Welding and fabrication of high temperature components for advanced power plant - Part 1

TWI Bulletin, January/February 1999

Trevor Gooch graduated in industrial metallurgy at the University of Birmingham, and continued to obtain the degrees of MSc(Eng) and PhD from the University of London. Dr Gooch joined TWI in 1965. In 1980, Dr Gooch became Head of the Materials Department, embracing the welding characteristics of virtually all metallic materials. In 1994, he was made Materials Technology Manager for TWI, involved with joining characteristics of all materials used for construction.

Within the power generation industry, in the UK and worldwide, many changes have been brought about, partly as a result of widespread deregulation, but primarily because of the drive for improved thermal efficiency and reduced emissions. As Adrienne Barnes, Richard Jones, David Abson and Trevor Gooch report, this has led to material developments, to meet the demands of higher operating temperatures and pressures, while ensuring the continued and safe operation of plant throughout its life.

Part 1 of this article concludes with the identification of a number of specific areas, in terms of welding behaviour, weldment properties, and service performance, where further work is required.

There are ever increasing demands on designers, fabricators and operators, for both economic and environmental reasons, to increase the thermal efficiency of power plant. This has been further promoted by the widespread deregulation within the power generation industry, which has given rise to significant competition among the generators. The net effect is an increase in operating temperatures and pressures, which in turn places higher demands on material and weldment performance to ensure the safe and reliable operation of the plant throughout its life. It is to meet such demands that we see the moves to combined cycle gas turbine plant and 'clean coal' technology.

Many of the independent power producers favour gas turbine based plant in new construction, as it is both cost-effective and can be built relatively quickly, allowing them to compete with the established utilities. Combined heat and power installations are also now an attractive proposition. However, it seems hard for many of us living in the western world to recognise that there is still a very high percentage of the world's population that do not have access to electricity. It is no surprise that there are huge plant construction programmes, either planned or indeed in progress, to meet the increasing supply demands, particularly in countries such as Korea, China and India. Clearly such programmes are constrained by environmental factors, particularly the need to minimise emissions to slow down global warming, but by far the most significant driving factor is cost. Therefore, at least in the short/medium term, efforts in these geographic areas are concentrated on the existing technology of pulverised fuel plant.

A large number of materials, section thicknesses and product forms is used in the construction of power plant, requiring a variety of welding techniques, and each accompanied by different fabrication and service problems. Further, an inevitability of the diversity of materials used within power plant is the frequent need for dissimilar metal joints, involving either ferritic to austenitic combinations or different ferritic alloys.

It is clearly impractical, here to cover all material and joining issues related to high temperature plant, so the principal emphasis will be on the family of Cr-Mo steels. This review is intended to provide a brief historical overview of materials and welding issues associated with both the fabrication and service operation of power plant, through to the present day. In light of problems that have arisen in the past, and the changes occurring in the operating environment, a projected view of future developments in both fabrication practice and the materials used will be presented. The article will broadly be divided into two sections, the first addressing many of the materials issues, and the second covering fabrication practice and welding processes in Part 2.

Metallurgy of low alloy creep resisting steels

The low alloy creep resisting steels typically contain between 0.5 and 9% chromium, to give good corrosion/oxidation resistance, rupture ductility and resistance to graphitisation. This, together with carbon and 0.5-1.0% molybdenum, contribute to the creep strength through secondary hardening. In addition, they contain small additions of carbide-forming elements such as Nb, V and Ti to achieve grain refinement and precipitation strengthening. Such materials are supplied in a number of heat treatment conditions, for example quenched and tempered, normalised and tempered, and annealed.

The microstructure will be determined by both the composition and the heat treatment condition; it can be a ferrite/pearlite or ferrite/bainite mix, fully bainitic, or martensitic, and the mechanical properties will vary accordingly. Ultimately it is the required balance of toughness and strength for a particular application that will govern the heat treatment and particularly the tempering temperature used. Broadly this family of steels can be divided into three categories, viz the Cr-Mo steels, the Cr-Mo-V steels and the modified Cr-Mo steels.

Application of welding to Cr-Mo steels for high temperature applications

In conventional, fossil-fired power plant in the UK, low alloy steels such as 0.5Cr 0.5Mo 0.25V, 1.25Cr-0.5Mo, 2.25Cr-1Mo, to name a few, have been used in a variety of applications, for example header pipes, steam pipework, boiler tubes, and turbine casings. The Cr-Mo-V and Ni-Cr-Mo-V steels have been used in the manufacture of rotors, although high temperature rotors, bolting, and the turbine blades have generally been made from 12%Cr stainless steels, with various Mo, W and V modifications.

The low alloy steels have traditionally been welded using the conventional arc welding processes. Although these materials, with their potentially high hardenability, are susceptible to hydrogen induced cracking, procedures have been developed incorporating appropriate preheat levels and often the use of temper bead techniques, to minimise the risk of such cracking. The occurrence of solidification cracking has generally not been a major problem, although isolated instances have been encountered.

In the following sections, some of the principal fabrication and service problems associated with power plant components will be highlighted. Where appropriate, the remedial action taken will be indicated, in terms of design changes or the use of alternative materials.

Reheat cracking

When a weld is reheated during stress relief or high temperature service, intergranular 'reheat' or 'stress relief' cracking can occur if either the weld metal or coarse grained HAZ have insufficient creep ductility, Fig.1. As illustrated by Middleton et al, ^[1] in such instances the relaxation of welding residual stresses occurs through either grain boundary cavitation, or intergranular cracking, as opposed to plastic flow. The susceptibility to cracking is dependent on the level of restraint, and hence joint geometry, as well as alloy composition (for example 0.5Cr 0.5Mo 0.25V showed a higher susceptibility than 2.25Cr 1Mo steel).

Historically there have been a number of widely reported instances of reheat cracking. A detailed review of reheat cracking was undertaken by Dhooge and Vinckier ^[2] . Their review considered the details of the cracking mechanism and factors affecting the cracking susceptibility. In the 1950s, significant cracking problems were encountered, particularly in high temperature steam piping made from type 347 stainless steel. ^[3] The 1960s brought cracking in 2CrMo and CrMoV steam pipework, and later constructional low alloy steels, including the quenched and tempered grades used in pressure vessel fabrication, were found to be potentially susceptible.

Since the 1970s, the problem has also been encountered in the form of 'underclad cracking' in nuclear pressure vessels. ^[4,5] The reheat cracking problems in low alloy steels largely resulted from the use of weld procedures that led to the development of a high proportion of coarse grained microstructure in the heat-affected zone (HAZ) with little refinement. The development of controlled deposition procedures, giving rise to significant weld metal and HAZ refinement, ^[6-10] together with tighter control of impurity levels, and optimisation of the postweld heat treatment procedures, have virtually eliminated the problem.

Fig.1 Optical micrograph showing reheat cracking in the coarse grained heat affected zone of Cr-Mo-V steel. 

Temper embrittlement

Temper embrittlement is a phenomenon that has plagued industry for many decades, and is primarily associated with tempered commercial alloy steels. A comprehensive and interpretive review of studies on temper embrittlement was carried out by McMahon. ^[11] Temper embrittlement is encountered particularly with the higher alloying levels required to achieve through-section hardening at large section thickness; the larger ingots involved give rise to greater segregation, and require slower cooling through the embrittlement temperature range of ~375-575°C.

Further, embrittlement can also occur during long-term service within this temperature range. Temper embrittlement only occurs when certain impurity elements are present within the steel ^[12] and segregate to the prior-austenite grain boundaries. The principal embrittling elements (in decreasing order of importance) are antimony, phosphorus, tin and arsenic, in trace quantities. In addition, larger quantities of Si and Mn have been found to have an embrittling effect. The susceptibility increases with increased levels of these elements, and exhibits a C-curve time-temperature dependence.

Temper embrittlement is a reversible phenomenon; the toughness of an embrittled material can be restored by heat treatment at a temperature above 600°C. The embrittlement is manifest in a number of ways, viz. the ductile/brittle transition temperature is raised, the fracture path is intergranular with respect to the prior austenite grains, and the prior austenite boundaries may show preferential etching. Whilst there have been many studies carried out to determine the details of the mechanism associated with temper embrittlement, this has yet to be established definitively. A number of compositional parameters have been defined, for both base material and weld metal, to allow an assessment to be made of the susceptibility to temper embrittlement. Examples of these are the Bruscato factor, X, ^[13] and the 'J' factor proposed by Watanabe. ^[14]

Ligament cracking in headers

The superheater headers of a boiler are used to transport high pressure steam from the boiler to the steam turbine. The temperatures encountered are typically in the region of 540°C, which is within the creep regime for the materials that were traditionally used namely 2.25Cr-1Mo and 1.25Cr-1Mo steel. As reported by Viswanathan et al, ^[15] and Middleton et al, ^[1] the headers were designed for base loading and not for the cyclic operation that was frequently encountered; the latter led to thermal stresses and cyclic loading. This, in combination with the high temperature operation, gave rise to premature cracking in the headers by a creep-fatigue mechanism widely referred to as 'ligament cracking'. The cracks typically occurred around tube bore holes, and often extended, linking adjacent holes. The cracking generally initiated on the inside of the header, and propagated radially into the ligament between holes, and axially into the tube-hole. A thick oxide layer has been frequently associated with the ligament cracks, arising from localised temperatures above those of the bulk material. The occurrence of ligament cracking was reportedly independent of material type or the age of the header, and depended primarily on the thermal/stress history.

When cracking was discovered in a header, the remedial action depended largely on the operating conditions and the extent of the cracking. In some instances, the header was deemed fit for continued service. For more severe cracking, weld repair was carried out or, in extreme cases, the header was replaced. Many of the low alloy steel headers have now been replaced with modified 9Cr1Mo (grade 91) headers. The use of this alloy allows considerable reduction in section thickness by virtue of its higher creep strength. This in turn leads to a decrease in the thermal gradients, reducing the cyclic loading experienced, and thereby reducing the propensity for ligament cracking.

In-service cracking

For the most part, the in-service problems encountered with the ferritic alloys widely adopted in high temperature plant have been few. However, many of these alloys, for example 0.5Cr-Mo-V, 2.25Cr-1Mo and modified 9Cr-1Mo (grade 91), are susceptible to Type IV cracking. This cracking occurs at the edge of the heat affected zone adjacent to parent material. ^[16] The occurrence of cracking is attributed to the development of a localised 'soft zone' in this region, from the weld thermal cycle and postweld heat treatment, giving rise to localised creep deformation under the action of bending stresses. Type IV cracking can result in the weldment creep performance being significantly poorer than that of the base steel. It has been shown ^[17] that some improvement in performance and reduced propensity for Type IV cracking can be achieved through careful heat treatment procedures, whereby the base material is supplied after only a partial tempering treatment, and the full component subjected to a post-weld stress relieving treatment. This is clearly an approach of limited practical application. An alternative approach is to incorporate a greater safety margin in terms of stress into the design, but this carries an economic penalty.

Dissimilar metal joints

In view of the wide variety of materials used in power plant construction, the need for dissimilar metal joints is inevitable, for example between low alloy steels and modified 9Cr1Mo steel, or between ferritic and stainless steels. However, there is uncertainty in the choice of weld procedure, consumable, and heat treatment schedule to optimise the mechanical properties of the joint. The recommendation given in AWS D1.118 is that the filler material should be selected to match the low alloy material and the joint heat treated as for the higher alloy material. However, because of the differences in the Cr content of the materials, carbon diffuses from the low Cr material into the high Cr material or weld metal, giving rise to the formation of a carbon-depleted zone in the low Cr steel, Fig.2 (and a carbon-enriched region in the neighbouring material), and the presence of such regions can result in premature failure due to strain concentration in service. The extent of the carbon migration will depend on the heat treatment temperature and time, and can really only be eliminated by the use of nickel-based weld metal. However, in principle the extent of carbon migration can be reduced by the introduction of a buttering layer onto the higher alloy materials, to allow a more gradual compositional gradient. For the welding of modified 9Cr-1Mo (grade 91) to 2.25Cr-1Mo steel, a commonly adopted procedure involves the buttering of the modified 9Cr-1Mo steel with conventional 9Cr weld metal, intermediate PWHT (to temper the hard, potentially brittle high alloy heat affected zone), completion of the main fill, again using conventional 9Cr weld metal, followed by a lower temperature PWHT.

Fig.2. Optical micrograph of the interface between 2.25Cr 1Mo base material and stainless steel weld metal (309LNb/347) showing clearly the carbon depleted zone in the Cr-Mo HAZ. Vickers hardness indents show hardening at the fusion line and softening in the decarburised region.

Material developments

With the drive for higher operating temperatures and pressures to improve the thermal efficiency of new plant, as well as the bid to extend the life of existing plant, there have been significant changes in the materials used. For operation up to 620°C, a new generation of ferritic steels with 9-13%Cr has been developed, containing additions of tungsten (1-2%) to give improved high temperature properties over the traditional modified 9Cr grades. These materials, like grade 91, are predominantly martensitic, with varying amounts of retained delta-ferrite. Such materials have been the subject of extensive research programmes, for example EPRI project 1403-5019 and COST 501. A number of these materials now have ASME code case approval, and indeed have now been introduced into actual plant for full-scale trials. However, further long term trials are required and, in view of the difficulty of matching the properties of the cast and PWHT weld metal microstructure with those of the base material, additional welding consumable development is needed before we can expect to see the widespread application of these grades.

The principal grades that have evolved to date are:

• E911- as studied in COST 501
• NF616, produced by Nippon Steel [A213 T92/A335 P92]
• HCM12A, produced by Sumitomo Metal Industries [A213 T122/A335 P122]
• TB12M, produced by Forgemasters Steel and Engineering Ltd

The compositional requirements of these grades, and typical creep strength values are detailed in the Table.

Table: Chemical composition of 'new' 9-13%Cr steels

These materials offer considerable advantages over conventional grade 91. The use of NF616 (code case 2179), for example, may allow a ~35% increase in allowable stress at 600°C. This in turn permits a decrease in section thickness, and thereby a reduction in weight. Consequently, the through wall temperature gradients will be lowered, giving a reduction in the thermal fatigue loading experienced.

The diagram given in Fig.3 compares the relative wall thickness for a pipe of 290mm internal diameter for operation at 557°C, 20MPa for grade 22, grade 91 and grade 122, although the benefits of the new advanced ferritic steel are even more pronounced at higher temperatures and pressures. ^[21]

Fig.3. Variation in wall thickness with material grade (with an internal diameter of 290mm) for service at 557°C 20MPa. ^[21]

On-going steel developments are now looking at non-austenitic steels that are suitable for service up to 650°C to give further improvement to the thermal efficiency of ultra supercritical power plant. ^[22] One of the new steels, NF12, designed for boiler application, contains ~12%Cr, ~2.5%W and ~2.5%Co, the addition of cobalt preventing the retention of delta-ferrite in the microstructure. ^[22,23] A rotor steel, HR1200, has also been developed, intended for use in ultra supercritical turbine rotors for service at temperatures of 620 and 650°C. This steel contains alloying additions of W, V, Nb, Co, and B and a low N content of ~200ppm. Data generated to date on the material have indicated that it exhibits excellent creep rupture strength, corresponding to that of the precipitation hardened austenitic alloy A286, but with a more favourable (lower) coefficient of thermal expansion

Further Reading

Welding and fabrication of high temperature components for advanced power plant - Part 2

References