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Subsea dissimilar joints: Failure mechanisms and opportunities for mitigation (March 2009)

V C M Beaugrand, L S Smith and M F Gittos

Granta Park
Great Abington, Cambridge, CB21 6AL, UK

Paper presented at Corrosion 2009, Atlanta, Georgia, USA, 22-26 March 2009. Paper #09305.


Dissimilar joints are a common part of subsea oil and gas systems where clad forged fittings are joined to clad or CRA pipe. Whilst the majority of these joints have given successful service, a small number have failed. In the current work, environmental testing was carried out on a dissimilar joint to an 8630M forging, buttered using alloy 625. Tests were monitored using a sensitive technique in order to capture the first moments of crack initiation and to determine the minimum threshold stress intensity factor Kth, below which there is no environmentally assisted cracking. Two distinct initiation mechanisms were identified; one dependent on hydrogen-charging in service (under cathodic protection) and one not. A detailed microstructural investigation, including chemistry, structure and strain, gave insight into the attributes of the 'partially mixed zone' (PMZ) immediately adjacent to the fusion boundary, on the weld side, in which one of the failure modes occurred. The chemical and microstructural characteristic of this PMZ are influenced by welding and postweld heat treatment (PWHT) parameters. Based on this investigation, the potential for mitigation of cracking susceptibility by materials selection or by modifying the welding/PWHT procedure is discussed.

Key words: Dissimilar joints, subsea, hydrogen embrittlement, cracking, microstructure.


Dissimilar joints are a common part of subsea oil and gas systems. Typically, forged and machined hubs or tees made from low alloy steel are buttered with a nickel alloy, such as alloy 625, and then joined to line pipe micro-alloyed steels. The main advantage of buttering is that it enables the forgings to be post weld heat treated prior to making the completion weld in the field, as the HAZ of the completion weld lies wholly in the buttered layers. The primary dissimilar interface of interest in the present article is that between the forging and the butter. As for any fusion welded dissimilar joint, a narrow partially mixed zone (PMZ) is observed at the dissimilar interface, where the composition is graded between that of the parent forging and the bulk weld deposit chemistry of the butter runs. In order to prevent corrosion of piping systems, subsea, cathodic protection is applied, primarily using sacrificial aluminium anodes giving an applied potential of approximately -1100mVSCE. Whilst this does inhibit corrosion it can result in a certain amount of hydrogen charging, at any exposed metal surfaces, resulting in the potential for hydrogen embrittlement when an undesirable combination of susceptible microstructure and high applied load is present. Whilst the vast majority of subsea joints have given successful service, a small number has failed.

Hydrogen induced cracking at dissimilar interfaces has been observed historically in vessels for high temperature hydrogen service, specifically at the interface between austenitic claddings and low alloy Cr-Mo steels. Here, rather than hydrogen charging through cathodic protection, thermal hydrogen charging is experienced and cracking is a risk upon cooling during shutdowns. Asami and Sakai first coined the concept of three types of cracking at various positions across the dissimilar interface.[1] Type I cracking was defined as occurring at the interface itself, Type II cracking was broadly intergranular, close to the interface in the PMZ in shallow penetrating grains, and Type III cracking in the steel HAZ. The susceptibility to hydrogen embrittlement of various austenitic claddings was investigated by Gittos and Gooch.[2] High temperature post weld heat treatment resulted in the migration of carbon from the steel to the austenitic cladding, due to the presence of an activity gradient, and the formation of virgin martensite. Similar activity/mobility gradients for hydrogen result in the formation of a hydrogen peak close to the interface and the combination of brittle microstructure and high hydrogen content gives a high risk of disbonding during cooling. Failure was predominantly Type I for nickel based claddings and a combination of Type I and II for austenitic stainless steels (i.e. cracking either at the interface or intergranularly along shallow penetrating grains). Whilst there are similarities in terms of the interface between the high temperature hydrogen service claddings and the butter layers of the subsea dissimilar joints, there are some obvious differences in parent metals employed, fabrication and, of course, hydrogen charging conditions. The primary aim of the present work was to understand the failure mechanism(s) that can operate in dissimilar joints in subsea service and its relationship with the microstructure and chemistry at the dissimilar boundary.

Experimental procedure


The butter weld consisted of a build up of alloy 625 on a low alloy steel forging (8630M), deposited using hot wire TIG welding with an applied preheat. The buttered forging was subjected to PWHT and a completion weld made, again using alloy 625, to a steel pipe.

Microstructural assessment

Transverse and longitudinal sections were taken through the dissimilar joint. Sections were examined using light and scanning electron microscopy. Semi-quantitative elemental analysis was conduced using EDX in order to determine the profile of Fe, Ni, Cr, Mo and Nb across the dissimilar interface. The carbon profile was qualitatively established using an electron probe micro-analyser (EPMA) with a dedicated spectrometer for carbon. The results were correlated with Vickers hardness/micro-hardness across the dissimilar interface, conducted with a 5kg/25g load. Electron backscattered diffraction (EBSD) in a high resolution field emission gun (FEG) scanning electron microscope (SEM) was employed to determine the phases present with a spatial resolution of approximately 0.05-0.1µm and to qualitatively assess the distribution of local strain around the dissimilar interface. This information was collected primarily to identify any trends in plastic strain in the partially mixed zone (PMZ) that may induce hardening and/or susceptibility to cracking.

Environmental testing and assessment

The main purpose of the environmental tests was to identify any initiation and propagation events. Square section 12 x 12mm single edge notched bend (SENB) specimens were machined and notched through-thickness by electro-discharge machining (EDM). These notches were positioned so that the notch tip sampled the dissimilar interface. Specimens were hydrogen pre-charged at -1100mVSCE in 3.5%NaCl for a minimum of 48 hours prior to testing. Incremental, dwell loading tests were carried out. Specifically, tests were conducted by loading samples in three point-bending to an initial stress intensity of 40MPam0.5, continued by further, cumulatively applied load increments of 2.5MPam0.5. These tests were performed at 3°C in 3.5%NaCl aqueous solution under cathodic protection at -1100mVSCE, applied by potentiostat.

The specimens were instrumented with a pair of clip gauges mounted on knife edges at the notch mouth. The knife edges were mounted on steel shims attached to the specimen, so that crack mouth opening displacement (CMOD) could be calculated at the crack tip. The tests were carried out, generally following BS7448 Parts 1, 2 and 4, where possible. The test was terminated once crack initiation was identified, by monitoring the response of either the clip gauges or the acoustic emission signal. The latter is a non-destructive method, which collects mechanical energies, partly in the form of elastic waves, released when a microscopic deformation or cracking event occurs. Initial testing was performed to significant crack extension (i.e. crack extension of a few millimetres) to calibrate the acoustic emission monitoring equipment (including noise emissions from the testing equipment) and to determine the sensitivity of the method in detecting cracking at an early stage, but subsequent tests were performed and stopped shortly after a significant signal was collected by acoustic emission (AE).

The AE procedure applied in the present programme is described in greater detail elsewhere.[3] After testing, samples were heat tinted and broken open in liquid nitrogen prior to fractographic examination. The AE events (including energy) were correlated with features observed on the fracture faces, where possible. The threshold stress intensity was determined and the data were used to estimate KI(EHI) (EHI being external hydrogen induced, i.e. from cathodic charging as opposed from hydrogen from, say, welding). In the present work KI(EHI) was calculated at the loading increment for which cracking was first detected, although clearly a conservative value would be lower than this. For comparison, a limited amount of testing was performed in air. One specimen was tested using a similar methodology to the environmental tests, but in air and without hydrogen pre-charging. This gave a nominal threshold stress intensity KI(IHI) (IHI standing for internal hydrogen induced hydrogen embrittlement).

A further conventional CTOD test was performed to produce comparative KQ data. Finally, two notched specimens were simply broken open at a similar test rate to reveal typical fractographic features. One of these was hydrogen pre-charged at -1100mVSCE in 3.5%NaCl for two weeks, whilst the other was tested in the as-fabricated condition. Post-test fractographic examination of the broken halves was undertaken to determine evidence of the location of the initiation event(s) and measurement of the extent of any cracking. Examination was performed using light and scanning electron microscopy, as appropriate. Fracture features were correlated with initiation modes at the dissimilar joint interface/partially mixed zone.


Microstructure and hardness

The first butter layer has been diluted by the low alloy steel and exhibited some partly-fused 'swirls' of steel-rich material between adjacent weld passes (i.e. at inter-run positions). This occurs since the melting temperature of the alloy 625 (~1350°C) is somewhat lower than that of the alloy steel (~1500°C). As well as these 'inter-pass swirls' there were also 'intra-pass swirls' that form along the length of the weld beads of the first layer of butter deposit, as a consequence of periodic variations in weld penetration.

A narrow (typically a few tens of microns wide) PMZ is present between the non-fused steel parent and the bulk of the weld deposit, and has a chemical composition that lies between that of the steel and that of the bulk of the first layer. After PWHT, it is anticipated (and substantiated in the present work) that carbon migration from the steel to the Cr/Mo/Nb rich butter deposit creates a decarburized zone in the alloy steel adjacent to the fusion boundary and a carbon-enriched zone in the immediately adjacent fused metal. The microstructures and chemistries of the inter- and intra-pass swirls were complex, with a gradation in composition across the swirl and the material immediately around it (Figure 1b). Whilst individual swirls showed broadly comparable features, each was relatively unique due to slight variations in geometry, dissolution and chemistry. In comparison, the mid-bead position of the interface was quite reproducible (Figure 1a).

Fig.1. Micrographs of the dissimilar interface
Fig.1. Micrographs of the dissimilar interface

a) at mid-bead position;

b) at an inter-pass position

The microstructure across the dissimilar interface was found to consist of six different morphological zones. Not all of these zones were found at all positions, but when present were, in order of low alloy steel HAZ to the bulk of the first butter layer, as follows:
  1. Parent material showing a fine ferritic microstructure, with occasional 'fingers' of highly diluted butter deposit penetrating into the steel, apparently at prior austenite grain boundaries;
  2. A narrow decarburized zone in the region of the forging immediately adjacent to the fusion boundary and containing occasional high atomic number particles (labelled Zone Δ);
  3. Immediately at the fusion boundary (mostly within the PMZ, but perhaps extending slightly into non-fused steel), iron-rich martensite-like laths extending a few microns into the PMZ, primarily at inter-pass positions and within inter-pass and intra-pass swirls of diluted steel penetrating into the butter deposit (labelled Zone M). No similar features were found at mid-bead positions away from intra- and inter-pass swirls;
  4. A featureless region on the butter side immediately adjacent to the fusion boundary and exhibiting an apparently homogeneous single-phase microstructure on a microscopic scale (labelled Zone Φ);
  5. A region containing numerous high atomic number particles (labelled Zone Π). These particles exhibited an inter-dendritic distribution, consistent with the expected segregation of Mo and Nb during solidification of this highly diluted region of the butter deposit;
  6. The bulk of the first deposit layer and subsequent layers with close to alloy 625 chemistry, and a less populous inter-dendritic distribution of Nb/Mo particles.

Zones 3 to 5 constituted the partially mixed zone (PMZ) within the weld metal, as determined from the extent of the chemistry gradient between the parent metal and the bulk of the first butter layer. This PMZ mainly consisted of Fe, Ni and Cr with proportions that were closer to that of the forging close to the interface, and closer to the alloy 625 weld deposit further away, with a width from a few tens to a few hundreds of microns (although mostly at the lower end of this scale).

Average results of the Vickers hardness tests are presented in Table 1. It is noted that Vickers hardness indents are too coarse to identify hardness changes across the dissimilar interface, but results indicate that forging approximately met the maximum ISO 15156 hardness limit of 250HV for steel in sour service. The micro-Vickers hardness traverse results localized to the interface are plotted in Figure 2 and showed a peak of hardness in the PMZ, adjacent to the fusion boundary, with a peak hardness of approximately 450µHV. These results are for a mid-bead position, but similar microhardness peaks were noted at inter-pass positions within the martensitic region. Note that the hardness indent size is much greater than a martensite lath, however, although not reported here, nano-hardness measurements show significantly greater hardness within the laths.

Table 1 - Average Vickers hardness (load 5 kg)

Region sampledForging, bulkForging, HAZButter, bulkClosure weld, bulk
Hardness, HV 236 251 268 230
Fig.2. Micro-Vickers hardness survey across the dissimilar interface a) Micrograph showing the interface;
Fig.2. Micro-Vickers hardness survey across the dissimilar interface a) Micrograph showing the interface;
b) Micro-hardness survey across the interface
b) Micro-hardness survey across the interface

Phase and strain distribution

At the mid-bead positions, the 'butter side' exhibited a face centred cubic (FCC, or 'austenitic') structure with large grains; while the 'steel side' showed a fine grained body centred cubic (BCC) (ferritic) structure (Figure 3a), with no discernible martensite. The particle-free region (Zone Φ) consisted of large and elongated (in the direction parallel to the fusion boundary) grains, often restricted to this zone, as opposed to continuing seamlessly into the bulk of the butter deposit. The grains of the bulk weld bead were elongated perpendicular to the fusion boundary. The swirls exhibited a mixture of fine lath grains and equiaxed grains, depending upon location, with the lath grains found predominantly in positions where the swirl appeared to be partially dissolved by the butter weld metal. Fully automated EBSD analysis could not differentiate between the BCC ferrite and the BCT martensite, but carefully set up detailed semi-automatic analysis clearly identified martensite in the swirl regions, corresponding to the laths identified metallographically (Figure 1b). The high atomic number inter-dendritic particles observed in Zone π were identified as having primarily a niobium carbide crystal structure, NbC.

A small plastic strain was found in the PMZ, adjacent to the interface at the mid-bead positions (Figure 3b), but this was lower than that found in the adjacent ferrite. Limited local plastic deformation was also identified within the particle-free region of the PMZ (Zone Φ, with a lattice distortion of up to 2° similar to that found at the mid-bead position). It was clear, however, that the amount of plastic strain in the PMZ and steel was relatively small, consistent with what would be expected from relief of residual stress during PHWT. Plastic strain levels were comparable in interpass regions, with the strain around the martensite no greater than that found locally elsewhere.

Fig.3. EBSD analysis at the mid-bead position a) Phase map showing the FCC in red and the BCC in green b) Kernel average misorientation (KAM) map showing the deformation distribution
Fig.3. EBSD analysis at the mid-bead position a) Phase map showing the FCC in red and the BCC in green b) Kernel average misorientation (KAM) map showing the deformation distribution

Composition across the dissimilar interface

At mid-bead positions, the featureless region (Zone Φ) of the PMZ, immediately adjacent to the fusion line at the mid-bead position, was Fe-based and enriched in Cr and Ni and typically comprised 20wt%Ni, 6-7wt%Cr, and approximately 0.5wt%Nb. The remainder of Zone Φ was nickel-based, enriched in iron (35wt% at mid-width) with significant Cr (14 to 15wt%), Mo (5 to 6wt%) and Nb (approximately 2.5wt%). The PMZ, as defined by the transient Fe composition, was up to 100µm wide, and the iron content 50µm away from the fusion boundary in the bulk of the first layer, was approximately 12wt%.

Fig.4. Electron probe micro-analyses across the dissimilar interface a) SEM image
Fig.4. Electron probe micro-analyses across the dissimilar interface a) SEM image
b) Qualitative C profile
b) Qualitative C profile

At inter-pass positions the coexistence of unmixed steel and butter deposit within swirls resulted in complicated variations in chemistry. The region of the low alloy steel immediately adjacent to the fusion boundary was slightly enriched in Cr (1.5wt%) and Ni (1.8wt%). The region of the PMZ immediately next to the fusion boundary was Fe-based, with an example of one such region containing 1.3wt%Mo, 4wt%Cr and 10 to 11wt% Ni. The chemistry of the swirls of diluted steel that penetrated into the butter were very similar to the forging, being slightly enriched in Ni (2wt%), Cr (1.7wt%) and Mo (0.8wt%). No niobium was identified in the swirl. Several swirls were examined and, by and large, were found to exhibit similar features. EPMA analysis across the dissimilar interface exhibited a peak of carbon near the fusion boundary in the featureless region of the PMZ (Zone Φ) as shown in Figure 4. These results are qualitative, but are indicative of a significant peak in carbon content.

Environmental tests

AE monitoring was found to be more sensitive for the detection of the initiation of cracking than monitoring the clip gauge and, during the initial test used for calibrating the AE, cracking was indicated 4 hours before any appreciable clip gauge response was detected. Two types of signal were identified from subcritical cracking mechanisms: (1) A signal emitting waves with amplitude between 50 and 60dB and an absolute energy lower than 2,000aJ; (2) A signal emitting waves with amplitude between 50 and 70dB and an absolute energy above 20,000aJ. These were correlated with two distinct fracture morphologies, as described later.

The lowest stress intensity factor for crack initiation during environmental testing and under cathodic protection, KI(EHI), was 55 MPam0.5. Note that this should not be viewed as a conservative threshold value since only a limited number of tests were performed and the value at which cracking was observed is reported. The stress intensity factor for crack initiation for a specimen tested in air (e.g. with only that hydrogen present from fabrication), KI(IHI), was 74MPam0.5. It is noted that the hydrogen content in the weld metal, measured by inert gas fusion, was 8ppm. For comparison, the stress intensity factor KQ, determined by conventional CTOD testing was 76 MPam0.5. Note that these represent a limited data set and are intended to provide indicative values only. Even so, KI(IHI) and KQ are very close to one another, illustrating that in this instance hydrogen from fabrication has less influence than hydrogen from the testing environment.

Post-test examination of the environmentally-tested SENB specimens (stopped shortly after identifying crack initiation) revealed plastic deformation around the edges of the samples, together with cracking along the dissimilar interface. Similar fractographic features were noted in the tests, namely:

  1. Some 'terraced' cleavage within the PMZ, comprising cleavage facets roughly in the plane of the dissimilar interface connected to one another by shorter cleavage facets, at a steeper angle to the dissimilar interface (Figure 5a). These features were within a zone of chemistry consistent with a nickel-rich region of the PMZ.
  2. Flat relatively featureless fracture at the interface, with no particularly discernible characteristic features other than a profile of cellular ridges. The chemistry of the fracture face showed minimal amounts of dilution, consistent with fracture at or very close to the dissimilar interface.
  3. A mostly flat dimpled fracture at the interface between the diluted steel swirls and the butter deposit (Figure 5b). These zones again showed minimal amounts of dilution, consistent with fracture at or very close to the dissimilar interface.
  4. Ductile fracture, principally within the low alloy steel.

When testing was stopped immediately after initiation, one or both of two fracture morphologies were identified. In some instances, a tiny 'patch' of cracking exhibiting terraced cleavage morphology was observed roughly in the middle of the specimen, initiating from the notch. Similarly, small patches showing a flat, dimpled fracture morphology were typically found near the edges of the SENB specimen. Where only one mode of initiation was identified, the fracture morphology could be correlated with the AE signal acquired during testing. The lower energy/amplitude signal corresponded to the dimpled fracture mode shown in Figure 5b, whereas the higher energy/amplitude signal was from the terraced cleavage shown in Figure 5a.

Fig.5. SEM micrographs showing the two types of fracture morphology in the initiating stages of cracking a) Terraced cleavage
Fig.5. SEM micrographs showing the two types of fracture morphology in the initiating stages of cracking a) Terraced cleavage
b) Type I disbonding
b) Type I disbonding

When tested in air (for the KI(IHE) test), the fracture faces exhibited mostly ductile features, and numerous isolated patches of flat, dimpled fracture located, once again, mainly at the edges of the samples. Similar results were observed for the conventional CTOD test specimen and for an SENB specimen that was hydrogen pre-charged for two weeks, heat tinted and then simply broken open. Metallographic sections taken through the fracture initiation features demonstrated that the flat, dimpled fracture occurred predominantly at the fusion boundary and that terraced cleavage features were found within the featureless region (Zone Φ) of the PMZ. In the present work, cleavage within the PMZ was only found in specimens tested under hydrogen charging conditions.


Microstructural features and their formation

The dissimilar joint between 8630M and the alloy 625 butter showed several microstructural zones on traversing across the dissimilar interface. Based on the examination summarised in the present paper, a microstructural model was established and is presented in Figure 6.

Fig.6. Microstructural model of the dissimilar interface of a 8630M buttered with alloy 625. An interpass swirl is shown on the right hand side, whereas the left hand side is comparable to a mid-bead position
Fig.6. Microstructural model of the dissimilar interface of a 8630M buttered with alloy 625. An interpass swirl is shown on the right hand side, whereas the left hand side is comparable to a mid-bead position

Macroscopically, 'swirls' of steel penetrating up to a few hundred microns into the first butter layer were present at inter-pass positions, and intra-pass locations where local perturbations in penetration had occurred. The inter-pass positions, including at partially dissolved steel swirls, also exhibited martensite as a consequence of the higher dilution in this zone. Fingers of butter weld metal penetrating into the steel were also present, akin to shallow liquid metal penetration along the austenite grain boundaries of the forging GCHAZ at high temperature. No cracking was observed in these 'fingers' and it is not believed that these features have any significance in terms of structural performance.

The PMZ, immediately adjacent to the fusion boundary, consisted of a featureless particle-free zone (zone Φ) and, further into the butter deposit, a region with an interdendritic distribution of primarily NbC (zone π). It is postulated that zone Φ forms as a consequence of the high Fe and C dilution from the parent steel, both reducing the proportion of Mo and Nb segregates and changing the solidification mode to planar (thereby avoiding interdendritic segregation completely).[4] The transition to zone π occurs when the solidification mode changes to dendritic.

During post weld heat treatment, diffusion of interstitial species, in particular carbon, across the dissimilar interface dominates the microstructural changes experienced. A narrow decarburized zone was present within the steel, immediately adjacent to the fusion boundary, accompanied by a peak of carbon in zone Φ of the PMZ, formed due to a 'pile-up' of carbon. Given that zone Φ has been enriched in carbon, diluted from the parent steel during welding, and then further enriched in carbon, as a consequence of its diffusion from the steel during PWHT, it is somewhat surprising that no carbides have formed, especially considering that the carbon solubility in nickel is typically lower than that in iron. However, it is believed that the microstructure in the featureless zone after PWHT consists of carbon in a supersaturated solid solution, perhaps accompanied by a fine dispersion of nano-sized carbides. The equilibrium solubility limit of C in ferrite (pure iron) is at most 0.1at% (0.02wt%). However, this represents equilibrium conditions and thus assumes adequate mobility of all elements. At the PWHT temperature the mobility of strong carbide formers is very low and thus assumptions of para-equilibrium conditions (i.e. considering only the diffusion of carbon) are more appropriate. This has been exploited for low temperature carburisation surface hardening of stainless steel,[5,6] where by control of temperature and carburising atmosphere, supersaturated solid solutions containing as much as 12at.%C (3wt%) can be achieved together with hardness levels in excess of 600HV.[7.8] The PWHT temperatures of relevance to the present work, although greater than that used to produce such massively supersaturated solid solutions, are still expected to permit quite significant supersaturation.

Environmental performance and failure mode

In the present work, environmental tests highlighted susceptibility to hydrogen embrittlement from cathodic charging. In the alpha BCC phase, hydrogen diffusion is fast but solubility is low, whereas in the gamma FCC phase, diffusion is slow, but solubility is high. Consequently, a hydrogen activity gradient drives hydrogen from the steel to the dissimilar interface where it forms a 'pile-up' in the PMZ due to the low diffusion rates in the FCC material. Cracking either at the interface or within the PMZ is favoured by both a concentration of hydrogen and the presence of a hard microstructure (susceptible to hydrogen embrittlement).

Different morphologies of fracture were highlighted by the environmental tests: principally terraced cleavage within zone Φ of the PMZ or flat, dimpled fracture at interpass interfacial regions. Both of these were often observed in test specimens, having initiated independently, even when testing was stopped upon detection of the earliest stage of cracking. On this basis, it would appear that whilst there are two mechanisms, there is no single microstructural 'weak link'.

Terraced cleavage was only observed when testing under hydrogen charging conditions and was found within the austenitic zone Φ. Hydrogen embrittlement of austenitic material is not particularly common, but can occur under severe conditions. Given the unusually high carbon content of this part of the PMZ, its high hardness and the tendency for hydrogen to build up in this zone under hydrogen charging conditions,[7] local conditions are clearly suitably extreme for embrittlement and cleavage to occur. This mode of cracking has not been reported for the high temperature service hydrogen vessels and may be peculiar to higher alloy deposits or cracking under cathodic protection.

The flat dimpled fracture was located predominantly at inter-pass and intra-pass positions and was associated with Type I disbonding directly at the dissimilar interface or the martensite in this region, but was observed on all tested specimens, including non-hydrogen pre-charged specimens, tested in air. No evidence of pre-existing disbonding flaws was found in any of the metallographic sections and the acoustic emission data supports the formation of disbonding during testing. The difference in threshold stress intensity factor between testing in air and testing under hydrogen charging was not particularly large, a difference of less than 20MPam0.5 was found between the internal (fabrication only) and external (from cathodic protection) hydrogen induced threshold stress intensity factors (respectively 55MPam0.5 and 74MPam0.5). However, conventional CTOD testing gave a KQ of 76MPam0.5 demonstrating that the internal hydrogen content may have little influence on performance and that the microstructure at the swirl interfaces represents a local brittle zone.

This finding suggests that, given sufficient installation stresses/strains, Type I disbonding cracks can initiate in the brittle martensite prior to installation subsea and behave as 'pre-existing' flaws for subsequent subsea service under hydrogen charging conditions (i.e. under cathodic protection).

Mitigation of hydrogen embrittlement of dissimilar joints

Clearly, the risks of hydrogen embrittlement can be dramatically reduced by preventing hydrogen charging. Hydrogen charging could theoretically be prevented by successful application of coatings (although the coating needs to be sufficiently robust to survive installation and service) or by reducing the current density and/or increasing the applied potential of the cathodic protection (CP) system (i.e. making it less electronegative). However, such 'solutions' whilst possible within a laboratory are not very pragmatic for field application.

Two crack initiation modes were observed in the present work, which were primarily observed in two different microstructural zones. Thus, modification of both of these features would be required to raise performance levels. Firstly, type I disbonding was observed directly at the dissimilar interface in the highly diluted chemistry (promoting martensite formation) at inter-pass and intra-pass positions, including swirls. Type I disbonding was also observed, albeit at higher stress intensities, in specimens that were tested without external hydrogen charging. Thus, minimising the formation of swirls should improve performance of the joints and at least reduce the risk of cracking during installation, under moderate plastic strain. Work on high temperature hydrogen service claddings has come to similar conclusions.[9-11] However, the complete elimination of swirls seems unlikely, thus moving to a lower carbon forging would seem to offer the most practical means of reducing the consequence of any martensite in swirl positions, by reducing hardness or ensuring a chemistry that is tempered during PWHT.

Secondly, hydrogen embrittlement of a region of the PMZ supersaturated in carbon was also observed. It is believed that this zone is susceptible to forming a hard carbon supersaturated zone during PWHT. Again, reducing the carbon content of the forging will be of benefit by reducing the peak carbon content in the PMZ (by lowering the activity gradient).

From the above, two potential means of avoiding dissimilar joint failure have been proposed:

  1. Employ lower carbon forging materials than 8630M
  2. Employ a low alloy steel butter or pup piece instead of alloy 625

Even so, appropriate qualification of fabrication procedures (on environmental performance) is necessary to give sufficient confidence in these solutions, since the relationship between chemistry, fabrication parameters, microstructure and performance is complex. Even item 2, whilst seemingly avoiding much of the problem by eliminating the PWHT of a dissimilar interface is not certain, since the carbon content of the steel adjacent to the alloy 625 completion weld will probably be greater than that typically encountered in line pipe (diminishing any comfort gained from the adequate performance to date of alloy 625 completion welds). Work is ongoing[12] to define acceptance criteria and provide guidance on demonstrated solutions and enable confident design of robust dissimilar joints for subsea service.


The dissimilar interface between an alloy 8630M forging buttered with alloy 625, postweld heat treated and completion welded to line pipe steel for subsea application was considered in the present work. The following conclusions are drawn:

  1. Dissimilar joints between low alloyed steel and alloy 625 are susceptible to hydrogen embrittlement under cathodic protection in subsea applications. In the present work a non-conservative threshold stress intensity factor, KI(EHI), of 55MPam0.5 was determined.
  2. Hydrogen embrittlement failure occurred by a combination of cleavage and Type I disbonding with both initiating independently at similar loads during the incremental dwell load KIH testing. Cleavage initiated within a hardened featureless region of the PMZ. Type I disbonding occurred at inter-pass and intra-pass positions, including steel swirls penetrating into the weld.
  3. Cleavage occurred in a narrow, 'non-segregated' portion of the PMZ that appeared to have solidified in a planar manner. Carbon diffusion resulted in the supersaturation of this zone in carbon together with consequent hardening. This mode of failure was only observed when testing under cathodic protection.
  4. Type I disbonding occurred at the fusion boundary within a local brittle zone of martensite that was most apparent at inter-pass and intra-pass locations (including swirls). Type I disbonding was observed in specimens subjected to incremental dwell load KIH testing under cathodic protection and in air, and in conventional air CTOD tests.
  5. A non-conservative conventional fracture toughness measured in air from a single specimen was KQ = 76MPam0.5 and failure occurred by a mixture of tearing and Type I disbonding, the latter at swirl positions. Similar fracture features were observed on a notched specimen broken open by bend testing at a fast rate.
  6. The threshold intensity factor of a single specimen subjected to incremental load KIH testing in air was KI(IHI) = 74MPam0.5. This value is very close to that recorded in item 5, suggesting minimal influence of hydrogen from fabrication.
  7. The partially mixed zone is divided into two regions: a planar solidification zone and a region exhibiting a dendritic solidification structure with numerous inter-dendritic NbC particles. Its width varied from a few tens of microns to a few hundreds of microns.
  8. Carbon migration during PWHT resulted in supersaturation of carbon in the PMZ, with resultant hardening. Longer PWHT promoted further carbon migration and a shift of peak hardnesses further into the PMZ.


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  2. M.F. Gittos, T. G. Gooch, 'Welding Journal', Vol. 71 No.11, 1992, pp. 461-72.
  3. C. Ennaceur and V. Beaugrand, 28th European conference on acoustic emission testing, EWGAE, Krakow, (September 17-19, 2008), pp. 64-69.
  4. C. Højerslev, N. Tiedje and J. Hald, Materials Science Forum Vol. 508, (2006), pp. 373-378.
  5. Y. Cao, F. Ernst, G.M. Michal, Acta Materialia 51 (2003), pp. 4171-4181.
  6. G.M. Michal, F. Ernst and A.H. Heuer, Metallurgical and materials transactions A 37A, (2006 June), p. 1819.
  7. M.F. Gittos, Corrosion 2008, NACE International, New Orleans, (2008).
  8. B. Alexandrov and J. Lippold, 3rd Stainless Steel World America Conference, Houston, USA, (October 2004).
  9. Z. Wang, B. Xu and C. Ye, Welding research supplement, (August 1993), p. 398-s.
  10. G. Gnriss, Conference IIW/IIS Ljubljana, (July 2001).
  11. M.F. Gittos, Welding in the World, 52 issue 3/4 (2008).
  12. V. Beaugrand TWI GSP/JIP 17434, report 17434/08/9, September 2008.

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