service, was attributed to incorrect heat treatment. Others in header end plate welds occurred after 36,500 h (with 469 hot and 72 cold starts). Damage in both cases was identified as Type IV cracking, Fig. 6 [4]. Nominal operating conditions for the headers were 568 C (1055 F), and 160 bar, and the failed weld had been inspected using manual UT and found clear of indications only 8663 h prior to failure. Following the failure, high sensitivity UT was carried out on three other Grade 91 headers on the affected unit and the equivalent headers on another unit. This led to the discovery of further damaged end plate welds. Failures have been
In general, the factors affecting the life of transition welds are understood. However, it is apparent from the examples given below that the understanding which exists is not always being applied to design and fabrication of components. Because this type of joint will continue to be necessary within pressure boundary components it is important to apply this understanding to new equipment and when remedial action is needed on existing plant. Potential opportunities for additional ways to solve problems continue to become available; these include the development and validation of new weld consumables as well as improvements to optimization of manufacturing process control. There are often requirements to join Grade 91 type steels to low alloy components. This type of joint creates questions regarding the consumables (normally a joint will use weld consumables which
Fig. 5. Cracking typically seen associated with saddle welded laterolets used in hot reheat systems. Significant Type IV cracking has been observed after around 35,000 h of service.
Please cite this article in press as: Parker J, In-service behaviour of creep strength enhanced ferritic steels Grade 91 and Grade 92 e Part 2 weld issues, International Journal of Pressure Vessels and Piping (2012), http://dx.doi.org/10.1016/j.ijpvp.2012.11.004
J. Parker / International Journal of Pressure Vessels and Piping xxx (2012) 1e12
5
Fig. 6. Schematic diagram showing the header end cap geometry and the crack location (a) and a metallographic section showing detail of the Type IV cracking (b) [4].
match the parent but when there are two different parents in a joint it is impossible to match both), the requirements for post weld heat treatment and the geometry. In many cases the best engineering solution is to fabricate a transition piece which will accommodate the changes in materials and geometry while minimizing fabrication problems. However, it is frequently the case in plant that these transitions have both changes in alloy and geometry at the same location. As shown in Fig. 7, when a weld between low alloy steel and P91 is fabricated with a Grade 91 type consumable care must be exercised over joint details. In the present example, the desire to ‘match’ thickness at the weld resulted in over stress conditions in the low alloy steel and relatively short term creep failure. Problems when the constituents are metallurgical very different are even more complicated. There have been a number of early service life failures of high energy steam piping dissimilar metal welds in combined cycle plants. Particular issues have been reported in joints which involve dissimilar metal girth welds made between P91 piping and austenitic stainless steel flow nozzles. These joints generally involve some combination of ENiCrFe-3 (Inconel 182) and ERNiCr-3 (Inconel 82) weld consumables [9]. Cracking has been found in relatively simple full thickness welds, and in joints made after application of Inconel 182 butter passes on the P91 side of the weld joint. The buttered layer is typically subjected to a shop post weld heat treatment before completion of the final weld. No additional PWHT is normally applied to these joints after completion of the weld fill passes. In other cases the entire final weld has been post weld heat treated. An example of an inservice crack is shown in Fig. 8.
A further complicating factor in assessing these components is the fact that almost no two flow elements have identical design and construction details. However, there are two broadly similar approaches used. Either the flow element is a short section of austenitic steel with an appropriate geometry (see Fig. 9) or the flow element is an austenitic steel insert within the P91 pipe spools. In the either case, an Inconel weld consumable is used to join the P91 to austenitic stainless steel flow element. Failures (leaks, severed pipe connections) of these welded connections have occurred after only 20,000 to 40,000 operating hours. 3. Factors affecting weld performance As far as the parent steel is concerned there is clearly a need to ensure that both the composition and all aspects of the heat treatment are performed in accordance with the latest standards. However, since in boilers there is a requirement to weld sections of parent, it is important that some consideration is given here to the effects of parent composition and heat treatment on the creep performance of welds. It is now well established that for subcritically heat treated welds the primary concern for long term creep life is Type IV cracking. The microstructural regions within a bead on plate weld were fully characterized by work performed at Marchwood Engineering laboratories by Alberry, Jones and co-workers [10e12]. The primary regions present have been represented based on the specific thermal cycles developed during welding, Fig. 10. The effect of the specific thermal cycles on final microstructure will depend on the
Fig. 7. In-service damage developed in a dissimilar weld between a low alloy steel casing and a Grade 91 pipe, a schematic diagram is presented in (a) with a photograph of the weld and the defect shown in (b). The cracking developed as a result of creep damage in the low alloy steel HAZ, (c).
Please cite this article in press as: Parker J, In-service behaviour of creep strength enhanced ferritic steels Grade 91 and Grade 92 e Part 2 weld issues, International Journal of Pressure Vessels and Piping (2012), http://dx.doi.org/10.1016/j.ijpvp.2012.11.004
6
J. Parker / International Journal of Pressure Vessels and Piping xxx (2012) 1e12
Fig. 8. Photograph showing a flow nozzle arrangement where the Austenitic steel flow element had been welded into the P91 piping system using Inconel filler. The inset picture shows damage development at the fusion boundary.
Fig. 9. Schematic diagram of an Austenitic Stainless Steel Flow nozzle welded into a Grade 91 steel pipe using Inconel 182 welds. Cracking was detected on the Grade 91 side of one of the welds [9].
details of the metallurgic system being considered. However, the general features in relation to the microstructures of martensitic steels are as follows: (i) Coarse grain region (CGHAZ): Material near the fusion boundary that reaches a temperature well above AC3 during welding. Precipitates which constitute the main obstacle to
Fig. 10. Schematic diagram illustrating how the thermal cycles during manufacture of a simple bead on plate type weld will influence the microstructure resulting in the formation of a Heat Affected Zone of complex gradients in metallurgical structure.
growth of the austenite grains must be dissolved, so that once austenite is formed growth of austenite grains is relatively rapid. In the 9e12 Cr steels, the region containing relatively coarse grained austenite transforms into martensite on cooling under typical conditions, Fig. 11a. (ii) Fine grain region (FGHAZ): As the distance from the fusion line increases the peak temperature achieved is above AC3 so that complete transformation to austenite occurs but this temperature is not sufficient to dissolve all of the precipitates present. Thus, austenite grain growth is limited by the incomplete dissolution of carbides, nitrides or carbo-nitrides. The fine grained austenite produced will mostly transform to martensite in the 9e12 Cr steels under typical cooling. However, in locations where the prior austenite grain size is very small the ability of these very fine grains to undergo a fully shear type transformation may be limited, Fig. 11b. (iii) Intercritical region (ICHAZ): There will be a region where the peak temperature exceeds the AC1 but not the AC3. Thus there will be some new austenite formed during the heating cycle. Moreover, some of the precipitates may dissolve but the combination of time and temperature is insufficient for complete dissolution of precipitates. There will thus be only a partial reversion to austenite on heating. This new austenite will predominantly transform to martensite on cooling, but it should be emphasized that the description ‘intercritical’ can be used to describe a wide range of conditions. Moreover, there will likely be some variation of specific microstructural features within an ‘intercritical’ region, Fig. 11c. (iv) Highly tempered region: There will be a region adjacent to the parent where the time; temperature combination will have
Please cite this article in press as: Parker J, In-service behaviour of creep strength enhanced ferritic steels Grade 91 and Grade 92 e Part 2 weld issues, International Journal of Pressure Vessels and Piping (2012), http://dx.doi.org/10.1016/j.ijpvp.2012.11.004
J. Parker / International Journal of Pressure Vessels and Piping xxx (2012) 1e12
7
Fig. 11. General appearance of the microstructure across the HAZ of a Grade 91 type steel weld, with optical micrographs detailing the structures in the relatively coarse grained region (a), the fine grained region (b) and part of the HAZ where incomplete transformation took place (c). This region is typically referred to as intercritical because the peak thermal cycles are between AC1 and AC3.
modified the local substructure. In CSEF steels this modification may relate to changes in precipitates or the locally high dislocation density or both. In either case these changes will likely not be easily resolved using optical metallography and will only be determined using advanced electron optics. Of course with multipass welds the relatively simple situation represented by bead on plate is significantly more complicated. Thus, depending on factors such as heat input, electrode size and angle, degree of overlap etc there can be regions of the parent which are exposed to multiple thermal cycles. Moreover, regions of deposited weld will also be thermally influenced and these locations greatly increase the possible variations. The final process associated with weld fabrication is normally PWHT. In 9e12% Cr steels this heat treatment tempers the new martensite introduced by the welding thermal cycles. However, gradients in the microstructure and mechanical properties, extending typically over a few millimetres from the fusion boundary are likely to be present even after this treatment. As the extent of the tempering during PWHT increases there is the potential that the transitions in properties will also be reduced. An example of the increase in tempering during PWHT on microhardness of a Grade 91 weld is illustrated in Fig. 12 [13]. As shown, the lowest degree of tempering, 2 h at 732 C, shows the highest hardness in the weld metal and the lowest hardness in the HAZ. As the level of tempering increases the hardness of the weld and parent both exhibit reductions. However, even at the highest degree of tempering, 40 h at 732 C the lowest hardness values continue to be measured in the HAZ. It is apparent from cross weld creep testing of CSEF steels that for conditions approaching in-service stresses the time to Type IV cracking falls below that of the parent [14]. This behaviour is illustrated for Grade 92 steel in Fig. 13. It is apparent that all CSEF steels are highly susceptible to Type IV cracking. This susceptibility comes about since while most of the component and weld have
excellent strength (conferred by the well controlled heat treatment which avoids deleterious effects) there are regions in the HAZ where the thermal cycles due to welding result in microstructures with low creep strength. Simply put, there are regions in the HAZ with a very fine grain size and no martensitic microstructure. Furthermore, it has also been has shown [14]that M23C6 precipitates and Laves phases form faster in the fine grain HAZ region in 9Cr martensitic type steels compared with the other regions of the weldment. This metallurgical effect further increases the vulnerability of the Type IV region. Since not only are matrix-strengthening elements such as Cr, Mo and W depleted but the Laves phase offers potential sites for the nucleation of creep voids. There are many other examples of Type IV cracking in other CSEF steels. The pictures in Fig. 14 are from a major European
Fig. 12. Microhardness traverse across a SA weldment produced at ORNL. Both the base plate and filler wire were of the modified 9Cr-lMo composition. Post weld heat treatments were for the times shown at 732 C [13].
Please cite this article in press as: Parker J, In-service behaviour of creep strength enhanced ferritic steels Grade 91 and Grade 92 e Part 2 weld issues, International Journal of Pressure Vessels and Piping (2012), http://dx.doi.org/10.1016/j.ijpvp.2012.11.004
8
J. Parker / International Journal of Pressure Vessels and Piping xxx (2012) 1e12
Fig. 13. Comparison of the Creep Rupture behaviour of Grade 92 steel parent and the results of Cross Weld Creep tests at 650 C [based on data from Ref. [14], micrograph from Ref. [4]].
Programme studying the behaviour of E911 [15]. These investigations again show that the most susceptible region of the HAZ is the fine grained location. The failure occurred where the HAZ microstructure consisted of very small equiaxed prior austenite grains where the grain boundaries were decorated with large precipitates. No martensitic lath structure was observed within this region. Creep voids mainly formed along prior austenite grain boundaries. Final fracture occurred by interconnection of single creep voids to form micro-cracks, which themselves finally coalesced to form macro-cracks. As shown in Fig. 14, high densities of creep voids are developed over the HAZ, with crack formation and final propagation occurring only very late in creep life.
There has been some disagreement over details of the formation of creep voids associated with Type IV cracking. In general, where loading is reasonably uniform across the section, i.e. in the absence of significant bending loads, it is now clear that creep voids nucleate relatively early in the creep life. These voids are initiated at locations below the component surface so that even a properly performed examination at the surface of a weld will often under estimate the damage present. These observations have been seen in both laboratory cross weld experiments and in post service examination of pressure boundary welds. The results shown in Figs. 14 and 15 are typical of laboratory generated information. When a relatively large specimen is tested under the appropriate conditions damage will typically be developed in both the HAZs within the gauge length. Thus, preparation of a sample allows assessment of both the fractured HAZ and the interface where damage levels are relatively close to fracture. This situation is demonstrated in Fig. 15. The HAZ to the right of the weld has failed in a classic Type IV manner and Microdamage is clearly present in the left hand HAZ. Detailed optical metallography, Fig. 15a, reveals that individual voids were associated with the damage at both HAZs. In addition to variations in metallographic preparation techniques a wide variety of methods have been used to examine and quantify the creep voids present. It is apparent that even with careful preparation there will be variation in the number of voids detected depending on the magnification used for examination. As a general rule the number density of voids will increase with the magnification used at least in the range 200 times to around 2000 times. There have been several comprehensive investigations of void development reported. Fig. 16 shows post test examination results from a set of long term Grade 91 cross weld tests which
Fig. 14. Illustration of the Type IV damage development in the CSEF steel Grade E911 in Cross Weld Creep tests performed at 600 C [15]. The examination using Electron Backscatter Diffraction allows high resolution, detailed mapping of features such as grains, subgrains and martensite laths.
Please cite this article in press as: Parker J, In-service behaviour of creep strength enhanced ferritic steels Grade 91 and Grade 92 e Part 2 weld issues, International Journal of Pressure Vessels and Piping (2012), http://dx.doi.org/10.1016/j.ijpvp.2012.11.004
J. Parker / International Journal of Pressure Vessels and Piping xxx (2012) 1e12
9
Fig. 15. Details of creep voids formed during a cross weld creep test which was interrupted prior to failure (a), and a montage showing details of Type IV cracking (b). The right hand HAZ shows through section cracking with the left hand HAZ showing microcracking.
Fig. 16. Post test examination showing the development of creep voids in long term cross weld creep tests with increasing life fraction [16].
were interrupted at known life fractions. It is clear there is a relatively high density of creep voids present in the Type IV region after a creep exposure of approximately 70%. Even at this level of creep life fraction void damage does not appear to extend to the sample surface and there is no clear evidence of cracking. This is only apparent after around 90% of life. The data presented in Fig. 17 compare the creep lives observed in relatively low strength parent, the data from bar 257, with results obtained from cross weld creep testing. It is apparent the cross weld lives observed in the bar 257 steel fall well below the expected scatter band for parent steel. Consideration of the behaviour of a second steel, identified as bar 817, with typical parent behaviour indicates that the cross weld life is similar to that of the welds in bar 257. It thus appears that variations in chemical composition and heat treatment have a greater effect on the creep strength of the parent than they do in the HAZ. The high aluminium, weak Grade 91 material bar 257, which substantially underperforms mean P91 parent creep data, showed a much lesser shortfall in its HAZ
Fig. 17. Comparison of the creep rupture lives of cross weld samples which failed by Type IV cracking in the HAZ with Parent behaviour, all data are for Grade 91 steels tested at 600 C. [Based on data from Ref. [4]].
Please cite this article in press as: Parker J, In-service behaviour of creep strength enhanced ferritic steels Grade 91 and Grade 92 e Part 2 weld issues, International Journal of Pressure Vessels and Piping (2012), http://dx.doi.org/10.1016/j.ijpvp.2012.11.004
10
J. Parker / International Journal of Pressure Vessels and Piping xxx (2012) 1e12
performance. It may be that high aluminium, which is known to combine with nitrogen and suppress the formation of key creep strengthening carbonitride precipitates, has a more harmful effect on parent properties than on those of the HAZ. It could therefore be that, whereas optimized precipitation strengthening is the key to optimized parent performance, such strengthening mechanisms become less effective in the non-optimal and lath-free fine grained HAZ microstructure. Other factors, for example simple solution strengthening, may perhaps become more important in determining HAZ performance. Moreover, as discussed later there could also be an influence on Type IV cracking susceptibility due to welding residual stresses. 4. Discussion It is apparent that welding needs to be carried out carefully using qualified staff and appropriate procedures, further background is provided in reference [17]. Assessment of the creep performance of welds is then frequently studied by programmes which include cross weld testing. In these programmes it is important to select test conditions and specimen geometries that produce failures where the damage mechanisms are relevant to long terms service. In general, it is now widely accepted that in creep tests at relatively high stress and temperature, the results are not typical of long term damage in component welds. Thus, results from tests under these conditions should not be used as a guide to in-service behaviour. The concept of a weld strength reduction factor (WSRF) has been introduced in design codes for high temperature nuclear power and for process piping. It is defined as a ratio of the stresses causing rupture at the same time in welded joint and base metal. Although no explicit employment has been made in design codes for non-nuclear plants so far, the introduction of a WSRF is regarded necessary to avoid failure at welded joints. In general, the procedure involves obtaining a welded joint “allowable stress” based on cross weld creep rupture data and then calculating WSRF as the ratio of this value to the allowable stress for parent [18]. Thus, the average stress value and 95% lower value of 105 h rupture stress were calculated at each temperature using the equations for lower stress region. Then, the welded joint “allowable stress” was given as minimum value of (the average 0.67) and (95% lower value 0.8). The calculated WSRF for Grade 91, Grade 92 and Grade 122 are listed in Table 1 [18]. Residual stresses are often introduced unintentionally during fabricationdfor example, during welding or heat treatment. A few elegant experiments illustrate how phase transformations interact with the buildup of residual stress with the results summarized in Fig. 18 [19]. Using bainitic, martensitic, and stable austenitic steels, Jones and Alberry [20,21] demonstrated that transformation plasticity during the cooling of a uniaxially constrained sample from the austenite phase field acts to relieve the buildup of thermal stress as
Fig. 18. Interpretation of experimental data showing how residual stresses develop on cooling for an austenitic steel (no transformation), a bainitic low alloy steel (relatively high temperature of transformation) and a martensitic steel (relatively low temperature of transformation). From Ref. [19] based on data from Refs. [20] and [21].
the sample cools. By contrast, the non transforming austenitic steel exhibited a continuous increase in residual stress with decreasing temperature, as might be expected from the thermal contraction of a constrained sample. When the steels were transformed to bainite or martensite, the transformation strain compensated for any thermal contraction strains that arose during cooling. Significant residual stresses were therefore found to build upbuildup only after transformation was completed and the specimens approached ambient temperature [22,23]. There have been considerable efforts to create welding consumables which on solid state phase transformation partly compensate for the stresses which develop when a constrained weld cools to ambient temperatures [24,25]. Many of these studies have considered on structural steels which are ferritic. There is however, some work reported which has used alloy design methods to create a stainless steel welding consumable which solidifies as ferrite, transforms almost entirely into austenite which then undergoes martensitic transformation at a low temperature of about 220 C. Based on consideration of the effect of martensitic transformation on welding residual stresses there now could be a further reason why 9 to 12%Cr CSEF steels are universally susceptible to Type IV cracking. As shown earlier, there will be a band within the HAZ where martensitic transformation does not take place. This could be in the fine grained location or where peak temperatures were in the intercritical region. In either case it appears that the lack of a martensitic transformation will lead to locally high residual stresses. These stresses will relax by strain, however, since martensitic steels are strain softening these relaxation processes could lead to further reductions in strength. In any event it is likely that the relaxation strain would be relatively high so that there would be effective reduction in subsequent creep life.
Table 1 Evaluation of weld strength reduction factors for Grade 91, Grade 92 and Grade 122 steels on the basis of a recent reevaluation of the results of long term creep testing [18]. Material
Gr.91
Gr.122
Gr.92
Weld strength reduction factor
KA-SCMV28 KA-STPA28 KA-SFVAF28 KA-SUS410J3 KA-SUS410J3TP KA-SUSF410J3 KA-STPA29 KA-SFVAF29
Note
525 C
550 C
575 C
600 C
625 C
650 C
1.00 1.00 1.00 1.00 1.00
0.90 0.90 0.90 0.84 0.84
0.82 0.74 0.74 0.68 0.60
0.79 0.67 0.68 0.57 0.50
0.79 0.65 0.65 0.50 0.50
0.79 0.65 0.65 0.50 0.50
1.00
1.00
0.74
0.62
0.53
0.53
<¼76 mm >76 mm
Please cite this article in press as: Parker J, In-service behaviour of creep strength enhanced ferritic steels Grade 91 and Grade 92 e Part 2 weld issues, International Journal of Pressure Vessels and Piping (2012), http://dx.doi.org/10.1016/j.ijpvp.2012.11.004
J. Parker / International Journal of Pressure Vessels and Piping xxx (2012) 1e12
11
Table 2 Summary of key variables being considered in the EPRI project assessing the factors affecting the performance of repair welds in Grade 91 steel. Weld
1A, 1B 2A, 2B 3A, 3B 4A, 4B 5A, 5B 6A, 6B 7A, 7B 8A, 8B 9A, 9B 10A, 10B
Base material
As-received Grade 91 (A) Re-normalized Grade 91 (B)
Weld metal AWS Desig.
Trade name
E9015-B9 H4 E9015-B9 H4 E9015-B9 H4 E9015-B9 H4 E8018-B8 E9015-G E9015-G E9018-B3 H4 EPRIP87 ENiCrFe-2
Thermanit Chromo Thermanit Chromo Thermanit Chromo Thermanit Chromo 9Cre1Mo Thermanit P23 Thermanit P23 Bohler E9018-B3 EPRIP87 INCO-WELD A
It is clear that Grade 91 type steels are widely used in boilers and piping. Thus, there will be an increasing need to perform remedial action of high temperature components manufactured from this steel. To optimize the processes associated with remediation, it is critical to establish key information on removal of damage and on approaches for repair welding. Selected variables and weld processes are under investigation in a current EPRI project. This work includes weldment testing linking high temperature performance to details of how the repair was carried out. This aspect is considered important since in contrast to traditional low alloy steels which are relatively ‘user friendly’ as far as repairs are concerned; Grade 91 steels introduce additional complications. Because the properties and performance of the base metal will be critically dependent on the original composition and the full heat treatment history, there have been concerns expressed that the reparability of Grade 91 steel will be limited by the condition of the base metal. Because there are a very large number of variables to be evaluated the project will be performed in two phases. Phase 1 will provide a Ranking of Repair performance. The following will be the primary factors for consideration: Discussion of Methods and Extent of Excavation, Weld Procedure Considerations (including consumables) and Heat Treatment Post Repair evaluation of microstructure, damage etc Specimen Geometry and Testing Conditions including development of test matrix Analysis to identify best option repairs e ‘best option’ to be based on factors such as speed of welding, initial quality, creep life The key considerations for the first Phase are summarized in Table 2. The general use of summary phrases such as “Normal” and “Temperbead” in Table 1 should not be taken to suggest that the details were the same for each weld fabricated. These specifics were customized as necessary for each weld produced. The parent sections where of the same composition. One set of welds, designated A, were manufactured in steel in the ex-service condition with the other set, designated B, in the same sections after full Renormalizing and tempering heat treatment. It is clearly the expectation that by understanding the interaction of thermal cycles, composition and microstructural development the repair methods developed will have the potential for welding without the need for post weld heat treatment. The second Phase of this project will involve analysis of the phase 1 results to identify best option repair methods. The preferred methods will then be applied to sections of an ex-service header which were removed after around 75,000 h of operation. The performance of this header was well documented in several publications (e.g. 6). The published work has been supplemented by extensive post service NDE and selective laboratory metallography. The main tasks
9V 9V 9V 9V
Mod. Mod. Mod. Mod.
Welding procedure
PWHT
Normal þ Rec’d. PWHT Normal þ Min. PWHT Temperbead Poor Practice Temperbead Temperbead Temperbead Normal þ Rec’d. PWHT Temperbead Temperbead Temperbead
1375 25 F/2 h 1260 10 F/2 h None None None None 1375 25 F/2 h None None None
involved in Phase 2 will be associated with agreeing methods and extent of excavation, specifics of the best option repair methods and performance assessment of the manufactured welds. 5. Concluding remarks Grade 91 steel is one of the family of creep strength enhanced ferritic steels (CSEF). This steel is now routinely installed in both fossil-fuel fired and combined cycle units. Recent in-service experience has demonstrated that cracking can occur in Grade 91 steel early in life [1]. The results also show that while differences in material heat treatment prior to welding clearly affect parent behaviour, they have remarkably little influence on HAZ performance. Thus, it is very likely that the primary damage location inservice will be Type IV cracking. Potential problems associated with this form of in-service damage include challenges over detection and appropriate methodologies for repair. Since repeated heat treatment leads to over tempering and a degradation of properties specific procedures for performing and then lifing repair welds are required. The following issues need to be considered when considering performance: 1. It is now widely accepted that in creep tests at relatively high stress and temperature the results of cross weld creep testing are not typical of long term damage in component welds. 2. Clearly then it is important to select test conditions and specimen geometries for laboratory test programs so as to produce failures where the damage mechanisms are relevant to long terms service. 3. Using these conditions it is apparent that failure occurs as a consequence of the nucleation, growth and link up of creep voids. It appears that the damage is significantly greater within the volume of the specimen where relatively high constraint conditions are developed. 4. The Type IV life is significantly below that of the parent under the same conditions. The effects of weld thermal cycle on microstructural are clearly important to the susceptibility for creep damage to occur. However, there is at least inferential evidence that transformation behaviour is effecting the distribution of weld residual stress as well as the local creep strength. The significance of residual stresses to the susceptibility for Type IV cracking in 9e12%Cr steels is one explanation for why the creep strength of the parent has only a second order effect on Type IV performance. Further work is in progress to examine Grade 91 welded samples which have been tested to different creep life fractions. Advanced characterization techniques will be used to establish further details of creep cavity nucleation and growth within the weld HAZ.
Please cite this article in press as: Parker J, In-service behaviour of creep strength enhanced ferritic steels Grade 91 and Grade 92 e Part 2 weld issues, International Journal of Pressure Vessels and Piping (2012), http://dx.doi.org/10.1016/j.ijpvp.2012.11.004
12
J. Parker / International Journal of Pressure Vessels and Piping xxx (2012) 1e12
Acknowledgements Thanks are due to the participants of the recent EPRI project “Life Management of CSEF steels” for much stimulating discussion on all aspects of the fabrication, installation and performance assessment of Grade 91 and Grade 92 steels. References [1] DiStefano JR, Sikka VK. Summary of modified 9Cr-1Mo steel development program, 1975e1985. ORNL-6303. Oak Ridge National Laboratory; October 1986. [2] Guidelines and specifications for high-reliability fossil power plants: best practise guideline for manufacturing and construction of Grade 91 steel components. Palo Alto, CA: EPRI; 2011. 1023199. [3] Service experience with Grade 91 components. Palo Alto, CA: EPRI; 2009. 1018151. [4] Brett SJ, Allen DJ, Buchanan LW. The Type IV creep strength of Grade 91 materials. In: Proceedings of the third international conference on integrity of high temperature welds. Institute of Materials; 2007. p. 409e20. [5] Brett SJ, Bates JS, Thomson RC. Aluminium nitride precipitation in low strength Grade 91 power plant steels. In: Proceedings of the fourth international conference on advances in materials technology for fossil power plants, October 25 to 28. Hilton Head Island, South Carolina: EPRI; 2005. p. 1183e97. [6] Brett SJ. Service experience with a retrofit modified 9Cr (Grade 91) steel header. In: Proceedings of EPRI conference on advanced materials, Paper 4B04; 2010. [7] Foldyna V, Kubon Z. Consideration of the role of Nb, Al and trace elements in creep resistance and embrittlement susceptibility of 9-12% Cr steel. In: Conference proceedings, York, UK; 1994. p. 175e87. [8] Parker JD, Parsons AWJ. High temperature deformation and fracture processes in 2 1/4Cr1Mo-1/2Cr 1/2Mo 1/4V weldments. International Journal of Pressure Vessel and Piping 1995;63(1):45e54. [9] Patterson S, Geanaldi R, Cronin M. Early service life cracking of steam piping welds in combined cycle power plants. In: Proceedings of the EPRI Conference on Boiler Assessment; 2010. [10] Alberry PJ, Jones WKC. Diagram for the prediction of weld heat affected zone microstructure. Materials Technology 1977;4:360e4. [11] Alberry PJ, Jones WKC. Comparison of mechanical properties of 2Cr-Mo and 0.5Cr-Mo-V simulated heat affected zones. Metals Technology 1977;4: 45e51.
[12] Alberry PJ, Brunnstrom RRL, Jones KE. Computer model for predicting heat affected zone structures in mechanized tungsten inert gas weld deposits. Metals Technology 1983;10:28e38. [13] King JF, Sikka VK, Santella ML, Turner JF, Pickering EW. Weldability of modified 9Cr-1Mo steel. ORNL-6299. Oak Ridge National Laboratory; September 1986. [14] Abe F. Critical issues for development of high-Cr ferritic steels for USC power plant at 650 C. In: Proceedings of CREEP8: eighth international conference on creep and fatigue at elevated temperatures July 22e26, 2007; 2007. San Antonio, Texas; ASME Paper CREEP 2007 e 26255. [15] Cerjak H, Holzer I, Mayr P, Pein C, Sonderegger B, Kozeschnik E. Application of a comprehensive R&D concept to improve long-term creep behaviour of martensitic 9-12% Cr steels. In: Proceedings of the Fifth International Conference on Advances in materials technology for fossil power plants, October 2007. Marco Island, Florida: Electric Power Research Institute; 2007. Paper 4B e 07. [16] Li Y, Hongo H, Tabuchi M, Takahashi Y, Monma Y. Evaluation of creep damage in heat affected zone of thick welded joint for Mod.9Cr-1Mo steel. International Journal of Pressure Vessel and Piping 2009;86:585e92. [17] Creep strength enhanced ferritic (CSEF) steel welding guide. Palo Alto, CA: EPRI; 2011. 1024713. [18] Yaguchi M, Matsumura T, Hoshino K. Evaluation Of Long-term Creep Strength of Welded Joints of ASME Grades 91, 92 and 122 type steels. In: Proceedings of the ASME 2012 Pressure Vessels & Piping Conference, Paper PVP; 2012. 78393. [19] Bhadeshia HKDH. Effect of materials and processing: material factors. Handbook of residual stress and deformation of steel. Ohio, USA: ASM International; 2002. [20] Jones WKC, Alberry PJ. Ferritic steels for fast reactor steam generators. British Nuclear Engineering Society; 1977. p. 1e4 and Metals Technology 11: 557e566. [21] Jones WKC, Alberry PJ. Residual stresses in welded constructions. The Welding Institute; 1977. Paper 2. [22] Deng D, Murakawa H. Prediction of welding residual stress in multi-pass buttwelded modified 9Cre1Mo steel pipe considering phase transformation effects. Computational Materials Science 2006;37:209e19. [23] Francis JA, Bhadeshia HKDH, Withers PJ. Welding residual stresses in ferritic power plant steels. Materials Science and Technology 2007;23(9):1009e20. [24] Shirzadi AA, Bhadeshia HKDH, Karlsson L, Withers PJ. Stainless steel weld metal designed to mitigate residual stresses. Science and Technology of Welding and Joining 2009;14(6). [25] Payares-Asprino MC, Katsumoto H, Liu S. Effect of martensite start and finish temperature on residual stress development in structural steel welds. Welding Journal 2008;87:279e89.
Please cite this article in press as: Parker J, In-service behaviour of creep strength enhanced ferritic steels Grade 91 and Grade 92 e Part 2 weld issues, International Journal of Pressure Vessels and Piping (2012), http://dx.doi.org/10.1016/j.ijpvp.2012.11.004