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Advances in welded creep resistant 9-12%Cr steels (October 2007)

   
D J Abson, J S Rothwell and B J Cane

Paper presented at 5th International Conference on Advances in Materials Technology for Fossil Power Plants, 2-5 October 2007.

Abstract

Research on high chromium ferritic materials for high temperature power plant components generally concentrates on the properties of the parent steel. Weldments, however, are often the weak link, leading to premature failures and associated forced outages and high maintenance spend. Clearly, consideration of the creep performance of weld metals and associated heat-affected zones (HAZs) in these materials is important. Despite this, relevant weldment creep rupture data are not commonly available, and weldment creep rupture 'strength reduction factors' are not always known. This paper provides comment on the available information on parent materials, and highlights the need for the assessment of the creep performance of weldments. Strategies for increasing HAZ creep rupture strength are reviewed, and some available weldment data are considered. Less conventional welding processes (GTA/TIG variants and EB welding) appear to provide improved creep performance of weldments. They therefore merit further study, and should be considered for welding the new steel grades, particularly in supercritical and ultra-supercritical applications.

Introduction

World electricity demand is increasing relentlessly, with many parts experiencing very rapid growth. At the same time, since power plants are a major source of greenhouse gases, there is growing pressure to reduce emissions per unit of electricity generated. Whilst nuclear power and alternative or renewable energy generating systems are likely to be increasingly introduced, IEA projections of new plant-builds indicate a reliance on fossil fuels over the next 20 to 30 years. Electricity generated from coal currently accounts for around 40% of the world's generating capacity and, because of its abundance in many of the world's developing countries, it is likely to remain the dominant fuel in growing economies such as China and India for the foreseeable future.

With this situation, there is a clear incentive to increase the efficiency of coal-fired plants. Sub-critical plants typically operate at efficiencies of 35%, whereas supercritical (SC) plants have efficiencies up to 46%, and ultra-supercritical (USC) plants achieve efficiencies around 50%. Such efficiency improvements provide significant reduction in greenhouse gas emissions. More than 240 high efficiency SC plants are in operation worldwide. World Coal Institute figures for 2005 indicate that China, for example, has 22 SC units, supplying around 14 GW of electricity. 24 USC plants also exist, with units in Denmark, Germany, the Netherlands, Japan and the U.S.A., and many are at the planning stage elsewhere. The concept of carbon dioxide capture is now adding further importance to the need for higher efficiencies, since the integration of carbon capture plant gives an efficiency penalty of around 10%. These SC and USC plants (with provision for integration of carbon-capture) are thus clear candidates for the next generation of coal-fired plants.

SC and USC plants operate at higher temperatures and pressures than conventional sub-critical plants, as shown in Table 1 [1] . High temperature materials technology has consequently become the major challenge, particularly for USC plant. The reality of USC plant has been made possible by major developments in Europe, U.S.A. and Japan in 9 to 12% Cr ferritic steels for operation up to 650°C. The evolution of steels for power plants is shown in Table 2 [2] , and also in Figure 1. [3]

Table 1: Temperatures and pressures for different types of coal-fired plant [1]

 Temperature (°C)Pressure (bar)
Sub-critical 538 167
Supercritical 540 - 566 250
Ultra-supercritical 580 - 620 270 - 285

Table 2: Evolution of creep rupture strength for ferritic steels [2]

YearsAlloy Modification600°C /10 5 h Creep Rupture Strength, MPaExample AlloysMaximum Metal Use Temp., °C
1960-70 Addition of Mo or Nb, V to simple 12Cr and 9Cr steels 60 EM12, HCM9M, HT9, Tempaloy F9, HT 91 565
1970-85 Optimization of C, Nb, V 100 HCM12, T91, HCM2S 593
1985-95 Partial substitution of W for Mo 140 1 P92, P122, P911 (NF 616, HCM12A) 620
Emerging Increase W and addition of Co 180 1 NF12, SAVE 12 650 2

1 In a later paper by Masuyama [3] , 140MPa was revised to 130MPa and 180MPa was revised to 150MPa.
2 Masuyama [3] added MARN and MARB2 steels (10 5 h creep rupture strength 150MPa).

Fig.1. Chart showing the progressive development of 9-12%Cr steels [3]

Fig.1. Chart showing the progressive development of 9-12%Cr steels [3]

The opportunities offered by the advanced steels in terms of USC plants, however, may not fully materialise if plant reliability suffers. Poor confidence in the long term performance of welds, particularly in components such as steam piping, headers and valves, is a major safety and economic concern. The phenomenon of type IV cracking is recognised as a primary failure mechanism that limits the life of components operating at elevated temperatures.

Type IV cracking has plagued the power industry for around 30 years. Unfortunately, with increasing precipitate-strengthening in the advanced ferritic steels, the effect on lifetime may be exacerbated. In other words, the weldment-parent creep strength ratio, WPCSR (commonly called the weld creep rupture 'strength reduction factor') can decrease, i.e. worsen, with increasing parent metal creep strength. Type IV cracking can thus erode the benefits achieved in parent metal creep strength. Unless the problem can be reduced or overcome, for example by placing welds in lower stress or lower temperature locations, there will need to be more conservative design philosophies and remnant life assessments based on parent metal strength, which realistically account for long-term WPCSR values.

This paper considers possible solutions to the problem of type IV cracking, including improvements in the welding process, and the need for long-term weldment data as a prerequisite for any new high temperature design approach using advanced ferritic steels.

Parent steel developments

Whilst substantial improvements have been made over the last four decades in the development of 9-12%Cr creep-resistant steels, [4] the steam temperatures in recent and projected thermal power stations are now approaching the upper temperature limit for ferritic steels. An understanding of the microstructural factors that control long term creep rupture strength is vital for further improvements in creep rupture strength. Two basic alloy design philosophies are possible. One is to adopt nominally 9%Cr as the base, recognising that the alloys will have inadequate steam oxidation resistance, and will require cladding in order to operate at temperatures of the order of 650°C. Alternatively, the poorer creep rupture strength of 10.5%Cr steels [5] might be improved through judicious choice of alloying. The improved oxidation resistance of higher chromium steels would eliminate the need for cladding. In all cases, it is essential to avoid the formation of delta ferrite, since its presence is detrimental to creep rupture strength. [6]

In designing further alloys based on a nominally 9%Cr composition in the European COST 536 programme, Hald [7] stated the basic design ideas as:

  • Avoid Z-phase formation (with low Cr contents)
  • Achieve Laves phase strengthening
  • Add other nitride formers (Ti, Ta, Zr and Hf)
  • Optimize the boron and nitrogen contents

Creep damage in power plant steels is associated with grain boundary precipitates. In most Cr-containing ferritic steels, grain boundary precipitates, such as M 23 C 6 , provide favourable nucleation sites for creep damage, in the form of grain boundary cavities and micro-cracks. An increase in the proportion of intragranular MX type particles, such as VN, is beneficial. Hald [7] reviewed the development of 9-12%Cr steels. He noted that the poor long term creep rupture strength of some of the recently-developed alloys demonstrated that the ideas behind the alloy design have been based on a lack of understanding of the long-term microstructure stability of this class of alloys. Microstructural explanations for the improved creep strength of the new 9-12% Cr steels should be found in mechanisms which retard the migration of dislocations and sub-grain boundaries, and thus delay the accumulation of creep strain with time.

Among recent developments, including some on experimental alloy systems, Yin and Faulkner [8] showed that ion implantation of hafnium into thin foils of a 9 wt-%Cr ferritic steel can markedly improve the creep properties of the material. The addition of hafnium markedly increased the volume fraction of MX precipitates, and completely eliminated the formation of the common grain boundary M 23 C 6 particles.

Semba and Abe [9] showed the beneficial effect on creep rupture strength of a progressive increase in boron up to 0.0139% in their 0.08%C-9%Cr-3%W-3%Co-0.2%V-0.05%Nb-0.008%N alloys. Steels with boron or nitrogen contents at higher levels lead to the detrimental formation of coarse BN precipitates. The solubility limit of BN at 1150°C is expressed by the equation:

log(%B) = - 2.45 log(%N) - 6.81    [10]

Figure 2 shows how this equation defines a boundary beyond which coarse precipitates form. Although nitrogen is an important 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 steel maker. Taneike et al, [11] employed a solution treatment temperature of 1,300°C to take titanium carbo-nitrides into solution, but 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. Optimizing the solution treatment temperature is an important aspect of producing modern B-containing steels.

Fig.2. Solubility limit for BN at 1150°C for high Cr steel, with the corresponding information for the Fe-B-N system from Fountain et al superimposed [9]

Fig.2. Solubility limit for BN at 1150°C for high Cr steel, with the corresponding information for the Fe-B-N system from Fountain et al superimposed [9]

Type IV failure and its consequences for creep rupture life

Creep rupture studies on modified 9-12%Cr steels have been largely focused on the rupture strength of the parent steels. Much less attention has been given to the behaviour of weldments. The failure of high chromium steels in power plant components is often associated with failure in the 'type IV' region. In short term, high stress laboratory tests on welded joints, failure generally occurs in the parent steel or weld metal. However, at lower stresses and after longer times, the failure location switches to the outer region of the visible HAZ. This has been classified as a type IV failure, and the location identified as the type IV region. The classification of the different failure locations has been discussed by Brear and Fleming [12] and Francis et al. [4] Type IV cracking is the most troublesome in-service failure classification facing the industry today.

In 9%Cr steels, Cerjak and Letofsky [13] found that the transition to type IV cracking occurred after test durations of ~10,000h. In their article on 91 grade steel weldments, Brear and Fleming [12] examined equations to describe parent steel and type IV creep rupture failures in grade 91 steels. The intersection of these two equations predicts the transition from parent metal failure (at high stresses and short times) to HAZ failure (at lower stresses and longer times). They found good agreement between experimental data and the equations describing HAZ behaviour due to Nath and Masuyama. [14]

Abe [15] carried out creep rupture tests on simulated HCM12A specimens that had been heated to temperatures in the range 800 to 1,000°C, to simulate different parts of the HAZ. He reported that, while hardness had its minimum value after heating to temperatures near the Ac 1 , the creep rupture time in tests carried out at 650°C had its minimum value in specimens heated to temperatures near the Ac 3 . From their results, it appears that at high stresses (140MPa), failure preferentially occurred in specimens heated slightly below the Ac 3 i.e. intercritically heated, whilst at lower stresses (<120MPa), the minima shifted to specimens heated to just beyond the Ac 3 i.e. the fully transformed fine grained specimens. This apparent shift in failure position with stress may explain why both regions of the outer HAZ have been associated with type IV failure.

It has been reported that type IV cracking of the weldments can reduce creep life, compared with parent design life, by as much as a factor of five. [16-18] When considering creep rupture strength, Nath and Masuyama [14] reported that, at 600°C, 10,000h extrapolated data for grade 91 weldments indicated a 37% disparity between the parent steel and the HAZ, i.e. a WPCSR of 0.63. However, it should be recognised that the extent of the disparity increases progressively with increasing log (exposure time), as shown by Sawaragi [19] for 2.25%Cr-1%Mo weldments. In creep rupture tests at 538°C on tubular modified 9%Cr-1%Mo specimens, a progressive reduction in WPCSR from 0.94 to 0.93 to 0.90 for rupture times of 10 1 , 10 3 and 10 5 h, respectively, was reported by Blass et al. [20] In a study of grades E911, 91, 92, and 122 weldments, in the UK-funded Fourcrack programme, it was concluded that the weldment creep rupture strength 'falls towards a floor value of about 60% of the parent strength in the longer term'. [21]

Parent steel and cross-weld creep rupture data for 92 grade steel, tested at 650°C, have been presented by Abe, [22] Figure 3. Similar data have also been presented by Abe and Tabuchi [23] for grade 122 steel. The important features of this figure are that, for long exposure times, the parent steel creep rupture strength falls below a linear extrapolation of the short term data, and that the disparity between the parent steel and the cross-weld creep rupture strength increases progressively with increasing life, i.e. the WPCSR decreases.

Fig.3. Creep rupture data of Abe [22] for parent steel and cross-weld specimens of 92 grade and a boron-containing steel tested at 650°C

Fig.3. Creep rupture data of Abe [22] for parent steel and cross-weld specimens of 92 grade and a boron-containing steel tested at 650°C

The first issue, i.e. the parent steel creep rupture strength, has been addressed by Japanese researchers, who incorporated an increased slope for a second straight line fitted to long term creep data, obtained at stress levels less than 50% of the 0.2% offset yield strength. [24,6] They have called this method 'region splitting'. It is clearly essential, when extrapolating short-term parent steel test data to projected long service lives, to ensure that appropriate allowance is made for the increasing fall-off in parent steel creep rupture strength with time.

Concerning the second issue, i.e. the weldment behaviour, WPCSR values estimated from the data in Figure 3, are plotted in Figure 4, along with similar data obtained for grades 91 and 122 from a variety of sources. The linear extrapolation suggests that the WPCSR appropriate for a 30 year design life at 650°C for this steel is of the order of 0.27! Furthermore a linear extrapolation may be optimistic. This is clearly at variance with, and much more alarming than, the conclusion of the Fourcrack programme cited above, even though 92 and 122 steel grades were included in both studies. The information in Figure 4 reveals that, in addition to measures needed for extrapolation of parent material data, allowance must also be made for the progressive decrease in WPCSR, when extrapolating short term data to longer times.

a) Larson-Miller constant of 36
a) Larson-Miller constant of 36
b) Larson-Miller constant of 20
b) Larson-Miller constant of 20

Fig.4. WPCSR values derived from creep rupture data of cross-weld or simulated fine-grain HAZ specimens tested at 650°C for grade 122 [23] , grade 92 [22] and grade 91. [25] A linear extrapolation is tentatively assumed, up to a 30 year life. A single data point for grade 9, from ASME code case N-253-14 [26] , is represented here as a 100,000h value. Equivalent times at 600°C are shown, using a Larson-Miller constant of

In all cases, it should also be recognised that, particularly for small diameter test specimens, at least part of the apparent fall-off in creep rupture strength at long exposure times will be attributable to the loss of cross-section through oxidation, unless a suitable protective atmosphere is provided. [27]

Masuyama [28] and Kimura et al [29] noted that changing the constant in the Larsen-Miller parameter has a strong influence on the value of the parameter, and thus on the estimated creep rupture life at a different temperature. Kimura et al [29] have shown how more accurate fitting of data can be achieved in the type IV region by using a different constant. Therefore, also included in Figure 4, on a second abscissa is the equivalent time at 600°C, estimated using the Larsen-Miller parameter with an assumed constant of 36 ( Figure 3a) and 20 ( Figure 3b). This demonstrates the strong effect of changing the Larsen-Miller constant on the lifetimes of weldments at temperatures different from that of the test data available. Clearly, the test temperature and the creep rupture life must be reported whenever WPCSR values are quoted. If predictions are made for temperatures other than the test temperature, there must be confidence in the applicability of the parametric constant chosen.

In agreement with this, data obtained by Kubon and Sobotka [30] for grade 91 steel show that the ratio was 0.77 after 10,000h at 600°C, but decreased with increasing time and temperature. Even with this modest diminution of extrapolated creep rupture strength, for a 91 grade parent steel creep rupture life of 10,000h, there is a dramatic reduction in the HAZ (type IV) rupture life to only 1,700h.

The sizing of component thicknesses in the ASME boiler and pressure vessel construction codes [31-33] , is normally based on closed form equations which include an allowable stress for the material of construction and an efficiency factor (E). The allowable stresses are supplied by ASME (2004), and are based on the parent material properties at the design temperature. Most commonly for creep service the allowable stress is taken to be 67% of the average creep rupture stress of the parent material at 100,000 h. The allowable stress does not consider the loss of strength due to the presence of weldments. The efficiency factor varies between the different ASME construction codes. However, Section I (2004) states that E is 1 for seamless and welded cylinders or the ligament efficiency. Hence, for components designed to ASME section I with no ligaments, there will be no allowance for weld creep rupture strength reduction.

It appears from the information presented above that premature weldment failure is likely wherever weldments in fabrications from 91 grade and similar steels operating at temperatures around 600°C are subjected to stresses that are of similar magnitude to the parent steel design stresses. It is likely that utilities that have used these steels in high pressure steam service at such temperatures with inadequate allowance for the presence of weldments are 'sitting on a time-bomb', and that monitoring and perhaps in-service stress measurements are required urgently to avoid catastrophic premature failures.

Improving the WPCSR and avoiding type IV failure

Through parent steel developments

In view of the importance of failure in the type IV region during long-term service, it is inevitable that increasing attention will be given to the creep rupture properties of the HAZ. One approach to improving the creep rupture strength of the type IV region is to modify the parent steel composition.

For 91 grade steel tested t 650°C, Iseda et al [34] reported an improvement in both the parent steel creep rupture strength and the WPCSR, as nitrogen was increased, for times up to 40,000h. However, lower nitrogen levels would clearly be required if the beneficial effects of increased boron levels are to be exploited.

The single most effective step towards the elimination of type IV failure and its negative effect on the WPCSR has been the addition of boron. As well as improving the parent steel creep properties, boron has been shown to improve the WPCSR substantially. [35,22,34] Tabuchi et al [35] added from 90 to 180ppm B to low nitrogen (<0.0017%), 9%Cr steels. They tested cross-weld specimens, and reported that 90-130ppm was the optimum level to improve both parent steel and HAZ creep rupture strength. A similar conclusion was reached by Kondo et al [36] , who studied welded joints and simulated HAZs. The beneficial effect of boron was attributed to it combining within M23C6, and suppressing its coarsening, thereby reducing the minimum creep rate. [37-38] It also prevented the formation of a fine-grain HAZ, which adds to the evidence pertaining to this region being the most susceptible to type IV cracking.

The substitution of tungsten for Mo has also been reported to be very effective in enhancing the creep rupture strength of the HAZ. [39] However, this route of enquiry does not appear to have been addressed by other investigators.

Through welding process development

A novel approach to improving the creep rupture strength of the type IV region, as noted by Bell [40] , is to carry out a partial tempering of the parent steel prior to welding, and completing the tempering treatment as the post-weld heat treatment, as described by Sikka. [41] This approach almost doubled the creep life, compared with conventionally heat treated 9%Cr steel weldments in the short term (<10,000 hours); [42] it would be interesting to see the effect over a longer test period. Other process variables suspected to have a strong influence on the type IV creep life of weldments have been highlighted by Francis et al. [43] Among the variables shown to influence type IV cracking, as determined by Baysian neural network analysis, were pre-heat temperature and PWHT time.

A further strategy, which could presumably be combined with the half tempering approach, is to join modified 9% Cr steels by welding processes that create a very narrow HAZ (and therefore a shorter thermal cycle, and less time for over ageing of precipitates). Abe [15] reported that electron beam welding, which produces a very narrow HAZ, had a beneficial effect on the creep rupture performance of the type IV region, with more recent data confirming the behaviour; [23] see Figure 5. The approximate HAZ widths of electron beam and GTA/TIG welds were 0.5mm and 2.5mm, respectively, and the cross-weld creep rupture life of EB welds was approximately twice that of GTA/TIG welds. [15,23] However, consideration should be given to any adverse effects on toughness of EB welds in such steels. [44] A number of alternative processes, including EB, narrow gap GTA/TIG, and friction welding are currently being investigated in one of the creep research programmes at TWI within the European COST programme.

Fig.5. Creep rupture data for grade 122 weldments produced with different welding processes [23]

Fig.5. Creep rupture data for grade 122 weldments produced with different welding processes [23]

Albert et al [45] have shown that simulated fine grain HAZ and real weldments behave significantly differently in terms their creep strain. 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'.

Weld metal development

Welding consumables have been developed with matching or near-matching properties, for welding 92 grade 9%Cr steels by manual metal arc and submerged arc welding [46-48] and flux-cored wire welding [47-48] However, where creep rupture properties are presented, they are generally marginally inferior to those of the 92 grade parent steel. In view of the lower creep rupture strength of the HAZ at long durations, the creep rupture strength of the weld metal is probably of little consequence.

For their series of experimental weld metal compositions, Vanderschaeghe et al [47] showed the detrimental effect of increasing oxygen and also boron content on impact toughness. However, few long-term creep rupture test data are available, and so it will be of interest to see how the various consumables perform in long-term tests.

Limited creep rupture data, from tests at 600°C have been presented for experimental rutile flux-cored wires, for which the optimum Ti addition was ~0.06%. [49] For welds containing 9%Cr-1% Co~0.06%Ti~0.003%B, the maximum test duration was ~9,000h, with a creep rupture strength ~40% greater than that of 92 grade parent steel. Despite the short test duration, it is a dramatic improvement in creep rupture strength that warrants further qualification. An unusual feature of the failure was the extensive ductility (~80% reduction of area) displayed. The reason for the substantial enhancement of creep rupture strength from the Ti addition has not yet been elucidated, but is assumed to result from a change in the character of some of the precipitating phases.

Whilst it would be undesirable to employ a weld metal that over-matched the parent steel creep rupture strength by such a large margin, it is nevertheless vital to press ahead with the development of weld metals, in order to be able to weld the higher creep strength parent steels that are currently emerging. The development of improved weld metals will become particularly important as the problem of type IV cracking diminishes.

Concluding remarks

In view of the decrease in WPCSR values with increasing creep rupture life, the design of existing plant that operate at temperatures close to 600°C should be reviewed, and appropriate action taken if the assumed WPCSR values were non-conservative. Where the deployment of new steel grades is concerned, long term (>10,000h) test results should be taken into account not only for the parent steel, but also for the weldments to be used. Extrapolation of short term data obtained outside the type IV regime is greatly misleading, and possibly dangerous. The change in WPCSR over time needs to be predicted accurately for the type IV region, so that appropriate life-time projections can be made. It is clear that the test temperature and the creep rupture life must be reported whenever WPCSR values are quoted, and that care must be exercised in manipulating data determined at different temperatures.

Research should be aimed at increasing the WPCSR and increasing further our understanding of the mechanisms by which type IV degradation occurs. To this effect, revised parent steel compositions will need to be combined with alternative welding processes to achieve optimal performance. Results for EB welds appear to give promising improvements in the HAZ creep rupture strength ratio, and should be explored further, together with other alternative welding processes. R&Dmp;D programmes to address these issues are being initiated at TWI and work is ongoing within the European COST program, with emphasis on commercial fabrication, as well as performance improvement.

As continued improvements are made to the creep strength of parent steels and of the HAZ region, it is clear that further development of welding consumables is required, to ensure that the weld metal is not the creep weak region in long-term service.

Conclusions

  1. WPCSR values (commonly called weld creep rupture 'strength reduction factors') are not constant, but decrease with increasing creep rupture life.
  2. There is evidence that type IV behaviour can be improved by incorporating boron.
  3. Type IV behaviour can be improved by welding with processes producing a short thermal cycle, thereby producing a narrow HAZ, together with attention to weld preparation, preheat and PWHT.

Acknowledgements

The Authors would like to acknowledge Paul Woollin and Ian Partridge of TWI for their contributions to the paper.

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