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Review of type IV cracking of weldments in 9– 12%Cr creepstrength enhanced ferritic steelsD. J. Abson and J. S. Rothwell* TWI Ltd

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

Abstract

The improvement of thermal efficiency of power plants has provided the incentive for the development of themartensitic–ferritic 9–12%Cr creep-resistant steels. Good progress has been made in developing such steels, whichare 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 ingeneration IV nuclear designs. The high temperature conditions that these steels operate under in fossil fuelled powestations induce type IV cracking. This type of cracking occurs in the intercritical or fine grain region of the heatedaffected zone via a creep mechanism, and results in fractures with relatively little total cross-weld strain. Despite theoccurrence of type IV cracking experienced in lower alloy predecessors, successor alloys have been introduced andwidely used with insufficient consideration given to the consequences of welding them. Unfortunately, the newersteels suffer from reduced cross-weld creep strength due to type IV cracking to a greater degree in the temperaturerange of operation expected of them, and thus many failures by this mechanism have occurred. The subject of type cracking has been an area of active research interest. This review aims to serve as an update, drawing selectively osome of the vast amount of literature that has been published over the last 30 years.

Introduction

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. Thetarget operating temperature for these steels is 650uC, with a common target design life of 100 000 h. Increasinglythe demand for efficiency is linked to efforts to reduce CO2 emissions, in order to meet environmental obligations anminimise any form of punitive environmental tax. Figure 1 illustrates how CO2 emissions vary with thermalefficiency. 1 Government agencies such as the Department of Energy in the USA are funding large, ambitious projecfor the development of ultra- supercritical (USC) power plant designed to operate at 760° C and conversions fromsupercritical to USC opera- tions at 700° C have been carried out in Japan.2Retrofitting of existing plant is a popularoption for improving output with minimal investment, since existing fittings can be used. The inclusion of CO2 captuand 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 operatingtemperatures 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 showpiecedemonstrates achievements in dissimilar metal joining necessary for minimising capital costs.3 Alloy development

programmes that involve the collaboration of many partners from industry and academia feed into the hightemperature plant initiatives and an example of this is the European Co-operation in Science and Technology (COSTprogramme that has developed alloys for use in the AD 700 initiative and other state-of-the-art fossil power plantsuch as the Neurath plant to be constructed in Germany.4 Many of these technically challenging new power plantsincorporate the latest ferritic alloy developments and an increasing amount of austenitic and nickel- based materialswhile more conservative constructions rely on established, coded steels that have been available for some time, sucas grades 91 and 92 that are popular pressure vessel materials, and that represent the current highest performingtypes in the 9%Cr family of steels. In the developing economies, such as that of China, where construction of newpower plant is very rapid, the demand for these kinds of materials has recently been beyond the rate of supply.5 Thhuge demand for such materials continues, despite significant problems experienced with them during fabrication anservice.

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

Table 1 and Fig. 3 illustrate how, over the years, good progress has been made in developing creep-resistant steelswhich are being used particularly in the forged form as tubes and pipes in boilers, but which are also being considerfor 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 rupturstrength is attractive to designers who want to take advantage of the higher hot tensile and creep strength to reducpipe wall thicknesses, and thereby minimise thermal stresses for a more reliable plant. This is increasingly importantfor today’s power plants, which are subject to temperature cycles in an effort to respond to the peaks and troughs odemand, and improve profitability. A schematic illustration of the reduction in pipe wall thickness possible through thuse 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- ingconsequences for the power industry. Also, incorrect heat treatment of grade 91 steel in particular and an initiallydetrimental 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 theheated affected zone (HAZ) of welded joints, particularly in high alloy precipitation- strengthened grades that arepopular today, such as grades 91 and 92. These grades have a tempered martensite microstructure that is solid

solution strengthened, but with much of the strength derived from fine

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

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

Years Alloy modification600° C/105 h creeprupturestrength/MPa

Example alloys Maximum metal us

temperature/° C

1960–1970Addition of Mo or Nb, Vto simple 12Cr and 9Crsteels

60 EM12, HCM9M, HT9, Tempaloy

F9, HT 91 565

1970–1985 Optimisation of C, Nb, V 100 HCM12, T91, HCM2S 593

1985–1995Partial substitution of Wfor Mo

130 P92, P122, P911 (NF 616,

HCM12A) 620

Emerging Increase W and addition

of Co, B and controlled N150

NF12, SAVE 12, MARN, MARB2,MARBN

650

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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 theknowledge 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

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4 Illustration showing possible reduction in wallthickness achievable through use of advanced9%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 andof the extent of long term creep strength degradation, and strategies for improved performance, would constitute a

useful addition to existing knowledge, and would help designers and operators to take the right decisions whendesigning, 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 discusse

Type IV cracking – characteristics and causes

Introduction

The problem of type IV cracking has existed for many years in power plant operated at elevated temperatures. Asearly as 1974, in tests carried out at 550° C, Schuller in Germany reported the low cross-weld strength of weldment

in DIN 10CrMo9-10 (equivalent to AISI A182 grade F22 – 2•25%Cr–1%Mo). 6 He observed that, in tests extending tmore than 104 h, the cross-weld strength fell more than 20% below that of the parent steel. In the UK, many powerstations built in the 1960s and 1970s incorporated a large amount of 0• 5%Cr–0• 5%Mo–0• 25%V steel (BS 3604grade 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 IVcracking. Meanwhile, continental Europe deployed the 12%Cr steel X20CrMoV121 (X20), with modest servicetemperatures in thinner section sizes, and encountered few problems, thanks to an inherent resistance to reheatcracking and a resistance to type IV cracking at the lower service temperatures.7 Grade 91, provided a cheaper andstronger 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 developedcreep-resistant steels to type IV cracking, falls in a more dramatic way than for their lower alloy predecessors. Earlydeterminations of the ratio of cross-weld creep strength to that of the parent steel revealed values substantially belounity. 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 listedvalues for many other materials in elevated temperature service, including X20.) A warning note was sounded by Brhl et al., who measured cross-weld creep strengths that were substantially lower than parent steel values.10,11 Thiand the following warning note of Middleton appear to have been largely to no avail: ‘The operation of super- criticaplant constructed from either grade 91 or X20 steels at y600uC would therefore involve a degree of risk of type IVcracking, the level depending on the medium to long term strength loss found for cross-weld testing, unless a modeoverdesign were applied.’12 From a comparison of known creep rupture strengths and observed failure times of powplant welds, Middleton also deduced the likely level of system stresses. He commented that ‘At 600uC the superiorrupture strength of P91 renders it more resistant to type IV failures than X20 beyond y30 kh. The presence of acreep-weak HAZ leads to a cross weld rupture strength that is severely reduced, however, leading to a much lowertolerance 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 metHAZ and parent steel are subjected to the same strain for circumferential welds, the lower HAZ creep strength of P9weldments is of concern for longitudinal welds, as they are subjected to the same stress as the parent steel (which itwice the cross-weld pressure stress to which the girth welds are subjected).


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Almost all testing in alloy development programmes is for parent steels, and extensive testing programmes have be

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 osimulated HAZs as long ago as 1990 by Middleton and Metcalfe, who carried out creep rupture testing on grade 91samples that had been heated to temperatures between Ac1 and Ac3 (intercritical and austenitic phase fields); seeFig. 5.12 Since then, similar studies have been carried out highlighting the poor creep strength of material heatedabove 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 theparent steel falling in the range from 0•5 to unity.9 However, as discussed in the section on ‘Design codes, publisheweld strength factors and extrapolation of data’, with a few exceptions, such information has been slow to find its winto 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 inJapan where extensive cross-weld testing or the imposition of simulated HAZ heat treatments, followed by creeprupture testing, have been carried out. European institutes that have particular interest in the problem of type IVcracking include the University of

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

ElementGrade91

NF616(grade 92)

HCM12A TB 12M Grade

911Grade122

FB2

C0.08-0.12

0.07-0.13 0.07-

0.140.10-0.15

0.09-0.13

0.07-0.14

0.13

Mn0.20-0.60

0.30-0.60 ≤0.70 0.40-

0.600.30-0.60

≤0.70 0.82

Si0.20-0.50

≤0.50 ≤0.50 ≤ 0.50 0.10-

0.50 ≤0.50 ...

S ≤0.010 ≤0.010 ≤0.010 ≤0.010 ≤0.010 ≤0.010 ...

P ≤0.020 ≤0.020 ≤0.020 ≤0.020 ≤0.020 ≤0.020 ...

Cr8.00-9.50

8.50-9.50 10.00-

12.5011.0-11.30

8.50-10.50

10.0-12.5

9.32

Mo0.85-1.05

0.30-0.60 0.25-

0.600.40-0.60

0.90-1.10

0.25-0.60

1.47

W ... 1.50-2.00 1.50-

2.501.60-1.90

0.90-1.10

1.50-2.50

...

Co ... ... ... ... ... ... 0.96

Ni ≤0.40 ≤0.40 ≤0.50 0.70-1.0 ≤0.40 ≤0.50 0.16

Cu ... ... 0.30-

1.70 ... ...

0.30-1.70

...

V0.18-0.25 0.15-0.25

0.15-0.30

0.15-0.25

0.18-0.25

0.15-0.30 0.20

Nb0.06-0.10

0.40-0.09 0.09-

0.100.04-0.09

0.06-0.10

0.04-0.10

0.05

N0.030-0.070

0.030-0.070

0.040-0.100

0.04-0.09

0.040-0.090

0.040-0.100

0.019

Al ≤0.04 ≤0.040 ≤0.040 ≤0.010 ≤0.04 ≤0.04 ...

B ... 0.001-

0.006 ≤0.005 ... ...

0.0005-0.005

0.0083


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Sn ... ... ... ≤0.010 ... ... ...

As ... ... ... ≤0.010 ... ... ...

Sb ... ... ... ≤0.005 ... ... ...

105 hcreep

600uC 94 (115)* (115) [150

(105 h)] (115) ... ...

rupturestrength

/MPa

650uC 50 (60)* (60) [80 (105

h)]

(65) ...

(): 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, plussubsequent PWHT at 570 or 600C on creep rupture 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 a760C/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 introducedin 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 arestarting to be incorporated into design codes to help prevent premature failures. Initiatives such as the EuropeanCreep Collaborative Committee (ECCC) and the state-funded analysis of data carried out at the National Institute oMaterial Science (NIMS) in Japan have been set up to collect data for design purposes. However, although reliableperfor- mance may now be designed in more easily by way of increased thickness of material, the poor creep strengof welds that leads to type IV cracking remains, and prevents the impressive properties of the parent alloys being


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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 thefabrication of a joint and may be solidification, hydrogen, reheat, temper embrittlement and occasion- allylong-term creep-related cracks. Type IV cracking forms towards the outer edge of the visible HAZ, beside the parentmetal (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 illustratethe different subzones that commonly exist in the HAZ. Figure 10 shows a section of the phase diagram for grade 91steel, and how different parts of

7 Rupture times for HAZ simulated coupons for ASME grade122 steel: open points are furnace heat treated, and closedpoints represent coupons treated in a weld simulator108


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8 Classification of cracking in weldments from Brear and Fleming,16 accordingto 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 cracking occurs towards the outer edge of the visible HAZ, the exact location, in terms of microstructure has beenreported 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

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 HAZmicrostructure 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 temperature–stressregimes may bring about differences in the observed failure location, although the failure location in long-term (lowstress) service is most commonly reported to be the fine-grained heat affected zone (FGHAZ), if stresses transverseto the weld are the most significant. It is worth noting that, in multiple pass welds, the fine-grained and intercriticalHAZ 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 this 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 thetriaxiality is highest.19 This observation clearly has implications for the use of surface replicas to investigate theextent of creep cavitation.

Type IIIa is another long-term creep failure mechanism, associated with carbon migration from the CGHAZ near thfusion 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 thweld and HAZ, which is often due to differences in the Cr concentration. It tends to be more prevalent in butt weldswhere less inspection is focused, and consequently more time is available for diffusion and development of cracks attypically longer durations than type IV cracks. The most commonly associated combination of materials affected

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10 Schematic representation of the subzones of the HAZ corresponding to thecalculated 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 intercriticzone 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 Asdiscussed in the section on ‘Technology gaps and future trends’, EPRI and Metrode Products Limited have developedfiller for the avoidance of this problem which has a lower carbon activity, and is especially useful for welding austenisteels 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 3

000 h.7,15 While much of the earliest type IV cracking detected occurred in heats with low N/Al ratios (1•5), manyexamples 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 proceto 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 heattreatment, 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 lowertemperatures, failures are more common in the weld metal or parent metal. As the temperature is increased and thestress reduced, there comes a point at which type IV creep cracking becomes the dominant mechanism, as illustrateschematically in Fig. 12 and for experimental steel FB2 in Fig. 13. This implies that, where the primary stress is acrothe weld, below a critical or threshold stress level, the creep strain in a weldment is localised, and creep failure occuin the type IV region of the HAZ much earlier than would be

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11 Heat affected zone region of a ruptured creep cross-weld specimen thatfailed 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



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a HAZ region on non-ruptured side. Etched in 1 : 5 v/v Hcl (35%) and distilled water with 0•5–1 g K 2S2O5, 2 g

NH4HF2 per 100 g of solution; b HAZ region beside rupture with an attempt at defining the different regions of theHAZ. Etched in 2•5% picric 2•5% Hcl v/v in ethanol

12 Schematic representation of the reduction in creep rupture strength beyond theonset 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 muchearlier 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 inASTM/ASME grade 91 steel, with a line distinguishing the boundary conditions for type IV cracking. An approximateequation 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 sea

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 historicfailures of pipes that resulted in fatalities. However, longitudinal seams in headers and other pressure vessels areinevitable.

From their experimental measurements and subsequent computation of the creep of an X20 vessel loaded at 610°Cwith an internal pressure of 200 bar, Steen et al. concluded that ‘the normally expected failure mode in a pipe unde

pure internal pressure will be HAZ cracking perpendicular to the hoop direction. 22 The lifetime of the weld joint will fbelow the lower bound of the base material scatter-band, although when no weld efficiency factor is adopted, the timsafety margin may be very small’.

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

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The experimentally observed dependence of type IV cracking on temperatureand 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. fracture

located mainly in the outer HAZ region and featuring limited ductility; those bounding the grey arearepresent failures located in the parent or weld metal, with higher ductility

14 Failure locations for cross-weld samples of grade 91 steel for differentrupture lives and temperatures describing the temperature dependence of typeIV cracking7

poor placement of welds, e.g. on tight radii of end caps, resulted in reduced failure times due to increased axial stres

across the welds.23 Girth welds that were subjected to pure pressure failed due to transverse cracking of the weldmetal or parent material rather than the HAZ region, and the importance of creep ductility in such situations washighlighted, 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. Thematching filler failed by axial cracking in the weld metal, while the higher ductility Nibased filler weld failed in theparent, 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 loadingcondition where bending predominates, and type IV creep strain accumulates in a localised region, failure throughdiscrete 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 twhich the girth welds are subjected. Moreover, the failure of a seam weld is likely to have much more seriousconsequences than the failure of a girth weld. However, the superposition of system stresses can and does allow typIV failure to occur in the HAZs of girth welds, for example those attaching header end caps. Multiaxial constraint isexpected to increase with specimen size and improve creep rupture times. It follows that high strength, low ductilityweld metals and high angle fusion lines should increase constraint and creep rupture times for transversely loaded


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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 penetrationbutt welds, ‘X’ type, perform worse than ‘U’ groove geometries for large-scale (30 mm2) uniaxial cross-weldspecimens, and when tested in hot pressure vessels (seam welds subject to hoop stresses) for grade 122 steel at650°C. The results presented also showed a large reduction in time to rupture for seam welds compared with thelarge uniaxial specimens, although no explanation for this was offered. In contrast, Yaguchi et al. observed similarcreep rupture lives for conventional creep rupture specimens and internally pressurised welded pipe, illustrating theconflicting findings that can found in the literature.19 Their study involved narrow-gap longitudinal seam welds in a shaped preparation in 9%Cr and 12%Cr steels. Type IV failure was reported as occurring in the FGHAZ. The planestress/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 influencof weld angle, and thus the inclination of the fusion boundary to the stress axis (between 0 and 45°), on creep lifewas modelled by Tanner et al., who showed up to a 60% reduction in a creep life, with the minimum at an inclinatio

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 consider several axial (end load) and hoop stress configurations for weld angles between 0 and 37•5° for a weld wit

uniform HAZ properties.29 The analysis suggested an interaction between applied stress state, weld angle and time 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 ahigher 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 weld

as they are subjected to the same stress as the parent steel.10,11 In short-term laboratory tests, the cross-weld andparent creep strengths are usually very 14 Failure locations for cross-weld samples of grade 91 steel for differentrupture lives and temperatures describing the temperature dependence of type IV cracking7 Abson and RothwellReview of type IV cracking of weldments in 9–12%Cr creep strength enhanced ferritic steels International MaterialsReviews 2013 VOL 58 NO 8 445 similar until the onset of type IV failure. With the onset of type IV failure comes adegradation of cross-weld creep performance compared with that of the parent steel under the same conditions. Onplot 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 stresthe difference between parent and cross-weld creep strength widens, so that a progressively greater discrepancyexists at lower stresses and longer durations. The difference between the weldment and parent creep strength at agiven 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 i

the creep rupture strength of cross-weld specimens at time t and temperature T.31 The terminology for describing treduced creep strength of weldments compared with that of the parent steel is sometimes confusing. It is desirable tachieve 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 streng

between heats.32 However, there is also a trend for the WSF to decrease with improved parent performance betweegrades, thereby strongly reducing the effective creep strength of a fabrication compared with a single piececomponent. In some cases, the improved creep strength of the parent material has been shown to be completelycountered 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 thatemploying a higher strength parent steel will bring no further advantage in terms of cross-weld creep performance athe 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 MPacompared with welded joints in grade 91 (that has a lower parent creep strength) that have a creep rupture strengtof 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 fullyaccounted for by design codes until many years after the introduction of steels such as grade 91 into the marketplace. Figure 15 shows creep rupture data for parent and cross-weld grade 92 specimens, and illustrates the pointthat if plant were designed to the allowable stress specified by ASME alone (before 2008), weldment failure could be


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expected beyond 20 000 h of service for welded grade 92 material, if the allowable stress for 100 000 h service wasemployed.34 Several of the standards that relate to out-ofcore nuclear components (ASME Section III Code Case N4the French Code RCC-MR and the Japanese code ETSDG) were reviewed by Bhoje and Chellapandi, who observed th

comprehensive creep SRFs for weldments are provided in the French code RCC-MR.35 Weldments are discussed inSection 4/5 of the R5 assessment procedure. In addition to stating that ‘A particular problem with weldments is thepresence of residual welding stresses’, there is a short section on ‘type IV cracking in 0•5 CrMoV weldments’, whichdiscusses crack growth in the type IV region of girth welds, with no explicit mention of low cross-weld creep strengtThe 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 thatthe 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 clethat this allowance was not sufficiently adequate. Until recently, no factor accounting for poor cross-weld creepstrength was incorporated into more widely used design codes such as the ASME code, where allowable stress isbased upon safety factors for parent material performance. The ASME code relating to high temperature pressurevessels 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 ofdrums and headers has not changed, but the definition of the efficiency factor E has been expanded to incorporate wfor the seamwelded condition

(4)

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 parentmaterial 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•77and post-weld heat treatment (PWHT) (0•5) are illustrated. The w multiplier was introduced into ASME

2008 code in recognition of the poor cross-weld creep performance of joints. 34 This factor effectivelyincreases 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 fac


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E for seam-welded components is equal to w, or the ligament efficiency, whichever is lower. In the case where theweld seam is penetrated by the openings forming the ligament, E is taken to be the product of w and the ligamentefficiency.

The value for w is provided in a table in the code, and is dependent on heat treatment of the component afterwelding, 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 thenormalised and tempered condition, the value falls from 0•86 to 0•77. The introduction of w effectively reduces themaximum stress (S6w), as indicated in Fig. 15, and increases the required thickness for seam-welded components, 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, thecross-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 adjoiningparent 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

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


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17 Creep rupture WSFs for grade 92,115 grade 122107 and grade 91.68 Valuesderived from creep rupture data of crossweld or simulated FGHAZ specimenstested 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, asdiscussed 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 specifiedlinear interpolation between values of 1•0 at 510°C and 0•5 at 815°C, for all materials. Since the 2008 issue (Table302•3•5), values of W range from 1 to 0•5 for different materials and service temperatures. For socalled CSEF steelin the post-weld heat treated condition, it is 0•5 for the temperature range from 510 to 649°C. For weldments in theN & 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 operatingtemperature range than it was for the older less alloyed steels, as indicated in Figs. 17 and 18, which shows anextrapolated 100 000 h WSF of y0•4 for grade 122 (12%Cr) steel at 650°C. Grades 92 and 122 show broadly similabehaviour, namely a stronger downward trend than that shown by grade 91.40 As noted above, this is supported bydata presented more

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 for9%CrMoV steels and 0•5 for 12%CrMoV steels.41 The graphs show clearly that the WSF is not constant, but ratherdecreases with increasing rupture life (decreasing applied stress), and with increasing temperature, reaching levelswell 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 bHolmstro¨m and Auerkari.42 Their plot of stress versus WSF for grade 91 steel at temperatures ranging from 575 to650°C, Fig. 19, shows WSF values falling from ~0•9 to ~0•6 as the stress decreases. Data points added to theoriginal graph are the WSFs given


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Stress versus WSF for grade 91 steel derived by Holmstro¨mand Auerkari42 from 2005 ECCC data, to which the followingdata 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 1

000 h with increasing service temperature. Similar behaviour is reflected in graphs presented by Laha et al.43 Theyshowed 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 WSFcannot be determined as a simple ratio from short-term creep test data. Rather, careful extrapolation must be madeto stresses and lifetimes that are relevant to the intended application. The implication of the poor cross-weld creepstrength and the weld SRFs introduced by ASME described above is that either operating temperatures or pressureswill 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 liketo feature in further standards in the future. Towards this end, reliable extrapolations of WSF are needed, preferablyup to 200 000 h. EN 13445-2:2002 (E) Issue 35 (2009-01), Annex C, ‘Procedure for determination of weld creepstrength reduction factor (WCSRF)’, which is based on VdTU V-Merkblatt 1153, requires testing at stresses selected give durations up to one-third of the creep design life at two test temperatures within a range of ± 30°C of the mea

design temperature.44,45 Hence, if steels are to have a design life of 20 years, then a 7-year testing programme wilbe required. The standard states that ‘if the failure is located in the HAZ extrapolation is not allowed without furthertesting at longer times showing no further apparent decrease (in the WSF)’; see Appendix 1. However, for materialsoperating at the high temperature end of their application ranges, there is ample evidence to show that theweldment/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. Thecriterion 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 an

Serco Assurance, and reviewed by Ainsworth, has a section (Volume 4/5) that provides two methods for calculatingcreep/fatigue crack growth. It incorporates ‘fatigue SRFs’ to modify the stress range, but it does not appear to treatexplicitly the lower cross-weld creep strength compared with that of the parent metal.46 Further discussion of theallowances made in national codes has been given by Ref. 47.

Microstructural degradation and failure mechanism


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Type IV failure is the result of localised strain and void formation in the outer region of the HAZ (FG/IC HAZ). Thiscreep-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 HAZspecimens. Voids nucleate on precipitates mainly at prior-austenite grain boundaries and, at an advanced stage, thevoids coalesce or ‘unzip’ to form a crack. Once this final stage begins, creep crack growth progresses rapidly to failuand immediate repair or replacement is required. Investigation of creep crack propagation in grade 91 steel willultimately help utilities to take decisions on the safest and most economical way to operate, once a type IV crack ha

been detected.48

Creep-resistant steels rely on alloy and precipitate strengthening, with the most important fraction of strengtheningunderstood 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, sometimesvia 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 itcontributed. 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 uponthe main strengthening phases in these steels, due to their low coarsening rates and their ability to pin grainboundaries and dislocations. NbC forms at high temperatures, where it can prevent grain growth and is thereforeimportant during the steelworking operation, as well as during service. VN forms during tempering, precipitatingthroughout 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 yearand is usually accompanied by a sharp fall in creep strength.51 Another precipitate that forms during service is theintermetallic 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 precipitatestrengthening from Laves phase may counteract this to some degree. The range and scale of the phases that exist a

any time are dependent on the temperature of operation and the sequence of precipitation reactions.52,53 Coarseninkinetics 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 steelsFB2, FB6 and FB8 as a function of time to rupture at 650C.FB6 and FB8 both suffered from shortterm onset of Z phasecausing 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 toform after extended creep service durations as the final equilibrium phase. Until recently, this phase was unknown icreep-resistant martensitic–ferritic steels, and none of the thermodynamic databases available even predicted itsoccurrence in these steels. The onset of Z-phase formation after long-term exposure results in a rapid reduction increep properties; see Figs. 20 and 21. Significant efforts towards developing a series of high Cr alloys, with improvecreep 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 developmentprogrammes 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 thoptimum distribution of the precipitates and their interparticle spacing, which the manufacturer generated by judicio


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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 oftensupported by minimum creep rupture strength data from simulated specimens. In this region, where temperatures atypically between 900 and 1100°C during welding,55 complete reversion to austenite occurs. Partial dissolution of thM23C6 phase is expected followed by rapid growth according to equilibrium and kintetics calculations and diagrams f

9–12%Cr steels. MX particles would also be expected to grow at a somewhat increased rate. An in-depth TEM studythe distribution and composition of precipitates within this and other regions of the HAZ via simulated thermal cyclein the as-welded, post-weld heat treated and creep exposed condition, confirmed that the dissolution of M 23C6precipitate 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 furthertemperature exposure during creep tests, the number density decreased rapidly, i.e. growth rates were comparativehigh, in the FGHAZ. Furthermore, they found that the MX particles contained relatively high amounts of Cr, leading tthe conclusion that Cr diffusion (rather than V) was controlling the growth rate of MX particles following greaterdissolution 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 approximately3%Co and 3%W and no boron. The reorganisation of austenite grains is also an important factor as it is the preferresite for M23C6 nucleation. During the welding thermal cycle, the boundaries change position, and the carbide locatio

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 creepresistant 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 backstressdue to Z-phase explains the poor creep strength of alloys that suffer fromprecipitation 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 thisregion, some austenite exists; some precipitates are dissolving, while others are coarsening at relatively higher ratescompared with the surrounding parent steel. Concentration gradients between austenite and a quickly temperingmartensite are formed and a less homogenous microstructure results on cooling. Lee et al. reported a substantially

higher growth rate for M23

C6

precipitates in the ICHAZ.57 The growth of creep cavities that nucleate on these partic

is assisted by the creep deformation occurring in this HAZ region. Francis et al. and Smith et al. quantified the numb

of creep cavities as a function of the fraction of rupture life.24,58 Kimura et al. attributed the low creep rupturestrength 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 recoveof martensite that contains less solid solution strengthening and less effective precipitates at boundaries defined bythe phase interfaces that existed during welding.

In summary, the microstructure of creep-resistant steels is a finely tuned dynamic system, with changes in thecharacter of the precipitates occurring alongside the movement of dislocations, subgrain boundaries and grainboundaries during PWHT and subsequent creep service. The weld thermal cycle imposed upon the HAZ creates asevere disturbance in the finely balanced microstructure, with a severe detrimental effect on creep strength. Althougsimulated specimens demonstrate fairly consistently that the FGHAZ has the poorest creep resistance across the HAin practice, it is difficult to differentiate between FG and ICHAZ when trying to define the exact microstructure that a


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type IV crack has nucleated in and propagated through in a real welded joint.

Hardness

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 measurementsmade 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 aproperly 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 notstraight forward. Some weldments do not display the typical HAZ characteristics as clearly as others. Relationshipsbetween 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 tesspecimen, when normalised by being divided by the hardness H0 of the (aged but not strained) grip region, is as

follows64

H/H 0=0•98-0•15 t/tr (5)

where t is the test duration and tr is the time to rupture. This equation gives an approximate indication of thehardness 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 an

determination of the hardness values H and H0

tr=0•15t/(0•98{H/H 0 ) (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/t r ) (7)

This equation can be written as

H/H0=1-0•14t=t r (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 othis technique in the field.’ 66

In a study of creep degradation in grade 91 steel weldments, Masuyama noted that the lowest hardness in the HAZwas approximately 10 HV below that of the parent steel, and that this hardness difference persisted throughoutalmost 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 bPWHT indicated that the lowest hardness following PWHT is in the intercritically heated specimens(Ac1>temperature


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Plot of Vickers hardness (open symbols) and creep rupturetime (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, theoverall fall in HAZ hardness during service may be used to estimate remaining life in a way similar to that in parentmaterial. 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 thereforethe former that is potentially at greater risk of type IV cracking and that poses the greatest risk to safety (althoughthe weld strength factor will be higher for the seam weld than for a girth weld in the same pipe). However, type IVfailures 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 attentionmoved to weldment performance, once the susceptibility of this steel to type IV cracking was recognised. Type IVfailures 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 IVcracking 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 IVcracking are known. It is also clear that older units designed for base-load operation and used in this capacity overmany years are very susceptible to component failure when they are then cycled regularly. One of the West Burtonfailures 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 bythe low hardness of the grade 91 steel. The steel contained a low nitrogen to aluminium ratio, which gave rise to AlNprecipitates, and reduced the amount of nitrogen available to give the fine MX intragranular precipitates, therebyreducing 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. Ona 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 sharpchange in section was a further contributory factor to an end cap failure after 36 000 h of service. The seriousness othe problem and its frequent occurrence is illustrated by the report of 100 instances of cracking during 2004 inoutages in power stations operated by Innogy/RWEnpower.61 Type IV cracking in grade 91 retrofit headers that havbeen 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 beenresponsible for poor performance, and incorrect heat treatment is also suspected to have contributed in some


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instances. However, as noted earlier, warnings of likely problems from failure in the type IV region, if adequateprovision for the system stresses was not made, were not heeded. Moreover, the construction codes do not yet allrequire that adequate provision be made for the reduced cross-weld creep strength compared with that of the parensteel, 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 includedconsideration of service experience of type IV cracking in girth welds and in seam welds, with separate considerationof steam lines and headers. Seam-welded pipe failures have occurred in the USA, for example the 1985 failure of a inch (760 mm diameter) steam reheat line of the Mohave power station that had been in service for 14 years. Thisfailure, 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 plantwere 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 onemould for the pressing, rather than two). Since the seam welds occur along the extrados and intrados, stresses on twelds arising from flexure of the pipe induced by any temperature fluctuations will be higher on the extrados than ifthe welds were situated on the neutral axis, and will be in addition to the pressure stress. Such a design of elbowrequires 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 91pipe. The best documented failure occurred during commissioning of unit 2 at Datong Power Station in 2006, andresulted in two deaths. The main stem line was not manufactured in grade 91 steel, as had been believed at the timof installation. Similar failures have occurred elsewhere, with an additional four deaths, and there are believed to bean additional 30 plants containing this pipe. Such short-term failures are unlikely to have been the result of type IVcracking; however, they do illustrate the potential danger arising from the use of seam-welded pipe in power plant.

Type IV creep life prediction

Introduction

The ability to predict accurately the creep strength of specimens/components subject to long-term exposure is highlydesirable, since it would avoid the need for longterm testing during the development of creep-resistant steels, andalso allow plant operators to make more confident decisions on maintenance schedules. Despite advances in methodof 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 studiesParent 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 notalways be valid, as different phases may exist, and different mechanisms may be operational. For these reasons, thuse of the Larson–Miller (L–M) parameter to predict performance at different temperatures is not encouraged, as itcan yield nonconservative lifetime predictions. Thus, the L–M parameter, although useful in some instances, do notalways provide a reliable method for the interpolation and extrapolation of data over long durations or interpolationbetween different conditions of temperature and stress. This is fundamentally due to the differences in microstructurdeveloped under different conditions, and over the lifetime of a specimen that can give differences in creep resistancand differences in fracture mechanism between short-term and long-term creep rupture tests or service, as discusse

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 empiricallyvalues very close to this have been found to be quite accurate for 9–12%Cr creep resistant steels. Using values of 3and above, as is sometimes done for Cr–Mo steels,


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23 Creep rupture test results for grade 122 showing adistinctive sharp inflexion in creep rupture stress. Thisinflexion is associated with the presence of delta ferrite andZ-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 inpredicted values of rupture stress from the Manson–Haferd (MH) analysis, than from the L–M or Dorn methods. TheMH parameter is given by

MH=(T-Ta)/(logtr-logta)

It is derived from plots of temperature, T versus logtr, at constant stress, where tr is the time to rupture, and Ta an

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

(10)

where DL is the lattice diffusion coefficient, μ is the shear modulus, b is the Burgers vector, σa is the effective stress

is Boltzmann’s constant, T is the absolute temperature and A and n are numerical constants determined byexperiment.

By splitting the data into regions, and assigning different constants to each part, time/temperature predictionsbecome more accurate and useful, but difficulties still exist when trying to predict behaviour at different temperaturewhere 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, whichoccurs at a stress that is approximately 50% of the 0•2% offset yield strength; see Fig. 23. Creep deformation in th

low stress regime is governed by diffusion-controlled phenomena, while in higher stress regimes, it is controlled bydislocation glide. For dual-phase steels, i.e. those containing δ-ferrite, such as grade P122, this inflection is morepronounced.


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

σ is creep rupture stress, σTS is tensile strength at test temperature, tf is time to failure, Q is activatio

energy for diffusion through the matrix (300 kJ mol

-1

), R is universal gas constant (8•314 J K21 mol

-1

),is temperature89

In their analysis of grade 122 (parent steel) creep rupture data, Wilshire and Scharning noted that, not only are longterm 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 showedthat, if creep rupture stress values for grade 122 steel are normalised by dividing by the tensile strength at the testtemperature, data derived at different temperatures can be superimposed using a function including time to rupturetemperature and the activation energy for matrix diffusion (300 kJ mol21); see Fig. 24. This appears to be a morereliable and fundamental way of relating data collected at different temperatures to each other than the usualapproach of applying the L–M parameter. Taking logarithms yields a straight line plot that allows simple extrapolatioof 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 anapproach 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 Complexpolynomial parameters are being refined using extensive datasets in a rigorous mathematical fashion to give a singlbest fit curve. However, such an exercise is particular to the steel grades involved, and seems to add little to theaccuracy of extrapolations when there is already a vast dataset available. Such a parametric analysis serves only togive the most probable lifetime estimates, based on existing data. Extensive numerical manipulation of data has beeperformed by organisations such as the ECCC. As a consequence of the availability of increased amounts of data andof more reliable extrapolation methods, greater confidence can increasingly be given to long-term creep strengthextrapolations for parent steels. It is important to establish reliable representations of parent steel creep behaviournot only for design purposes, but also to assist in the derivation of more reliable extrapolations of weld strengthfactors; see the section on ‘Design codes, published weld strength factors and extrapolation of data’.

Weldments

The inflection observed in parent steels, as mentioned above for dual phase steels, is relatively minor compared witthe more pronounced change in performance seen at the type IV threshold stress for welded specimens. As notedearlier, 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 interms of time, temperature and stress have been created for grade 91 steel by various authors. Bell gave an equatiothat defined the transition from parent metal to type IV failure of cross-weld creep specimens in 9%Cr steel; see Fig

25a94

(11)


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

represent their experimental data95

(12)

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

(13)

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

Modelling

The limitations of parametric methods have meant that more fundamental physical models based on microstructuraevolution 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 propertiesspecified times and temperatures. The fusion zone itself in multiple pass welds is recognised as an extremelyinhomogeneous entity. This is reflected in the studies conducted by Hyde and co-workers, for example Hyde and Su

whose models reflect the presence of columnar as-deposited and equiaxed reheated weld metal regions. 96 The studof these authors include the determination of longitudinal and transverse all-weld metal creep data for incorporationinto their model.

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


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25 Creep rupture data and the lower bound creep rupture strength for parentsteel 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 tonarrow region in the HAZ, they devised a novel finite element model of the behaviour of a creep-weak layer beside asingle-V weld (simulating the type IV region) inclined to the stress axis. Their model gave good agreement betweentheory and experimental observations when they allowed for relaxation of constraint via the sliding of adjacentelements. They concluded that ‘the creep-weak type IV layer experiences no measurable constraint from the adjacematerial’. In a later experimental programme, Smith and co-workers concluded that cavitation in the type IV region a consequence of grain boundary sliding, leading to relaxation of constraint and multiaxial rupture governed by the


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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 themaximum 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 proportioof a the equation due to Bell;94 b the equation due to Brear and Fleming16 25 Creep rupture data and the lowerbound creep rupture strength for parent steel and type IV cracking, according to two different equations, from Brearand Fleming16 Abson and Rothwell Review of type IV cracking of weldments in 9–12%Cr creep strength enhancedferritic steels International Materials Reviews 2013 VOL 58 NO 8 455 cavitated boundaries, it is likely to be beneficiato keep the HAZ as narrow as possible, which may be one of the reasons for the better cross-weld performance of Ewelds 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 temperaturecreep 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 weldedge angle (0–22°) for longitudinally welded pipes in E911 steel. Their modelling has indicated that using anundermatchi

review of type iv cracking of weldments in 9– 12%cr.pdf

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