A novel approach to characterising the mechanical properties of supermartensitic 13 Cr stainless steel welds

A novel approach to characterising the mechanical properties of supermartensitic 13 Cr stainless steel welds

Materials Science and Engineering A 384 (2004) 83–91 A novel approach to characterising the mechanical properties of supermartensitic 13 Cr stainless...

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Materials Science and Engineering A 384 (2004) 83–91

A novel approach to characterising the mechanical properties of supermartensitic 13 Cr stainless steel welds A. Griffiths, W. Nimmo, B. Roebuck, G. Hinds, A. Turnbull∗ Materials Centre, National Physical Laboratory, Queens Road, Teddington, Middlesex, TW11 0LW, UK Received 2 February 2004; received in revised form 17 May 2004

Abstract Novel measurements of the tensile properties of the weld metal, heat-affected zone (HAZ) and parent plate of supermartensitic 13 Cr stainless steel welds have been made by taking sections from each region and testing with a miniaturised tensile testing machine. Measurements were made on welds produced with a super-duplex stainless steel and a matching filler. The tensile tests were conducted at various temperatures and emphasised the need to account for the temperature sensitivity of the mechanical properties of the individual regions while conducting 4-point bending on welded specimens. Scanning microhardness measurements were also made and significant sub-surface hotspots were highlighted, which could have been inadequately characterised by conventional line-scan measurements. Residual stress measurements on a specimen taken from a super 13 Cr weld with a super-duplex filler indicated local regions of high compressive stress, up to 350 MPa in magnitude. This needs to be taken into account while interpreting the hardness results. © 2004 Elsevier B.V. All rights reserved. Keywords: 13 Cr stainless steel; Welds; Mechanical properties; Hardness; Residual stress

1. Introduction A range of low carbon 13 Cr martensitic stainless steels has been developed recently for welded pipelines transmitting oil and gas. A key element is resistance to environment assisted cracking. The environment inside the pipe contains chloride, CO2 and H2 S, which may result in stress corrosion cracking due to either absorption of hydrogen, generated by corrosion, or anodic dissolution of sensitised regions. The external surface of the pipe will be under cathodic protection and hydrogen embrittlement may occur due to absorption of hydrogen. To provide a framework for evaluating the resistance of super 13 Cr steels to environment-assisted cracking under these varied environmental conditions, the mechanical properties of two types of super 13 Cr weld have been characterised. Novel measurements of the tensile properties of the weld metal, heat-affected zone (HAZ) and parent plate ∗ Corresponding author. Tel.: +44-20-8943-7115; fax: +44-20-8614-0436. E-mail address: [email protected] (A. Turnbull).

0921-5093/$ – see front matter © 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2004.06.015

have been made using a miniaturised tensile testing machine. The primary purpose of these measurements was to provide input data for finite element modelling of the strain distribution in welded 4-point bend specimens [1]. The microhardness distribution was measured using the scanning microhardness technique to provide information about sub-surface hard spots. Selective measurements of residual stress were also made.

2. Experimental method 2.1. Material The material tested was a 12 Cr–5 Ni–2 Mo steel supplied in the form of a pipe with an external diameter of 203 mm and a wall thickness of 11.1 mm. The chemical composition is shown in Table 1. The 0.2% proof stress at 23 ◦ C was measured in a conventional tensile test as 701 MPa and the UTS was 843 MPa. Girth welds were produced by tungsten inert gas welding with a super-duplex stainless steel filler and

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Table 1 Chemical composition of supermartensitic stainless steel (mass %) C Cr Mo Ni Mn Si Cu Nb N V Al S P

0.009 12.13 2.06 5.3 0.62 0.14 0.02 <0.003 0.0064 0.044 0.003 <0.001 0.019

with a matching filler. Some of the welds were then post-weld heat treated at 650 ◦ C for 5 min followed by an air cool. The heat input/travel speed in welding was 0.6–1.1 kJ/min; the welding position was 3G; the shielding gas was argon; the filler material was about 2 mm diameter and typically, about five passes were made. No prior heat treatment of the superduplex stainless steel weld metal was undertaken. 2.2. Miniaturised testing A miniature mechanical testing system developed at NPL [2] was used to obtain stress–strain data for the weld metal, HAZ and parent material over a range of temperatures. Specimens of 2 mm × 1 mm cross-section and 40 mm length were machined from the parent plate, HAZ and weld metal, with the orientation of the specimens such that the 40 mm length of the specimen was tangential to the circumference of the girth weld. The specimens taken from the weld metal and HAZ of the matching weld were removed from the central 7 mm of the pipe wall. Ideally, the precise location of the specimens machined from the weld with super-duplex filler should have been recorded, but this was not done. However, to achieve such a specimen length given the curvature of the pipe, the mid-section of the specimen, at which the key measurements are made, would be in the middle region of the pipe wall or close to the inner surface of the pipe (root of the weld). The specimens from the HAZ were removed at the mid-point of the HAZ. For these welds, the width of the HAZ, as determined by hardness measurements, was about 2 mm. The miniature mechanical testing system uses a computercontrolled DC power supply to heat the specimen, with the temperature controlled with respect to a thermocouple located centrally on the specimen. The grips are maintained at room temperature resulting in a parabolic temperature distribution along the specimen. However, the temperature is reasonably constant (typically within ± 2.5 ◦ C) over the central 2–3 mm of the testpiece, and the strain is measured in this region. Measurement of strain on specimens of these dimensions is a challenge. Typically extensometers are used on gauge

lengths of 10–25 mm but are not appropriate for smaller testpieces. An alternative method of strain measurement, based on change in resistance, is used with the NPL miniature mechanical testing system. The resistance changes significantly during plastic deformation due to a change in the crosssectional area. The strain is determined by measuring the resistance using thin, 0.1 mm diameter, spot-welded Pt-13% Rh wires as potential contacts. The wires are located 2–3 mm apart in the central region of the specimen. Treatment of the elastic behaviour is described later. The load is measured using a load cell with 0.2 N resolution. Tests were conducted at temperatures of 20, 60 and 130 ◦ C at a loading rate of 10 N s−1 . 2.3. Microhardness scanning Microhardness scanning was carried out on specimens taken through the cross-section of girth welds. The specimens were prepared to a 3 ␮m diamond finish, and twodimensional microhardness maps of the weld, HAZ and adjacent parent plate were produced using the NPL scanning microhardness facility [3]. This machine operates by automatically indenting the specimen over an area defined by a grid of measurement points. At each location, indentation depth and load are measured as the load is applied. Measurements were made using a Vickers indenter with a load of 0.3 N at spatial intervals of 0.1 mm over a 220 × 150-point grid. A typical indent diameter would be about 15 ␮m with a typical indent depth of 1.5 ␮m. The loading rate was approximately 0.15 N s−1 . For these welded specimens, the time to complete a scan is large, about 55 h, but the system is fully automated. Microhardness is determined using the principles outlined in the ISO standard for Instrumented Hardness Tests, ISO 14577. The position of the surface is identified by fitting a polynomial to the initial 5% of the loading curve. The displacement at maximum load incorporates elastic and plastic deformation. A linear fit to the top 20% of the unloading curve (representing elastic unloading) is extrapolated to zero load to determine the depth of material in contact with the indenter at maximum load (hc ). This is a measure of the plastic deformation at maximum load. The microhardness, H, (force per unit area associated with plastic deformation) is then calculated using: H=

F , 24.5 × h2c

(1)

where F is the maximum load and 24.5 × h2c is the area of the indenter in contact with the material at maximum load. To provide a comparison between the scanning hardness measurements and conventional Vickers hardness measurements (differences arise because of different indenters, and the residual indenter depth after load removal is measured in the Vickers measurement), microhardness profiling was carried out in the plane of the weld section using a load of

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9.806 N. Indentations were made along a line parallel to the edge of the pipe close to its mid-thickness. The distance between indentations was 0.5 mm. 2.4. Residual stress Measurement of residual stress was conducted on a crosssection through the girth weld of the super-duplex filler weld. The specimen was prepared to a 1 ␮m diamond metallographic finish, and the residual stress was measured using X-ray diffraction. The area sampled for each measurement was 5 mm × 1.5 mm and the depth of penetration of the Xray beam about 10–15 ␮m. In addition, an attempt was made to measure the residual stress in the miniaturised test specimens. Although significant residual stress relaxation was expected in the preparation of these specimens, the possibility of a retained longitudinal component could not be overlooked. In determining the residual stress, we need to be able to follow the angular position of a diffraction peak coming from a specific (hkl) plane. In practice, the diffraction peak appeared as a composite of a number of overlapping peaks. Although analysis yielded a relatively small value for the residual stress of about 39 MPa, the confidence level was very low. 2.5. Austenite content The austenite content was measured in two specimens taken from the parent plate. The 50 mm × 50 mm × 10 mm specimens were prepared to a 1 ␮m diamond metallographic finish. In addition, one of the samples was chemically cleaned to confirm that the polishing procedure did not cause transformation of retained austenite. The austenite content was then measured using X-ray diffraction according to ASTM E-975 [4].

3. Results

Fig. 1. Stress–strain curves for the super-duplex stainless steel weld at 130 ◦ C.

where σ is the stress, Em is the apparent elastic modulus in the miniaturised tests, and Ec is the elastic modulus measured in conventional tensile tests. The elastic modulus measured for the parent material was used for the parent material, HAZ and weld metal of both types of weld. We would not expect any significant impact of the microstructure on the modulus, and the decrease of the modulus with temperature, determined from conventional tensile tests on the parent material, was only 3%. 3.1.1. Super-duplex stainless steel filler weld The stress–strain curves measured for the parent material, HAZ and weld of a super 13 Cr weld produced with a superduplex filler are shown in Figs. 1–3. Occasional outliers are obtained with the resistance measurement technique, which explains the scatter in Fig. 2 for example. The values of the 0.2% proof stress are plotted versus temperature in Fig. 4. The values of 0.2% proof stress were calculated from an analytical curve-fit assuming a Ramberg–Osgood model, which included a parameter to accommodate any offset from the origin. Relevant fitting parameters are given in an appendix.

3.1. Miniaturised tensile testing The plastic strain measured using the resistance probes, εm , was calculated from the resistance during deformation, Rt using the following Eq. (2):  Rt , (2) εm = ln Rs where Rs is the resistance prior to deformation. However, this analysis does not account for the effects of elastic behaviour. The true strain ε was calculated from the strain measured in the miniaturised tests, εm , using the following equation:   1 1 ε = εm − σ , (3) − Em Ec

Fig. 2. Stress–strain curves for the super-duplex stainless steel weld at 60 ◦ C.

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Fig. 3. Stress–strain curves for the super-duplex stainless steel weld at 20 ◦ C.

Fig. 5. Stress–strain curves for the matching weld at 130 ◦ C.

To provide confidence in the data, the values of the 0.2% proof stress for the parent material obtained from the miniaturised testing were compared with those measured in conventional tests (Fig. 4). There is good agreement with a maximum difference of 20 MPa. This would suggest that the methodology associated with the miniaturised testing is satisfactory. It can be seen from Fig. 4 that the proof stress of the parent material decreases slightly with temperature. However, the proof stress of the super-duplex stainless steel weld metal is more strongly dependent on temperature. At 20 ◦ C, the proof stress is similar to that of the parent material, but as the temperature increases to 130 ◦ C, the value for the weld metal is 80 MPa lower. This relative decrease with temperature is consistent with other published data for super-duplex stainless steel [5]. The proof stress of the HAZ is similar to the parent material at 20 ◦ C and at 130 ◦ C but greater at 60 ◦ C. However, as the result at 60 ◦ C is based on a single specimen, this temperature response of the proof stress may be a consequence of sampling a region of the HAZ of higher strength rather than an effect of temperature on the properties of the

HAZ per se. Some variation in the mechanical properties of the HAZ and weld metal might be expected due to differences in the thermal history associated with individual weld passes.

Fig. 4. Variation of the 0.2% proof stress with temperature for super 13 Cr steel with super-duplex stainless steel weld filler.

3.1.2. Matching weld The stress–strain curves measured for the parent material, HAZ and weld of a super 13 Cr weld produced with a matching filler are shown in Figs. 5–7. The values of the proof stress are plotted versus temperature in Fig. 8. At 130 ◦ C, all three materials have similar strength levels. At 60 ◦ C, the parent material and weld metal have similar strength levels with the HAZ having lower strength. At 20 ◦ C, there is a large variation in the proof stress values with the HAZ having the lowest value and the parent material having the highest value. In view of the similarity in results for the weld metal and parent metal at elevated temperature, the difference at ambient temperature is surprising. It is feasible that experimental variability could account for some variation. However, the HAZ would seem to be consistently lower in yield strength. The microstructure showed a much larger grain size, which may

Fig. 6. Stress–strain curves for the matching weld at 60 ◦ C.

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Fig. 7. Stress–strain curves for the matching weld at 20 ◦ C.

contribute to the reduced proof stress. The austenite content would also be different from the parent material, although not measured specifically in the HAZ. 3.2. Scanning microhardness measurements Microhardness scans for welds with both matching and super-duplex stainless steel fillers with and without post weld heat treatment (PWHT) are shown in Figs. 9–12. The hardness is quoted as universal hardness. To relate this to Vickers hardness, a conventional microhardness line scan close to the mid-thickness of the sample is compared to the corresponding section from the microhardness scan for the super-duplex filler weld. The results are shown in Fig. 13. The hardness values in Figs. 9–12 scale linearly with Vickers hardness values and are higher by about a factor of 2 . It should be noted that the hard regions observed very close to the boundary of the specimen, such as that just below X marked on Fig. 9, are an artefact of the measurement technique as it approaches the free surface and are not genuine values.

Fig. 8. Variation of the 0.2% proof stress with temperature for the super 13 Cr weld with matching filler.

Fig. 9. Microhardness of super 13 Cr weld with super-duplex filler.

A microhardness scan of a super-duplex filler weld is shown in Fig. 9. Generally, the HAZ is harder than the parent material. However, there is considerable variation in the hardness through the HAZ, probably induced by the varying thermal history associated with individual weld passes. The weld metal generally has a hardness intermediate between that of the HAZ and the parent material, although there are local hard spots. In addition, there is a large hard region at the root of the weld. Fig. 10 shows the same scan for a super-duplex filler weld from a pipe subjected to PWHT. It must be emphasised that Figs. 9 and 10 represent specimens taken from different pipes and as such, direct comparison is not ideal. In that context, however, a modest decrease in hardness after PWHT is apparent, notably at the root of the weld.

Fig. 10. Microhardness of super 13 Cr weld with super-duplex filler after PWHT.

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A. Griffiths et al. / Materials Science and Engineering A 384 (2004) 83–91 Table 2 Residual stress measurements in the transverse direction Position Residual Error stress (MPa) (MPa) 1 2 3

Fig. 11. Microhardness of super 13 Cr weld with matching filler.

−111.4 110.4 −139.6

7 14.2 17.4

Residual shear Error stress (MPa) (MPa)

Uncertainty (MPa)

−6.8 39.4 −14.6

10.36 20.9 24.5

1 3.4 3

A microhardness scan for the matching weld is shown in Fig. 11. The dotted lines show the estimated location of the fusion line based on the weld geometry. The pattern of hardness in the weld with matching filler is very different from that in the weld with super-duplex filler. In the matching weld, the hardness is high in the entire weld and HAZ. There is some variation in hardness in the weld region associated with individual weld passes. A similar scan for a matching weld specimen after PWHT is shown in Fig. 12. Here, the effect of PWHT appears more pronounced than for the super-duplex filler weld, as exemplified by the lower and more uniform hardness values. 3.3. Residual stress

Fig. 12. Microhardness of super 13 Cr weld with matching filler after PWHT.

The residual stress in the weld with super-duplex filler was measured using X-ray diffraction (no measurements on the matching weld have as yet been undertaken). The area sampled was 5 mm × 1.5 mm, and the location of the measurements for longitudinal residual stress is shown by the boxes in Fig. 9 marked 1, 2 and 3. The area sampled for the measurement of transverse residual stress was centred at the same location but rotated through 90◦ . The results are given in Tables 2 and 3. The longitudinal direction is transverse to the weld (i.e. horizontal in Fig. 9), and the transverse direction is through the thickness of the weld (i.e. vertical in Fig. 9). The residual shear stress values are perpendicular to the direction of the main residual stress being measured and in the plane of the specimen. The error values quoted in Tables 3 and 4 are based on the error in fitting the data, and the uncertainty value is based on the repeatability of the measurements. At position 2, the transverse stress is tensile. Otherwise, all other stresses measured at the three locations are compressive and in the longitudinal direction, relatively large. The values reported are average values and therefore, the maximum values of the residual stress may be significantly higher than those in Tables 2 and 3. Table 3 Residual stress measurements in the longitudinal direction Position Residual Error stress (MPa) (MPa)

Fig. 13. Comparison of universal hardness values with Vickers hardness values along a line close to mid-thickness of the super-duplex filler weld.

1 2 3

−219.6 −253.0 −353.8

17.6 13.8 21.2

Residual shear Error stress (MPa) (MPa)

Uncertainty (MPa)

1.2 −1 −20.2

26.8 24.8 32

4.4 2 4

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3.4. Austenite content The austenite content measured for the parent material specimen that had been polished to a 1 ␮m diamond finish was 12.2 vol.%, with a standard deviation of 1.6 vol.%. The value for the specimen that had been chemically cleaned was 12.1 vol.%, with a standard deviation of 2.8 vol.%. The similarity of these results confirms that the polishing procedure did not cause transformation of retained austenite.

4. Discussion 4.1. Relationship between tensile properties, hardness, residual stress and microstructure In assessing the significance of the hardness, the effect of residual stress on hardness measurement needs to be taken into account [6]. Conceptually, compressive residual stresses would be expected to result in an apparent increase in hardness, while tensile residual stresses would result in an apparent decrease in hardness. The magnitude of the effect is much debated, but in a recent paper, Gibmeier and Scholtes [6] suggest that the magnitude can be up to 17%. In the present work, the residual stresses in a sample of the weld with a superduplex stainless steel filler were measured at three points, which had high hardness values. At all three locations, the residual stresses were compressive in the longitudinal direction and significant in magnitude, up to 50% proof stress. Of course, the underlying origin of both high hardness and residual stress may be the same, but it is the possible impact of the residual stress on the hardness value that needs to be resolved. We have not established that herein but would simply indicate that the compressive stress could contribute to the magnitude of the apparent hardness in the HAZ, compared to the parent material. It is usually assumed that a region of material with a high hardness value will be more susceptible to cracking, certainly by hydrogen embrittlement. However, if this increased hardness value reflects compressive residual stresses to an extent, rather than a more susceptible material, the sensitivity may be overestimated. The more concerning situation is where tensile residual stresses may lower the apparent hardness but make the material more susceptible to cracking. If hardness is used as an acceptance criterion, this could be misleading. However, the impact of residual tensile stress will depend on the extent to which the material work-hardens. Further work is required to provide quantitative insight. 4.2. Difference between duplex and matching filler welds In the weld with super-duplex stainless steel filler, there are significant variations in the tensile properties and hardness in the weld and the HAZ. However, in broad terms, the parent metal, weld and HAZ have similar proof stress values at 20 ◦ C but as the temperature increases to 130 ◦ C, the proof stress

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of the super-duplex weld decreases much more rapidly than the HAZ and the parent metal. In contrast, for the matching weld filler, the proof stress values for the parent material, weld metal and HAZ are broadly similar at 130 ◦ C, but at lower temperatures, the proof stress of the HAZ would appear consistently lower. The difference between the hardness patterns for the two types of weld is striking. In the super-duplex stainless steel weld, the HAZ is generally hard but with local variations, and the high hardness zones are predominantly near the root. In contrast, the high hardness zones in the matching weld are generally closer to the cap region. Conclusions about differences in the relative magnitude of the hardness in the two types of welds would be premature without more detailed examination at different locations on each pipe. When examined, the microstructure of the HAZ appeared to be superficially similar for the two types of weld. This might be expected, as it is likely to be dependent on the heat input during welding but less on the composition of the filler metal. However, the proof stress of the HAZ is consistently lower for the matching weld and that is not easy to rationalize. If it were simply a sampling issue in relation to the location of removal of the specimen from the HAZ, a more random variation might be expected. However, it may be systematic insofar as the specimens are removed at the same location. This may correspond to a local “soft” region where otherwise the yield strength of the HAZ could be greater. It highlights the primary limitation of reading too much into discrete measurement in testing of welds. It is comparable to relating a single line scan with the scanning microhardness. 4.3. Implications for 4-point bend testing of welds In undertaking this investigation, a key objective was to provide an improved framework for 4-point testing with respect to both the preparation of specimens from a weld with non-uniform properties and the distribution of stress and strain in a welded specimen under load at the test temperature. The hardness scans provide a basis for making informed decisions about the optimum specimen thickness to test, although qualified with the need to ensure that the hardness is intrinsic and not an artefact of compressive residual stress. For example, in the super-duplex stainless steel weld, there is a significant hard region centred at a distance of about 2 mm from the surface of the parent material at the weld root. If the specimens were prepared with the root surface under tension, the stress at this distance from the surface would be considerably lower than that on the external surface of the specimen, with the exact value depending on the thickness of the bend specimen. In that context, assuming that this sub-surface hard region is susceptible to hydrogen embrittlement, it would be prudent to use as thick a specimen as is practically possible to increase the stress at this point. Alternatively, if testing fully machined specimens as distinct from retaining the as-welded

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surface, it would be advisable to prepare the specimens so that the hard zone was close to the surface. In the same context of failure by hydrogen embrittlement, the test duration for the as-welded test might be chosen to ensure that there is sufficient time for hydrogen to diffuse from the external surface of the specimen to the sub-surface hard region. In this specific case, with the sub-surface hard region in the root, this may be less of an issue for corrosion resistant alloys exposed to the internal fluid in the pipe as general corrosion is very unlikely. Localised hydrogen uptake at pitting corrosion sites would be more important, but the hydrogen from such a localised source would be unlikely to accumulate significantly in this apparent hard-spot region. A sub-surface hard zone is more of an issue for the external surface of the pipe when it is cathodically protected. In this case, for the super-duplex stainless steel weld, the hardness hotspots happen to extend to the surface and should be accessible without significant time delay with respect to hydrogen uptake. The hardness of the matching weld is relatively uniform through the thickness of the weld and therefore, the thickness of the specimen is less of a concern. However, it can be seen (Fig. 11) that there is a step difference in height of 0.7 mm across the weld and this will affect the distribution of stress and strain in a 4-point bend test. Such steps are relatively common on welded pipe and add to the variability of data when testing as-welded material. The presence of residual stresses, up to 50% of the proof stress, could clearly affect the applied stresses in any 4point bend test depending on the location at which the specimens are removed. For the weld with super-duplex filler, the residual stresses in the longitudinal direction are compressive and consequently, the applied stress during testing will be considerably less than the nominal applied stress. A knowledge of the distribution of residual stress in 4-point bend specimens is important in interpreting the results of 4-point bend tests. The problem is that the residual stress can vary quite markedly around the pipe and there is reluctance for cost reasons to measure the residual stress in all sections. The effect of residual stress can be eliminated by testing stress-relieved specimens but the heat treatment affects the microstructure of the material. For example, it has been demonstrated [7] that heat treatment at 650 ◦ C for 5 min eliminates the Cr-depleted zones around carbide precipitates in the HAZ of super 13 Cr welds. Four-point-bend constant-displacement specimens are often stressed at ambient temperature and then heated to the test temperature. There is uncertainty about the effect this will have on the strain distribution in the specimen. The tensile data indicate that the proof stress of the parent material, HAZ and weld metal of the weld with the matching filler vary by less than 10% as the temperature increases from 20 to 130 ◦ C and therefore, the effect of increasing the temperature on the strain distribution would not be expected to be significant. However, the proof stress of the duplex weld metal decreases

significantly with temperature. The required deflection (typically to achieve 100% of σ 0.2 at temperature, for oil and gas applications [8] is determined by strain gauges located on the parent plate adjacent to the weld. However, to achieve this deflection at ambient temperature prior to immersion, the material may have to be deformed beyond yield. This will result in work-hardening of the material. It is simply a limitation of adopting constant displacement testing. Finite element analysis is being used to quantify the magnitude of the effect. It emphasises the importance of characterising the mechanical properties at temperature.

5. Conclusions Novel miniaturised tensile testing has provided unique local tensile data for the parent metal, HAZ and weld metal of super 13 Cr welds with matching filler and super-duplex filler. The tensile tests were conducted at various temperatures and emphasised the need to account for the temperature sensitivity of the mechanical properties of the individual regions when conducting 4-point bend tests on welded specimens. In the specific case of the super-duplex stainless steel weld metal, a marked decrease in proof stress was observed with increasing temperature, which would cause a high degree of strain localisation. Microhardness scans have identified sub-surface hard-spots, which may be missed by a line-scan. However, the location of these high hardness zones correlated with regions of significant compressive residual stresses (up to 350 MPa), and these may have contributed to the apparent high hardness values. The location of local hard-spots and regions of residual stress needs to be considered when manufacturing 4-point test specimens from welds.

Acknowledgements This work was conducted as part of the “Life Performance of Materials” programme, a joint venture between the United Kingdom Department of Trade and Industry and an Industrial Group comprising BP, Shell Exploration, KBR, PhillipsConoco, Expert Metallurgical Services, HSE and ChevronTexaco.

Appendix A The stress–strain data from the miniaturised testing were fitted using a Ramberg–Osgood relationship:  σ P3 σ ε = 100 × + , (A.1) P1 P2 where ε = strain in % and σ = stress in MPa. The 0.2% proof stress is then given by σ0.2 = P2 × 10(log 0.2/P3) .

(A.2)

A. Griffiths et al. / Materials Science and Engineering A 384 (2004) 83–91 Table 4 Ramberg–Osgood fitting parameters for weld with duplex stainless steel filler (see Eqs. (A.1) and (A.2)) Material

Temperature (◦ C)

P1 (MPa)

P2 (MPa)

P3

Parent Parent Parent Weld Weld Weld HAZ HAZ HAZ

20 60 130 20 60 130 20 60 130

2 × 105 2 × 105 2 × 105 2 × 105 2 × 105 2 × 105 2 × 105 2 × 105 2 × 105

770 762 704 760 690 623 746 810 749

15 31 18 16 13 16 16 33 22

Table 5 Ramberg–Osgood fitting parameters for weld with matching filler (see Eqs. (A.1) and (A.2)) Material

Temperature (◦ C)

P1 (MPa)

P2 (MPa)

P3

Parent Parent Parent Weld Weld Weld HAZ HAZ HAZ

20 60 130 20 60 130 20 60 130

2 × 105 2 × 105 2 × 105 2 × 105 2 × 105 2 × 105 2 × 105 2 × 105 2 × 105

770 762 704 708 742 707 682 688 680

15 31 18 15 30 28 13 18 22

The fitted parameters, P1, P2, P3, for the weld with the duplex stainless steel filler and for the matching filler are given in Tables 4 and 5.

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References [1] W. Nimmo, A. Griffiths, L. Crocker, R. Shaw, A. Turnbull, NPL Report, MATC(A) 153, 2003. [2] B. Roebuck, J.D. Lord, L.P. Orkney, in: M.A. Sokolov, J.D. Landes, G.E. Lucas (Eds.), ASTM STP 1418, 4th Symposium on Small Specimen Test Techniques, Reno, Nevada, 2001, p. 234. [3] B. Roebuck, M. Stewart, R. Morrell, M.G. Gee, G. Plint, Surf. Eng. 17 (2001) 447. [4] ASTM E-975-00, Standard Practice for X-Ray Determination of Retained Austenite in Steel with Near Random Crystallographic Orientation. [5] J. Charles, Proceedings of Duplex Stainless Steels ’91, Beaune, Les Editions de Physique, vol. 1, 1991, p. 3. [6] J. Gibmeier, B. Scholtes, Proceedings of the 6th European Conference on Residual Stresses, Portugal, July, 2002, Materials Science Forum, vols. 404–407, 2002. p. 349. [7] E. Ladanova, J. K. Solberg, T. Rogne, Supermartensitic Stainless Steels 2002. KCI Publishing BV, Zutphen, Netherlands, 2002. Paper 028. [8] Corrosion Resistant Alloys for Oil and Gas Production: Guidance on General Requirements and Test Methods for H2S Service, European Federation of Corrosion, Publication no. 17, 2nd edition, The Institute of Materials, London, 2002.