Weld Joint Cracking in Subsea Oil and Gas Systems Pt I

Department of Engineering, University of Leicester Leicester, UK.

M. Milititsky, R. P. Barnett, V. F. Marques, and M.F. Gittos
TWI Ltd. Great Abington, Cambridge, UK

Presented at Proceedings of the ASME 2012 31st International Conference on Ocean, Offshore and Arctic Engineering OMAE2012 July 1-6, 2012, Rio de Janeiro, Brazil

Abstract

In subsea oil and gas systems, steel linepipes (typically API X65) have been welded to 8630M low alloy steel hubs, using a buttering technique and nickel alloy filler materials (e.g. Alloy 625). After buttering, postweld heat treatment (PWHT) is carried out to temper the heat-affected zone of the low alloy steel. However, while the vast majority of subsea joints provide successful service, a small number of in-service failures has had significant ramifications. During PWHT of the buttered joints, carbon migrates from the ferritic steel to the high alloy ‘austenitic’ weld metal. The resulting microstructures and consequent localised hardening cause sensitivity to hydrogen induced cracking.

Subsea systems employ cathodic protection (CP) using aluminium based anodes in order to prevent corrosion. However, while cathodic protection has proved successful as a means of preventing corrosion of the steel components, complications can arise due to hydrogen charging.

This paper describes the detailed characterisation of such dissimilar interfaces from several joint types, including examples recovered from the seabed after exposure under CP and newly fabricated joints, using electron microscopy, microanalysis and nanohardness testing.

1. Introduction

1.1 Background

Gas tungsten arc (GTA) nickel alloy buttering deposits are commonly deposited on low alloy steel forgings, facilitating in-shop PWHT of an assembly to remove residual stress and temper hard heat-affected zones (HAZs), formed in the as-welded condition.  During PWHT of the buttered joints, carbon migrates from the ferritic steel to the high alloy ‘austenitic’ weld metal [1], causing hardening by martensite formation, carbide precipitation [2] and carbon supersaturation [3].  After machining a beveled weld preparation in the buttered layer, a closure weld, typically using shielded metal arc or GTA welding with a nickel alloy consumable, is made to join the assembly to a micro-alloyed steel pipeline. The bevel is positioned such that the HAZ of the closure weld lies wholly within the austenitic buttering layers, so that no further heat treatment is required.

Subsea, the dissimilar joint is subject to CP via aluminium based sacrificial anodes. The majority of joints of this type have been in service without complication, however, a small number have failed at the interface between the buttering and the low alloy steel.

The work presented in this paper is part of an ongoing programme being conducted by TWI Ltd and the University of Leicester on dissimilar joints. The emphasis of part I is characterisation of the joints using both scanning and transmission electron microscopy, as well as nanohardness testing. Part II of the paper will focus on slow strain rate mechanical testing, fracture morphologies and fracture mechanisms from samples tested in a chloride environment under cathodic protection.

1.2 Previous work on the microstructure of low alloy steel - 625 buttering interfaces

As with all dissimilar welds, there are interfacial zones with chemical gradients between the weld metal and the parent metal. At the edge of the melted zone is a region, known as the partially-mixed zone (PMZ), which has been the focus of a number of studies in joints of this kind, [4] and [5]. A concensus on nomenclature has not been reached between authors on the subject, so the various regions, illustrated schematically in Fig. 1, are described below:

 

1. The body centred cubic (BCC) ferritic microstructure of the parent low alloy steel forging which contains inclusions and carbides.
2. A mainly ferritic, decarburised, grain-coarsened heat-affected zone (GCHAZ) immediately adjacent to the dissimilar interface with fingers of fused weld metal that have penetrated into prior austenite grain boundaries.
3. A martensitic region at bead overlap positions (where ‘swirls’ normally occur) that extends into the PMZ. This zone exhibits a body centred tetragonal structure. The formation of martensite is thought to be promoted with lower welding arc energies [5].
4. An austenitic FCC and apparently ‘particle free’ supersaturated carbon solid solution in a region that appears to have solidified in a planar manner. This region has been termed the ‘Featureless zone’, the ‘White Etching Zone’ or ‘Zone Φ’.
5. A region within the PMZ that contains higher atomic number interdendritic particles. This region is austenitic.
6. The bulk metal further away from the fusion zone that exhibits an FCC matrix with a less diluted chemistry.

1.3 Typical Failure modes

When bend specimens are notched at the fusion line (using electro discharge machining) and subject to slow strain rate testing under cathodic protection, a number of fracture morphologies can be witnessed: 

1. In the planar solidification region of PMZs a ‘terraced cleavage’ morphology is found.
2. Regions with a morphology which has an apparently interdendritic character are seen at bead overlap positions. Cracking in these regions appears to correlate with the presence of martensite.

2. Materials and sample Fabrication

Samples from buttered joints with three parent steels were selected: 8630 low alloy steel, F6NM (13Wt%Cr 4Wt%Ni stainless steel) and F65 (C-Mn steel). Each of these parent steels had been TIG-welded with ERNiCrMo-3 (Alloy 625) wire. The nominal chemical compositions of the parent steels and welding consumable are given in Table 1. F65 is a C-Mn steel which does not usually require PWHT; F6NM is a low carbon martensitic stainless steel requiring PWHT to temper hardened HAZ microstructures and the 8630 grade is a high carbon air hardenable material, which will form high hardness HAZs and consequently will require PWHT. Although the exact thermal history of the retrieved subsea joints is unknown, typical industrial heat treatments for 8630 forgings buttered with Alloy 625 are 650°C for 5-10 hours. The F6NM steel was supplied in the quenched and tempered condition. After buttering using a layering weld sequence, it was given an industry standard PWHT.

Table1 – Nominal compositions of welding consumable (Alloy 625) and parent metals.

	ERNiCrMo-3 (Alloy 625) Nominal Wt%	AISI 8630 Nominal Wt%	F65 Nominal Wt%	ASTM A182 F6NM Nominal Wt%
Fe	0.8	Balance	Balance	Balance
Ni	65	0.85	-	4
C	0.015	0.33	0.15	0.02
Mn	-	1	1.2	0.7
Cr	22	1	-	13
Si	0.1	0.25	-	-
S	0.005	0.01	0.005	-
P	0.005	0.01	0.01	-
Mo	9	0.4	-	0.55
Nb	3.5	-	-	-
V	-	0.1	0.1	-
Cu	-	-	0.25	-

Table 2 – History and origin of dissimilar weld specimens.

Sample Reference	Joint Combination (Parent Metal - Buttering)	Origin
Sample 1	F6NM - Alloy 625	Joint fabricated for test work
Sample 2	F65 - Alloy 625	Joint fabricated for test work
Sample 27	8630 - Alloy 625	Recovered from service (9 months)
Sample 35	8630 - Alloy 625	Recovered from service (9 months)
TEM sample	8630 - Alloy 625	Recovered from service

3. Experimental Approach

3.1 Metallurgy across the Dissimilar Interfaces

Transmission electron microscopy (TEM) specimen creation: An 8630 joint buttered with Alloy 625 was selected for in-depth characterisation by TEM. An FEI quanta 3D dual beam focussed ion beam-scanning electron microscope (FIB-SEM) with Omniprobe micromanipulator was used in-situ to create TEM samples approximately 60nm thick. The key steps in TEM specimen preparation using a dual beam system are described briefly below: 

• The ion beam is used to mill either side of a deposited protective platinum strip. The trenches created allow the wafer face to be seen when tilted about eucentric height.
• The ion beam cuts through the thickness of the wafer leaving the side opposite the micromanipulator needle attached (Figure 3a).
• The needle is placed in contact with the wafer and secured by a small platinum deposit. The wafer is then cut free from the bulk sample (Figure 3b).
• The sample is secured to a copper grid using deposited platinum tabs. It is then detached from the needle.
• An electron transparent window is created by thinning front and back faces using a small ion beam current. The sample is measured to ensure transparency and uniformity.

TEM sample analysis: Jeol JEM 2010 and Jeol 2100 microscopes coupled with Oxford Instruments energy-dispersive X-ray (EDX) microanalysis equipment were used to image and perform phase identification from the samples created in the dual beam system. When appropriate, elemental mapping was used.

Microstructure and Chemistry: Sections from each joint combination cut transverse to the weld and the buttering interface were prepared using standard polishing techniques and examined using light and scanning electron microscopy. Semi-quantitative analyses were conducted using EDX in order to determine profiles for Fe, Ni, Cr, Mo and Nb.

Nanohardness Indentation: Due to the fine scale of the various microstructures near the fusion line, nanohardness testing was selected to measure the variation in hardness across the steel-buttering interface. A Micromaterials Nanotester, equipped with a Berkovich tip, was used at room temperature to determine nanohardness. For the F65 and F6NM samples, 5 x 20 arrays were used with an indent spacing of 5µm. For the 8630 buttered with 625, angled arrays of 10 x 10 indents were used, providing a higher point-to-point resolution. Measurements extended up to 15 and 60μm into the steel and nickel alloy, respectively. Indentations were performed at a maximum load of 1.5mN using load/unload rates of 0.05mN/s and hardness and elastic modulus were calculated by Oliver and Pharr analysis [6] .

4. Results

4.1 Metallurgy across the welding interface

SEM Microstructure and Chemistry: All joints examined exhibited typical microstructures as described in Figure 1 . However, joint 27 was also found to contain localised cracks along the interface, Fig. 4. These cracks appeared to have formed during the seabed exposure of the sample and were typical of cracks produced by testing under CP to simulate subsea interfacial failures. It was notable that these cracks were all subsurface.

 

Typically, the width of the PMZ in 8630-625 welded samples was 40µm. The planar solidification region (‘featureless zone’) was approximately 20µm wide and was identified as being the distance between the dissimilar interface and the regions where interdendritic niobium-molybdenum based carbide occurs, as seen in TEM sample 2 (Fig. 5).

TEM Microstructure and Chemistry: An approximately mid-wall thickness position in a continuous (i.e. mid-weld bead position) PMZ was selected for TEM analysis. Three extractions were made, two across the interface and one towards the outer edge of the planar solidification region, (Fig. 6). TEM sample 1 was extracted from a region towards the edge of a weld bead, i.e. close to a discontinuous PMZ (bead overlap swirl). Sample 3 was selected from a mid-bead position.

Along the interface of sample 1, lath-like martensitic grains were seen running vertically through the wafer. Immediately adjacent to this is a transition zone, typified as the white, carbide free region near to the protective platinum cap. The nickel alloy, seen on the right hand side of Fig. 7, appeared featureless in the SEM but, at this magnification, nanoprecipitates were visible as dappled specks surrounded by an FCC matrix. Point electron diffraction analysis was inconclusive but it is suggested they were M₇C₃ carbides, as they were consistent in size and shape with the precipitates documented in [7] and [8].

At higher magnification, tilted images (Fig. 8) revealed the lath structure more clearly and also a narrow band of particles which are shown in the EDX elemental maps of sample 3, given in Fig. 9b, to be rich in chromium. These particles are potentially M₂₃C₆ carbides where M is principally Cr. Dark globular particles, approximately 200nm x 50nm located on a grain boundary obliquely crossing the so-called ‘featureless zone’ in Figs. 7 and 8 are also thought to be carbides, although point electron diffraction did not confirm this. Sample 3, further away from the discontinuous PMZ at the bead overlap position, had a considerably narrower lath martensite zone but displayed a more pronounced carbide free region in the Alloy 625 next to the interface (Fig. 9a).

4.2 Nano-indentation Testing

By setting the nano-indentation array at a shallow angle to the fusion line, indent spacing limitations were overcome. An effective spacing of less than 0.5μm perpendicular to the interface was created without interference from the plastic zones of neighbouring tests, Fig. 10. A calibrated FEGSEM was used to measure the perpendicular (±1°) distance from the centre of each indent to the fusion line. A typical array of indentations is shown in Fig. 11. 

The results from the F65 and F6NM nanohardness tests have been discussed at length in [9][6. Discounting the high nanohardness values that appear to follow the dendrites away from the interface, the F65 to Alloy 625 joint (no PWHT) showed no hardness peak on the Alloy 625 side. Occasionally, near-interface peaks in hardness were seen in the F65 parent. Figure 10c shows nanohardness of the F6NM to Alloy 625 interface. Higher values were seen in the Alloy 625 side, with lower hardness in the F6NM. Nanohardness measurements showed a slight decrease in hardness immediately adjacent to the interface on the Alloy 625 side. It should be noted that indentation spacing may have not been fine enough to characterise some near-interface features.

In an attempt to maximise resolution, for the 8630 parent material joints (specimen numbers 27 and 35) the indents were aligned at an acute angle of between 5° and 15° to the interface (Fig. 11). Perpendicular measurements to the interface allowed hardness profile results to be plotted as a function of distance from the apparent fusion line. Each array was accompanied by an SEM-EDX linescan from the interface towards the bulk weld metal, between the middle columns of indents.

The plots in Fig. 12 and Fig. 13 for 8630/625 interfaces are largely consistent with the earlier hardness maps for F65/625 and F6NM/625 in that the highest hardness was measured deeper into the buttering, rather than immediately at the interface. The EDX traces showed an apparent difference in the chemical gradients at the interfaces of the two different 8630 samples, despite the fact that the bulk dilution of the beads seemed to be leveling out at a similar composition. Unlike the F65 and F6NM interfaces, however, three out of four plots for 8630 did show a hardness peak just inside the melted zone. One of the arrays from the sample which contained cracks, 27B, gave the clearest indication of a hardness peak close to the fusion line.

5. Discussions

5.1 Nanohardness Testing

Nanohardness testing indicated a different material response for the 8630-Alloy 625 interface when compared to the F65 and F6NM joints buttered with the same filler. In the Alloy 8630 joints, significant hardening was found close to the fusion line. This is an important finding and may be consistent with the failures observed in the 8630-nickel alloy combinations in the field ([1] and [10]). The 8630-625 dissimilar interface is known to have inferior fracture toughness values, compared to other interfaces tested [11]. In previous work, it has been suggested that, if nanohardness profiles could be correlated with fracture toughness of dissimilar interfaces measured under CP, the technique might be used to qualify welds of this type in the future [9].

 

However, some refinement of the nanohardness testing technique may be necessary before definitive conclusions can be made. Consideration must be given to the unknown service conditions of each individual joint. There is also significant point-to-point variability in the nanohardness tests shown in this paper, despite the arrays being positioned in very similar regions and similar weld procedures being used throughout. TEM work has demonstrated that there can be very large microstructural variations experienced in distances much less than the hardness indentation spacing: this is likely to create nanohardness profile fluctuations. Large disparities in hardness may arise from either grain misorientation between successive indents or different phases, e.g. carbides, being sampled.

The cracks found in sample 27, which had been retrieved from the seabed after being subjected to CP for 9 months, are significant because they yield important clues about the development of interfacial cracking in service. Thus, the indications are that cracks develop progressively, rather than being generated in a single large event, and are capable of initiating below the surfaces of components. Similar cracking to that seen in sample 27 has been observed at low alloy steel-625 interfaces tested at slow strain rates in chloride solution with CP. Cracks have been observed forming ahead of the main crack front, away from the area of maximum bending stress. Fracture toughness data for joints of this kind will be presented in part II of this paper.

6. Conclusions

The following conclusions can be drawn from this work:

 

1. The planar solidification region in 8630/625 buttering interfaces is decorated by fine precipitates that are potentially M₇C₃, the formation of which is likely attributable to carbon diffusion during PWHT.
2. Hardness peaks at the interface between 8630 steel and Alloy 625 weld metal may arise from either martensite or carbide particles, or a combination of both. Peaks in nanohardness at the 8630/625 interfaces were not replicated at those with F65 and F6NM.
3. Characterisation of dissimilar joints by nanohardness arrays requires an effective indentation spacing perpendicular to the interface of much less than 0.5µm. Parallel to the interface, indentations should be spaced as close together as possible, within the limitations imposed by the plastic zones of surrounding indentations. The described FIB-SEM lift-out procedure has great potential for analysing individual indents by TEM. Combined with the nano-indentation material response, TEM analysis would describe unequivocally the microstructures responsible for peak hardness near the interface.
4. Cracks found in an 8630-625 joint, retrieved from the seabed, indicate a progressive cracking mechanism for interface cracking which can initiate subsurface.

Acknowledgments

This work was joint funded by TWI Ltd., Cambridge, and Mintweld, a European Commission Seventh Framework Programme (FP7) sponsored project. The authors acknowledge the University of Loughborough and Oxford Materials for use of dual beam and TEM systems, respectively. Hiroto Kitaguchi, Ashley Spencer and Geoff West are thanked for their contributions.

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