Deepwater Corrosion Fatigue Performance of Welded Risers

Paul Woollin, Richard J Pargeter and Stephen J Maddox

Paper presented at Corrosion 2004, March 28 - April 1, 2004, New Orleans, Paper no. 04144, Symposium 04-STG-32(2).

Abstract

Steel catenary and top tension risers for deepwater oil and gas field developments are subject to corrosive environments on both internal and external surfaces and to fatigue loading. The girth welds control the fatigue life of such structures. Therefore, there is a need to quantify the fatigue performance of girth welds under the influence of the external seawater, typically with cathodic protection, and the internal environment, which may contain water, CO₂, H₂S and chloride and bicarbonate ions. The effect of temperature must also be taken into account.

This paper presents a review of published data to illustrate the effects of seawater, sweet (CO₂) and sour (H₂S) exposure on corrosion fatigue of welds in carbon steels and corrosion resistant alloys (CRAs). In view of the dominance of the fatigue crack propagation process in the fatigue lives of girth welds, particular attention is paid to fatigue crack propagation data. The data indicate the range of predominantly detrimental effects of the various environments on the fatigue performance of welded carbon steel and CRAs and illustrate the effects of weld microstructure. Limitations of existing data are discussed. Challenges remain with respect to design of welded risers against corrosion fatigue and the data required to meet these challenges are discussed.

Introduction

The environmental limits of application of pipeline materials, which are primarily subjected to static loading, are fairly well defined, at least for C-Mn steels. However, under cyclic loading the environmental limits are not well defined. There is a range of candidate materials for risers carrying corrosive produced fluids including C-Mn steel, various stainless steels, nickel alloys and titanium, but few relevant data exist on which to base materials selection. Also, it is possible that compositional and microstructural effects will influence corrosion fatigue behaviour, whereas behaviour in air is dominated by geometric effects. Thus, although preferred welding processes for risers have been developed based on testing in air, there is a need to establish preferred processes and consumable selection for corrosive applications. The paper gives a description of relevant corrosion fatigue phenomena, summarises existing corrosion fatigue data for C-Mn steels and stainless steels (i.e. supermartensitic, duplex/superduplex and austenitic stainless steels, with the last of these being used as cladding), which represent the most economically attractive candidate materials, and discusses the research required to assist future materials and welding process selection.

Corrosion fatigue phenomena associated with riser girth welds

The fatigue behaviour in air of girth welded pipes is primarily controlled by the weld toe geometry ^[1] . Hence efforts to modify weld geometry via the welding procedure have been beneficial for the fatigue life of steel catenary risers. However, fatigue in corrosive environments may be affected also by material-environment interactions, leading to reduced fatigue lives compared to operation in air. Of particular concern in the initial stages is preferential weld area corrosion (PWC) and localised pitting, which can provide stress concentrations that maybe are more severe than the weld toe, but under different corrosive conditions sharp cracks and features may be blunted by corrosion.

Once crack propagation is underway, interactions may be of various types and may be related either to the applied cyclic stress (Δ σ)/stress intensity factor range (ΔK) or the absolute stress/stress intensity factor (σ or K respectively). When cyclic modes of crack extension only are operating, behaviour is classified as 'true corrosion fatigue', whilst the presence of static modes gives rise to 'stress corrosion fatigue', i.e. fatigue with a contribution from stress corrosion cracking, which occurs when σ _max max or K _max exceeds σ _SCC or K _ISCC . The relative contributions to crack development from purely mechanical fatigue, true corrosion fatigue and stress corrosion cracking will vary (i) with the particular material-environment combination in question and (ii) during the propagation of a particular crack, due to the change in K _max and ΔK as the crack extends. It should also be recognised that the corrosive conditions at the crack-tip will be dependent on transport of species down the crack, which will depend on crack length and tightness, the presence and nature of corrosion products in the crack, and any pumping action due to cyclic opening and closing. A feature which is often regarded as an experimental inconvenience, but which may also occur in service, is crack closure effects caused by build up of corrosion products within the crack. (Having said this, any benefit from such crack closure seen in fatigue tests is rarely assumed in design on the basis that corrosion products cannot be relied upon to remain intact in service).

Thus, the nature and extent of material-environment interactions is microstructure and environment dependent, and the microstructures developed at a welded joint will contribute significantly to the corrosion fatigue properties of that joint. There is a need therefore for appropriate weld microstructure control, in addition to geometry control, in order to maximise corrosion fatigue performance of riser girth welds. Fatigue cracks developed at a weld joint will typically encounter HAZ and/or weld metal microstructures for most of their lives. With respect to hydrogen-related effects, the hardness is likely to be the dominant parameter, at least for C-Mn steels, but local grain size and composition may also be significant particularly for CRAs. The primary environmental factors that are likely to be important include the pH and partial pressure of H₂S in particular, but also Cl- content, temperature and partial pressure of CO₂. Since risers made from the less corrosion resistant materials will be inhibited, the effects of inhibitor, and of occasional lapses in inhibition, are also important.

In summary, mechanisms that may contribute to corrosion fatigue include: (i) localised corrosion, e.g. pitting or preferential weld metal/HAZ corrosion, leading to formation of geometric stress concentrators, which will tend to enhance crack initiation; (ii) crack-tip dissolution, which may increase crack propagation rates by physically removing material, or potentially decrease them by crack-tip blunting; (iii) crack-wedging, caused by formation of corrosion product/scale within the crack, which will increase crack closure and (iv) hydrogen embrittlement of material ahead of the crack-tip, as a result of pick-up of atomic hydrogen from corrosion or cathodic protection. Local hydrogen charging due to corrosion at an active crack-tip may be more significant, and more severe, than general corrosive hydrogen charging.

Both true corrosion fatigue and stress corrosion fatigue mechanisms are time dependent. Consequently, it is essential that testing to establish corrosion fatigue behaviour is performed at a frequency representative of service. In general, both true corrosion fatigue and stress corrosion fatigue crack propagation rates (in terms of extension per cycle, da/dN) will continue to increase with decreasing frequency, and no single frequency exists beyond which no further increase in crack propagation per cycle will occur. Fatigue testing of carbon steels under cathodic protection has demonstrated that crack propagation rates continue to increase approximately linearly with reduction infrequency to at least 0.01Hz ³ .

Review of existing data

C-Mn Steels

A large body of fatigue crack propagation data exists for C-Mn steels in seawater, both with and without cathodic protection. ^[1-4] Together with fatigue endurance data obtained from welded joints, these have formed the basis of current design rules. ^[6-8] However, these rules are specifically aimed at structures subject to wave loading and typical near-surface seawater temperatures of 6-8°C. Risers may be subject to different loading frequencies and may be hot, hence the design guidance for other offshore structures might not be directly applicable. Nevertheless, it is apparent that at low ΔK (and hence, generally low K _max ) cathodic polarisation reduces propagation rate, whereas there is a region of higher ΔK where propagation rates are faster with cathodic polarisation, and higher at more negative potentials. This suggests a hydrogen effect under these conditions.

For hardened weld HAZs and high strength materials, the detrimental effect of cathodic protection on crack propagation is higher than for softer parent steels. This has been linked to the introduction of static modes of hydrogen embrittlement crack propagation. ^[5] Intergranular fracture was observed at fairly high hardness levels of the order of 350HV in a simulated HAZ, giving a dramatic increase in crack propagation rate, particularly at low cycling frequencies. For a HAZ hardness of250HV, the effect was lower but an appreciable increase in crack propagation rate was still recorded. Most of these data were at recorded stress intensity ranges of above 300N/mm ^3/2 and the trends suggested that much smaller effects were present in the near-threshold regime.

Some fatigue data exist for welded joints in both sour ^[9-14] ( Figures 1-3) and sweet ^[15] ( Figure 4) produced fluids but more data are required particularly for sweet fluids, where only one reference was found. Both types of environment are apparently detrimental. Crack propagation data indicate little effect of room temperature sour environments at near-threshold stress intensity ranges but an increasingly detrimental effect is found at higher ΔK or K _max levels, Figure 2. The effect of H ₂ S depends on the corrosivity of the test environment, but a substantial effect is observed even in dry H ₂ S or oil with H ₂ S ^[11] . Cyclic frequency was demonstrated to be important, Figure 3, with crack propagation rates increasing as cyclic frequency decreased. ^[9,13] Aqueous sour environments are more damaging, as are high levels of H ₂ S, Figure 2. On the basis of fatigue endurance tests on representative riser girth welds, Buitrago and Weir reported a factor of 10-20 reduction of life for when tested in NACE TM0177 solution B with 0.5 and 1psi H ₂ S partial pressure at a frequency of 0.33Hz ^[10] , Figure 1. An effect of temperature may be anticipated, especially in relation to the kinetics of any hydrogen effects, as for sulfide stress cracking.

One study of fatigue of welds in a sweet environment (3 bar CO₂ with a solution of 10%NaCl and 10%CaCl₂ at 95°C) has been reported by Szklarz. ^[15] He found that crack propagation was faster than in air, believed to be primarily due to crack-tip dissolution, and that there was a strong dependence of propagation rate on loading frequency, Figure 4. This was found at moderately high ΔK (over 600N/mm ^3/2 ). He observed preferential propagation along the weld fusion line and preferential corrosion of the fusion line/HAZ area. The application of inhibitor from the start of test reduced the environmental effect on crack propagation at low ΔK but had little effect if applied after a crack had started to propagate.

Supermartensitic Stainless Steel

Few fatigue data exist for supermartensitic grades, although there are more data for the more conventional, higher carbon martensitic stainless steel grades. ^[16-19] The limited data suggest that fatigue behaviour of welded joints in air is very similar to that of C-Mn steels. Testing of conventional grades in seawater has been performed and gave reduced life compared to tests in air butthis was apparently dominated by the highly corrosive nature of seawater with respect to 13%Cr steel, which gave significant surface corrosion. ^[16,17] Reducing the frequency gave increased crack propagation rates. Produced fluids are much less corrosive, as they are not oxygenated. Crack-propagation rates measured in water and steam are also higher than in air. ^[18,19] Pit formation has been shown to be a critical factor controlling fatigue crack development in martensitic stainless steels, which is time dependent. Hence, any meaningful testing to establish fatigue lives must allow time for any such time-dependent corrosion processes. No data are available for behaviour of supermartensitic or conventional martensitic stainless steels in seawater with cathodic protection or in produced fluids.

Duplex and Superduplex Stainless Steels

The fatigue performance of welded joints in ferritic-austenitic stainless steels in air is broadly similar to that of C-Mn structural steels. ^[20,21] When exposed to seawater, both with and without cathodic protection, higher fatigue crack propagation rates were observed when K_max was above around 700-1000N/mm ^{3/2 [20,22-27]} . Below this level, especially in the near threshold regime, crack propagation behaviour was similar to that in air, as was the fatigue threshold. The increase in propagation rate was greater under cathodic protection. This behaviour was associated with the appearance of a cleavage-like fracture morphology, suggesting an effect of hydrogen embrittlement in both cases. ^{[22,23,26,27]} Without cathodic protection, the active hydrogen presumably arises from corrosion at the crack-tip. Crack propagation rates increased with decreasing frequency. It was found that corrosion fatigue in seawater was sensitive to microstructure, with a grain coarsened heat affected zone with high ferrite content showing higher propagation rates than parent steel. ^[27] Overall, the data are consistent with a stress corrosion fatigue mechanism, where the static mode of crack propagation is hydrogen embrittlement stress cracking. It is likely, however, that dynamic straining of the crack-tip encourages local corrosion and hydrogen generation, so that the stress corrosion component becomes active under less severe environmental conditions than might be expected from static load data. No published data were found for testing in sour or sweet media.

Austenitic Stainless Steels

Austenitic stainless steel welds have generally similar performance to C-Mn steel welds and it has been proposed that the same design S-N curves can be used for both providing there is no influence of the environment. ^[20] Testing in seawater gave a reduction of fatigue life by a factor of approximately two when compared to air. ^[20,21] Overall, few fatigue data are available for austenitic stainless steel welds in corrosive environments ^[21] and no data from tests in produced fluids have been reported. However, materials with austenitic microstructure, including austenitic stainless steels and nickel alloys, are expected to be less susceptible to hydrogen embrittlement in a weld HAZ than the competing supermartensitic and duplex/superduplex grades.

Discussion

For C-Mn steels, it has been demonstrated that sour and sweet produced fluids may reduce fatigue resistance of welded joints substantially compared to air. Hence, the environmental limitations of welded C-Mn steel are greater undercyclic load than under static load and there is a need to establish the cyclic limits. There is less information on the application limits of the various candidate CRAs under cyclic load, but, again, the indications are quite clearly that corrosive environments reduce fatigue resistance and/or increase fatigue crack propagation rate.

The work which has been carried out to date on corrosion fatigue of material/environment combinations relevant to SCRs has concentrated on generating fatigue data (propagation rates and endurances) and there is very limited information on mechanisms. Nevertheless it is evident that there are at least two primary mechanisms which are operating to affect crack propagation, namely crack-tip dissolution and hydrogen embrittlement, in addition to any wedging/closure effects. An improved understanding of these, and possibly other, effects is likely to explain many current uncertainties. For example, where hydrogen is being generated it is usually by corrosion, so both crack-tip dissolution and hydrogen absorption are taking place simultaneously. The overall effect, however, will depend on a range of material and loading factors, which may vary throughout the test duration. A simple example is that crack extension due to hydrogen embrittlement may not occur below a threshold K value, and thus any such crack acceleration will be dependent on K _max , while crack-tip dissolution will be active over a wider loading range. Even confining attention to crack-tip dissolution and hydrogen effects provides a number of different scenarios. For example, dissolution may either accelerate crack propagation, or blunt a crack and thereby retard it, and hydrogen may embrittle a material but can also affect slip, and thus affect fatigue crack propagation.

In principal, it may be possible to develop a predictive system that allows the contributions from the various mechanisms to be added together, such that the results of a matrix of crack propagation tests examining each of the contributory modes of cracking separately could be used to predict behaviour in a wide range of service environments. Such an approach has been proposed previously for crack propagation rates in C-Mn steel welds, ^[5,9] with hardness being used as a descriptor of HAZ microstructural sensitivity to corrosion fatigue under cathodic protection. Such an approach would allow a degree of extrapolation that use of S-N curves would not easily allow. However, this approach requires a comprehensive understanding of the various contributing factors and a quantitative relationship to be developed for each one. Also, the testing required to allow such an approach to be adopted with confidence would be substantial and at present it is more appropriate to undertake case by case testing to establish performance in specific service environments. Furthermore there is an underlying need to establish how environmental factors and weld area microstructure contribute to corrosion fatigue behaviour. For example, it is known that hardness would only be an appropriate descriptor for hydrogen assisted mechanisms, and that grain size and phase balance are important factors in duplex/superduplex stainless steels. Such an understanding will be essential if a prediction system is to be produced, and will ensure that lessons learned from previous experience are applied appropriately.

In the absence of a clear understanding of all the mechanisms contributing to corrosion fatigue, it would still be possible to derive a series of design S-N curves for sour and sweet environments. The advantage of such an approach is that it is empirical and full understanding of the processes occurring is not required. These would need to be based on fatigue tests performed under relevant conditions. Clearly, for such an approach to be robust all the controlling variables, relating to both the weld and environment, must be identified and the test data must be obtained under appropriate conditions with respect to each variable. However, the essential variables controlling weld performance and suitable allowable ranges for procedure qualification are not defined and, hence, there is a need to establish them.

For sour environments, it is anticipated that weld area hardness (especially for C-Mn steels), pH and H ₂ S partial pressure will be dominant, as for sulfide stress cracking under static load. However, this needs to be confirmed. Effects of frequency and any tendency for the environment to form a protective surface corrosion product in the form of a scale (which depends on temperature and will affect crack closure), need to be established.

For sweet environments, the controlling variables are less clear. However, the tendency for preferential weld metal/HAZ corrosion in C-Mn steel weldments is likely to be a significant concern, (as indeed it may be in some sour service). Preferential weld corrosion in sweet conditions has been found to be controlled by a combination of weld metal/HAZ hardness and grain size and the levels of minor alloying elements in the weld metal; with low hardness, fine grain size and low Ni and Si levels in the weld metal being beneficial. ^[28] Inhibitor performance has been highlighted as critical for control of preferential weld corrosion ^[28] but its ability to influence subsequent corrosion fatigue crack propagation has been the subject of only one study ^[15] and more work is needed. Scaling tendency and conductivity of the environment are controlling environmental factors, with low conductivity and intermediate scaling behaviour giving greatest risk of preferential weld corrosion. However, in common with sour environments, atomic hydrogen will be generated in sweet environments and, if absorbed by the material, this could lead to a contribution to crack extension from hydrogen embrittlement, as found for duplex stainless steels in seawater. This may be particularly significant for CRA materials, for which the passive film acts as a barrier to corrosion and subsequent localised hydrogen embrittlement, under static load. At a fatigue crack-tip, the passive film will be repeatedly broken and will not be protective, so some hydrogen generation is expected.

In the absence of either a predictive system, or a body of S-N curve data, there will be a need to design girth welded risers based on project-specific corrosion fatigue test data, measured under conservative test conditions. The difficulty of this approach is that there is currently only limited understanding of the critical variables and hence how to design an appropriate test. Suitable specimen geometry and stressing have been considered by Buitrago and Weir, ^[10] but choice of an appropriate test environment remains subjective. For example, testing at the maximum service temperature might not be conservative, as high temperatures may encourage scaling and/or crack blunting via extensive corrosion, which will reduce crack propagation rates. Too low a pH might also encourage crack blunting. The development of data applicable to the service loading frequency remains problematic due to the time required to perform such testing. Extrapolation to account for differences between the test and service frequencies may be appropriate but the feasibility of such an approach remains unproven.

One solution to the problem of endurance testing at low frequencies is to employ crack propagation rate testing, and then either factor the S-N curve in proportion to the change in crack propagation rate, or use fracture mechanics calculations to predict the S-N curve. However, without a full understanding of mechanisms, there are risks in relying entirely on propagation rate testing. In the first place, as indicated above, K _max needs to be at an appropriate level to ensure that any effects of hydrogen embrittlement are correctly accounted for. Secondly, data generated for deep cracks will not necessarily be representative of the early stages of the fatigue life of an actual girth weld, when the crack is still very small, due to the different environmental conditions in deep and shallow cracks. Finally, the very early stages of crack propagation and/orinitiation may be affected by surface corrosion effects; pitting, preferential weld corrosion, or general corrosion leading to crack-tip removal or blunting. Such effects will not be reproduced by crack propagation rate testing. Thus, at this stage it is advisable for any corrosion fatigue life predictions based on crack propagation rate testing to be confirmed by endurance testing.

Conclusions

• Design of girth welded SCRs for corrosion fatigue service currently requires development of project-specific data.
• Resistance of welded C-Mn steels to fatigue in both sweet and sour produced fluids is substantially lower than in air. The HAZ is typically more susceptible than parent steel, and the same is expected for weld metal. The environmental limitations of welded C-Mn steel under cyclic loading are apparently greater than under static load.
• Welded supermartensitic and duplex/superduplex stainless steels have reduced resistance to fatigue of the HAZ in seawater under CP, compared to air, and reduced properties in sour produced fluids may be anticipated. The HAZ is more susceptible than parent steel.
• There is an absence of relevant data for austenitic stainless steel and nickel alloys (which may be used in clad pipe). However, the corrosion fatigue performance of austenitic stainless steel and nickel based alloy weld HAZs may be expected to be less influenced by hydrogen embrittlement than competing martensitic and duplex stainless steel weld HAZs.
• There is a need to establish the variables and to understand the mechanisms controlling the corrosion fatigue behaviour of all candidate riser materials, with emphasis on the behaviour of the welds, i.e. HAZ and weld metal. A robust approach to design against such corrosion fatigue may be developed by generating a series of design S-N curves. However, this may require very extensive testing and an alternative predictive approach based on summation of the various contributions to crack propagation might have merit.
• Crack propagation rate measurements coupled with fracture mechanics calculations provide an attractive approach to predicting lives at low frequencies. Sufficient uncertainties remain about material-environment interactions early in life and the relative environmental conditions in deep and shallow cracks, however, that endurance testing will continue to be necessary for the time being.
• There is a need for further testing to identify preferred materials and welding processes for welded risers carrying sweet and sour produced fluids.

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