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Type IV Cracking of Weldments in Enhanced Ferritic Steels


Review of Type IV Cracking of Weldments in 9-12%Cr Creep Strength Enhanced Ferritic Steels


D. J. Abson and J. S. Rothwell*

International Materials Reviews, November 2013, 58(8), 437-473.


The improvement of thermal efficiency of power plants has provided the incentive for the development of the martensitic–ferritic 9–12%Cr creep-resistant steels. Good progress has been made in developing such steels, which are being used particularly in the wrought form as tubes and pipes for fossil fuelled power stations. They are also finding use in high temperature process plant within the oil and gas sector, and are being considered for use in generation IV nuclear designs. The high temperature conditions that these steels operate under in fossil fuelled power stations induce type IV cracking. This type of cracking occurs in the intercritical or fine grain region of the heated affected zone via a creep mechanism, and results in fractures with relatively little total cross-weld strain. Despite the occurrence of type IV cracking experienced in lower alloy predecessors, successor alloys have been introduced and widely used with insufficient consideration given to the consequences of welding them. Unfortunately, the newer steels suffer from reduced cross-weld creep strength due to type IV cracking to a greater degree in the temperature range of operation expected of them, and thus many failures by this mechanism have occurred. The subject of type IV cracking has been an area of active research interest. This review aims to serve as an update, drawing selectively on some of the vast amount of literature that has been published over the last 30 years.


Improved thermal efficiency of power plant has been the main driver for the development of ferritic–martensitic 9– 12%Cr creep-resistant steels that are also commonly known as creep strength enhanced ferritic (CSEF) steels. The target operating temperature for these steels is 650uC, with a common target design life of 100 000 h. Increasingly, the demand for efficiency is linked to efforts to reduce CO2 emissions, in order to meet environmental obligations and minimise any form of punitive environ- mental tax. Figure 1 illustrates how CO2 emissions vary with thermal efficiency. 1 Government agencies such as the Department of Energy in the USA are funding large, ambitious projects for the development of ultra- supercritical (USC) power plant designed to operate at 760° C and conversions from supercritical to USC opera- tions at 700° C have been carried out in Japan.2Retrofitting of existing plant is a popular option for improving output with minimal investment, since existing fittings can be used. The inclusion of CO2 capture and storage technology for fossil fuel power plants represents a significant advance in terms of combating CO2. However, it also reduces overall output efficiency and therefore there is an additional incentive for raising operating temperatures and pressures in order to maintain the same amount of saleable energy per power station.

An example of the materials used in these state-of-the- art projects is illustrated in Fig. 2. This Alstom showpiece demonstrates achievements in dissimilar metal joining necessary for minimising capital costs.3 Alloy develop- ment programmes that involve the collaboration of many partners from industry and academia feed into the high temperature plant initiatives and an example of this is the European Co-operation in Science and Technology (COST) programme that has developed alloys for use in the AD 700 initiative and other state-of-the-art fossil power plant such as the Neurath plant to be constructed in Germany.4 Many of these technically challenging new power plants incorporate the latest ferritic alloy developments and an increasing amount of austenitic and nickel- based materials, while more conservative constructions rely on established, coded steels that have been available for some time, such as grades 91 and 92 that are popular pressure vessel materials, and that represent the current highest performing types in the 9%Cr family of steels. In the developing economies, such as that of China, where construction of new power plant is very rapid, the demand for these kinds of materials has recently been beyond the rate of supply.5 This huge demand for such materials continues, despite significant problems experienced with them during fabrication and service.

1 Plot showing the expected reduction in CO2 emissions by increased plant efficiency per unit energy1
1 Plot showing the expected reduction in CO2 emissions by increased plant efficiency per unit energy1

Table 1 and Fig. 3 illustrate how, over the years, good progress has been made in developing creep-resistant steels which are being used particularly in the forged form as tubes and pipes in boilers, but which are also being considered for rotors and their cast variants for turbine casings. Such steels are also being considered for the new breed of nuclear reactors (generation IV), and beyond the power industry, the oil refinery industry has become  more  interested  in  this  family  of  steels  for pressure vessels and piping. The improvement of tensile and creep rupture strength is attractive to designers who want to take advantage of the higher hot tensile and creep strength to reduce pipe wall thicknesses, and thereby minimise thermal stresses for a more reliable plant. This is increasingly important for today’s power plants, which are subject to temperature cycles in an effort to respond to the peaks and troughs of demand, and improve profitability. A schematic illustration of the reduction in pipe wall thickness possible through the use of more advanced materials is presented in Fig. 4.

Much less attention was paid initially to cross-weld testing of the 9%Cr steels, not without quite discomfort- ing consequences for the power industry. Also, incorrect heat treatment of grade 91 steel in particular and an initially detrimental allowable aluminium level have contributed to poor creep performance. However, the most significant, weld-related problem that has plagued the power industry for over 40 years and that still has not been overcome, despite alloy development, has been the phenomenon of type IV cracking. This is creep failure that occurs in the heated affected zone (HAZ) of welded joints, particularly in high alloy precipitation- strengthened grades that are popular today, such as grades 91 and 92. These grades have a tempered martensite microstructure that is solid solution strength- ened, but with much of the strength derived from fine

2 Demonstration showpiece illustrating the materials and joining techniques that are under consideration for the new USC plants3
2 Demonstration showpiece illustrating the materials and joining techniques that are under consideration for the new USC plants3

Table 1 Development of 9–12%Cr steels67,178

Years Alloy modification 600° C/105 h creep rupture strength/MPa Example alloys Maximum metal use temperature/° C


Addition of Mo or Nb, V to simple 12Cr and 9Cr steels


EM12, HCM9M, HT9, Tempaloy F9, HT 91



Optimisation of C, Nb, V


HCM12, T91, HCM2S



Partial substitution of W for Mo


P92, P122, P911 (NF 616, HCM12A)



Increase W and addition of Co, B and controlled N




3 Chart showing the progressive development of 9–12%Cr steels64,67
3 Chart showing the progressive development of 9–12%Cr steels64,67

MX precipitates such as vanadium nitride (VN). The composition of popular creep-resistant steels is given in Table  2.162,163

Safety is, of course, of primary concern in the design and manufacture of plant intended for steam service. Improvements in design, weld integrity and inspection techniques have all contributed to improved safety, with the knowledge gained being incorporated into relevant codes and standards. An increased understanding of the role of defects in failure has been reflected in assessment codes, with an improved ability to detect and size defects

4 Illustration showing possible reduction in wall thickness achievable through use of advanced 9%Cr parent alloys: temperature5600uC; pressure530 N mm22; 100 000 h life94
4 Illustration showing possible reduction in wall thickness achievable through use of advanced 9%Cr parent alloys: temperature5600uC; pressure530 N mm22; 100 000 h life94

contributing to improved safety arising from their use. An understanding of the phenomenon of type IV cracking and of the extent of long term creep strength degradation, and strategies for improved performance, would constitute a useful addition to existing knowl- edge, and would help designers and operators to take the right decisions when designing, operating or repairing plant, to maximise overall safety, operating efficiency and plant availability.

In this document, the phenomenon of type IV cracking, particularly with respect to weldments in 9– 12%Cr steels, has been reviewed, and associated technology gaps that exist in current materials and practices have been discussed.

Type IV cracking – characteristics and causes


The problem of type IV cracking has existed for many years in power plant operated at elevated temperatures. As early as 1974, in tests carried out at 550° C, Schuller in Germany reported the low cross-weld strength of weldments in DIN 10CrMo9-10 (equivalent to AISI A182 grade F22 – 2•25%Cr–1%Mo).6 He observed that, in tests extending to more than 104 h, the cross-weld strength fell more than 20% below that of the parent steel. In the UK, many power stations built in the 1960s and 1970s incorporated a large amount of 0• 5%Cr–0• 5%Mo–0• 25%V steel (BS 3604 grade 660) that allowed a substantial rise in the operating temperature (from around 540°C up to 565°C) for coal- fired plant.7 Problems with fabrication cracking were followed by long-term failure mechanisms such as type IV cracking. Meanwhile, continental Europe deployed the 12%Cr steel X20CrMoV121 (X20), with modest service temperatures in thinner section sizes, and encountered few problems, thanks to an inherent resistance to reheat cracking and a resistance to type IV cracking at the lower service temperatures.7 Grade 91, provided a cheaper and stronger alternative to X20 that could be deployed in higher temperature plant with target temperatures of 600°C. Unfortunately, at these higher temperatures, the resistance of both X20, grade 91, and of subsequently developed creep-resistant steels to type IV cracking, falls in a more dramatic way than for their lower alloy predecessors. Early determinations of the ratio of cross-weld creep strength to that of the parent steel revealed values substantially below unity. Townley gave a 10 000 h/570°C value of 0•65, and Etienne and Heerings gave values at 550, 600 and 650°C, determined by extrapolation to the design stresses of 1•0, 0•68 and 0•70 respectively.8,9 (These authors also listed values for many other materials in elevated temperature service, including X20.) A warning note was sounded by Bru¨ hl et al., who measured cross-weld creep strengths that were substantially lower than parent steel values.10,11 This and the following warning note of Middleton appear to have been largely to no avail: ‘The operation of super- critical plant constructed from either grade 91 or X20 steels at y600uC would therefore involve a degree of risk of type IV cracking, the level depending on the medium to long term strength loss found for cross-weld testing, unless a modest overdesign were applied.’12 From a comparison of known creep rupture strengths and observed failure times of power plant welds, Middleton also deduced the likely level of system stresses. He commented that ‘At 600uC the superior rupture strength of P91 renders it more resistant to type IV failures than X20 beyond y30 kh. The presence of a creep-weak HAZ leads to a cross weld rupture strength that is severely reduced, however, leading to a much lower tolerance of system stress, to the extent that P91 systems operated at 600uC are predicted not to sustain system stresses > 16 MPa without risk of failure for a 40% strength loss’. Bru¨ hl et al.10,11 also noted that, while weld metal, HAZ and parent steel are subjected to the same strain for circumferential welds, the lower HAZ creep strength of P91 weldments is of concern for longitudinal welds, as they are subjected to the same stress as the parent steel (which is twice the cross-weld pressure stress to which the girth welds are subjected).

Almost all testing in alloy development programmes is for parent steels, and extensive testing programmes have been used to optimise parent steel compositions.13 Any cross-weld testing appears commonly to be an after- thought, although the presence of a creep-weak HAZ in a 9%Cr steel was reported, as discussed above, and demonstrated on simulated HAZs as long ago as 1990 by Middleton and Metcalfe, who carried out creep rupture testing on grade 91 samples that had been heated to temperatures between Ac1 and Ac3 (inter- critical and austenitic phase fields); see Fig. 5.12 Since then, similar studies have been carried out highlighting the poor creep strength of material heated above the Ac1 temperature; see Figs. 6 and 7. A growing body of data, reviewed in 1993 by Etienne and Heerings, revealed the lower creep strength of weldments, with the ratio of cross-weld creep rupture strength to that of the parent steel falling in the range from 0•5 to unity.9 However, as discussed in the section on ‘Design codes, published weld strength factors and extrapolation of data’, with a few exceptions, such information has been slow to find its way into design codes.

Interest in type IV cracking has intensified in recent years. Service experience with grade 91 steel components has revealed many early failures in the HAZ.14,15 The need to assess HAZ creep rupture strength has been recognised in Japan where extensive cross-weld testing or the imposition of simulated HAZ heat treatments, followed by creep rupture testing, have been carried out. European institutes that have particular interest in the problem of type IV cracking include the University of

Table 2 Chemical composition of new 9–12%Cr steels and experimental European steel FB2179



Grade 91

NF616 (grade 92)


TB 12M

Grade 911

Grade 122













































































































































































105 h creep





[150 (105 h)]




rupture strength /MPa





[80 (105 h)]




(): estimated; ...: not specified.
*Robertson and Holdsworth give 113 MPa for 100 000 h at 600°C and 56 MPa for 100 000 h at 650°C.180

5 Influence of a brief hold at different austenitising temperatures, plus subsequent PWHT at 570 or 600C on creep rup- ture life for grade 91 steel12
5 Influence of a brief hold at different austenitising temperatures, plus subsequent PWHT at 570 or 600C on creep rup- ture life for grade 91 steel12
6 Plot of creep rupture stress displaying data for P91 parent material, simulated fine-grained heat affected zone and welded joints: all specimens were subject to a 760uC/2 h PWHT and were tested at a temperature of 600C53
6 Plot of creep rupture stress displaying data for P91 parent material, simulated fine-grained heat affected zone and welded joints: all specimens were subject to a 760C/2 h PWHT and were tested at a temperature of 600C53

Nottingham, the University of Loughborough, Imperial College, MPA Stuttgart and the Technical University of Graz.

The  poor  creep  strength  of  welded  joints  greatly undermines the advantages gained through alloy devel-
opments, and has resulted in many unexpected repairs, several large-scale failures and a subsequent loss of confidence in the new materials. Particularly in the UK, existing plants into which grade 91 steel has been introduced in replacement components, commonly headers, or power plant constructed in recent decades may contain under-designed components that need careful monitoring and/or repair. It is only recently that valuable long-term data are starting to be incorporated into design codes to help prevent premature failures. Initiatives such as the European Creep Collaborative Committee (ECCC) and the state-funded analysis  of data carried out at the National Institute of Material Science (NIMS) in Japan have been set up to collect data for design purposes. However, although reliable perfor- mance may now be designed in more easily by way of increased thickness of material, the poor creep strength of welds that leads to type IV cracking remains, and prevents the impressive properties of the parent alloys being taken advantage of to their full extent.

Failure location

The classification of service cracking in weldments devised by Schu¨ ller et al., which describes the location of cracking relative to the weld, is still used today and is shown in Fig. 8.6,16 Types I–III are associated with the fabrication   of   a   joint   and   may   be   solidification, hydrogen, reheat, temper embrittlement and occasion- ally long-term creep-related cracks. Type IV cracking forms towards the outer edge of the visible HAZ, beside the parent metal (Fig. 9), and is exclusively a creep cracking mechanism that occurs after long durations.

To explain the location of type IV cracks more precisely in relation to the HAZ microstructure, it is useful to illustrate the different subzones that commonly exist in the HAZ. Figure 10 shows a section of the phase diagram for grade 91 steel, and how different parts of

 7 Rupture times for HAZ simulated coupons for ASME grade 122 steel: open points are furnace heat treated, and closed points represent coupons treated in a weld simulator108
7 Rupture times for HAZ simulated coupons for ASME grade 122 steel: open points are furnace heat treated, and closed points represent coupons treated in a weld simulator108
8 Classification of cracking in weldments from Brear and Fleming,16 according to Schu¨ ller et al.6
8 Classification of cracking in weldments from Brear and Fleming,16 according to Schu¨ ller et al.6

the phase diagram relate to the different microstructural regions of the HAZ.17 While it is well established that type IV cracking occurs towards the outer edge of the visible HAZ, the exact location, in terms of micro- structure has been reported to occur in both the intercritically (IC) heated  HAZ  and  the  grain-refined or fine grain (FG) HAZ. Thus, there is not a single HAZ microstructural region associated with type IV cracking.

9 Type IV cracking in electron beam welded 9%Cr–1%Mo steel
9 Type IV cracking in electron beam welded 9%Cr–1%Mo steel

a macro; b micro image. mm scale shown

The distinction between the two zones is not a simple one to observe in practice. Figure 11 shows the HAZ microstructure of a recently developed boron-containing 9%Cr steel (FB2), and illustrates the difficulty of distinguishing the different parts of the HAZ.18 Different materials, welding thermal cycles and tem- perature–stress regimes may bring about differences in the observed failure location, although the  failure location in long-term (low stress) service is most commonly reported to be the fine-grained heat affected zone (FGHAZ), if stresses transverse to the weld are the most significant. It is worth noting that, in multiple pass welds, the fine-grained and intercritical HAZ regions are essentially almost continuous through thickness, unlike the coarse-grained heat affected zone (CGHAZ), which may exist in disconnected regions beside the fusion boundary. This provides a continuous region that is susceptible to the phenomenon. It should be noted that the number density of cavities is not uniform across the section. The highest density of cavities is expected to occur at midsection, as observed by Yaguchi et al., where the triaxiality is  highest.19  This  observation clearly has implications for the use of surface replicas to investigate the extent of creep cavitation.

Type IIIa is another long-term creep failure mechan- ism, associated with carbon migration from the CGHAZ near the fusion boundary over the lifetime of the weldment, resulting in a carbide-depleted zone with lowered creep strength that subsequently leads to cracking.20 This type of cracking is associated with differences in the carbon activity of the weld and HAZ, which is often due to differences in the Cr concentration. It tends to be more prevalent in butt welds where less inspection is focused, and consequently more time is available for diffusion and development of cracks at typically longer durations than type IV cracks. The most commonly associated combination of materials affected

10 Schematic representation of the subzones of the HAZ corresponding to the calculated phase diagram of X10CrMoVNb9-1 (grade 91 type) steel.17
10 Schematic representation of the subzones of the HAZ corresponding to the calculated phase diagram of X10CrMoVNb9-1 (grade 91 type) steel.17

The HAZ associated with a single weld bead is shown. Type IV cracking is associated with the intercritical zone or the adjacent fine grain region

by type IIIa cracking are 1/2Cr1/2MoV steel welded with a 2¼Cr1Mo weld metal. However, denuded CGHAZs are of much concern for 9Cr joints involving lower alloy bainitic or higher alloyed martensitic and austenitic steels.21 As discussed in the section on ‘Technology gaps and future trends’, EPRI and Metrode Products Limited have developed a filler for the avoidance of this problem which has a lower carbon activity, and is especially useful for welding austenitic steels to ferritic or martensitic steels.

Circumstances giving rise to type IV cracking

Service failure times for type IV cracking have been reported to be typically 6–10 years, but have been as short as 35 000 h.7,15 While much of the earliest type IV cracking detected occurred in heats with low N/Al ratios (1•5), many examples were found for which cracking was observed after 58 000 h service in heats with much higher N/Al ratios (some > 10, at nitrogen levels > 0•05%). Consequently, Brett commented that ‘Because

the cracking observed to date has occurred at such an early stage, there is unfortunately ample scope for this process to continue with even material with quite high N/Al ratio potentially cracking within the design life (typically 150,000 hours).’ Whether or not type IV cracking will occur is influenced by the microstructure developed by the quality heat treatment, the operating temperature and most importantly the stress state of the component. For transverse tests, the mechanism is favoured at higher temperatures and lower stress levels. At higher stresses and lower temperatures, failures are more common in the weld metal or parent metal. As the temperature is increased and the stress reduced, there comes a point at which type IV creep cracking becomes the dominant mechanism, as illustrated schematically in Fig. 12 and for experimental steel FB2 in Fig. 13. This implies that, where the primary stress is across the weld, below a critical or threshold stress level, the creep strain in a weldment is localised, and creep failure occurs in the type IV region of the HAZ much earlier than would be

11 Heat affected zone region of a ruptured creep cross-weld specimen that failed by the type IV mechanism.
11 Heat affected zone region of a ruptured creep cross-weld specimen that failed by the type IV mechanism.

The specimen was extracted from a flux cored arc weld made in a piece of FB2 rotor steel. The test conditions were 625°C at 110 MPa.18 The rupture time was 9664 h

a HAZ region on non-ruptured side. Etched in 1 : 5 v/v Hcl (35%) and distilled water with 0•5–1 g K2S2O5, 2 g NH4HF2 per 100 g of solution; b HAZ region beside rupture with an attempt at defining the different regions of the HAZ. Etched in 2•5% picric 2•5% Hcl v/v in ethanol

12 Schematic representation of the reduction in creep rupture strength beyond the onset of type IV cracking for weldments at a fixed temperature.
12 Schematic representation of the reduction in creep rupture strength beyond the onset of type IV cracking for weldments at a fixed temperature.

The expected fracture location for a cross-weld specimen is indicated. Type IV cracking occurs much earlier than for a parent or weld metal specimen under the same test conditions

expected for the parent steel. Figure 14 shows the dependence of creep life on temperature for weldments in ASTM/ASME grade 91 steel, with a line distinguishing the boundary conditions for type IV cracking. An approximate equation for this boundary is

log10(tf )=0•0235|(733-T) (1)

where log10(tf) is the logarith to base 10 of the time to rupture in hours, and T is the temperature in °C.

Similar behaviour is observed in girth welds subject to high system stresses that increase the axial loading, and seam welds subject to pressure loading with hoop stress as the main component.12 It was noted in the section on ‘Introduction’ that longitudinally welded pipe is not specified in USA and most European countries, due to historic failures of pipes that resulted in fatalities. However, longitudinal seams in headers and other pressure vessels are inevitable.

From their experimental measurements and subsequent computation of the creep of an X20 vessel loaded at 610°C with an internal pressure of 200 bar, Steen et al. concluded that ‘the normally expected failure mode in a pipe under pure internal pressure will be HAZ cracking perpendicular to the hoop direction.22 The lifetime of the weld joint will fall below the lower bound of the base material scatter-band, although when no weld efficiency factor is adopted, the time safety margin may be very small’.

Coussement and de Witte investigated these effects in X20 and P91 pressurised vessel tests, and found that

The experimentally observed dependence of type IV cracking on temperature and stress for weldments made by various processes in FB2 steel.18
The experimentally observed dependence of type IV cracking on temperature and stress for weldments made by various processes in FB2 steel.18

The points bounding the darker shaded red area to the upper left represent type IV failures, i.e. fractures located mainly in the outer HAZ region and featuring limited ductility; those bounding the grey area represent failures located in the parent or weld metal, with higher ductility

14 Failure locations for cross-weld samples of grade 91 steel for different rupture lives and temperatures describing the temperature dependence of type IV cracking7
14 Failure locations for cross-weld samples of grade 91 steel for different rupture lives and temperatures describing the temperature dependence of type IV cracking7

poor placement of welds, e.g. on tight radii of end caps, resulted in reduced failure times due to increased axial stress across the welds.23 Girth welds that were subjected to pure pressure failed due to transverse cracking of the weld metal or parent material rather than the HAZ region, and the importance of creep ductility in such situations was highlighted, i.e. the capacity to strain, offload and redistribute stresses over the weld. Evidence of the importance of creep ductility was highlighted when comparing girth welds in X20 made with matching and high Ni fillers. The matching filler failed by axial cracking in the weld metal, while the higher ductility Nibased filler weld failed in the parent, again by axial cracking, but with a 54% improvement in time to rupture; cracking did not occur in the HAZ. Similar conclusions were drawn by Smith et al., who reported that if welded components are subjected to a loading condition where bending predominates, and type IV creep strain accumulates in a localised region, failure through discrete crack growth will prevail, whereas with less constrained loading, longer plant lifetimes will result, with the tendency for net section creep strain to accumulate in the type IV zone.24

For welds under pure pressure, the principal stress for the seam welds, which is in the hoop direction, is twice that to which the girth welds are subjected. Moreover, the failure of a seam weld is likely to have much more serious consequences than the failure of a girth weld. However, the superposition of system stresses can and does allow type IV failure to occur in the HAZs of girth welds, for example those attaching header end caps. Multiaxial constraint is expected to increase with specimen size and improve creep rupture times. It follows that high strength, low ductility weld metals and high angle fusion lines should increase constraint and creep rupture times for transversely loaded crossweld specimens. These effects have been demonstrated by Masuyama, who considered the effects of specimen size, shape and loading configuration in high chromium martensitic steels.25 It was reported that double penetration butt welds, ‘X’ type, perform worse than ‘U’ groove geometries for large-scale (30 mm2) uniaxial cross-weld specimens, and when tested in hot pressure vessels (seam welds subject to hoop stresses) for grade 122 steel at 650°C. The results presented also showed a large reduction in time to rupture for seam welds compared with the large uniaxial specimens, although no explanation for this was offered. In contrast, Yaguchi et al. observed similar creep rupture lives for conventional creep rupture specimens and internally pressurised welded pipe, illustrating the conflicting findings that can found in the literature.19 Their study involved narrow-gap longitudinal seam welds in a U-shaped preparation in 9%Cr and 12%Cr steels. Type IV failure was reported as occurring in the FGHAZ. The plane stress/plain strain finite element modelling of Kimmins and Smith indicated only a small change in the strain in the type IV region when its inclination to the stress axis, initially perpendicular, was changed by 10–15°.26 The influence of weld angle, and thus the inclination of the fusion boundary to the stress axis (between 0 and 45°), on creep life was modelled by Tanner et al., who showed up to a 60% reduction in a creep life, with the minimum at an inclination of ~25°. They also compared their computed values with the experimental data of Francis et al..27,28

Hyde et al. have also used element modelling of biaxial loading on 2¼ Cr–1Mo girth welds in ½ Cr– ½ Mo–¼ V pipe to consider several axial (end load) and hoop stress configurations for weld angles between 0 and 37•5° for a weld with uniform HAZ properties.29 The analysis suggested an interaction between applied stress state, weld angle and time to rupture. At a low axial to hoop stress ratio, a high weld angle gave improved creep rupture life. A cross-over point, where no benefit was derived, was identified (at an axial to hoop stress ratio of y0?6), and the trend was reversed at higher ratios, but with a lesser difference between the predicted creep lives. Further similar studies have since been made for grade 91 steel, highlighting the effect of stress state.30

Weld strength factor (WSF) and strength reduction factor (SRF)

As noted earlier, Bruhl et al. pointed out that the lower HAZ creep strength of P91 is of concern for longitudinal welds, as they are subjected to the same stress as the parent steel.10,11 In short-term laboratory tests, the cross-weld and parent creep strengths are usually very 14 Failure locations for cross-weld samples of grade 91 steel for different rupture lives and temperatures describing the temperature dependence of type IV cracking7 Abson and Rothwell Review of type IV cracking of weldments in 9–12%Cr creep strength enhanced ferritic steels International Materials Reviews 2013 VOL 58 NO 8 445 similar until the onset of type IV failure. With the onset of type IV failure comes a degradation of cross-weld creep performance compared with that of the parent steel under the same conditions. On a plot of creep rupture stress versus duration, an inflection is observed for cross-weld specimens, as illustrated in Fig. 12. This inflection occurs at a threshold stress level for a given temperature and material. Below the threshold stress, the difference between parent and cross-weld creep strength widens, so that a progressively greater discrepancy exists at lower stresses and longer durations. The difference between the weldment and parent creep strength at a given duration is expressed in a number of ways, but commonly as the WSF or the SRF. These are often defined as

WSF(t, T)=Ru(w)/t/T/Ru/t/T (2)

SRF(t, T)=(R/ut/T - Ru(w)/t/T)/Ru/t/T (3)

where Ru/t/T is the creep rupture strength of parent material specimens at time t and temperature T, and Ru(w)/t/T is the creep rupture strength of cross-weld specimens at time t and temperature T.31 The terminology for describing the reduced creep strength of weldments compared with that of the parent steel is sometimes confusing. It is desirable to achieve a high WSF and a low SRF. Examples of some of the alternative parameters are given in Table 3. Allen et al. identified that there is a broad (but by no means perfect) correlation between HAZ and parent material creep strength between heats.32 However, there is also a trend for the WSF to decrease with improved parent performance between grades, thereby strongly reducing the effective creep strength of a fabrication compared with a single piece component. In some cases, the improved creep strength of the parent material has been shown to be completely countered by the detrimental effect of welding, when compared with lower alloy predecessors. For example, in their in-depth study of data from parent and cross-weld specimens of the 9–12%Cr family, Kimura et al. indicated that employing a higher strength parent steel will bring no further advantage in terms of cross-weld creep performance at the higher temperature ranges desired for operation.33 They demonstrated that for grade 92 (which was the strongest of the coded CSEF steels) at 625°C, the estimated 105 h cross-weld creep rupture strength is 46•2 MPa compared with welded joints in grade 91 (that has a lower parent creep strength) that have a creep rupture strength of 49•9 MPa for the same temperature and duration. Such findings have made the further development of creep-resistant steels less attractive unless susceptibility to type IV cracking is eliminated.

Design codes, published WSFs and extrapolation of data

The relatively poor creep strength of weldments at long durations was not understood, and generally not fully accounted for by design codes until many years after the introduction of steels such as grade 91 into the market place. Figure 15 shows creep rupture data for parent and cross-weld grade 92 specimens, and illustrates the point that if plant were designed to the allowable stress specified by ASME alone (before 2008), weldment failure could be expected beyond 20 000 h of service for welded grade 92 material, if the allowable stress for 100 000 h service was employed.34 Several of the standards that relate to out-ofcore nuclear components (ASME Section III Code Case N47, the French Code RCC-MR and the Japanese code ETSDG) were reviewed by Bhoje and Chellapandi, who observed that comprehensive creep SRFs for weldments are provided in the French code RCC-MR.35 Weldments are discussed in Section 4/5 of the R5 assessment procedure. In addition to stating that ‘A particular problem with weldments is the presence of residual welding stresses’, there is a short section on ‘type IV cracking in 0•5 CrMoV weldments’, which discusses crack growth in the type IV region of girth welds, with no explicit mention of low cross-weld creep strength. The time to rupture of the complete weldment is taken to be the lowest of the times to rupture for the constituent zones.36

As noted in the section on ‘Introduction’, there was a comparatively early awareness in Germany and in the UK that the cross-weld creep strength can be substantially lower than that of the parent steel. This awareness was reflected as early as 1988 in the German code.37 As noted by Gomes et al., this code incorporated a factor in the stress of 80% for fully loaded longitudinal seam welds.38 However, in the light of the data now available for grade 91 steel, it is clear that this allowance was not sufficiently adequate. Until recently, no factor accounting for poor cross-weld creep strength was incorporated into more widely used design codes such as the ASME code, where allowable stress is based upon safety factors for parent material performance. The ASME code relating to high temperature pressure vessels was, however, updated in 2008 to reflect cross-weld performance. More precisely ASME I introduced a range of ‘weld SRFs’ w used for calculating component thicknesses.34 The general equation for calculating the thickness of drums and headers has not changed, but the definition of the efficiency factor E has been expanded to incorporate w for the seamwelded condition


where P is the pressure, D is the outside diameter, S is the maximum allowable stress value, E is the efficiency, y is

Table 3 Some examples of terminology used for WSF or SRF

Description given by authors Designation according to equations (2) and (3) Author(s)
Weld reduction factor SRF(t, T) Middleton et al., 20017
SRF WSF(t, T)] ASME B31.3
Weldment parent creep strength ratio WSF(t, T)] Abson et al., 200783
Weld SRF * Kimura et al., 200833
Weld creep SRF WSF(t, T)) EN 13445-2, 200244

*Weld strength reduction factor (WSRF)50?86cross-weld creep strength/allowable stress33

15 Creep rupture plot showing the performance of ASTM/ASME grade 92 parent material and weldments.
15 Creep rupture plot showing the performance of ASTM/ASME grade 92 parent material and weldments.

The allowable stress S at 10 000 h for P92 seamless pipe, and S multiplied by weld SRF w for N+T (0•77) and post-weld heat treatment (PWHT) (0•5) are illustrated. The w multiplier was introduced into ASME I 2008 code in recognition of the poor cross-weld creep performance of joints.34 This factor effectively increases the thicknesses required by the code for seam welded pressure components

the temperature coefficient and C is the minimum allowance for threading and structural stability. The efficiency factor E for seam-welded components is equal to w, or the ligament efficiency, whichever is lower. In the case where the weld seam is penetrated by the openings forming the ligament, E is taken to be the product of w and the ligament efficiency.

The value for w is provided in a table in the code, and is dependent on heat treatment of the component after welding, the steel type and temperature range of operation. For grade 91 and similar steels in the PWHT condition, between 510 and 649°C, the value of w is 0•5. Over the same temperature range, for the same steels in the normalised and tempered condition, the value falls from 0•86 to 0•77. The introduction of w effectively reduces the maximum stress (S6w), as indicated in Fig. 15, and increases the required thickness for seam-welded components, as indicated in Fig. 16.

No specific guidance is given in the code for other types of welds, such as girth welds. However, for girth welds, the cross-weld stress arising from the internal pressure is half that to which a longitudinal weld is subjected. Moreover, the deformation of a girth weld in response to the pressure stress is limited by the strain produced in the adjoining parent steel. Hence, girth welds generally present a substantially lower risk of

16 Schematic showing the calculated pipe thickness according to ASME I 2007 and 2008
16 Schematic showing the calculated pipe thickness according to ASME I 2007 and 2008

highlighting the difference that a weld SRF E makes for seam-welded pipe: P55 MPa, T5572°C, y50•7, C50, D5864 mm (Ref. 181)

17 Creep rupture WSFs for grade 92,115 grade 122107 and grade 91.68 Values derived from creep rupture data of crossweld or simulated FGHAZ specimens tested at 650uC. A linear extrapolation up to 300 000 h has been made
17 Creep rupture WSFs for grade 92,115 grade 122107 and grade 91.68 Values derived from creep rupture data of crossweld or simulated FGHAZ specimens tested at 650uC. A linear extrapolation up to 300 000 h has been made

suffering type IV cracking, although s°Ch welds could be affected adversely by local stress concentrations and, as discussed by Middleton, by high system stresses.12 A weld joint SRF W was introd°Ced into ANSI B31•3 in 2004 [clause 302•3•5(e)], as discussed by Becht IV.39 With a general dearth of experimental cross-weld data, it specified linear interpolation between values of 1•0 at 510°C and 0•5 at 815°C, for all materials. Since the 2008 issue (Table 302•3•5), values of W range from 1 to 0•5 for different materials and service temperatures. For socalled CSEF steels in the post-weld heat treated condition, it is 0•5 for the temperature range from 510 to 649°C. For weldments in the N & T condition, it decreases from 1 to 0•77 over the same temperature range, varying linearly with temperature.

Unfortunately, the WSF is worse for the newer high chromium creep-resistant materials at their target operating temperature range than it was for the older less alloyed steels, as indicated in Figs. 17 and 18, which shows an extrapolated 100 000 h WSF of y0•4 for grade 122 (12%Cr) steel at 650°C. Grades 92 and 122 show broadly similar behaviour, namely a stronger downward trend than that shown by grade 91.40 As noted above, this is supported by data presented more

18 Weld strength factors predicted for various power plant steels after 100 000 h41
18 Weld strength factors predicted for various power plant steels after 100 000 h41

recently by Kimura et al.33 Schubert et al.41 show 600°C/ 100 000 h WSF values of approximately 0•75 for 9%CrMoV steels and 0•5 for 12%CrMoV steels.41 The graphs show clearly that the WSF is not constant, but rather decreases with increasing rupture life (decreasing applied stress), and with increasing temperature, reaching levels well below unity for exposure times of the order of 100 000 h at temperatures ≥600°C.

This behaviour is also illustrated by the analysis of ECCC cross-weld data for grade 91 and E911 steels carried out by Holmstro¨m and Auerkari.42 Their plot of stress versus WSF for grade 91 steel at temperatures ranging from 575 to 650°C, Fig. 19, shows WSF values falling from ~0•9 to ~0•6 as the stress decreases. Data points added to the original graph are the WSFs given

Stress versus WSF for grade 91 steel derived by Holmstro¨m and Auerkari42 from 2005 ECCC data, to which the following data points have been added:
Stress versus WSF for grade 91 steel derived by Holmstro¨m and Auerkari42 from 2005 ECCC data, to which the following data points have been added:

sun = 100 000 h WSF values from Schubert et al.,41 which indicate associated strength values; star = 2009 ECCC 100 000 h strength values, with assumed WSF plotted in accordance with underlying data

by Schubert et al.,41 which fit the underlying data set well, and 2009 ECCC 100 000 h strength values, from which, relevant WSF values can be estimated. These added data points reveal clearly the progressive decrease in WSF at 100 000 h with increasing service temperature. Similar behaviour is reflected in graphs presented by Laha et al.43 They showed the SRF (of equation (3)), derived from data extrapolated to 100 000 h, increasing up to approximately 30% at 600°C and to approximately 45% at 650°C. In the light of this decrease in WSF, it is clear that values of WSF cannot be determined as a simple ratio from short-term creep test data. Rather, careful extrapolation must be made to stresses and lifetimes that are relevant to the intended application. The implication of the poor cross-weld creep strength and the weld SRFs introduced by ASME described above is that either operating temperatures or pressures will have to be lowered or the pipe wall thicknesses will need to be increased, unless the problem of type IV and associated poor cross-weld creep strength is resolved.34

Allowance for the cross-weld creep strength at elevated temperature being lower than that of the parent steel is likely to feature in further standards in the future. Towards this end, reliable extrapolations of WSF are needed, preferably up to 200 000 h. EN 13445-2:2002 (E) Issue 35 (2009-01), Annex C, ‘Procedure for determination of weld creep strength reduction factor (WCSRF)’, which is based on VdTU V-Merkblatt 1153, requires testing at stresses selected to give durations up to one-third of the creep design life at two test temperatures within a range of ± 30°C of the mean design temperature.44,45 Hence, if steels are to have a design life of 20 years, then a 7-year testing programme will be required. The standard states that ‘if the failure is located in the HAZ extrapolation is not allowed without further testing at longer times showing no further apparent decrease (in the WSF)’; see Appendix 1. However, for materials operating at the high temperature end of their application ranges, there is ample evidence to show that the weldment/parent creep strength ratio continues to decrease with increasing time (and hence decreasing stress). Hence, test durations for qualification to this code are likely to be similar to that of the intended service life!

The guidance incorporated in the Volume 5 Part IIb of the ECCC Recommendations (2001) is more realistic. The criterion for allowing extrapolation of the creep rupture strength Ru(W)/t/T by a factor of 3 on life, beyond the life of the longest test, tu(W),max, is:

If ΔWSF(t, T)>0•1. WSF(t, T) between 0•86xtu(W),max

and tu(W),max then Ru(W)/t/T may be extrapolated to 36tu(W),max If DWSF(t, T) .0•1. WSF(t, T) between 0•86 x tu(W),max and tu(W),max then extrapolation is not advisable. While the R5 procedure, developed by British Energy and Serco Assurance, and reviewed by Ainsworth, has a section (Volume 4/5) that provides two methods for calculating creep/fatigue crack growth. It incorporates ‘fatigue SRFs’ to modify the stress range, but it does not appear to treat explicitly the lower cross-weld creep strength compared with that of the parent metal.46 Further discussion of the allowances made in national codes has been given by Ref. 47.

Microstructural degradation and failure mechanism

Type IV failure is the result of localised strain and void formation in the outer region of the HAZ (FG/IC HAZ). This creep-weak region accommodates the vast majority of strain detectable in the creep specimen and, as such, specimens fail with markedly low total strain to failure compared with parent steel specimens or simulated HAZ specimens. Voids nucleate on precipitates mainly at prior-austenite grain boundaries and, at an advanced stage, the voids coalesce or ‘unzip’ to form a crack. Once this final stage begins, creep crack growth progresses rapidly to failure, and immediate repair or replacement is required. Investigation of creep crack propagation in grade 91 steel will ultimately help utilities to take decisions on the safest and most economical way to operate, once a type IV crack has been detected.48

Creep-resistant steels rely on alloy and precipitate strengthening, with the most important fraction of strengthening understood to be due to fine MX precipitates such as VN and NbC. During the life of the steel, many intermediate, metastable precipitates exist which, during tempering, PWHT or service, transform to equilibrium phases, sometimes via another intermediate precipitate. M23C6 accounts for a large volume fraction of the precipitates present early in the life of such steels, but it is known to coarsen quickly, thereby removing any precipitate strengthening it contributed. Similarly, M2X precipitates grow quickly and also have a passing contribution to strengthening. The dislocation density reduces during tempering, PWHT and service, and the martensitic laths break down into softer equiaxed sub-grains as the martensite tempers further to ferrite.49 The MX precipitates NbC and VN are relied upon as the main strengthening phases in these steels, due to their low coarsening rates and their ability to pin grain boundaries and dislocations. NbC forms at high temperatures, where it can prevent grain growth and is therefore important during the steelworking operation, as well as during service. VN forms during tempering, precipitating throughout the matrix. MX precipitates are obstacles to dislocation movement, but possibly of similar significance is their role in pinning sub-grain boundaries and preventing loss of strength through grain recovery and growth.50

A high chromium content can prevent MX precipitates being the equilibrium phases, and instead Z-phase can form. Dissolution of the strengthening MX phases into a complex nitride Cr(V,Nb)N known as Z-phase can take many years, and is usually accompanied by a sharp fall in creep strength.51 Another precipitate that forms during service is the intermetallic known as Laves phase, Fe2(W,Mo). This phase is found in Mo- and Wcontaining steels, and is generally thought to weaken the steel through reduction in solid solution strengtheners, but a degree of precipitate strengthening from Laves phase may counteract this to some degree. The range and scale of the phases that exist at any time are dependent on the temperature of operation and the sequence of precipitation reactions.52,53 Coarsening kinetics lead to sharp changes in creep behaviour, for example, dual phase steels (containing retained δ-ferrite) demonstrate a sudden drop off in creep strength at long durations, associated with the accelerated growth of

20 Creep rupture strength of the COST development steels FB2, FB6 and FB8 as a function of time to rupture at 650C. FB6 and FB8 both suffered from shortterm onset of Z phase causing the creep rupture strength to fall dramatically40
20 Creep rupture strength of the COST development steels FB2, FB6 and FB8 as a function of time to rupture at 650C. FB6 and FB8 both suffered from shortterm onset of Z phase causing the creep rupture strength to fall dramatically40

precipitates at the delta ferrite/martensite interface. Also Z-phase, precipitated in various high Cr steels, is known to form after extended creep service durations as the final equilibrium phase. Until recently, this phase was unknown in creep-resistant martensitic–ferritic steels, and none of the thermodynamic databases available even predicted its occurrence in these steels. The onset of Z-phase formation after long-term exposure results in a rapid reduction in creep properties; see Figs. 20 and 21. Significant efforts towards developing a series of high Cr alloys, with improved creep resistance, have been thwarted due to the inability to predict the existence and onset of Z-phase formation. This example shows how the debilitating effects of unexpected phases can invalidate predictions made from short-term data or data acquired under different operating conditions. It is for just such reasons that development programmes should always employ extensive long-term testing.

The thermal cycle imposed by welding that produces the FG/ICHAZ type IV region has a strong adverse effect on the optimum distribution of the precipitates and their interparticle spacing, which the manufacturer generated by judicious alloying and careful processing, and which are vital for high creep strength.54

Commonly, the FGHAZ is reported as the most creep-weak region and location for type IV cracking. This is often supported by minimum creep rupture strength data from simulated specimens. In this region, where temperatures are typically between 900 and 1100°C during welding,55 complete reversion to austenite occurs. Partial dissolution of the M23C6 phase is expected followed by rapid growth according to equilibrium and kintetics calculations and diagrams for 9–12%Cr steels. MX particles would also be expected to grow at a somewhat increased rate. An in-depth TEM study of the distribution and composition of precipitates within this and other regions of the HAZ via simulated thermal cycles, in the as-welded, post-weld heat treated and creep exposed condition, confirmed that the dissolution of M23C6 precipitate in the as-welded state was more evident in the FGHAZ compared with ICHAZ and parent specimens. The study showed that precipitate distributions in the PWHT state were similar for all specimens, but after further temperature exposure during creep tests, the number density decreased rapidly, i.e. growth rates were comparatively high, in the FGHAZ. Furthermore, they found that the MX particles contained relatively high amounts of Cr, leading to the conclusion that Cr diffusion (rather than V) was controlling the growth rate of MX particles following greater dissolution of Cr containing M23C6 particles during simulation heat treatment. Hence, the faster growth rate and demise in coherency and creep resistance of the FGHAZ. It should be noted that the steel contained approximately 3%Co and 3%W and no boron. The reorganisation of austenite grains is also an important factor as it is the preferred site for M23C6 nucleation. During the welding thermal cycle, the boundaries change position, and the carbide locations remain the same, and can no longer provide grain boundary hardening benefits.56

As mentioned previously, the ICHAZ region has also been suggested, albeit in fewer references, as the least creep resistant and therefore most likely type IV

21 Predicted reduction in backstress with time associated with the growth of different precipitates for COST alloy CB8. The marked reduction in backstress due to Z-phase explains the poor creep strength of alloys that suffer from precipitation of thi
21 Predicted reduction in backstress with time associated with the growth of different precipitates for COST alloy CB8. The marked reduction in backstress due to Z-phase explains the poor creep strength of alloys that suffer from precipitation of this phase during extended service101

location. The appearance of M23C6 in this region has been found to be much more spheroidised compared with those in the FGHAZ or parent steel, making for a striking difference between the TEM images of the regions.68 In this region, some austenite exists; some precipitates are dissolving, while others are coarsening at relatively higher rates compared with the surrounding parent steel. Concentration gradients between austenite and a quickly tempering martensite are formed and a less homogenous microstructure results on cooling. Lee et al. reported a substantially higher growth rate for M23C6 precipitates in the ICHAZ.57 The growth of creep cavities that nucleate on these particles is assisted by the creep deformation occurring in this HAZ region. Francis et al. and Smith et al. quantified the number of creep cavities as a function of the fraction of rupture life.24,58 Kimura et al. attributed the low creep rupture strength of grade 122 steel cross-weld specimens to the concentration differences arising between austenite and martensite existing in the intercritical range.59,60 This increases the driving force for diffusion, and promotes recovery of martensite that contains less solid solution strengthening and less effective precipitates at boundaries defined by the phase interfaces that existed during welding.

In summary, the microstructure of creep-resistant steels is a finely tuned dynamic system, with changes in the character of the precipitates occurring alongside the movement of dislocations, subgrain boundaries and grain boundaries during PWHT and subsequent creep service. The weld thermal cycle imposed upon the HAZ creates a severe disturbance in the finely balanced microstructure, with a severe detrimental effect on creep strength. Although simulated specimens demonstrate fairly consistently that the FGHAZ has the poorest creep resistance across the HAZ, in practice, it is difficult to differentiate between FG and ICHAZ when trying to define the exact microstructure that a type IV crack has nucleated in and propagated through in a real welded joint.


Hardness is often used as a quick quality assessment for parent steel, before and during service. Measuring the hardness in different parts of the HAZ and linking this to performance is more difficult.61 Hardness measurements made on polished samples will reveal changes quite well. For grade 91 parent steel, it has been asserted that, if the correct heat treatment has been received, then hardness results should fall in the range 200– 270 VHN.62 For a properly heat treated grade 91 weldment, Cohn et al. indicated that, the normal hardness range is 200–295 HV, i.e. quite similar to the parent steel range.63 Hardness generally decreases across the weld from the fusion line towards parent material with respect to the peak temperature of the welding thermal cycle that the region has experienced. However, defining the particular region of the HAZ (CGHAZ, ICHAZ, etc.) that the indent is made in is often not straight forward. Some weldments do not display the typical HAZ characteristics as clearly as others. Relationships between room temperature hardness and the time to rupture have been reported for grade 91 parent steel. Masuyama demonstrated that the instantaneous (room temperature) hardness H of the gauge portion of a creep test specimen, when normalised by being divided by the hardness H0 of the (aged but not strained) grip region, is as follows64

H/H0=0•98-0•15t/tr (5)

where t is the test duration and tr is the time to rupture. This equation gives an approximate indication of the hardness change that corresponds to end of life, which occurs when the hardness has fallen by approximately 17% (when t/tr=1 and H/H0=0•83). The creep rupture life tr can be predicted from a knowledge of the test duration t and determination of the hardness values H and H0

tr=0•15t/(0•98{H/H0) (6)

Endo et al. have given a further example of an equation for the (room temperature) hardness of grade 91 steel

H=207{29•4(t/tr) (7)

This equation can be written as

H/H0=1-0•14t=tr (8)

with the value of H0 (207 HV for their particular steel) being the initial hardness, with end of life corresponding to a 14% drop in hardness from its initial value.65 Sposito et al. state that ‘Vickers hardness correlates very well with creep life if care is taken to prepare the surface before inspection, but large measuring errors limit the applicability of this technique in the field.’66

In a study of creep degradation in grade 91 steel weldments, Masuyama noted that the lowest hardness in the HAZ was approximately 10 HV below that of the parent steel, and that this hardness difference persisted throughout almost the whole of the creep life (for creep life fractions from 0•2 to 0•9), as both parent steel and HAZ softened, displaying a linear relationship with the creep life fraction (t/tr).64,67

Hardness testing of specimens subjected to thermal cycles representative of different regions in the HAZ, followed by PWHT indicated that the lowest hardness following PWHT is in the intercritically heated specimens (Ac1>temperature<Ac3) corresponding well to the type IV location in the same test specimens.68 However, the creep rupture life of simulated HAZ specimens has not always been shown to correspond with the hardness minima. It has been shown that the test stress has an influence on the ranking of the effect of heat treatment with respect to rupture time.

For grade 122 steel subjected to a simulated HAZ thermal cycle plus PWHT, it has been demonstrated that not only does increasing the stress reduce the overall time to rupture, but it shifts the shortest time to rupture from specimens simulating the FGHAZ to those simulating the ICHAZ. Thus, the higher stressed specimens rupture sooner in lower temperature simulated specimens of low initial hardness, as seen in Fig. 22.69,70

This may go some way to explaining why type IV failure has been reported in both regions, but it follows that the FGHAZ is the weakest in creep for the low loads expected during service conditions. Albert et al. also observed creep rupture minima in specimens heated above AC3 for grade 122 steel, and detected hardness minima at intercritically heated temperatures.71 By contrast, Laha et al. found that, after PWHT of simulated grade 91 specimens, the creep-weak region

Plot of Vickers hardness (open symbols) and creep rupture time (closed symbols) for grade 122 HAZ peak temperature simulated specimens after PWHT70
Plot of Vickers hardness (open symbols) and creep rupture time (closed symbols) for grade 122 HAZ peak temperature simulated specimens after PWHT70

corresponded to the lowest hardness and an intercritical heat treatment.68

In summary, hardness minima across the HAZ do not always represent the most creep-weak region. However, the overall fall in HAZ hardness during service may be used to estimate remaining life in a way similar to that in parent material. The estimation of remaining life is discussed in Appendix 2.

Reported examples of failures, mainly by type IV cracking

As noted above, the cross-weld pressure stress in longitudinal seam welds is twice that in a girth weld. It is therefore the former that is potentially at greater risk of type IV cracking and that poses the greatest risk to safety (although the weld strength factor will be higher for the seam weld than for a girth weld in the same pipe). However, type IV failures have often occurred in other weld configurations, including at least one example of a weld attaching an end cap to a header.72

The deployment of grade 91 steel in utility power boilers as thick section components began in the UK in the late 1980s on the basis of the performance of the parent metal strength.72–74 Only since the late 1990s has attention moved to weldment performance, once the susceptibility of this steel to type IV cracking was recognised. Type IV failures in grade 91 components have been more common in the UK than in the USA, as this steel grade has been in service longer in the UK,62 although it has found widespread use in power plant in the USA, where several failures have occurred, although these have largely been in lower alloy steel grades.62,75 At least one instance of type IV cracking in a grade 91 longitudinal seam weld is known to have occurred in Japan.

Reports of at least six failures at the West Burton power plant in the UK and over a hundred instances of type IV cracking are known. It is also clear that older units designed for base-load operation and used in this capacity over many years are very susceptible to component failure when they are then cycled regularly. One of the West Burton failures occurred in a forged header end cap, and four occurred in bottle type joints, the first after only 20 000 h of service at 565°C. A contributory factor was reported to be the comparative strength of the parent steel, implied by the low hardness of the grade 91 steel. The steel contained a low nitrogen to aluminium ratio, which gave rise to AlN precipitates, and reduced the amount of nitrogen available to give the fine MX intragranular precipitates, thereby reducing the creep strength of the steel. The specification for the aluminium content has since been revised from #8804;0•04 to #8804;0•02% in the ASTM material standards, but not yet in the equivalent European standards. On a retrofit header installed on a 500 MW unit in 1992, and operating for 58 000 h at 568uC, over 100 stub welds are reported to be affected by type IV cracking. Again a low nitrogen to aluminium ratio was implicated.76–78 A sharp change in section was a further contributory factor to an end cap failure after 36 000 h of service. The seriousness of the problem and its frequent occurrence is illustrated by the report of 100 instances of cracking during 2004 in outages in power stations operated by Innogy/RWEnpower.61 Type IV cracking in grade 91 retrofit headers that have been fabricated to the highest standards have been reported more recently, with the predominant location being the stub welds.15,79 Clearly, in some instances, the composition and geometry of steel constructions have been responsible for poor performance, and incorrect heat treatment is also suspected to have contributed in some instances. However, as noted earlier, warnings of likely problems from failure in the type IV region, if adequate provision for the system stresses was not made, were not heeded. Moreover, the construction codes do not yet all require that adequate provision be made for the reduced cross-weld creep strength compared with that of the parent steel, and hence systems in steam service have suffered failures, and others may still be at risk of type IV failure, even though they were code-compliant when fabricated.

A comprehensive review on type IV cracking has been presented by Ellis and Viswanathan.80 They included consideration of service experience of type IV cracking in girth welds and in seam welds, with separate consideration of steam lines and headers. Seam-welded pipe failures have occurred in the USA, for example the 1985 failure of a 30 inch (760 mm diameter) steam reheat line of the Mohave power station that had been in service for 14 years. This failure, however, was not type IV cracking, but rather occurred because of the poorer creep performance of the weld metal compared with that of the parent pipe. However, the third and fourth failures at the Mount Storm power plant were attributed to type IV cracking.80 While seam-welded pipe is not in common use in the UK, ‘clam-shell’ seam-welded grade 91 elbows have been installed in at least one UK power station (presumably because this required one mould for the pressing, rather than two). Since the seam welds occur along the extrados and intrados, stresses on the welds arising from flexure of the pipe induced by any temperature fluctuations will be higher on the extrados than if the welds were situated on the neutral axis, and will be in addition to the pressure stress. Such a design of elbow requires careful evaluation to ensure that it does not give an unacceptably short life. Several failures due to creep rupture were reported in the USA as early as 1996,81 and the midspan of the weldment on the extrados has been identified as a region for the early formation of creep cavities.82

In China, there have been a total of six deaths in three separate incidents due to failures of seam welds in grade 91 pipe. The best documented failure occurred during commissioning of unit 2 at Datong Power Station in 2006, and resulted in two deaths. The main stem line was not manufactured in grade 91 steel, as had been believed at the time of installation. Similar failures have occurred elsewhere, with an additional four deaths, and there are believed to be an additional 30 plants containing this pipe. Such short-term failures are unlikely to have been the result of type IV cracking; however, they do illustrate the potential danger arising from the use of seam-welded pipe in power plant.

Type IV creep life prediction


The ability to predict accurately the creep strength of specimens/components subject to long-term exposure is highly desirable, since it would avoid the need for longterm testing during the development of creep-resistant steels, and also allow plant operators to make more confident decisions on maintenance schedules. Despite advances in methods of prediction, confidence in a particular alloy is only built up after a degree of verification through testing. Clearly, adequate appropriately long-term cross-weld testing is vital before the acceptance of any new creep-resistant alloy.

Parametric studies

Parent steel

The equilibrium phases that exist in a steel and the kinetics of formation vary according to the composition, temperature and stress, which in turn affect creep performance and may cause significant changes in performance. Therefore, extrapolating results obtained at one temperature to predict results at a different temperature will not always be valid, as different phases may exist, and different mechanisms may be operational. For these reasons, the use of the Larson–Miller (L–M) parameter to predict performance at different temperatures is not encouraged, as it can yield nonconservative lifetime predictions. Thus, the L–M parameter, although useful in some instances, do not always provide a reliable method for the interpolation and extrapolation of data over long durations or interpolation between different conditions of temperature and stress. This is fundamentally due to the differences in microstructure developed under different conditions, and over the lifetime of a specimen that can give differences in creep resistance and differences in fracture mechanism between short-term and long-term creep rupture tests or service, as discussed below. However, the L–M parameter is widely used

L-M=T(C+logt)/1000 (9)

where T is the temperature (K), t is the time (h) and C is the material constant, commonly set at 20, as empirically values very close to this have been found to be quite accurate for 9–12%Cr creep resistant steels. Using values of 30 and above, as is sometimes done for Cr–Mo steels,

23 Creep rupture test results for grade 122 showing a distinctive sharp inflexion in creep rupture stress. This inflexion is associated with the presence of delta ferrite and Z-phase88
23 Creep rupture test results for grade 122 showing a distinctive sharp inflexion in creep rupture stress. This inflexion is associated with the presence of delta ferrite and Z-phase88

can introduce large differences in the estimated value of equivalent times at different temperatures.83

In his discussion of parametric methods for the extrapolation of high temperature data, Goldhoff favoured determination of the L–M constant from the dataset being analysed.84 However, he obtained smaller errors in predicted values of rupture stress from the Manson–Haferd (MH) analysis, than from the L–M or Dorn methods. The MH parameter is given by


It is derived from plots of temperature, T versus logtr, at constant stress, where tr is the time to rupture, and Ta and logta are material constants which describe the point of convergence of the isostress lines.85 One form of the Dorn equation, relating strain rate and stress, has been given by Evans86


where DL is the lattice diffusion coefficient, μ is the shear modulus, b is the Burgers vector, σa is the effective stress, k is Boltzmann’s constant, T is the absolute temperature and A and n are numerical constants determined by experiment.

By splitting the data into regions, and assigning different constants to each part, time/temperature predictions become more accurate and useful, but difficulties still exist when trying to predict behaviour at different temperatures where different microstructural regimes are operating. Kimura and researchers have coined the term ‘region splitting analysis’ for one of the very first methods used to separate out different regions of a creep rupture plot.59,60,87,88 They observed that, when creep rupture strength is plotted against time, there is an inflection in the curve, which occurs at a stress that is approximately 50% of the 0•2% offset yield strength; see Fig. 23. Creep deformation in the low stress regime is governed by diffusion-controlled phenomena, while in higher stress regimes, it is controlled by dislocation glide. For dual-phase steels, i.e. those containing δ-ferrite, such as grade P122, this inflection is more pronounced.

24 Normalised rupture stress as a function of time for failure for grade 122.
24 Normalised rupture stress as a function of time for failure for grade 122.

σ is creep rupture stress, σTS is tensile strength at test temperature, tf is time to failure, Q is activation energy for diffusion through the matrix (300 kJ mol-1), R is universal gas constant (8•314 J K21 mol-1), T is temperature89

In their analysis of grade 122 (parent steel) creep rupture data, Wilshire and Scharning noted that, not only are long-term tests required in order to provide extrapolated design data, but that the allowable (parent material) stress estimates have been reduced progressively as longer-term measurements have become available. 89 They showed that, if creep rupture stress values for grade 122 steel are normalised by dividing by the tensile strength at the test temperature, data derived at different temperatures can be superimposed using a function including time to rupture, temperature and the activation energy for matrix diffusion (300 kJ mol21); see Fig. 24. This appears to be a more reliable and fundamental way of relating data collected at different temperatures to each other than the usual approach of applying the L–M parameter. Taking logarithms yields a straight line plot that allows simple extrapolation of data. For grade 91 parent steel creep rupture data, they produced extrapolations from 5000 h and from 30 000 h out to times up to 200 000 h, and presented comparisons with measured data at 550, 600 and 650°C.90–92 Such an approach could potentially reduce the number of creep rupture tests required to characterise a parent steel, and give greater confidence in extrapolations to longer lifetimes. A similar analysis, with stress normalised by dividing by (initial) Vickers hardness, has been applied to weldments of 2•25%Cr–1%Mo steels by Brear et al.93 Complex polynomial parameters are being refined using extensive datasets in a rigorous mathematical fashion to give a single best fit curve. However, such an exercise is particular to the steel grades involved, and seems to add little to the accuracy of extrapolations when there is already a vast dataset available. Such a parametric analysis serves only to give the most probable lifetime estimates, based on existing data. Extensive numerical manipulation of data has been performed by organisations such as the ECCC. As a consequence of the availability of increased amounts of data and of more reliable extrapolation methods, greater confidence can increasingly be given to long-term creep strength extrapolations for parent steels. It is important to establish reliable representations of parent steel creep behaviour not only for design purposes, but also to assist in the derivation of more reliable extrapolations of weld strength factors; see the section on ‘Design codes, published weld strength factors and extrapolation of data’.


The inflection observed in parent steels, as mentioned above for dual phase steels, is relatively minor compared with the more pronounced change in performance seen at the type IV threshold stress for welded specimens. As noted earlier, a graph showing the rupture life at which the transition occurs in grade 91 steel weldments as a function of temperature has been presented by Middleton et al.; see Fig. 14.7 Equations defining the transition to type IV in terms of time, temperature and stress have been created for grade 91 steel by various authors. Bell gave an equation that defined the transition from parent metal to type IV failure of cross-weld creep specimens in 9%Cr steel; see Fig. 25a94


Nath and Masuyama, having had sight of Bell’s review before its publication, derived a broadly similar expression to represent their experimental data95


where tf is the rupture life in h, Tk is the temperature in K and s is the (uniaxial) creep rupture strength in MPa. The reported application ranges are 843–1005 K (570– 732°C), 40–75 MPa and 4–11 600 h. The equation was subsequently recast in a form similar to that of Bell’s original equation by Brear and Fleming;16 see Fig. 25b


As Fig. 25b shows, equation (12) gives a better fit to the specific experimental data presented than does Bell’s original equation. Nevertherless, Bell’s simple version developed using many data points from different weld types is often used as a good rough guide for studies of type IV cracking, with good reason.


The limitations of parametric methods have meant that more fundamental physical models based on microstructural evolution are being developed for lifetime prediction. Various institutions have been involved in the development of physical microstructural models for creep-resistant steels that predict the evolution of microstructure and properties at specified times and temperatures. The fusion zone itself in multiple pass welds is recognised as an extremely inhomogeneous entity. This is reflected in the studies conducted by Hyde and co-workers, for example Hyde and Sun, whose models reflect the presence of columnar as-deposited and equiaxed reheated weld metal regions.96 The studies of these authors include the determination of longitudinal and transverse all-weld metal creep data for incorporation into their model.

Much of this review has concentrated on the lower creep strength of the HAZ, and the consequent

25 Creep rupture data and the lower bound creep rupture strength for parent steel and type IV cracking
25 Creep rupture data and the lower bound creep rupture strength for parent steel and type IV cracking

a the equation due to Bell;94 b the equation due to Brear and Fleming16

premature type IV failures. Hence, modelling studies that address this issue are of particular relevance. Kimmins and Smith reviewed experimental studies and finite element models.26 Recognising that the creep damage is confined to a narrow region in the HAZ, they devised a novel finite element model of the behaviour of a creep-weak layer beside a single-V weld (simulating the type IV region) inclined to the stress axis. Their model gave good agreement between theory and experimental observations when they allowed for relaxation of constraint via the sliding of adjacent elements. They concluded that ‘the creep-weak type IV layer experiences no measurable constraint from the adjacent material’. In a later experimental programme, Smith and co-workers concluded that cavitation in the type IV region is a consequence of grain boundary sliding, leading to relaxation of constraint and multiaxial rupture governed by the von Mises stress.24 Hence, ‘rather than using conventional continuum damage models in finite element analyses, alternative models involving mechanisms of grain boundary sliding require development’. They found that the maximum fraction of cavitated boundaries in the type IV region, at times close to failure, was only approximately 1% (4000 cavities per mm2); this proportion must therefore increase rapidly before rupture. With such a small proportion of a the equation due to Bell;94 b the equation due to Brear and Fleming16 25 Creep rupture data and the lower bound creep rupture strength for parent steel and type IV cracking, according to two different equations, from Brear and Fleming16 Abson and Rothwell Review of type IV cracking of weldments in 9–12%Cr creep strength enhanced ferritic steels International Materials Reviews 2013 VOL 58 NO 8 455 cavitated boundaries, it is likely to be beneficial to keep the HAZ as narrow as possible, which may be one of the reasons for the better cross-weld performance of EB welds compared with tungsten inert gas (TIG) welds, as noted in the section on ‘Improved performance through welding procedure’.

The finite element-based creep continuum damage mechanics (CDM) method for modelling the high temperature creep damage initiation, evolution and crack growth behaviour of cross-weld specimens has also been used by Hayhurst and co-workers and by Hyde.30,97,98 Bauer et al.99 have been modelling the use of matching, overmatching and undermatching filler metal, together with the effects of pipe wall thickness (42–47 mm) and weld edge angle (0–22°) for longitudinally welded pipes in E911 steel. Their modelling has indicated that using an undermatching weld metal is beneficial by reducing the multiaxiality of the stress state, and thus the extent of creep damage. Thus, modelling studies are now capable of describing several aspects of cross-weld behaviour.

Continuum damage mechanics can be combined with the microstructural models to predict creep strength and other properties. The University of Loughborough and the Technical University of Graz appear to be leading the development of such models, which use thermodynamic and kinetic models to predict the growth and distribution of precipitates considered most important for performance. TU Graz have gone one step further by the application of back stress calculations to their CDM model.101 Clearly, progress is being made in modelling the growth of precipitates which, in the fine grain region of the HAZ is a precursor to void formation and linkage to form cracks. Hence, it should eventually be possible to introduce into the models parameters derived from the characteristics of a particular material, and to predict its type IV behaviour. Among the complexities that do not yet appear to have been incorporated into the models is the change in carbon content beside the fusion boundary that occurs during elevated temperature exposure of weldments with differences in alloy content, as a consequence of the migration of carbon across the fusion boundary.

Neural network analyses

If a sufficiently large dataset exists, a neural network analysis can be carried out, to predict the value of one or more parameters (depending on the software) if the value of all the other parameters is specified. Additionally, by fixing the value of all but one of the input variables, the way in which the output variable changes as the remaining variable are changed can be explored. The quality and applicability of the data-base being drawn on is of great importance in order to be able to make firm conclusions.

This approach has been adopted by Francis et al., whose analysis indicated a strong influence on predicted creep rupture strength of welded joints, not only of the test temperature and time, but also of the normalising and tempering temperatures, the tungsten content, the preheat temperature and the PWHT temperature.58,102 Experimental support for some of this work is discussed in Appendix 1. A subsequent study demonstrated the beneficial effects of increasing the preheat temperature and having a steep side to the weld preparation, with the influence of heat input, in the range 0•8–2•4 kJ mm-1, being small.28 It is, of course, recognised that a range of each of the input variables is required in order for trends to be revealed. As more data become available, this approach will be capable of demonstrating the effects of the variables with greater clarity.

Since grade 91 steel is strongly resistant to tempering, a PWHT temperature of 760°C is generally used for piping and pressure vessels to reduce the level of residual stress and to improve weld metal toughness. However, with a lesser improvement in weld metal toughness, PWHT at 680°C has been used for its derivatives in exploratory studies when welding turbine rotors. While this gives an initial benefit in terms of increased shortterm creep rupture stress in laboratory tests, it is likely that at least some of the benefit will have been lost when subjected to lower stresses in long-term service.

Despite difficulties in predicting long-term behaviour of creep properties, it is reasonable to assume from the available evidence that, as long as only type IV data are used, the extrapolation of these data and ignoring other short-term data should provide a simple and effective way of predicting failure to longer durations. Data demonstrating a clear departure from parent material creep performance should be used for a single particular temperature, if such a simplification is to be employed. Data from tests somewhat in excess of 100 h at 650°C or 300 h at 625°C will display the characteristic departure from parent material data for grade 91 steel; see Fig. 14.

Mitigation of type IV cracking


There are four principal ways of improving the crossweld creep rupture strength of creep-resistant alloys:

(i) normalising and tempering the weld
(ii) adjusting the welding procedure
(iii) modifying the tempering treatment
(iv) alloy development to increase inherent resistance to type IV mechanisms.

Normalising and tempering

This heat treatment may be carried out before incorporation into the main piping network and it effectively eliminates theHAZ of the weld, causing the microstructure to revert back to a state similar to that of the parent steel. The same regime of temperatures and times used in the production of the parent material may be employed and, as with parent manufacture, care in controlling the heat treatment is essential. Because the entire welded component, which may take various shapes, needs to be normalised to prevent an HAZ-like microstructure in any part of the component, there are significant difficulties in carrying out a successful treatment, and various companies have invested heavily in the optimisation of the process. One such example is Mitsubishi Heavy Industries in Japan who have reported the successful normalising and tempering (NzT) of narrow-gap seamwelded P91. The reported results show little or no difference between cross-weld and parent steel creep test results at durations up to 30 000 h and temperatures of up to 700°C.103 Figure 26 shows the basic details of the process used and the results obtained. There are also companies in Germany known to be practising NzT for grade 91 pipe that is too thick to be produced as seamless. The technology is also applicable to thin-walled large diameter pipes. Since ASME introduced

26 Heat treatment, welding details and cross-weld creep results for normalised and tempered P91 weldments103
26 Heat treatment, welding details and cross-weld creep results for normalised and tempered P91 weldments103

the ‘weld strength reduction factor’ mentioned in the section on ‘Weld strength factor (WSF) and strength reduction factor (SRF)’, there has been greater interest in NzT treatments, because this allows a strength factor of 0•8 instead of 0•5, and thus reduces the thickness of material needed. Obviously, the advantages of thinner pipe need to be weighed against the extra time and cost of theNzT process. It is not generally practical to normalise the attachment and site welds. However, such welds are generally circumferential welds; hence, there will still be welds withHAZs in the systemthat could suffer fromtype IV cracking, and could therefore compromise plant integrity. The pressure stresses on such welds are half those experienced by the longitudinal welds. Therefore, unless inadequate allowance has been made for system stresses or local stress concentrations, they should present a lower type IV cracking risk.

In some instances, it may be possible to butter the ends of grade 91 pipe work with a weld metal of higher creep strength and then carry out a PWHT, as noted by Coleman and Hainsworth.104 If the welding is carried out in a fabrication shop, it may even be possible to normalise the whole component. The subsequent welding would then leave the site weldment with its HAZ lying within the buttering layer of weld metal.

Improved performance through welding procedure

Although an HAZ susceptible to type IV cracking will always be produced by a welding thermal cycle in currently used creep-resistant steels, there appear to be opportunities to improve the performance by adjustments to the welding procedure.

A novel approach to improving the creep rupture strength of the type IV region, as noted by Bell, is to carry out a partial tempering of the parent steel before welding and to complete the tempering treatment as the PWHT, as described by Sikka.94,105 This approach almost doubled the short-term (< 10 000 h) creep life, compared with conventionally heat-treated 9%Cr steel weldments.106 It would be interesting to see the effect over a longer test period. Although it seems logical that partial tempering of the parent during steelmaking would reduce the size of precipitates in the steel, this approach is unlikely to be adopted widely. It would require code changes, and it would be necessary to persuade steelmakers to change the tempering treatment, and accept a change to the specification requirements. Unless only short lengths of parent steel were involved, the final tempering, in the field, would presumably require a mobile induction heating unit,

27 Plot showing improved cross-weld creep performance
27 Plot showing improved cross-weld creep performance

and thus this approach would have limited application. However, it may prove of benefit where the creep rupture life cannot be increased by any other means. Other process variables suspected to have a strong influence on the type IV creep life of weldments have been highlighted by Francis et al. with their Baysian neural network analysis.28 These were preheat temperature and PWHT time. The effect of PWHT on interparticle spacing has also been modelled, showing a negative effect on creep properties.

A further strategy, which could presumably be combined with the half tempering approach, is to join modified 9%Cr steels by welding processes that can create a very narrowHAZ (and therefore a shorter thermal cycle and less time for overageing of precipitates). Abe69 reported that EB welding which produces a very narrow HAZ, had a beneficial effect on the creep rupture performance of the type IV region for HCM12A steel, with more recent data confirming the behaviour; see Fig. 27.107 The approximate HAZ widths of EB and TIG welds were 0• 5 and 2• 5 mm respectively and the cross-weld creep rupture life of EB welds was found to be approximately twice that of TIG welds. Since reduced pressure EB welding eliminates the need for the fabrication to be housed in a vacuum chamber, this welding process appears worthy of consideration for the field welding of 9%Cr steels.

Albert et al. have shown that simulated FGHAZ and real weldments behave significantly differently in terms of their creep strain.108 This observation was attributed to the triaxiality introduced by the different creep properties in the various regions of the HAZ. They pointed out that changing the weld preparation angle can alter the stress state of the joint, and influence creep results significantly. They concluded that ‘by reducing HAZ width or the groove angle of the joint, this stress state can be altered to achieve significant improvement in rupture life of the weld joints’. In their report of studies of the HAZ in grade 91 steel, Laha et al. discussed measures to improve type IV cracking resistance, namely heat treatment, changing the HAZ width and modifying the steel composition.68

Modification of the weld metal properties alone has been shown to be ineffective in avoiding type IV failure in grade E911 steel.109 Below a certain stress, all failures were type IV. Nevertheless, there are a number of considerations when selecting a weld metal for use with creepresistant steels. Similar steam oxidation resistance is required to that of the parent steel. However, since the creep-week region is likely to be the type IV zone produced by the welding, the creep strength of the weld metal need only to be comfortably in excess of the cross-weld strength. Santella reported the elimination of a hardness minimum in surveys across the HAZ when the tempering treatment was carried out at 621°C, followed by welding and a conventional 760°C PWHT.110 In subsequent creep testing, failure occurred in the parent steel. While this gives some hope that the type IV zone may have been largely eliminated as a result of the revised tempering treatment, the creep test conditions were not disclosed fully, and it appears possible that the testing was of too short a duration for type IV cracking to develop. It is, however, possible that the extent of precipitates nucleation and growth during the tempering treatment was so small that the weld thermal cycle plus subsequent PWHT did not cause them to coarsen greatly beyond their optimum size. This approach clearly has some similarities with the partial tempering proposed by Sikka, which was discussed above.105

Alloy development


In view of the severe degradation in creep rupture strength that welding introduces in currently available steels, it is desirable to develop steels in which such degradation is much diminished. One approach to improving the creep rupture strength of the type IV region is to modify the parent steel composition. The most promising step towards the elimination of poor cross-weld creep strength has been the addition of boron as an alloying element, shown to improve creep strength in general for 9%Cr steels; low concentrations are present in coded steel grade 92, for example (the specified range is 0•001–0•006 wt-%).111 Higher concentrations have been shown to enhance the creep strength, with the TAF steels developed in the late 1990s by Fujita at NIMS (0•04%B) still being unsurpassed.112,113 Boron is reported to combine with M23C6 precipitates and prevent their coarsening and the onset of tertiary creep.114 Abe explained that theoretically, boron will segregate to grain boundaries during austenitising at 1100°C.115,116 During tempering at ~800°C, boron at the boundaries is incorporated into the precipitating carbides, giving M23(C,B)6 precipitates. Confirmation of the segregation of boron to grain boundaries and its presence in the carbides was obtained by autoradiography.114 Das et al. compared a grade 91 steel and a boron-containing variant.117 In addition to a uniform grain size across the HAZ for the boron-containing steel and a smaller trough in the hardness profile, they claimed that the consequent reduction in the ‘metallurgical notch’ contributed to the avoidance of type IV failure.

Unfortunately, TAF steels were not conducive to welding or forming, but recent experimental compositions appear to have overcome these difficulties. These new experimental alloys include MARBN steels developed at NIMS and a similar version – NPM1 – produced at TU

28 Creep rupture data of Abe for parent steel and crossweld specimens of 92 grade and a boron-containing steel tested at 650C 115
28 Creep rupture data of Abe for parent steel and crossweld specimens of 92 grade and a boron-containing steel tested at 650C 115

Graz.109 The most beneficial aspect of these steels, which contain 3%Co, 3%W, high boron and low nitrogen, is their apparent immunity to type IV cracking. The creep strength of the weldment in boron-containing steel appears similar to that of the parent steel, as illustrated in Fig. 28 by Abe and by Kondo et al., who found that no type IV failure occurred in the high boron steel.115,118 These steels have carefully controlled boron and nitrogen additions and have been subject to a normalising treatment at 1150°C. These measures prevent the formation of coarse boron nitride precipitates that are known to be detrimental for creep properties. Also, the HAZ region is not refined and appears virtually unchanged from its state before welding, i.e. almost the exact parent microstructure that existed before welding is present post-welding down to the location of prior austenite grain boundaries and martensitic laths; see Figs. 29 and 30.119 The propensity for this behaviour is not well understood but is thought to depend on the relative concentrations of boron and nitrogen. Figure 31 shows the solubility line for boron nitrides, as determined by observation of the fracture faces of 9%Cr steels with varying amounts of boron and nitrogen, after tensile testing.120 It shows that, for a given boron concentration, the nitrogen must be kept below a level defined by the line in order to avoid the formation of boron nitrides

29 Microstructure in steel ‘NPM1’ before and after welding thermal cycle simulation showing that the location and orientation of features remain the same119
29 Microstructure in steel 'NPM1' before and after welding thermal cycle simulation showing that the location and orientation of features remain the same119
30 Heated affected zone microstructure in steel ‘NPM1’ after PWHT showing a clear resemblance to the parent steel microstructure 109
30 Heated affected zone microstructure in steel 'NPM1' after PWHT showing a clear resemblance to the parent steel microstructure 109
31 Solubility limit for BN at 1150C for high Cr steel, with the corresponding information for the Fe–B–N system, derived by Fountain and Chipman 120 superimposed
31 Solubility limit for BN at 1150C for high Cr steel, with the corresponding information for the Fe–B–N system, derived by Fountain and Chipman 120 superimposed

Also shown is the composition specification for P92. The figure has been adapted from Sakuraya et al.121

and thus the formation of a refined HAZ. The solubility limit of BN at 1050–1100°C was expressed by the following equation121

log(%B)={2•45log(%N)6•81 (14)

However, the true position of this line is not well defined and alloys with relatively high levels of boron and nitrogen have been found to exhibit the same advantageous qualities as those with compositions on the left side of the diagram. Mayr et al.122 have attempted to further characterise the solubility of BN for different austenitising temperatures using thermodynamic calculations and examination of additional melts. Furthermore, the effect of varying the welding thermal cycle has not been investigated in these, rather special, boron-containing steels. Nevertheless, there is much excitement in industry surrounding the development of these steels because of their potential to eliminate problems with type IV cracking through much improved cross-weld creep strength. The European COST group are planning to scale up a cast of NPM1 (a 9%Cr steel containing B and N, and high levels of Co and W), and to also adjust B and N levels in the front running steels FB2 (a 9%Cr steel with B and N and lower levels of Co and W) and CB2 (the cast equivalent of FB2 that is of essentially the same chemical composition), in an effort to take advantage of the positive effects of controlling these elements. It appears that the production of large heats of steels with these modified compositions is likely. Information on the performance of welded joints in these materials is therefore clearly required. Unfortunately, the presence of boron in 9%Cr steels precludes their use in nuclear applications.

9–12%Cr steels strengthened by Z-phase

A separate parent steel development involves the formation of fine-scale precipitates of the equilibrium phase (Z-phase), which is the phase that, when present as coarse precipitates, is normally associated with the rapid degradation of creep strength at long durations in creep-resistant steels. Previously, a 12%Cr steel would have been expected to suffer from the reduction in creep strength associated with dissolution of finer-scale strengthening precipitates and the formation of coarse Z-phase precipitates during long-term service, as noted earlier; now it is possible to precipitate Z-phase very quickly, and in a fine stable dispersion, with higher Cr levels actually serve to aid this process. In their European patent, Danielsen and Hald described the production of a high Cr steel in which creep strengthening is achieved via a fine dispersion of Z-phase particles, with an average size of < 400 nm.123 This is an exciting development in which the most stable phase is precipitated during a tempering treatment to give a small interparticle spacing. This means that the steel cannot suffer from the competitive growth of different phases during its lifetime, and only the coarsening kinetics of the very stable Z-phase can cause the long-term degradation of creep properties. Another advantage of this type of steel is that higher Cr levels can be tolerated without limiting the creep life. Higher Cr levels give the additional advantage of improved steam oxidation resistance. Information on the cross-weld creep performance of such

32 Creep rupture strength at 600C of some experimental flux-cored wire deposits after PWHT, containing varying amounts of Ti and Nb, showing improved creep strength at Ti levels of ~0.06%Ti. Adapted from Abson et al. 83
32 Creep rupture strength at 600C of some experimental flux-cored wire deposits after PWHT, containing varying amounts of Ti and Nb, showing improved creep strength at Ti levels of ~0.06%Ti. Adapted from Abson et al. 83

steel has not been published at the time of writing, but it will clearly be of great interest to establish the level of susceptibility to type IV cracking and to confirm their expected good steam oxidation resistance. In view of the stability of the fine-scale Z-phase precipitates, it is anticipated that this new type of steel will show less degradation of creep strength in the type IV region that other 9–12%Cr steels, but this appears not yet to have been demonstrated.

Weld metal development

Weld metal compositions usually follow parent steel compositions. However, often the Mn and Ni contents are raised to improve strength and toughness.124 These elements can depress the eutectoid temperature. It is necessary for the PWHT temperature to be below the Ac1 temperature for the weld metal, and therefore certain elements that affect the Ac1 temperature need to be controlled. As Mn and Ni depress the lower transformation temperature, the sum of Mn+Ni should not exceed 1•5%, or the PWHT temperature range would need to be restricted.125 These elements also depress the Ms temperature, and this increases the risk of generating retained austenite.

According to Swanekamp,62 the ASME II Task Group has proposed new PWHT limits on grade 91 components based on Ni plus Mn content, specifically:

(i) PWHT temperature range must be 1350–1425°F (732–774 °C) if the chemical composition of the filler metal is not known precisely
(ii) if the chemical composition of matching filler metal is known precisely, the maximum PWHT temperature can be increased to 1470 °F (800 °C) (if Ni+Mn,1•0%) or 1450 °F (788 °C) (if Ni+Mn is between 1•0 and 1•5%).

In view of the low creep strength of the HAZ compared with the parent steel in most creep-resistant steels, the creep strength of the weld metal is of secondary importance. As noted above, weld metal compositions usually follow parent steel compositions; however, weld metals do not display type IV cracking. In the light of the promising results displayed by the new generation of boron-containing steels that appear to have a reduced susceptibility to type IV cracking, the creep strength of the weld metal could become the limiting factor for cross-weld creep strength. Further development of welding consumables must await the commercial realisation of newer steel compositions, but hopefully will not lag too far behind.

Although titanium is potentially a useful addition for the formation of fine MX precipitates, which enhance creep strength, high solution temperatures are required during steel processing, leading to extra costs for the steelmaker, making it unattractive for parent steels. Taneike et al. and Abe et al. employed a solution treatment temperature of 1300°C to take titanium carbo-nitrides into solution and observed a substantial reduction in the minimum creep rate at 650°C in 8•4%Cr steels containing 0•047%Ti, 0•13%C and 0•006%N.126,127

While it may be impractical to employ such a high solution treatment temperature for parent steels, due to cost implications, welds are effectively small-scale castings that undergo rapid solidification, so the incorporation of small additions of Ti into 9%Cr weld metals is more straightforward. Sperko suggested that Ti is an effective substitute for Nb.125 While the recovery of Ti in weld metal is difficult to control, partial substitution for Nb may be beneficial. Ti has been shown to be beneficial for the creep strength of weld metal; see Fig. 32. However, the potential detrimental effect of Ti on toughness must be recognised and the problem addressed.83

Repair of type IV damage and avoidance of PWHT

On a basic level, repairing a type IV crack in a weldment can be dealt with in a similar way to any other weld

33 Effect of PWHT on weld metal toughness demonstrated with Chromet 9-B9 129
33 Effect of PWHT on weld metal toughness demonstrated with Chromet 9-B9 129

repair, i.e. the damaged area is generally removed and either a fill weld is used or a new piece of material is welded in. Matching fillers and PWHT are standard practice, and procedures used during fabrication are used for repair of service-aged plant. However, there are a number of complications and issues with creepresistant steels that require consideration. The introduction of a weld and the associated HAZ represents a disturbed part of the microstructure where temperatures above the Ac1 and Ac3 have resulted in a redistribution of elements, including full or partial dissolution of precipitates and the formation of new martensite. One purpose of PWHT is to temper martensite in the weld metal and HAZ. However, if a PWHT is imposed, the microstructure developed in the parent steel during the tempering stage of manufacture can be modified and softened by overheating or prolonged PWHT, and this is a common cause for concern.

The effect on cross-weld creep properties of an (inadvertent) incorrect PWHT during fabrication is neither well understood, nor well documented. Most attention has been given to the weld metal toughness, which is important during start-up. The effect of PWHT on HAZ toughness and cross-weld creep strength has not been thoroughly investigated, although Shiga et al.119 presented limited data for grade 91 steel showing that the creep rupture strength after PWHT is lower than that in the as-welded condition.128

The improved toughness imparted by a higher PWHT temperature or longer PWHT time for grade 91 weld metal is shown in Fig. 33.129 While this information relates to the behaviour of one particular weld metal, it is particularly notable that the toughness is approximately doubled by imposing a 760°C/2 h PWHT rather than a PWHT in the AWS range. The improved absorbed energy levels imply that the dislocation density is reducing, carbon is coming out of solution (tempering the martensite) and precipitates are forming and coarsening. While PWHT is commonly considered necessary to precipitate fine-scaleMX, this might not always be the best approach for optimising the long-term creep properties, as further tempering occurs at service temperatures. After nucleation, a prolonged heat treatment may reduce the interparticle spacing, which is fundamental in influencing creep strength.100 In parts of the HAZ where the temperature has not been high enough to dissolve precipitates (more particularly the fine VN precipitates, which dissolve at temperatures above 1050°C), the effect of PWHT on creep life may be considered to be a negative one. However, cross-weld creep testing (where the fracturemode was type IV cracking) of weldments subject to several different subcritical heat treatments (< Ac1), considered to cover most scenarios arising in practice) with a total hold time of no more than 9 h, has shown negligible difference in terms of rupture life.130,131 It should, however, be recognised that somewhat longer hold times at similar subcritical heat treatment temperatures have been shown to affect the precipitates and creep performance significantly in parent materials and are therefore likely to also affect weldment strength.132–134 The same studies detailed the deleterious effects of treatments above the Ac1 temperature.

In aged material where precipitates have already undergone some coarsening, PWHT will not help to improve the creep strength of the HAZ, and is more likely to degrade it, thereby tending to increase the risk of type IV cracking. Since the creep life of an aged material may be somewhat degraded, it might be possible to relax the creep properties of the filler material used for repair, so it is essential to have an understanding of the creep properties of the aged material to ensure adequate strength in the weld metal. While it is the type IV region of the HAZ and not the weld metal that is the weak link, at least for the current generation of steels, the weld metal selected must have sufficient creep strength to support any offloading that may occur during service.

Since the toughness of a component is likely to degrade during elevated temperature service, it is not only during start-up after fabrication that there is a risk of brittle fracture. For repairs during fabrication or early service, where the toughness properties of the joint before and after PWHT are known, a decision can be taken as to whether it is safe to omit PWHT. Residual stresses and the changes in toughness during prolonged service and due to repair are generally not known, and so PWHT after a repair in aged material is generally considered necessary.

In view of the difficulties associated with PWHT in the field and the likely detrimental effect on cross-weld creep properties, a weld repair method for 9–12%Cr steels that does not require PWHT would be welcomed, as it would save valuable hours during maintenance shutdown periods. Such techniques are well established for lower alloy steels, as discussed by Friedman,126 who presented a four-layer MMA technique, and by Mitchell and Tolaini,127 who described the use of a low carbon Cr–Mo flux-cored wire for the repair of type IV cracking.135,136 For low alloy steels, care is required to ensure that stress corrosion cracking does not occur before some tempering of the hard weld metal and HAZ has been effected by exposure to service temperatures. 137–139 The use of a consumable that gives good ductility and toughness in the as-welded weld metal will be necessary for repairs in aged material without PWHT. Hence, the 9%Cr nickel-based consumable EPRI P87, which is discussed in the next section, is a natural candidate.

Controlled deposition techniques developed for lower alloy creep-resistant steels to avoid the need for PWHT, e.g. those discussed in the various articles in Welding Research Council Bulletin 412, particularly Friedman,126 have not been widely used for the newer, higher Cr steels, even though there are likely to be advantages if suitable techniques can be developed.135,140 Maximising the refinement and in particular the tempering of the CGHAZ and weld metal microstructure through the application of controlled deposition procedures may be a way to avoid the use of PWHT in some situations. However, devising such procedures presents a considerable challenge, since martensite of high hardness (> 400 HV) forms in both the weld metal and the HAZ at all normal cooling rates, and grade 91, having good creep strength, is very resistant to tempering, A fine grain size in the HAZ is required (for good toughness), and thus a low heat input must be used; the temperature must be reduced from 200°C (the usual preheat temperature) to 100uC (to allow transformation to martensite) in between each layer. Further layers need to introduce sufficient heat to effect tempering, and yet the steel is resistant to tempering.

Furthermore, since stress corrosion cracking has been known to occur in weldments awaiting PWHT, the hardness must be reduced to a sufficiently low level (which is yet to be defined) to avoid its occurrence. One benefit does arise from the low martensite transformation temperature range, namely that the accompanying expansion partially offsets the build-up of residual stress, resulting in lower levels of residual stress than those arising in 2•25%Cr–1%Mo steel.141,142 Extensive weld procedure qualification testing will be required before this desirable goal is achieved. This is an active area of research in several research centres, including TWI.

The necessary procedures can be time-consuming, difficult to carry out and rely heavily on the skill of the welder. Vekeman and Huysmans used a carefully constructed, controlled deposition procedure for cold weld repair of new T91, using a combination of the half and temper bead techniques, with a 2•2%Cr–1%Mo filler.138 Brett commented on weld repair without PWHT of a grade 91 header: ‘Prior to the 2006 outage, a weld repair procedure to be used without PWHT, capable of maintaining the header in operation between the 2006 and 2008 outages, was approved.’15 Details of this procedure, which involved the use of 9%Cr–1%Mo welding electrodes (with carbon at the lower end of the available range), have now been revealed.143 They carried out CTOD tests on the as-welded repair, and creep tests following a heat treatment that simulated service, and concluded that a repair had adequate toughness, and adequate creep strength to serve for the 4-year interval between planed outages.

A novel approach to the improvement of the crossweld creep strength of seam welds in low alloy steel pipe has been described by Coleman and Coleman and Gandy.144,145 Using flux-cored arc welding, they deposited a wide three-layer controlled deposition band straddling the whole of the seam weld, thereby strengthening the weldment and putting the potentially creep-weak HAZ into compression, and thus extending the creep life considerably. A demonstration repair to a 2•25%Cr–1%Mo elbow was welded with a 232° C preheat and no PWHT.144 In their studies of repairs in 2•25%Cr–1%Mo and 9%Cr–1%Mo steels, Bhaduri et al. used half bead and temper bead repairs separately, preferring the latter for grade 91 steel.146 TWI is currently extending the approach for service-aged tube material. As noted above, a most important consideration is that the interpass temperature must be low enough to allow complete transformation to martensite in the weld metal and HAZ (~100° C) between passes (or between layers) compared with the normally high interpass temperature used (~200° C), so that the martensite transformation proceeds to near completion, and subsequent passes are able to temper the underlying microstructure effectively, and effect some relief of residual stresses. The difference in creep properties between controlled deposition welds and heat-treated welds appears not to have been investigated. Other methods for the avoidance of PWHT may be worth exploring, including autogenous remelting to temper the weld beads and the HAZ.

There appears to be little published information on the creep properties of as-deposited weld metal and asformed HAZ for grade 91 steels. In indentation creep tests, the creep strength of as-deposited grade 91 weld metal was shown to be much higher than that of the parent steel.147

The service performance of several weld repairs for high temperature power plant applications has been discussed by Klenk et al.148 Existing methods for the weld repair of low alloy heat-resisting ferritic steels have been reviewed by Issler et al.149 The comment is made that PWHT is the most problematic step in practice, and that Ni-based consumables can be used to repair without PWHT, at least for temporary repairs. The advantages and disadvantages of using Ni-based weld metal have been reviewed by Brett et al. and are listed in Table 4.150 One difficulty is that dissimilar welds created in such a way are difficult to inspect nondestructively, and so are typically replaced with matching consumables at the next available opportunity. Ultrasonic inspection of dissimilar joints is receiving much attention at TWI. If successful methods are developed, this could lead to more widespread use of Ni-based filler materials in creep-resistant steel fabrications for the power industry.

Clearly, a large proportion of repairs are likely to be carried out following the occurrence of some type IV cracking. Any repair welding is likely to introduce a type IV region of lower creep strength than the service-aged parent steel. Hence, careful consideration must be given to the consequences of carrying out such a repair. The issues that should be addressed in a method statement relating to the repair welding of 9%Cr steels have been described by Henry and Bezzant, who stated that such repairs can be successfully performed if:

(i) there is uncompromising attention to detail
(ii) the unique metallurgy of the material is fully understood
(iii) the current condition of the material is fully understood
(iv) the reason for the repair is clearly understood (v) a detailed repair plan is developed by a competent technical specialist and all aspects of the plan are successfully executed by the end user or by the mechanical contractor acting on the end user’s behalf.151

Weld repairs carried out in Swedish power plant have been reviewed by Storesund and Samuelson.152 Most of the repairs were to 0•5Cr–0•5Mo–0•25V steel, but some involved repairs related to service-induced damage in dissimilar welds between low alloy and 12%Cr steels. No repairs were found in pipework of 9–12%Cr steels. They gave the following recommendations relating to repair welding:

(i) avoid welding procedures that may cause a strongly creep soft HAZ. The HAZ will always have a creep-soft part (the fine-grained and intercritical parts). That part, as well as the whole HAZ, will be wider with increasing heat input. A wider HAZ decreases the creep strength of the weld. Coarsening of carbides may be more pronounced during the weld thermal cycle with higher heat input, particularly in the HAZ. The coarsening of carbides results in lower creep strength, so the use of a heat input as low as practically possible is recommended. To avoid wide HAZs, it is also recommended to use interpass temperatures of #300uC. (Higher than 250uC would be more appropriate for 9%Cr steels.)
(ii) select design solutions that minimise system stresses as (even) small system stresses reduce the creep life of weldments significantly. If the repair was not dictated by the presence of appreciable system stresses, it is very important to reduce or eliminate these
(iii) select materials for weld repair that are somewhat overmatched (creep strong) in relation to the remaining service-exposedmaterial. Hardness testing, replica testing or the extraction of boat samples, service data and chemical analysis give understanding of the creep strain resistance in the aged material. New weld metal should be tested corresponding to an ordinary weld test. These investigations will help selection of weld repair material and procedure. Typically a new weld metal of the same material as the service-exposed ones will give a strongly overmatched repair
(iv) wide and medium deep geometry of the excavation is optimal for the lifetime of a weld repair. Deeper or full repairs are, however, necessary if cracks, microcracks or creep cavities are present deep in the material. Repair of at least the whole width of the original weld, including the HAZs is always recommended, even if the damage is local, such as cracks in the HAZ only at one side of the weld.

Table 4 Summary of the advantages and disadvantages of the use of ferritic and Ni-based consumables for repair

Ferritic Nickel-based
Advantages Advantages
Minimal material discontinuity Lower residual stress
Ultrasonic inspectability Intrinsic resistance to hydrogen-assisted cracking
Long-term integrity Better fracture toughness No specialised weld deposition techniques required
Disadvantages Disadvantages
Higher residual stress Transition joint limitations
Risk of hydrogen-assisted cracking Problems of inspectability
Poor fracture toughness  
Specialised weld deposition techniques essential  

Technology gaps and future trends

The use of creep-resistant steels in existing plant and their popular application in current builds mean that there will be a risk of continued recurrence of type IV cracking for the foreseeable future, and therefore, in addition to the usual monitoring strategies, detection and repair strategies will be required for at least the next 30 years. In China and India, where (in the light of the rapid expansion in electric power generation in these countries) hundreds of new power plants have recently been built using creep-resistant steels, there is likely to be considerable scope for the detection, monitoring and repair of type IV cracking. Worldwide standards need to reflect the knowledge that now exists. There remains a real danger arising from fabrications built to old code, including seam-welded elbows, with apparently no allowance for the diminished cross-weld creep strength. While the cross-weld stress to which girth welds are subjected is half that to which seam welds are subjected, failure to limit system stresses can, as discussed earlier, make girth weld vulnerable to failure by type IV cracking. Clearly, there is an urgent need to increase awareness of the severe degradation in creep strength that welding introduces, and to reflect this information

in standards. When improvements are effected in parent steels through boron and possibly other alloying additions, and cross-weld strength in the type IV region increases significantly, improved weld metals will be required, to avoid the weld metal becoming the creepweak region. It is therefore now time to undertake an intensive programme of weld metal development, involving long-term creep rupture testing.

There is growing awareness of type IV issues in Europe, the USA and Japan, where measures are being taken to mitigate the risks, and research to develop resistant steels is under way. Following further development, boron-containing steels are likely to be put into widespread use in fossil-fuel plant at temperatures up to 650°C. Long-term cross-weld data are therefore required for these newer steels.

In weldments in the 9%Cr steels, martensite forms both in the weld metal and in the CG and FG HAZs. The transformation from austenite to martensite occurs at a sufficiently low temperature that the associated expansion will partially offset the contraction occurring during weldment cooling. The effect of this expansion is to reduce the level of residual stress in this class of steels. While this is a topic that has not received extensive attention, it is unlikely to be an area of concern in fabrications subjected to a PWHT. However, if strategies are ever devised for carrying out repair welding without a PWHT, residual stresses will assume greater significance.

In the near future, N+T treatments to minimise the reduction in cross-weld strength could become more commonplace, and may also come to be used for seamwelded pipe. This may bring a different set of problems associated with temperature control and the effect of such treatments on weld metals in particular. Electron beam welding (which removes the complication of filler metals) may offer productivity improvements for large repetitive manufacture of some parts. One area where EB welding may be of use, especially when coupled with an N+T treatment, may be for the manufacture of large dished ends of pressure vessels by joining forged petals together. However, further research to prove such technology is required.

An interesting recent development is the revised normalising and (double) tempering treatment for grade 92 steel that was first modelled and then carried out by Yin and co-workers and Morris and co-workers.153,154 These later authors showed that a new heat treatment schedule (normalising at 1150°C and double tempering at 660°C for 3 h+3 h) extended the (parent steel) creep life by more than a factor of 3, with a further substantial increase arising from a reduction in the C/N ratio. (This demonstration of the strong influence of normalising temperature provides experimental support for the neural network prediction of Francis et al., mentioned in the section on ‘Neural network analyses’.102) It will clearly be of interest to establish if the improved creep performance is sustained for times appropriate to power plant service. So far, no cross-weld data have been reported, but it will also be of interest to establish the weld strength factor appropriate to 105 h tests.

To meet future energy needs, base-load power is most likely to be met by nuclear stations, with renewable energies contributing. In countries where coal is abundant, coal-fired stations will continue to make a major contribution, particularly as carbon capture and storage becomes viable. For coal-fired stations and combined cycle gas turbine units, the degradation of alloy steels in high temperature service as a result of combined creep and fatigue is likely to become increasingly important. Continued effort is required with data needed not only on grades 91 and 92 steels, but also on the emerging alloys.

If the pilot plant operating at > 700°C shows good progress, then more dissimilar welds between ferritic and nickel-based alloys will become commonplace, and this may bring new problems, including the cross-weld creep strength and non-destructive testing difficulties. The welding of dissimilar metal joints poses special problems in the selection of filler metals and PWHT temperatures. During PWHT and early service, carbon will diffuse to the higher alloy steel from the lower alloy steel, thereby creating a carbon-depleted zone in the latter, which is creep-weak. The issue has been researched by several authors, including Roman et al., who obtained the lowest cross-weld creep strength of a variety of dissimilar grade 22 to grade 91 joints where the grade 22 was buttered with grade 91 weld metal.155 In his extensive tests, which included elevated temperature exposure simulating service, Allen concluded that grade 22 weld metal was the most appropriate consumable for such joints.156 A table giving recommended weld metal compositions for dissimilar metal joints is given in AWS D10•8-96.157 A newly developed versatile welding consumable, EPRI P87, a nickel-based consumable alloyed with 9%Cr in order to minimise the carbon diffusion that occurs during welding, PWHT and service of dissimilar welds, has been reported.104,158–160 When joining Cr–Mo steel components with buttered joint faces that have been subjected to PWHT to other similarly prepared Cr–Mo steels or to stainless steels, the completed joint can be left as-welded.

In the light of the comments made above, important areas for future activity, including research, appear likely to be the following:

(i) further changes in standards to reflect the low cross-weld creep strength of creep-resistant steels, and the implications for fabrications, particularly those in steam service
(ii) scaling up the use of boron-containing steels, in the light of research findings of the apparent absence of the ‘type IV’ degradation in crossweld creep strength
(iii) new steels strengthened by fine Z-phase dispersions should be further investigated, especially the effect of welding them
(iv) the consequences of NzT treatments following welding merit further study, as does the EB welding of thick section units that can be fabricated offsite
(v) further investigation of the use of higher normalising temperatures than those currently in use should be carried out, together with revised tempering treatments to investigate whether improvements can be achieved in long-term creep strength
(vi) the consequences for creep life of thermal cycling merit further study, particularly for grades 91 and 92 steels and emerging steels, to Abson and Rothwell Review of type IV cracking of weldments in 9–12%Cr creep strength enhanced ferritic steels International Materials Reviews 2013 VOL 58 NO 8 465 ensure that any detrimental effects are understood and allowed for in design
(vii) dissimilar welds between martensitic steels and Ni alloys are coming into increasing focus in light of the target 700uC operating temperatures. Increasing use appears likely to be made of the new Ni-based consumable EPRI P87. Continued research is therefore required to gain a more complete understanding of the benefits that it provides. The prospect of using it for the repair of creep-resistant steels without PWHT merits investigation.


From a review of the literature particularly with respect to type IV cracking in 9–12%Cr creep-resistant steels, the following conclusions have been drawn.

  1. Type IV cracking has become a significant problem for steel in creep service. Laboratory tests have shown that it has generally increased in severity as the creep strength of parent steel grades has increased over the last 40 years.
  2. The mechanism of type IV failure is strongly linked to the relative overageing of precipitates in the intercritical and fine grain regions of the HAZ. The welding thermal cycle causes a degree of premature aging before service and some softening of this HAZ region. These fine grain regions of HAZ microstructure promote localised diffusion and expedite precipitate growth during service. The resulting creep weak region accumulates strain in a continuous narrow band of the HAZ, approximately parallel to the fusion boundary, leading to void nucleation, cracking and finally to low ductility fracture.
  3. Despite the occurrence of type IV failures in lower alloy steels, grade 91 steel and other creepresistant steels were employed with inadequate attention paid to cross-weld properties. Hence, many instances of type IV cracking have been detected in high temperature plant. Failures have occurred, and continue to occur, as a result of inadequate allowance for the presence of a creep-weak HAZ region and the influence of service stresses.
  4. The extent of the disparity in creep strength between the parent steel and the type IV region of the HAZ increases as the creep life increases, i.e. as the applied stress decreases. Hence, the extrapolation of results from laboratory tests to likely service lifetimes must be carried out with due care and attention. The inclusion of data collected from long term cross-weld creep specimens that fail by the type IV mechanism are essential before any extrapolation is carried out for the purposes of lifetime predictions.
  5. The introduction of new ASME rules in 2008 addressing the poor creep strength of welded boiler components should help to eliminate the premature occurrence of type IV damage in new fabrications.34 However, the recent code changes require the use of thicker components that are more expensive to manufacture and they may suffer from higher residual, system and thermally induced stress levels. The increased thickness to mitigate type IV cracking over component lifetimes and to comply with code requirements, negates many of the advantages that these steels were designed to offer. As a consequence, at the time of writing, the current grades of creep-resistant steels available are no longer perceived to offer such an advantage over older, low alloy counterparts.
  6. The introduction of new alloys with controlled boron and nitrogen levels appears to be the most promising way forward in terms of mitigating type IV failure in the long term. In the short term, normalising and tempering treatment of welded components will become more widespread, with final site welds either in areas of low system stress or of greater section size.
  7. The most desirable goal in the repair of creepresistant steels and type IV damage in them is the removal of the need for PWHT. Solutions to this problem may include controlled deposition techniques derived from those used for lower alloy steels, and also the use of Ni-based consumables. The criterion for the need for PWHT and the effect on creep properties is not clear.

Appendix 1

Digest from EN 13445-2:2002 (E) Issue 35 (2009-01) Unfired pressure vessels. Materials

Weld creep strength reduction factor (WCSRF): factor to account for possible creep strength reduction at the weldment

4•2•4•2 Creep properties of weldments: Creep properties of weld joints subjected to stresses normal to the weld may differ significantly from those of the base material.

For the design of vessels in the creep range, this is taken into account in EN 13445-3 by making use of a weld creep strength reduction factor zc obtained from tests on weldments. If no data are available, a default value of zc is used.

An acceptable method to determine zc by cross-weld tests is given in Annex C (see also Ref. 161).

Annex C (informative). Procedure for determination of the weld creep strength reduction factor (WCSRF)

The WCSRF will be taken as 1 when all the following conditions are fulfilled by the steel manufacturer:

  1. Stress rupture tests on weldments made on specimens of the same steel products as used in the vessel and which are comparable as regards consumable shall be carried out according to the European Creep Collaborative Committee (ECCC) Recommendations.162
  2. Two test temperatures shall be selected within a range of ±30°C about the mean design temperature. At each of these temperatures, creep tests shall be carried out at stresses selected to give durations up to one-third of the creep design life (typically 1000, 3000, 10 000, 30 000, 60 000, 100 000 h, etc.). It has to be shown that the lower limit of the achieved creep values of the welded joint are not lower than the lower accepted scatter band (-20%) of specified mean values of the creep strength of the base material according to the materials standard. However if the failure is located in the heat affected zone (HAZ), extrapolation is not allowed without further testing at longer times showing no further apparent decrease. In this case extrapolation may be made by a factor equivalent to the factor Abson and Rothwell Review of type IV cracking of weldments in 9–12%Cr creep strength enhanced ferritic steels 466 International Materials Reviews 2013 VOL 58 NO 8 showing stabilised conditions used in these longer tests.
  3. When no cracking in the HAZ has been found in the tests prescribed above, an additional set of tests at a higher temperature shall be made with the value of the Larson–Miller parameter (LMP) equal to or greater than that at the extrapolation point This testing shall be made to confirm that the location of the failure does not change from the base material to HAZ. The temperature shall ideally be no more than 50°C greater than the higher temperature test in C.2 (in order to avoid an unacceptable modification of the microstructure). The stress shall lead to a minimum testing time of 10kh. The temperature and testing time shall be selected so that the creep time–temperature parameter (TTP), e.g. LMP in these tests is at least the value at the extrapolation point (time and temperature). A minimum of three samples shall be tested. The fracture location of the creep specimens shall be checked by microscopic examination.
  4. If fracture location of the creep specimens in C.3 is within the base material, the WCSRF may be taken as unity for a time equal to the time achieved in the tests in C.2 multiplied by a maximum of three.
  5. When the creep strength properties of cross weld specimens fall below the minimum value given in the scatter band a specific weld reduction factor can be used based on the ratio of the average value of the creep strength compared to 80% of the mean value of the base material.

Appendix 2

Estimation of remaining life


Similar to other engineering assessment work, the first step in high temperature component life assessment is to determine the likely damage mechanisms involved in the deterioration of the material. The creep of most components is characterised by continuum damage (creep rupture or creep strain). However, in some cases fracture mechanics dominates the total life of the component; examples of this scenario could be thick components such as headers and high pressure/temperature steam piping. Once the applicable damage mechanisms are identified for each part of the system (pure creep rupture, creep fatigue, creep crack growth, etc.), an appropriate assessment procedure, including nondestructive testing, continuous monitoring, fitness-forservice, etc. can be adopted to ensure the integrity of component in a given operating interval. Welds that are subjected to cross-weld service stress or to high system stresses are potentially at risk of failure by type IV cracking. Hence, consideration or the potential to type IV cracking will figure in any life assessment of systems in creep service.

The following initial screening inspections are recommended, and these can be used in addition to and before other techniques. (Note, however, that each method may be usable only over a particular temperature range.)

(i) visual survey for bulging, sagging and general deformation, particularly in weld region
(ii) dimensional checks – ultrasonic thickness measurements, tube strapping (for diametric growth of tubes)
(iii) crack detection – for surface-breaking flaws at accessible welds, header ligaments, etc., using magnetic particle inspection, dye penetrant and/ or electromagnetic techniques to determine defect size. Ultrasonic inspection can be used for suspected internal cracks
(iv) thermographic inspection through peep holes to check for hot spots on tube banks/boiler sections, to focus inspection sample locations
(v) examination of thermocouple data from process operators to check for hot regions, to focus inspection and sampling.

A variety of techniques are available for further examination, including the preparation and examination of acetate replicas to establish the extent of any creep cavitation; hardness measurement, small-scale punch creep testing on extracted samples; creep rupture testing on extracted samples; the continuous monitoring of creep strain; intermittent measurement of creep strain and microfocus X-ray. The comparatively recent development of equipment for the removal by electrodischarge machining of a small slab of material facilitates damage assessment and remaining life assessment, as the sample can be used for acetate replica preparation and hardness, small punch creep, conventional creep (albeit with a miniature sample) and ultrasonic testing.163 A critical review of the assessment of creep damage in steels employed in the power generation industry has been carried out by Sposito et al.66 In the light of shortcomings of the use of replicas for damage assessment in the secondary and tertiary stages of creep, they concluded that ultrasonic and potential drop techniques appear to be most promising, but that further research is needed before they are fully mature for deployment in the field.

Several techniques are available for monitoring creep strain, and thus estimating remaining life. Hardness measurements, discussed below, are not sufficiently sensitive to allow prediction of the creep behaviour of weldments, and thus the risk of type IV cracking. The remaining techniques, if deployed in weldment regions or across the HAZ, have the potential to do so.

Remaining life assessment from hardness measurements

With the caveat noted in the section on ‘Hardness’, namely that hardness minima across the HAZ do not always represent the most creep-weak region, the remaining creep life can be estimated from hardness readings. However, it should be stated that using hardness values to determine remaining creep life is extremely specific to each individual heat of steel, and should therefore be used as an continuing monitoring technique from one inspection to the next. If the material hardness is determined at the start of its service life, the approach discussed in the section on ‘Hardness’ (equation (8)) can be applied to steel in elevated temperature service, with hardness being measured at room temperature during a shutdown. Allen and Fenton and Allen have developed an approach for estimating remaining life from hardness measurements.164,165 The model requires the determination of creep rupture strength data that is normalised by dividing by the high temperature flow stress.

Damage assessment using acetate replicas

In view of the need to have an acetate sheet in contact with the metal surface, this technique can only be used on high temperature plant during a shutdown. The preparation of acetate replicas and their examination (usually in an optical microscope), is a well-established technique. However, care must be exercised to ensure that the surface region being sampled is representative of the underlying bulk material and subsurface damage. This principle is most likely to be violated where it is desired to establish the extent of creep cavitation in a highly localised region, such as the HAZ. Particularly for a double V weld, the presence of creep cavitation near midthickness may not be monitored readily from the surface, as there will not normally be a similar extent of creep cavitation at the outer surface.

According to Jaske and Viswanathan, the guidance given by the classification of Neubauer and Wedel listed below, which was derived from extensive observations on steam pipes in German power plants, is used by many plant operators in their steam pipe assessment programmes. 166,167

(i) class A: isolated cavities, requires no remedial action
(ii) class B: oriented cavities, requires reinspection within 1•5–3 years
(iii) class C: microcracks, requires repair or replacement within 6 months
(iv) class D: macrocracks, requires immediate repair or replacement.

Remaining life assessment from creep measurements

There are two primary methods by which creep measurements can be made, both of which require extracted samples. Indentation creep measurements can be made, typically on a small slab of the test material.168 Data generated are compared, for example on a Larson–Miller plot, with established data for the test material to determine a converted stress. A similar approach is taken if sufficient material is extracted to permit the machining and testing of uniaxial creep rupture specimens.169

Continuous monitoring of creep strain

Continuous monitoring requires the attachment of strain gauges. Online displacement measurements at elevated temperature are typically undertaken using capacitance devices, which usually monitor the strain in the parent steel.170 The trend in strain accumulation under multidirectional conditions is similar to uniaxial, i.e. the rate decreases during primary creep before accelerating in a tertiary stage.

Intermittent measurement of creep strain

An initial image of the surface of the material of interest is taken, and a detailed comparison made with later images, permitting the microstrain to be determined. Image correlation methods are more suitable for monitoring the HAZ in a weld.171 Morris et al. discussed the autoreference creep management and control (ARCMAC) system used by E.ON UK.172 They report that using a light-emitting diode light source and appropriate software, the equipment can capture strains as low as 500 microstrain which, they state, is more than twice the expected strain on steam pipe after 2 years’ service. A similar device, speckle image correlation analysis (SPICA), which requires a light source and a video camera, has been used extensively by KEMA over several years. The method involves making an optical fingerprint of a given surface in order to compare it with another image recorded later. Local and integral strain can be calculated. Evaluation criteria are based on the results of tests conducted on test specimens; for HAZ, the criteria are based on the elongation of this particular area. By using electron speckle pattern interferometry, flaws are detected by examining interference patterns.173 This technique is reported to be able to detect deformation and displacement, and allow strain measurement down to 1025.171

Laser profilometry is used for measuring the inner profile of piping and accurate measurements of tube diameters are possible, allowing the detection of incipient creep damage by measuring a 2% diameter increase.174

Ultrasonic testing

For life prediction of super heater/reheater coils, the metal service temperature can be estimated by in situ measurement of the thickness of steam-side oxide scale, which is effected by measuring time-of-flight of high frequency ultrasonic waves.175 The data are used in the software program developed by BHEL for life prediction. 176 Bhattacharya mentioned the use of A-scan for a similar purpose in boiler components in Indian industry.177

Ultrasonic or radiographic techniques may be able to detect the physical damage due to creep (cavities or cracks), and thus detect advanced stages of type IV cracking. For improved imaging, phased array ultrasonics or time-of-flight deflection may be preferred or microfocus X-ray equipment with digital screens may be used. Acknowledgements The authors are grateful to Dr A. K. Bhaduri and Dr K. Laha, for supplying data used in plotting Fig. 17 and to colleagues at TWI, including B. J. Cane, C. J. Ablitt and J. R. Rudlin for their help and advice. The contributions of the referees, in suggesting many improvements to the manuscript, are also gratefully acknowledged.


The authors are grateful to Dr A. K. Bhaduri and Dr K. Laha, for supplying data used in plotting Fig. 17 and to colleagues at TWI, including B. J. Cane, C. J. Ablitt and J. R. Rudlin for their help and advice. The contributions of the referees, in suggesting many improvements to the manuscript, are also gratefully acknowledged.


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