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The Influence of Ti, Al and Nb on the toughness and creep rupture strength of grade 92 steel weld metal

   
D J Abson, J S Rothwell and P Woollin

Paper presented at Conference on the Integrity of High Temperature Welds, IOM3, London, 24-26 April 2007.

Abstract

Whilst cored wires are being used increasingly, to achieve higher productivity, the deposits generally contain higher levels of Ti and Al than MMA welds. The present project was undertaken to determine the influence of Ti and Al contents on the toughness and creep rupture strength of FCAW deposits for 9%Cr steels.

Butt welds were produced with six 1.2mm diameter cored wires in which the Al and Ti contents were varied. Mechanical testing of the panels comprised hardness testing, room temperature tensile testing, Charpy and CTOD testing, and creep rupture testing at 600°C.

Weld 0.2% proof strength values ranged from 627 to 757MPa, and tensile strengths from 753 to 871MPa. The medium Ti weld (W2, ~0.06%Ti) with the low level (0.01%) of Al, had the best toughness. The welds with high Al and high Ti (W3)and with high Al and medium Ti (W4), which contain the highest levels of oxygen, had the poorest toughness levels, with 40J temperatures ≥95°C and 0.1mm CTOD temperatures >140°C.

Measured or extrapolated 10,000h creep rupture strengths were typically ~150MPa at 600°C. However, for test specimens from two of the welds [one specimen from the root of weld W4 (high Al & medium Ti) and one specimen from the root of weld W2 (medium Ti)], it was of the order of 200MPa. At a stress of 216MPa, the base-line weld W1 and the low Ti weld W5 survived for only ~100h. Longer rupture lives at this stress level were obtained for weld W3 (~1,000h)and weld W4 (~1,800h). However, the medium Ti weld (W2) survived for a remarkable 9,000h.

1 Introduction

The operators of fossil fuel power plant have achieved improvements in operating efficiency by increasing operating temperatures. This increase has been made possible by using alloy steels with improved high temperature performance, Research on 9 to 13%Cr parent steels has shown that changes in the level of the minor elements can have a strong influence on creep behaviour. Aluminium is claimed to have a detrimental effect on microstructural stability, [4] and high (~0.03%) levels in parent steels have been reported to induce large Al nitride precipitates, thereby impairing the creep ductility. [3] Titanium is beneficial, as it contributes to the formation of fine-scale MX precipitates, thereby delaying dislocation recovery. [8] Whilst deliberate additions of Nb are made to enhance creep strength, increases in Nb in the range 0.1 to 0.5% have also been shown to have a detrimental effect on the creep strength of 9%Cr parent steels. [4]

The development of weld metals for welding such steels has lagged behind the parent steel. In recent work at TWI, additions of 1% Co have been found to be beneficial both to toughness and to creep rupture strength. [1] It is understood that a 1%Co addition is now made to a commercial 92 grade cored wire.

Semi-automatic welding is becoming more attractive for increasing productivity, with cored wire welding [10] assuming greater importance. However, differences between cored wire and MMA deposits include higher levels of Al and Ti in the former. The use of rutile also leads to the associated addition of Nb, which is an impurity in natural rutile. In previous work on modified 9%Cr weld metals, [9] Nb has been shown to be detrimental to toughness, after a 760°C/8h post weld heat treatment (PWHT).

This paper describes a study examining the influence of variations in the levels of Al, Ti and Nb on toughness and creep rupture behaviour of weld metal deposited with cored wires suitable for welding grade 92 steels.

2 Experimental details

2.1 Material and welding

The chemical composition of the 30mm thick grade 91 steel plate used is given in Table 1. A backing bar for each panel was buttered with the consumable that was to be used for the subsequent weld. The root face was 2mm, and the included angle was 55°. The panels were restrained by welded strong-backs.

Table 1 Chemical analysis of the parent plate.

Element, wt%
CSiMnPSCrMoNiAlCoCuNbVN
0.096 0.30 0.36 0.019 <0.002 8.8 0.94 0.14 0.02 0.01 0.07 0.08 0.20 0.040

Four 1.2mm diameter flux-cored wires, with the flux system based on a commercial rutile flux-cored wire (Supercore F92) were used. A metal-cored variant was used in order to allow low levels of Ti to be achieved for the remaining two consumables, Tables 2 and 3.

Welding was carried out using a DC electrode +ve current. Each panel was traversed beneath a fixed welding torch, depositing stringer beads. The pre-heat and minimum interpass temperature was 200°C, and the maximum inter pass temperature 300°C. The welding parameters are summarised in Table 3.

Table 2 Target baseline weld metal compositions. (Target B=0.003)

Element, wt%
CSiMnCrMoNiAlNbTiVWN
0.10 0.30 0.70 9.00 0.45 0.50 0.005 0.03 0.038 0.20 1.70 0.05

Table 3 Target weld compositions, actual welding conditions, and creep data.

Weld codeElement examinedTarget composition
(Wt %)
Welding parametersExtrapolated rupture stress
(MPa)
Reduction of area
(%)
Weld region
AlNbTiTravel speed
(mm/min)
Wire feed speed
(m/min)
Heat input
(kJ/min)
1,000h10,000h
W1 Base-line 0.005 0.03 0.038 240-400 9.3 0.90-1.0 185 150 6
78
Cap
Root
W2 Medium Ti 0.005 0.03 0.06 240-400 9.3 0.73-1.0 190
248
150
212
69 & 78
82
Cap
Root
W3 High Al & high Ti 0.03 0.03 0.10 240-400 9.3 0.70-1.2 210 170 21 & 51
25
Cap
Root
W4 High Al & medium Ti 0.03 0.03 0.06 240-360 8.5 0.78-1.2 240 193 35 & 70
27
Cap
Root
W5 Low Ti* 0.005 0.03 0.01 200-280 8.0 1.2-1.7 185 150 32
79
Cap
Root
W6 Low Ti & high Nb* 0.005 0.05 0.01 200-280 8.0 1.0-1.4 188 150 7
75
Cap
Root
*Metal cored wire

Transverse weld sections cut from each panel were retained for examination in the as-welded condition. The remaining portion of each panel was subjected to a PWHT in air at 760°C for 2h, followed by furnace cooling.

2.3 Mechanical testing

A Vickers hardness traverse, with a 10kg load, was performed down the centre of each of the as-welded and PWHT transverse metallographic sections. One all-weld-metal tensile test specimen, with a nominal gauge diameter of 10mm, was machined from the root region of each weld after PWHT. The tensile specimens were tested at room temperature, with machining and testing according to BS EN10002-1: 2001.

Charpy specimens were machined from both the upper ('cap') and lower ('root') regions of each weld. The specimens were tested over a range of temperatures, according to BS EN 10045-1: 1990. Single edge notched bend (SENB) CTOD specimens were machined, transverse to the weld, from each panel. They were notched through-thickness on the weld central plane, to BS 7448: Part 1: 1991 and BS 7448: Part 2: 1997.

Creep rupture testing was carried out at 600°C on cylindrical all-weld-metal specimens machined from the root and cap regions of each weld. These were M12 specimens, with a thread diameter of 12mm. The specimen gauge length was~110mm, and the gauge diameter was 7mm. The machining and testing were to EN 10291: 2003. Specimens taken from the weld root were all tested at 216MPa; those extracted from the weld cap were tested at both higher and lower stresses.

Fractured toughness and creep specimens were examined in a scanning electron microscope (SEM), and features were analysed by energy-dispersive X-ray analysis (EDX), and by wavelength-dispersive X-ray analysis in an electron probe micro-analyser (EPMA).

3 Results and discussion

3.1 Weld metal characteristics

The weld metal analyses, Table 4, show that reasonable success was achieved in meeting the target weld metal compositions. There were, however, slightly higher carbon levels in welds W1 to W4, and a slightly high nickel level in weld W4 and lower than target levels in welds W5 and W6.

Table 4 Weld metal chemical analyses

Weld code Element examined
W1 Cap Base-line
W1 Root Base-line
W2 Cap Medium Ti
W2 Root Medium Ti
W3 Cap High Al & high Ti
W3 Root High Al & high Ti
W4 Cap High Al & medium Ti
W4 Root High Al & medium Ti
W5 Cap Low Ti
W5 Root Low Ti
W6 Cap Low Ti & high Nb
W6 Root Low Ti & high Nb
As <0.01; Sn ≤0.012; B ~ 0.003, except W1 root (B = 0.0023) and W5 root (B = 0.0025)

Table 4 Weld metal chemical analyses - Continued

Weld codeElement, wt%
CSiMnPSCrMoNiAl
W1 Cap 0.13 0.29 0.58 0.019 0.007 9.7 0.40 0.51 0.010
W1 Root 0.11 0.29 0.56 0.019 0.005 9.5 0.43 0.37 0.011
W2 Cap 0.12 0.33 0.58 0.020 0.006 9.9 0.35 0.42 0.010
W2 Root 0.13 0.34 0.67 0.018 0.005 9.6 0.46 0.44 0.011
W3 Cap 0.13 0.35 0.70 0.019 0.006 9.9 0.40 0.48 0.018
W3 Root 0.13 0.33 0.61 0.020 0.006 9.5 0.61 0.41 0.018
W4 Cap 0.12 0.29 0.62 0.019 0.008 9.6 0.44 0.55 0.028
W4 Root 0.13 0.31 0.62 0.019 0.007 9.6 0.42 0.53 0.038
W5 Cap 0.10 0.29 0.75 0.018 0.008 9.6 0.45 0.34 0.012
W5 Root 0.10 0.31 0.74 0.017 0.006 9.5 0.46 0.32 0.013
W6 Cap 0.10 0.30 0.76 0.018 0.008 9.6 0.46 0.34 0.013
W6 Root 0.10 0.31 0.72 0.017 0.007 9.5 0.47 0.31 0.013
As <0.01; Sn ≤0.012; B ~ 0.003, except W1 root (B = 0.0023) and W5 root (B = 0.0025)

Table 4 Weld metal chemical analyses - Continued

Weld codeElement, wt%
CoCuNbTiVWNO
W1 Cap 0.02 0.04 0.03 0.050 0.22 1.64 0.0414 0.0684
W1 Root 0.02 0.05 0.04 0.031 0.22 1.14 ND ND
W2 Cap 0.02 0.04 0.03 0.055 0.22 1.21 0.0389 0.0651
W2 Root 0.02 0.04 0.04 0.059 0.26 1.39 ND ND
W3 Cap 0.02 0.04 0.04 0.077 0.24 1.44 0.0425 0.0753
W3 Root 0.02 0.05 0.05 0.097 0.24 1.28 ND ND
W4 Cap 0.02 0.04 0.04 0.056 0.23 1.80 0.0384 0.0923
W4 Root 0.02 0.04 0.04 0.063 0.23 1.90 ND ND
W5 Cap 0.02 0.04 0.03 <0.005 0.24 1.69 0.0505 0.0653
W5 Root 0.02 0.05 0.04 <0.005 0.24 1.38 ND ND
W6 Cap 0.02 0.05 0.05 <0.005 0.24 1.69 0.0463 0.0690
W6 Root 0.02 0.05 0.05 <0.005 0.24 1.49 ND ND
As <0.01; Sn ≤0.012; B ~ 0.003, except W1 root (B = 0.0023) and W5 root (B = 0.0025)

Welds W1 to W4 all had a similar shallow profile, and the reheated regions did not show clearly, Figure 1. By contrast, for the metal-cored welds W5 and W6, the beads were deeper, and the reheated regions showed up quite clearly.

Fig.1. Transverse sections typical of as-welded panels: a) W1
Fig.1. Transverse sections typical of as-welded panels: a) W1
b) W6
b) W6

The microstructures of all deposits were predominantly martensitic, with a high density of non-metallic inclusions. There were some differences in the appearance of the prior delta ferrite grain structure, which appeared narrowest in weld W2, Figure 2. 'Blocky' angular grains of delta ferrite were non-uniformly distributed within welds, being present in slightly higher proportion in the capping layer, Figure 3.

Fig.2. As-deposited, as-welded microstructure in the capping layer of weld W2, showing the narrow columnar prior-delta ferrite grains
Fig.2. As-deposited, as-welded microstructure in the capping layer of weld W2, showing the narrow columnar prior-delta ferrite grains
Fig.3. As-deposited, as-welded microstructure in the capping layer of weld W5, showing blocky regions of delta ferrite
Fig.3. As-deposited, as-welded microstructure in the capping layer of weld W5, showing blocky regions of delta ferrite

Close to the boundaries reheated by subsequent passes, a further morphology of retained delta ferrite at the prior-austenite grain boundaries was seen in welds W1, W2, W5 and W6, Figure 4. This figure reveals precipitates, assumed on the basis of the EDX and EPMA studies to be carbides, and narrow bands of delta ferrite at some of the prior-austenite boundaries and triple points. The presence of such narrow regions of delta ferrite was explained by Barnes [2] as resulting from the near-completion of the transformation from delta ferrite to austenite during initial cooling, with subsequent carbide precipitation during PWHT.

Fig.4. Reheated weld metal region in the as-welded transverse section from weld W5 showing dark-etching prior-delta ferrite boundaries
Fig.4. Reheated weld metal region in the as-welded transverse section from weld W5 showing dark-etching prior-delta ferrite boundaries

The weld microstructures are largely insensitive to weld cooling rate, and all were predominantly tempered martensite. However, the cooling rate could have influenced the amount of delta ferrite.

The PWHT tempered the as-welded microstructures, with alloy carbides being precipitated, some of which were coarse enough to be resolved optically. Figure 5 shows the microstructure in the capping layer of W5 after PWHT.

Fig.5. As-deposited microstructure in the capping layer of weld W5, after PWHT
Fig.5. As-deposited microstructure in the capping layer of weld W5, after PWHT

The Vickers hardness surveys gave average values between 400 and 440 HV10, as-welded, and between 239 and 262HV10 after PWHT. Tensile strength ranged from 627 to 757MPa for 0.2% proof strength and 753 to 871MPa for ultimate tensile strength. Elongation ranged from 15 to 19%.

3.3 Toughness data

The Charpy and CTOD data, Figures 6 and 7, show a clear difference between welds W3 and W4 on the one hand, which have poor toughness, and the remaining welds on the other hand, which show better toughness. Linear regression analysis was carried out to investigate the possibility of correlation between both the 0.1mm CTOD temperature T 0.1 (°C) and the temperature for an absorbed energy of 40J, T 40J (°C) to the Al, Ti, oxygen, tungsten and Nb contents. The resulting equations are given below:

T 40J = - 21 + 2149 Al + 428 Ti + 1023 Nb

T 0.1 = 42 + 2897Al + 497Ti

Fig.6. Weld metal Charpy data from the weld cap
Fig.6. Weld metal Charpy data from the weld cap
Fig.7. Cross weld CTOD data
Fig.7. Cross weld CTOD data

Both equations state that Al has the strongest detrimental effect, with Ti having a smaller effect. Thus, the difference in Al and Ti between the cap and root regions may have contributed to observed differences in toughness between the cap and root Charpy toughness specimens. Niobium up to 0.05% appears not to influence cleavage fracture initiation, although it did feature in the equation for Charpy toughness.

3.4 Examination of selected Charpy and CTOD specimens

Examination of the sectioned specimens in the optical microscope revealed differences between the high and low toughness specimens. In Weld W5, the delta ferrite grains of the solidification structure did not feature predominantly. By contrast, this delta ferrite grain structure was clearly visible in the two lower toughness specimens. Detailed examination revealed small islands of delta ferrite along the original solidification cell boundaries.

All CTOD specimens displayed mixed mode, predominantly cleavage fracture. There were few clear differences between the welds, although the lower toughness of welds W3 and W4 was revealed by the larger number of patches of cleavage fracture that formed during fatigue precracking.

4 Creep data

The creep rupture data, for all-weld-metal creep specimens, are presented in Figure 8, together with data from ECCC data sheets [7] for grade 92 parent steel and data for 92 grade Thermanit 616 submerged arc welds. [6,5] Several of the present welds gave actual or extrapolated 10,000h rupture stresses of ~150MPa, which is similar to the creep rupture strength of the 92 grade parent steel. These were weld W1, the two specimens from the cap of weld W2, weld W5 and weld W6. Detailed fractographic examination using an SEM revealed similar microvoid coalescence in all these specimens.

Fig.8. Creep rupture data for all-weld metal test specimens, determined at 600°C. (100h = 4.2 days; 10,000h = 1.1 years). All specimens tested at 216MPa were taken from the weld root; all other specimens were takenfrom the weld cap
Fig.8. Creep rupture data for all-weld metal test specimens, determined at 600°C. (100h = 4.2 days; 10,000h = 1.1 years). All specimens tested at 216MPa were taken from the weld root; all other specimens were takenfrom the weld cap

Specimens containing a medium amount (~0.06wt%) of Ti, ie root and cap specimens from W4 and the root specimen from W2, show a clear improvement over the rest of the data, with extrapolated 10,000h creep rupture stresses ~30% and ~40%, respectively, higher than the 92 grade steel. Approximate 1,000h and 10,000h creep rupture stresses are summarised in Table 3. The data were interpolated, or extrapolated by extending the lines shown, or assuming the same slope as the 92 grade steel, as appropriate. However, it is recognised that there is a lack of data, especially for W2above 100h, and thus further testing would be beneficial. The next best performing specimens were those taken from weld W3, which contained higher levels of titanium.

A surprising feature of the creep rupture data is the extensive ductility displayed, Table 3, particularly by the two specimens that showed the best creep rupture strengths (W2 root, 8,988h, reduction of area (RA) = 82%, and W4 cap, 2,418h, RA = 70%). They failed by what appeared to be ductile tearing, rather than by creep cavitation on the grain boundaries. The reason for this behaviour has not yet been elucidated, but is presumed to be linked to the presence of a fine distribution of titanium carbo-nitrides, and perhaps to finer-scale alloy carbides on grain boundaries than would be the case in the absence of Ti. Metallographic examination of the longitudinally sectioned creep rupture specimens revealed, as expected, alternate bands of as-deposited and reheated weld microstructure. Detailed examination of the microstructures revealed that precipitation had occurred on the lath boundaries during creep testing, to a similar extent in root and cap specimens. The main difference between root and cap specimens appeared to be the greater density of second phase particles in the former, which were revealed more clearly in the delta ferrite regions. Such a difference is as expected, given the longer test period for the specimen from the root region (9,000h, compared with 2.2h for the specimen from the weld cap).

5 Conclusions and recommendations

The low Ti weld (W5, <0.005%Ti), with a low level (0.013%) of Al, had the best toughness. The welds with high Al and high Ti and with high Al and medium Ti, which contain the highest levels of oxygen, have the poorest toughness levels.

The adverse effect of increased Al and Ti on toughness is attributed to their influence both in increasing the weld metal oxygen content, and thus the inclusion content, and also possibly to an indirect adverse effect on the detailed character of the solidification structure.

Several of the creep test specimens, including those from the low Ti weld, displayed creep rupture behaviour that was broadly similar to that of 92 grade parent steel.

The highest creep rupture strength was obtained in welds containing the medium level (~0.06%) of Ti (welds W2 and W4). For test specimens from two of the welds [one specimen from the cap of weld W4 (high Al & medium Ti) and one specimen from the root of weld W2 (medium Ti)], the extrapolated 10,000h creep rupture strength was 212MPa, which is ~60MPa (~40%) higher than the grade 92 parent steel. Creep rupture ductilities were often unusually high, with failure occurring by ductile tearing.

Where creep rupture strength comparable with that of 92 grade parent steel is acceptable, and high toughness is required, the Al and Ti levels in flux-cored wire weld metals should be controlled to ≤0.010%Al and ≤0.005% Ti. However, Ti at a suggested level of ~0.06%, should be incorporated for optimum creep properties.

6 Acknowledgements

The work was funded by Industrial Members of TWI.

7 References

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  3. Cantini S, Ghidini A and Di Gianfrancesco A: 'Innovative Cr-Mo-Co-B steel grades for cast turbine components - Preliminary results of product characterisation'. Proc. Sixth Charles Parsons Turbine Conf. on Engineering issues in turbine machinery, power plant and renewables, 16-18 Sept. 2003, Dublin, Ireland, Ed. A Strang et al., Maney, 2003, 397-410.
  4. Foldyna V and Kubon Z: 'Consideration of the role of Nb, Al and trace elements in creep resistance and embrittlement susceptibility of 9-12%Cr steel'. Proc. Conf. on Performance of bolting materials in high temperature plant applications, 16-17 June 1994, York, U.K., Ed. A Strang, The Institute of Materials, 1995, 175-187.
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  7. ECCC creep data sheets 2005.
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  10. Zhang Z, Marshall A W and Holloway G B: 'Flux-cored arc welding: the high productivity welding process for P91 steels'. Proc. 3rd EPRI Conf. on Advances in materials technology for fossil power plants, Swansea, April 2001, Ed. R Viswanathan.

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