Jean Kittel
email: jean.kittel@ifp.fr

Véronique Smanio

Laurence Garnier

Xavier Lefebvre

Hydrogen induced cracking (HIC) testing of low alloy steel in sour environment -Impact of time of exposure on the extent of damage

Keywords: Hydrogen induced cracking", "HIC", "hydrogen permeation", "hydrogen sulphide", "H 2 S"

The selection of materials for application in the oil and gas industry in environments containing hydrogen sulphide (H 2 S) is an important matter for the safety of equipments. In the presence of H 2 S, hydrogen atoms originating from the cathodic corrosion reaction can diffuse into the steel, resulting in cracking mechanisms such as hydrogen-induced cracking (HIC) or sulphide stress cracking (SSC). For the SSC cracking mode, a diagram defining different regions of environmental severity was established in the 1990's. The requirements for SSC testing of steels in a given environment are well defined and accepted by the industry. On the other hand, fit-for-purpose (FFP) testing for HIC remains a subject of debate. The most common practice consists in 96 hours exposure under 1 bar H 2 S. But this test environment is extremely severe, and is often considered to be too conservative for the selection of steels for use in mildly sour environments. However, a universally applicable FFP test method for HIC has not yet been established. In addition to pH and H 2 S partial pressure, test duration is another major parameter for HIC testing. Therefore, this experimental study was launched to investigate the impact of time of exposure on HIC. Experiments consisted of HIC exposure tests based on the NACE TM0284-2003 standard method. Experiments were conducted in different nodes of the pH -pH 2 S diagram respectively between pH 6.5 and 3.5, and between 10 mbar and 1 bar of H 2 S. For each condition, tests with different times of exposure were realised, from only a few hours and up to three months. Characterisation of HIC consisted in ultrasonic testing. Permeation experiments were also performed in selected pH -pH 2 S conditions. Permeation transients were analysed for calculation of diffusion coefficients. Modelling of hydrogen profiles in HIC specimens was then realised, and compared with the results of HIC immersion tests. For each test condition, it was possible to draw the evolution with time of the extent of HIC. Depending on the environment, the minimum duration necessary to initiate HIC was extremely variable, from a few hours in the most severe conditions to several weeks in low severity conditions. The extent of cracking at equilibrium also seemed to vary with pH and H 2 S. A good correlation was found between HIC and permeation results, confirming that HIC is strongly linked with the increase of hydrogen concentration in the steel. These results confirm that low severity conditions require longer exposures to reach equilibrium levels of internal hydrogen. This must be taken into account for FFP HIC testing in mildly sour environment. With that aim, permeation experiments could represent an interesting technique to assess the time to reach a certain level of hydrogen in the material.

INTRODUCTION

Hydrogen damage is one of the major causes of steel equipment failures in the oil & gas industry [START_REF] Bruno | History and development of TM0284[END_REF][START_REF] Hara | Conditions of hydrogen-induced corrosion occurrence of X65 grade line pipe steels in sour environments[END_REF][START_REF] Kane | NACE committee report 8X294: review of published litterature on wet H2S cracking[END_REF][START_REF] Mahmoud | Overview of hydrogen induced cracking (HIC) of pressure vessels in upstream operations[END_REF][START_REF] Moore | Hydrogen-induced damage in sour, wet crude pipelines[END_REF][START_REF] Smith | An overview of european federation of corrosion documents EFC16 and EFC17[END_REF][START_REF] Smith | Corrosion of carbon steel by H2S in CO2 containing oilfield environments[END_REF]. It is caused by the dissolution of hydrogen in the steel, as a result of the cathodic reduction of H + which accompanies the anodic oxidation of iron in acid media. The reduction of water molecules by cathodic overprotection is another potential source of reduced H + . Several failure modes can occur in service, for example:

• hydrogen induced cracking (HIC), corresponding to internal cracks generated by the recombination of hydrogen to gaseous molecules at certain appropriate traps in the steel, like MnS inclusions of pearlite bands. This failure mode is strictly internal, and does not require an external stress [START_REF] Hara | Conditions of hydrogen-induced corrosion occurrence of X65 grade line pipe steels in sour environments[END_REF][START_REF] Bosch | Fit-for-purpose HIC assessment of large diameter pipes for sour service application[END_REF][START_REF] Hara | The condition of HIC occurence of X65 linepipe in wet H2S environments[END_REF][START_REF] Hermann | HIC assessment of low alloy steel line pipe for sour service application -Literature survey[END_REF][START_REF] Kittel | Hydrogen induced cracking (HIC) -Laboratory testing assessment of low alloy steel linepipe[END_REF][START_REF] Revie | Hydrogen-Induced Cracking of Linepipe Steels .1. Threshold Hydrogen Concentration and Ph[END_REF][START_REF] Revie | Hydrogen-induced cracking of line pipe steels used in sour service[END_REF][START_REF] Tiwari | A study of internal hydrogen embrittlement of steels[END_REF]. • Hydrogen stress cracking (HSC), also designed as sulphide stress cracking (SSC) when H 2 S is present: this cracking mode is generated from the surface of the steel and requires an applied stress [START_REF] Asahi | Prediction of sulfide stress cracking in high-strength tubulars[END_REF][START_REF] Koh | Effect of line pipe steel microstructure on susceptibility to sulfide stress cracking[END_REF][START_REF] Leyer | SSC resistance of a 125 KSI steel grade in slightly sour environments[END_REF][START_REF] Liao | Effect of Molybdenum on Sulfide Stress Cracking Resistance of Low-Alloy Steels[END_REF][START_REF] Marsh | Comparing hydrogen permeation rates, corrosion rates ans sulphide stress cracking resistance for C-110 and P-110 casing steel[END_REF][START_REF] Mendez | Application of hydrogen permeation for monitoring sulfide stress cracking susceptibility[END_REF][START_REF] Mougin | Sulfide stress cracking and corrosion fatigue of steels dedicated to bottom hole assembly components[END_REF][START_REF] Szklarz | Sulfide stress cracking of a pipeline weld in sour gas service[END_REF][START_REF] Tsay | Sulfide stress corrosion cracking and fatigue crack growth of welded TMCP API 5L X65 pipe-line steel[END_REF]. The presence of H 2 S is known to have an aggravating influence.

The selection of steels for oil and gas environments containing H 2 S thus requires appropriate testing, to ensure resistance to cracking in the field conditions. NACE MR-0175/ISO 15156 Part 2 [START_REF]Petroleum and natural gas industries -Materials for use in H2S containing environments in oil and gas production -Part 2: Cracking-resistant carbon and low alloy steel, and the use of cast iron[END_REF] describes most of the testing methods used for SSC and HIC evaluation. For SSC, the pH -pH 2 S severity regions were established on an empirical basis, and are well accepted by the industry. On the other hand, a similar diagram for HIC resistance does not exist yet. In addition, NACE MR0175/ISO 15156 Part 2 does not place any lower H 2 S limit below which an assessment of HIC resistance is unnecessary. For severe conditions (i.e. 1 bar H 2 S), HIC testing methodology and acceptance criteria are the subject of a large consensus. On the other hand, mildly sour testing is still under debate, especially concerning the time of exposure under moderate H 2 S partial pressure. A recent study by the European Pipeline Research Group (EPRG) Corrosion Committee [START_REF] Kittel | Hydrogen induced cracking (HIC) -Laboratory testing assessment of low alloy steel linepipe[END_REF] proposed a preliminary HIC severity diagram, based on testing five different sweet service steels. One of the main conclusions of this study was the necessity to use long term (2 to 4 weeks) immersion tests for mildly sour conditions. However, testing for durations up to several weeks is a major drawback for a rapid fit-forpurpose assessment. Permeation measurements could represent an interesting alternative in addition to standard HIC immersion tests. As illustrated recently [START_REF] Kittel | Effect of membrane thickness on hydrogen permeation in steels during wet H2S exposure[END_REF], the diffusion parameters (subsurface concentration 0 C and diffusion coefficient D ) for hydrogen in steels exposed to sour environment could be obtained from permeation experiments through thick steel membranes.

Starting from these insights, an experimental study was launched with the objective of studying the time dependence of HIC for different pH -pH 2 S conditions. The test methodology was in accordance with NACE TM0284-2003, and several tests with different durations were performed for each condition. HIC cracking was evaluated using ultrasonic testing. Additionally, hydrogen permeation experiments were realised, using 10 mm thick membranes made of the same steel as for the HIC tests. From the permeation transients, the diffusion coefficient D and the hydrogen subsurface concentration 0 C were determined for different pH -pH 2 S. These parameters were then used to calculate the concentration profiles in standard HIC specimens, and the results were compared with HIC tests.

EXPERIMENTAL

Tested Material

The steel used for this study was an API X65 TMCP steel, coming from a sweet service LSAW pipe 42 inches diameter and 1 inch (25.4 mm) thick. Its chemical composition (Table 1) was analysed by optical emission spectrometry (OES), except for carbon and sulphur, analysed by a chemical method. The microstructure consisted mainly of ferrite with a portion of pearlite, with a banded structure particularly pronounced in the centre line. This material is the same as steel "P" of ref. [START_REF] Kittel | Hydrogen induced cracking (HIC) -Laboratory testing assessment of low alloy steel linepipe[END_REF]. Some results of this previous study were used in this paper. The same steel was also used for the permeation study reported in ref. [START_REF] Kittel | Effect of membrane thickness on hydrogen permeation in steels during wet H2S exposure[END_REF]. The sample dimensions were 100 mm long, 20 mm wide and 24 mm thick (0.5 to 1 mm were removed from the inner and outer surfaces). All faces were machined to obtain a finish equivalent to that achieved using 320 grit SiC paper.

Specimens for permeation experiments were 70 mm x 70 mm plates, with a thickness of 10 mm or 0.5 mm. The surface finish was the same as that of HIC samples.

Test solutions and exposure conditions

All solutions were prepared according to EFC16 document [START_REF]Guidelines on materials requirements for carbon and low alloy steels for H2S-containing environments in Oil and Gas production[END_REF], solution A, and thus contained 5% sodium chloride and 0.4% sodium acetate. The bubbling gas is either pure H 2 S, or a mixture of CO 2 with H 2 S (1 or 10%). After a thorough deoxygenating step, the solution was saturated by bubbling gas at ambient pressure. The pH was then adjusted to desired level through addition of deaerated 1N hydrochloric acid or deaerated 1N sodium hydroxide NaOH. The solution was then poured into the test cell.

For HIC experiments, standard glass vessels were used. All HIC tests were performed following the general NACE TM0284-2003 standard, including three replicates for each condition and a volume / surface ratio greater than 3mL.cm -2 . pH measurements of the test solution were performed with a period of 48 hours or less, and adjustments were made if the value was not within -0.1 / + 0.3 pH units of the desired level.

After each test, all specimens were inspected using ultrasonic method. The analysis was performed with a ULTRAPAC system (Euro Physical Acoustics SA) and a 15MHz transducer (1/4" diameter). The system was calibrated according to a well-established specification. For each specimen, the total area of the defects was calculated from the UT scans, and the crack area ratio (CAR) was determined as the ratio between the area of the defects and the total area of the specimen in the short transverse plan. The values reported in this paper are the average of three replicates.

Permeation experiments were conducted with a Devanathan-Stachurski [START_REF] Devanathan | The mechanism of hydrogen evolution on iron in acid solutions by determination of permeation rates[END_REF] cell (Figure 1) that was designed specially for thick specimens. The exposed surface area was 19.6 cm 2 . The experimental set-up consisted of two identical electrolytic cells separated by the steel membrane. The charging surface was left at the corrosion potential. The exit surface was coated with Pd [START_REF] Manolatos | Necessity of A Palladium Coating to Ensure Hydrogen Oxidation During Electrochemical Permeation Measurements on Iron[END_REF][START_REF] Manolatos | The electrochemical permeation of hydrogen in steels without palladium coating. Part I: Interpretation difficulties[END_REF][START_REF] Manolatos | A thin palladium coating on iron for hydrogen permeation studies[END_REF] and held in a 0.1 M NaOH solution at a potential of +300 mV vs. the saturated calomel electrode (SCE). Thus all the hydrogen atoms diffusing through the membrane were oxidised, this oxidation current providing a measure of the hydrogen permeation flux.

RESULTS AND DISCUSSION

Permeation data

As described in details in previous papers [START_REF] Kittel | Effect of membrane thickness on hydrogen permeation in steels during wet H2S exposure[END_REF][START_REF] Crolet | Revisiting hydrogen in steel, part I: theoretical aspects of charging, stress cracking and permeation[END_REF], great care must be taken for the interpretation of hydrogen permeation measurements in H 2 S environments. For too thin steel membranes, the rate determining step is the charging mechanism at the entry face, and the permeation flux becomes independent of the membrane thickness. For a correct evaluation of the hydrogen sub-surface concentration and of the diffusion coefficient, it was shown for the steel of this study [START_REF] Kittel | Effect of membrane thickness on hydrogen permeation in steels during wet H2S exposure[END_REF] that membranes thicker than 3 mm were necessary.

Three experiments were performed on 10 mm thick steel membranes exposed to 900 mbar CO 2 and 100 mbar H 2 S at pH 3.5, 4.5 and 5.5. All the results were then analysed in order to determine the diffusion properties of hydrogen in this steel. Considering one-dimensional diffusion though a membrane of thickness L , the diffusion coefficient was determined from the analysis of the permeation transients ( ) J t . Under these conditions, the ratio between the permeation transient ( )

J t and the steady-state flux ss J corresponding to the second Fick's law depends on the diffusion coefficient D and on the membrane thickness L [START_REF] Crank | The mathematics of diffusion[END_REF][START_REF] Astm G | Standard practice for evaluation of hydrogen uptake, permeation, and transport in metals by an electrochemical technique[END_REF] :

( ) ( ) 2 2 2 1 1 2 1 exp n n ss J t D t n J L π ∞ = × = + - - (1)
Once the steady-state is reached, the hydrogen sub-surface concentration at the entry face 0 C can be determined with the first Fick's law of diffusion from ss J and the diffusion coefficient D :

0 ss J D C L = × (2)
An example of experimental result is given in Figure 2. An excellent agreement is found between the experimental data and calculations with the Fick's laws of diffusion. Table 2 reports the values of diffusion coefficient and subsurface concentration calculated from the experimental results under 100 mbar H 2 S and for different pH. When pH decreases, hydrogen diffusion becomes faster (increase of diffusion coefficient), and the equilibrium hydrogen concentration increases. However, it is worth pointing out that the orders of magnitude of changes of diffusion parameters are far from the changes in hydrogen concentration in the test solution: from pH 5.5 to 4.5, hydrogen concentration in the test solution increases tenfold, while hydrogen subsurface concentration and diffusion coefficient are only multiplied by 1.4 and 2 respectively. Another set of interesting permeation transients was obtained at pH 5.5 and with varying H 2 S partial pressure (Figure 3). Unfortunately, these experiments were performed with thin (0.5 mm) membranes, and it was therefore impossible to extract diffusion parameters. Nevertheless, an interesting trend is observed: under 1 bar H 2 S, the permeation current decreases rapidly after the initial rise following immersion. After a few hours, the level of permeation current is 4 times less under 1 bar H 2 S than under 100 mbar H 2 S. Similar trends have already been observed by Kimura et al. [34]. These authors observed that above a certain H 2 S partial pressure, corrosion rate of steel was decreasing, thus leading to a decrease in permeation rate (Figure 4). This decrease in corrosion rate might be caused by a more stable and protective iron sulphide layer.

According to this experimental result we might expect that exposure to pH 5.5 and 1 bar H 2 S should be less severe than pH 5.5 and 100 mbar H 2 S. This will be verified through long term HIC experiments described in the next part of this paper.

HIC experiments

A great number of HIC experiments were performed. Experimental parameters were pH (varied from 5.5 to 3.5), H 2 S partial pressure (10 mbar to 1 bar) and time of exposure (5 hours to 1080 hours). The range of experimental parameters is presented in Table 3. From all these experimental results, the evolution of CAR with time was analysed. For this purpose, two different representations were adopted. Figure 6 presents CAR evolution at different pH and at constant H 2 S partial pressure. Figure 7 presents CAR evolution at different H 2 S partial pressure and at constant pH. As already illustrated in a previous paper [START_REF] Kittel | Hydrogen induced cracking (HIC) -Laboratory testing assessment of low alloy steel linepipe[END_REF], pH and H 2 S partial pressure have a significant impact on the time necessary for HIC to initiate. At constant H 2 S partial pressure (Figure 6), increasing the pH leads to a delayed HIC initiation. This effect is particularly marked under 1 bar H 2 S: at pH 4.5, cracks are observed after only a few hours immersion, whereas more than 200 hours are necessary at pH 5.5. The same tendency is observed for the results under 100 mbar H 2 S. It also seems that H 2 S partial pressure has a more pronounced influence than pH on the maximum extent of HIC: indeed, under 100 mbar H 2 S, the same CAR values are observed for the longer immersion times for the tests at pH 5.5 to 3.5. However, this conclusion requires further work on a range of sweet service materials.

Another interesting result was found for the experiments at pH 5.5 and under 0.1 and 1 bar H 2 S. HIC experiments revealed that exposure under 1 bar H 2 S at pH 5.5 was far less severe than 0.1 bar H 2 S at pH 4.5 (Figure 7). Under 1 bar H 2 S, more than 200 hours exposures were necessary before cracks were observed, while under 0.1 bar H 2 S, cracks were detected after only 96 hours exposure. This behaviour is quite unexpected: indeed, it is usually agreed that pH and H 2 S partial pressure have a monotone impact on severity, i.e. increase of severity when the pH decreases and H 2 S partial pressure increases.

DISCUSSION: HYDROGEN DIFFUSION VS. HIC

The objectives of the discussion are to check how permeation data could help for the understanding of HIC evolution results. The first part of the discussion focuses on experiments under 100 mbar H 2 S, for which both permeation on thick membranes and HIC tests are available. Diffusion coefficients determined from permeation experiments were used for modelling hydrogen profiles in HIC specimens, which could then be compared to the results of HIC exposure tests.

The second part of the discussion concerns the unexpected results at pH 5.5, for which cracking occurred much more rapidly under 100 mbar H 2 S than under 1 bar H 2 S. An explanation to this result could be proposed from permeation data.

Modelling of hydrogen profiles in HIC samples

For the particular set of experiments under 100 mbar H 2 S, permeation measurements allowed us to determine the diffusion coefficients of hydrogen in the steel (Table 2). For the comparison with HIC results, these diffusion parameters were used for the determination of hydrogen concentration profiles in HIC samples at different immersion times. Those calculations were performed using a finite element model software (Comsol multiphysics). The geometry of the samples was considered parallelepipedic (20 mm wide and 24 mm thick) with an infinite length. Hydrogen diffusion was modelled by a pure Fickian behaviour using experimental 0 C and D values. Hydrogen was supposed to enter from all four lateral faces of the samples. Figure 8 illustrates the evolution of hydrogen concentration with time in the section of a HIC sample at pH 5.5 and with 100 mbar H 2 S. Similar calculations were performed for the diffusion parameter determined from the experiments at pH 4.5 and 3.5 (Table 2).

Then, since all ultrasonic evaluations of HIC samples revealed that cracking was located near the mid-thickness, the analysis of hydrogen profile was restricted to the centre line, i.e. at a depth of 12 mm. Evolutions with time of concentration profiles in the centre line are presented in Figure 9. Obviously, hydrogen concentration increases more rapidly at pH 3.5 than at pH 5.5, and the final equilibrium level is also higher.

Finally, for the analysis of this modelling towards HIC, we considered a critical hydrogen concentration, above which cracking might occur. From a previous study using the same steel [START_REF] Kittel | Hydrogen induced cracking (HIC) -Laboratory testing assessment of low alloy steel linepipe[END_REF], this threshold hydrogen concentration was found to lie between 0.9 and 1 ppm. This threshold value was reported in Figure 9. For each time increment, we also calculated the average hydrogen concentration in the centre line. The corresponding results are plotted in Figure 10. (time increment between each profile = 2 hours).

For the discussion of time to initiate HIC, comparison between Figure 10 and Figure 6 (results at 100 mbar) is extremely interesting. Modelling of hydrogen diffusion in the centre line of HIC specimens (Figure 10) shows that the average concentration reaches the critical cracking level (0.95 ppm) after 3 hours at pH 3.5, after 7 hours at pH 4.5, and after more than 40 hours at pH 5.5. On the other hand, the evolution of HIC extent with time (Figure 6) reveals that cracks are observed after only a few hours immersion under 0.1 bar H 2 S at pH 3.5 and 4.5, while it takes between 24 and 96 hours at pH 5.5. These values are in very close agreement with the timescales obtained from permeation data and diffusion modelling.

This result suggests that permeation measurements might represent an interesting method for the optimisation of HIC testing procedures. From permeation data, it is possible to determine diffusion coefficients, which can then be used for calculations of hydrogen profiles in HIC samples, and thus give an indicative value of the minimum time of exposure necessary to reach a critical hydrogen concentration level. Additionally, when permeation data shows hydrogen subsurface concentration below the critical level, it might become useless to perform HIC tests.

Evaluation of severity from permeation experiments

The second interesting comparison between HIC and permeation tests lies in the experiments at pH 5.5 and under 100 mbar and 1 bar H 2 S. HIC experiments revealed that exposure under 1 bar H 2 S at pH 5.5 was far less severe than 0.1 bar H 2 S at pH 4.5 (Figure 7). Under 1 bar H 2 S, more than 200 hours exposures were necessary before cracks were observed, while under 100 mbar H 2 S, cracks were detected after only 96 hours exposure. At the same time, permeation current measured at pH 5.5 under 1 bar H 2 S rapidly decreased well below the values under 100 mbar H 2 S (Figure 3). This decrease might probably be explained by the build-up of a more protective iron sulphide layer at 1 bar H 2 S, and a reduced corrosion rate and hydrogen reduction reaction. Consequently, hydrogen charging in the metal takes more time under 1 bar H 2 S, and the time to reach the critical hydrogen concentration is strongly delayed.

This result suggests that permeation measurements could be used in order to define pH -pH 2 S severity regions as already suggested by Duval et al. [START_REF] Duval | Hydrogen permeation through ARMCO iron membranes in sour media[END_REF].

CONCLUSIONS

These results confirm that HIC assessment of sweet service linepipe steels in mild sour environments requires precautions. For the lower H 2 S partial pressure and high pH combination, more than 4 weeks might be necessary before HIC initiation. Therefore, it might become necessary to adopt long term immersion tests for the qualification of sweet service steels in terms of HIC resistance. For this purpose, permeation measurements might be an interesting and complementary experimental technique. The determination of diffusion coefficients can be used for calculations of hydrogen profiles in HIC samples, and give a correct evaluation of the minimum time of exposure necessary to reach a critical hydrogen concentration level. Additionally, when permeation data shows hydrogen subsurface concentration below the critical level, it might become useless to perform HIC tests.

Another important result is the fact that pH and H 2 S partial pressure do not always have a monotone impact on severity, i.e. increase of severity when the pH decreases and H 2 S partial pressure increases. Indeed, some results of this study showed faster cracking under 100 mbar H 2 S than under 1 bar H 2 S (at pH 5.5). This is probably linked with the build-up of a more protective iron sulphide layer under 1 bar H 2 S. Once again, this result which required more than 2 weeks HIC testing could have been predicted from 15 hours permeation measurements in both environments. Permeation technique could therefore represent an interesting means to investigate the impact of environmental parameters on HIC.

Figure 1 :

 1 Figure 1: Devanathan-Stachurski cell for permeation measurements on thick specimens.

Figure 2 :

 2 Figure 2 : Hydrogen permeation transient measured for a 10 mm steel membrane exposed to test solution at pH 5.5 under 900 mbar CO 2 and 100 mbar H 2 S. Comparison with ideal Fickian behaviour.

Figure 3 :

 3 Figure 3 : Hydrogen permeation transient measured for 0.5 mm steel membranes exposed to test solution at pH 5.5 and 1 bar or pH 5.5 and 100 mbar H 2 S.

Figure 4 :

 4 Figure 4 : Hydrogen permeation rate through 10 mm thick API X70 steel membranes in synthetic seawater and at different H 2 S partial pressure [34].

Figure 5

 5 Figure 5 illustrates ultrasonic evaluations (time-of-flight representation) obtained on HIC triplicates after different time of exposure under 1 bar H 2 S and at pH 4.5. HIC extent clearly increases with the time of exposure. We can also notice that most of the cracks are located near the centre line, where the banded pearlite structure was the more pronounced.

Figure 5 :

 5 Figure 5 : US-scans after different times of exposure under 1 bar H 2 S at pH 4.5.

Figure 6 : 5 Figure 7 :

 657 Figure 6 : Evolution of HIC extent with time of exposure. Impact of pH at constant pH 2 S.

Figure 8 :

 8 Figure 8 : Calculated hydrogen concentration profile in the 20 mm x 24 mm section of a HIC sample exposed at pH 5.5 and 100 mbar H 2 S (D = 7.5 x 10-6 cm².s-1 ; C0 = 0.97 ppm).

5 Figure 9 :

 59 Figure 9 : Modelling of hydrogen concentration profiles in the centre line (Y = 12 mm) of HIC specimens in test solution under 100 mbar H 2 S and at different pH, using diffusion constants of Table 2 (time increment between each profile = 2 hours).

Figure 10 :

 10 Figure 10 : Average hydrogen concentration in the centre line (Y = 12 mm) of HIC specimens in test solution under 100 mbar H 2 S and at different pH, using diffusion constants of Table 2(time increment between each profile = 2 hours).

Table 1 :

 1 Chemical compositions (wt %) of tested steel.

	C	Mn	Si	P	S	Cr	Ni	Mo	Cu	Nb	V
	0.09	1.56	0.28 0.014 0.001 0.05	0.03	0.01	0.02	0.04	0.05
	Specimens for HIC experiments were sampled following the NACE TM0284-2003 standard
	test method.

Table 2 :

 2 Model parameters obtained by fitting experimental results (10 mm thick membranes at 100 mbar H 2 S) with the first and second Fick's laws of diffusion.

Table 3 :

 3 Range of test durations (hours) of HIC experiments at different pH and pH 2 S.

		10 mbar	100 mbar	1 bar
	pH 5.5	96 -336	24 -96 -336	24 -96 -120 192 -336
	pH 4.5	96 -336	24 -48 -96 -336	5 -8 -14.5 -24 44 -96 -120 -336
	pH 3.5	96 -336 -720 -1080	24 -48 -96 -336	24 -96

ACKNOWLEDGEMENTS

The authors would like to thank G. Parrain, who performed all HIC experiments and participated in the permeation measurements.


Véronique Smanio
email: veronique.smanio@ifp.fr

Marion Frégonèse

Jean Kittel

Thierry Cassagne

François Ropital

Bernard Normand

Wet Hydrogen Sulfide Cracking Monitoring by Acoustic Emission

Equipments which operate in wet hydrogen sulfide (H 2 S) environment can be subjected to Hydrogen Embrittlement (HE) damages like Hydrogen Induced Cracking (HIC), Stress Oriented Hydrogen Induced Cracking (SOHIC), or Sulfide Stress Cracking (SSC). The resistance of steels to HIC and SSC is usually evaluated through standard tests (NACE TM0177 and NACE TM0284). These tests give operators and steel suppliers a pass/fail response for a given application. However, the detailed mechanisms of wet H 2 S cracking are not fully understood yet, although a lot of work has been carried out over the years on this subject. In particular it is of interest to better understand the initiation and propagation phases of HIC and SSC. For this purpose, some laboratory studies have used the acoustic emission (AE) technique to monitor the materials behaviour during HIC and/or SSC tests. Most of the results were promising, as some correlations were found between AE parameters and the cracking susceptibility of different materials. However, the most important gap stands in the analysis of AE data, which is always considered as a bulk, without any attempt to separate the contribution of other acoustic sources (such as general corrosion or hydrogen bubbling in the solution) from the cracking itself. A methodology was developed in this work to discriminate HIC signals from other sources. The aim was to isolate the cracking contribution, and compare HIC propagation mode for different non sour service steels. In addition tensile tests monitored by AE on two sour service steels are also presented. Difference between SOHIC and SSC mechanisms could be evidenced by AE.

Introduction

Wet H 2 S cracking of steels concerns a broad range of service applications in petroleum industry including oil and gas production, refining and gas processing. Several kinds of damage may be caused by H 2 S in contact with low alloyed steels [1]. In low strength plates and pipe steels, the main cracking process is Hydrogen Induced Cracking (HIC). HIC is caused by the accumulation of hydrogen at internal traps, such as non-metallic inclusions, regions of anomalous microstructures (e.g. pearlitic structure) produced by segregation of impurities and alloying elements, and other weak interfaces. When the concentration of hydrogen at the defects reaches a sufficient value, recombination to gaseous hydrogen occurs, with a build-up of high pressure leading to internal cracking. Under conditions of applied or residual tensile stresses, HIC cracks can become aligned in the through thickness direction resulting in 'ladder like' crack array referred as Stress Oriented Hydrogen Induced Cracking (SOHIC). This phenomenon is usually associated with lowstrength ferritic pipe and pressure vessel steels. The other kind of wet H 2 S cracking mode studied in this work is Sulphide Stress Cracking (SSC). SSC leads to the early failure of the metal due to the combined effect of both hydrogen penetration and tensile stress (residual and/or applied). High strength metallic materials and hard welded zones are prone to SSC. Crack initiation of SSC is always located at the surface of the metal. The mechanisms of HIC and SOHIC are generally well accepted. On the other hand, the SSC mechanism is still controversial. Some authors [2,[START_REF] Kane | NACE International Conference[END_REF] make the hypothesis that SSC is the result of hydrogen trapping, in a similar way to HIC. Traps would be located at high internal stress sites (i.e. grain boundaries, inclusions and regions of triaxial stress at notches). In the presence of tensile stress, these locations would become sites for embrittlement and could favour the initiation of brittle fracture. Other authors [START_REF] Crolet | [END_REF] suggest that SSC mechanism is caused by the powerful implantation of protons in the very first layers of metal which expands the crystal lattice. This expansion would generate a permanent surface tension. In addition to the working stress, this tension may then induce delayed cracking.

The main purpose of our study is to improve the understanding of cracking in wet H 2 S media. In the past, several studies [5][START_REF] Cayard | NACE International Conference[END_REF][START_REF] Gingell | NACE International Conference[END_REF][START_REF] Weng | [END_REF][9][START_REF] Amami | Environmental Degradation of Engineering Materials[END_REF] have used Acoustic Emission (AE) technique to monitor Hydrogen Embrittlement (HE) in sour media. Weng et al. [5] found a correlation between HIC damage evaluated by Crack Length Ratio (CLR) [START_REF]NACE International Standard Test Method NACE TM[END_REF] and the AE energy level. Most of the other studies were related to SSC and used AE monitoring during standard SSC tests described in NACE TM0177-96 [START_REF]NACE Standard Test Method NACE[END_REF]. Cayard et al. [START_REF] Cayard | NACE International Conference[END_REF] and also Gingell et al. [START_REF] Gingell | NACE International Conference[END_REF] showed that during SSC tests leading to the failure of the specimen, a higher AE energy build-up rate was recorded compared to no-failure tests. This critical energy build-up rate associated with SSC failure remained dependant on the tested steels and the applied stress level. Moreover, Gingell et al. [START_REF] Gingell | NACE International Conference[END_REF] showed that AE could be used to discriminate initiation and propagation stages of SOHIC (Stress Oriented Hydrogen Induced Cracking) development.

From these studies, it appears that AE is a powerful tool for H 2 S cracking evaluation. Some correlations can be found between the AE energy and the extent of cracking. However, for most of the past studies, the cracking mechanism was not always clearly established or documented, especially between SOHIC and SSC. Furthermore, AE data were treated in a global manner, i.e. considering simultaneously all the AE sources involved during the test. Considering this last point, a step forward was proposed recently by the authors [START_REF] Smanio-Renaud | Proceedings Eurocorr 2007[END_REF][START_REF] Smanio | NACE International Conference[END_REF]. A methodology was described, enabling to separate AE contributions from H 2 evolution due to acid corrosion, FeS layer build-up and evolution, and HIC. A correlation chart representing the absolute energy of the signals as a function of their duration was used for the discrimination (Figure 1). Thus, it was possible to separate AE events related to HIC from the background noise, allowing a better analysis.

In the present work, this methodology was applied in order to compare HIC mechanisms of two different steels. It was also used for tests under tensile load allowing distinguishing SOHIC and SSC rupture modes. Experimental procedure

Experimental set up

The experimental set up (Figure 2) was based on the standard tensile test described in NACE TM0177-96 (method A) [START_REF]NACE Standard Test Method NACE[END_REF]. All tests were conducted with the same experimental set up. For HIC tests, no load was applied to the test specimens. For SSC or SOHIC tests, a constant strain was applied with the proof ring. Cylindrical tensile specimens were used for both study (HIC and SSC). Only the gauge section of the specimen was immersed in the corrosive solution. Tests were conducted in buffered solution prepared according to EFC 16 [START_REF] Eliassen | Oil and gas production 2 nd Edition European Federation of Corrosion[END_REF] (5 wt% NaCl and 0.4% CH 3 COONa in distilled water) and under 1 bar H 2 S. De-aeration of the solution was made by nitrogen bubbling ensuring less than 15 ppb dissolved oxygen before H 2 S introduction.

A circulation of the corrosive fluid (sour media) was set-up between the test cell and a vessel, in which gas bubbling was maintained. This procedure allowed to avoid any interfering AE noise detection (e.g. gas bubbling in the test cell would result in AE) and to maintain constant the electrochemical conditions in the cell during the test (easier pH control). The pH was measured with a pH meter located in this second vessel and adjusted every day with HCl or NaOH additions. Signal duration (µs) Energy (10 -18 J)

Back AE monitoring was performed with two AE sensors placed at the top and bottom of the specimen (Figure 2). AE instrumentation consisted of a transducer, a preamplifier and a data acquisition system. The acquisition system was completely computer controlled. For each detected AE signal, online analysis was performed and several parameters were collected such as: o Events number: number of AE signals detected (discontinuous emission), o Amplitude: maximal amplitude of the considered AE event (peak), o Rise time: time between the first overshoot of the defined threshold and the peak, o Counts number: number of times that the threshold is overshot for a given AE event, o Duration: time between the first and the last overshoot of the define threshold, o Energy: envelop of the signal.

After HIC tests, damages in the specimen were evaluated by ultrasonic testing. From US scans, it is possible, with a software, to obtain an isosurface picture of the defects in the sample. In addition, metallographic examinations were made in three equidistant cuts along the length of the samples. After SSC and SOHIC tests, fracture surfaces of SSC samples were examined by SEM (Scanning Electron Microscopy). A 3D reconstruction of the fracture surfaces was made using a Hirox 3D microscope. This microscope is motorized in z direction allowing taking several pictures of the fracture surface with different focus. Then, it is possible to reconstruct a full 3D picture of the fracture surface and to extract some profiles.

Materials

Four different steels were studied: three of them were low alloy steel plates for line pipe applications and the last one was a well tubing material. Two grades were Sweet Service (SwS) and two grades were Sour Service (SS). The chemical compositions of all steels were analyzed by optical emission spectroscopy (OES), except for carbon and sulphur, analyzed by a chemical method (Table 1). X65 (SwS) 0,09 1,56 0,28 0,014 0,001 0,05 0,03 0,01 0,02 0,040 X65 (SS) 0,046 1,36 0,322 0,008 0,001 0,041 0,036 0,008 0,047 0,045 X70 (SwS) 0,060 1,25 0,227 0,007 0,003 0,092 0,194 0,017 0,028 0,038 C110 (SS) 0,309 0,394 0,343 0,015 0,002 0,964 0,037 0,834 0,018 0,033

The line pipe steels X65 SwS, X70 SwS and X65 SS exhibit a low carbon content (<0,1%) compared to the C110 casing steel, which presents higher mechanical properties (Table 2). The C110 grade also contains alloying elements, such as chromium and molybdenum, providing a better resistance to hydrogen embrittlement.

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The tensile properties of the steels in the longitudinal direction are given in Table 2. The microstructure of each steel was observed in the long transverse direction (LT) using an optical microscope, and after polishing down to 0.05µm diamond suspension and etching with a Nital solution (2%vol. nitric acid and ethanol). The different microstructures at the mid wall location are presented in Figure 3 to Figure 6. X65 SwS, X70 SwS and X65 SS has a ferrite pearlite microstructure. Microstructures of X65 SwS, X70 SwS are highly oriented in the rolling direction, which can affect their HIC resistance (Figures 3 and5). Pearlite bands in X65 SwS are located on parallel planes at different width whereas X70 SwS steel exhibits only one pearlite band located in the middle plane of the sample.

For the X65 SS grade, the microstructure is equiaxed and homogenous, which explains its good HIC resistance (Figure 4). The microstructure of C110 grade is a tempered martensite with macrosegregation (Figure 6). This microstructure is usually resistant to HIC [START_REF] Galland | Prévention et lutte contre la corrosion B[END_REF].

Results and discussion

Acoustic emission was applied to compare HIC propagation in two sweet service steels (X65 SwS and X70 SwS grades) and then to identify the rupture mode under tensile stress, SSC or SOHIC, of both sour service steels (X65 SS and C110 SS grades).

HIC cracking behaviour of sweet service steels HIC tests were performed on X70 SwS and X65 SwS at pH 4.5 under 1 bar H 2 S and for 90 hours exposure.

All the AE data recorded during the tests were then filtered as described in [START_REF] Smanio-Renaud | Proceedings Eurocorr 2007[END_REF][START_REF] Smanio | NACE International Conference[END_REF], in order to keep only the events related to HIC. This procedure is illustrated in figure 7 for the experiment on the X70 SwS steel, and in figure 8 for the X65 SwS steel. Each point on these graphs represents an AE event. For this experimental configuration and AE equipment, the HIC events are characterized by durations above 1500 µs [START_REF] Smanio-Renaud | Proceedings Eurocorr 2007[END_REF][START_REF] Smanio | NACE International Conference[END_REF]. Signal duration (µs) Energy (10 -18 J)

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Once the filtering of AE data was performed, the evolution with time of HIC related signals was examined for X70 SwS steel (Figure 9) and for X65 SwS steel (Figure 10). The evolution of the AE energy related to HIC is clearly different for X70 SwS and X65 SwS.

For the first steel, only one high energy jump is observed after 40 hours (Figure 9). On the other hand, AE signals related to HIC in X65 SwS proceed by successive steps between 43 hours and 68 hours. This difference may be explained by the number of cracks, by the difference in propagation path, or by different types of hydrogen traps in the steels.

After each test, ultrasonic testing was performed on the two samples. 3D reconstructions of the samples were made using US scans of X65 SwS (Figure 11) and of X70 SwS (Figure 12). These representations give a global view of the damage state in the samples. For the X70 SwS, only one large crack is observed in the middle of the sample. On the other hand, 6 distinct cracks can be identified in the X65 SwS sample.

These results are confirmed by metallographic cross sectioning. Figure 13 shows the observation of cracks in L plan for X70 SwS and Figure 14 presents the longitudinal cross section for X65 SwS. One main fine crack is identified in X70 SwS sample. Four large cracks can be observed by metallographic observations (Figure 14) inside the X65 SwS steel.

The difference in cracking damage in the two steels suggests that AE cumulative energy of HIC signals is correlated with HIC propagation. For the X70 SwS steel, one single crack propagates in the sample giving rise to one high AE energy jump. For the X65 SwS steel, several cracks propagate and result in several successive steps in cumulative energy evolution with time.

This comparison shows that AE provides real time information, not only on HIC kinetics but also on the details of the cracking. In this example, two types of cracking behaviours were identified. HIC in the X70 grade appears to be a single cracking event whereas X65 SwS seems to be multi-cracking. This difference in cracking might be explained by the difference in hydrogen trapping. The difference of cracking propagation seems to be related to the microstructure. X65 SwS steel presents multiple pearlite bands (Figure 3) whereas X70 SwS grade exhibits just one pearlite band in the middle of the sample (Figure5). In both steels, pearlite bands appear to be the main sites of hydrogen trapping leading to HIC.

In conclusion, AE seems to be a sensitive technique for the evaluation of HIC. The evolution of cumulative energy gives a foot print of the cracking process and information on HIC kinetics. Real-time information on the initiation, propagation and arrest of HIC can be obtained. In addition cracking modes (single or multiple cracking) can be identified.

Discrimination between SSC and SOHIC Tests performed on X65 SS and C110 SS steels showed that these steels are not susceptible to HIC (no applied stress) at pH 3.5 under 1 bar H 2 S. Then, tensile tests monitored by AE were made on these two sour service steels under the same conditions at 90% of actual YS. AE results of the test conducted on X65 SS steel are represented in the correlation chart that was used for HIC discrimination, i.e. absolute energy versus duration of the signals (Figure 15). SEM observation of the fracture surface of X65 SS is shown in Figure 16. Several internal cracks with a HIC appearance can be identified. Ductile microvoids can also be observed at higher magnification on the fracture surface (figure 20). Signal duration (µs) Energy (10 -18 J)

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A profile of the fracture surface was extracted from a 3D reconstruction of the surface using a 3D microscope (Figure 17). The 45° rupture observed on this profile is typical of ductile fracture. Moreover, no perpendicular brittle flat fracture plane is observed. From these observations, the fracture of X65 SS sample cannot be identified as SSC. The observed plastic deformation and the presence of internal cracks perpendicular to the fracture surface on one hand and the detection of HIC signals by AE on the other, strongly suggest a SOHIC mechanism in X65 SS under these test conditions.

Figure 18 shows the fracture surface of the C110 SS steel observed by SEM. Two areas can be identified. Fracture seems to initiate from the surface of the sample, and perpendicular to the loading direction (area 1). A trangranular fracture is observed in area 1 (figure 21) while microvoids are identified in area 2 (figure 22).

A profile of the fracture surface can be extracted from a 3D reconstruction of the surface using a 3D microscope (Figure 19). This profile is very different from the profile of the X65 SS sample. Area 1 which is brittle-like corresponds to a flat surface perpendicular to the direction of applied load. Area 2 identified as ductile corresponds to the 45° region. This profile is typical of a SSC rupture. Sulfide stress cracking corresponds to the brittle area, then, as the section of the loaded part is reduced, the applied stress reaches the UTS and the sample fails by ductile fracture. Two kinds of fracture have been presented. Fracture of X65 SS and C110 SS were identified as SOHIC and SSC, respectively. AE results of this two tests are different since in one case HIC signals were detected (X65 SS) and for the other one no HIC signals were detected (C110 SS). These results show that SSC and SOHIC fracture can be discriminated by AE. Moreover, SSC does not give rise to HIC signals and initiate on the specimen surface so it indicates that SSC is not related to subsurface trapping of hydrogen.

Sulfide stress cracking signals (not presented in this paper) exhibit a lower energy than HIC signals. Some indications to explain the difference observed between HIC and SSC signals can be inferred from the literature on stress corrosion cracking. Without forgetting that AE is dependent of each system it is possible to make a qualitative comparison of AE energy from different works. Mc Intyre and Green [START_REF] Mcintyre | [END_REF] showed that intergranular fracture gives rise to higher energy AE events than transgranular fracture during stress corrosion cracking of steels. In the line pipe steel tested in this study, the HIC propagation mode at the ferrite/pearlite interface can be assimilated to an intergranular crack propagation, whereas for the C110 steel SSC leads to a transgranular fracture. This could explain the higher AE energy of HIC/SOHIC compared to SSC. In addition, it is usually considered that the ability of crack detection is a function of crack velocity [START_REF] Rothéa | Corrosion sous contraintes : phénoménologie et mécanismes[END_REF]. Higher velocity cracking such as cleavage give rise to higher energy AE signals. Therefore, a difference in crack velocity could also explain the difference in AE cracking detection between SSC and HIC/SOHIC. In summary, the AE analysis confirms that different mechanisms of crack propagation are involved for HIC/SOHIC and SSC.

Conclusions

AE monitoring was applied to detect and analyse wet H 2 S induced cracking modes in different steels in sour media, i.e. HIC, SSC and SOHIC. First, the discrimination of the different active AE sources during an HIC test such as H 2 evolution, FeS layer, and cracking was made. Then, this method was applied to study the HIC susceptibility of different steels. This method provides new information on the cracking kinetics but also helps to distinguish HIC cracking mode (single or multiple cracking).

Tests conducted under tensile stress monitored by AE technique were carried out in order to study SSC. It was found that SSC and SOHIC can be differentiated by AE analysis. SSC signals cannot be discriminated easily since they are less energetic than those associated with HIC. The AE analysis confirms that the SSC and HIC/SOHIC mechanisms are different.

In the future, HIC tests under mildly sour conditions will be conducted in order to better understand the influence of environmental parameters on the kinetics and extent of HIC. SSC initiation and propagation will also be investigated using AE.

Figure 1 :

 1 Figure 1: Correlation chart representing absolute energy as a function of signals duration for a test conducted in the EFC 16 solution at pH 4.5 under 1bar H 2 S, on a X65 SwS specimen.

Figure 2 :

 2 Figure 2: Experimental device.

Figure 3 :

 3 Figure 3: Microstructure of X65 SwS grade (LT plane).

Figure 4 :

 4 Figure 4: Microstructure of X65SS grade (LT plane).

Figure 5 :

 5 Figure 5: Microstructure of X70 SwS grade (LT plane).Figure 6: Microstructure of C110 SS grade (LT plane).

Figure 6 :

 6 Figure 5: Microstructure of X70 SwS grade (LT plane).Figure 6: Microstructure of C110 SS grade (LT plane).

Figure 7 :

 7 Figure 7: Correlation chart representing absolute energy as a function of signals duration for a test conducted in the EFC 16 solution at pH 4.5 under 1bar H 2 S, on a X70 SwS specimen.

Figure 8 :

 8 Figure 8: Correlation chart representing absolute energy as a function of signals duration for a test conducted in the EFC 16 solution at pH 4.5 under 1bar H 2 S, on a X65 SwS specimen.The comparison between Figure7and Figure8shows a difference in AE related to HIC for the two steels. More HIC signals were detected during the test performed on X65 SwS steel compared to the test done with the X70 SwS grade.

Figure 9 :

 9 Figure 9: Evolution of cumulative energy of HIC class during a test performed on the X70 SwS steel grade at pH 4.5 under 1 bar H 2 S.

Figure 10 :

 10 Figure 10: Evolution of cumulative energy of HIC class during a test performed on the X65 SwS steel grade at pH 4.5 under 1 bar H 2 S.

Figure 11 :

 11 Figure 11: 3D picture of cracking in X70 SwS steel (reconstruction from US scan).

Figure 12 :

 12 Figure 12: 3D picture of cracking in X65 SwS steel (reconstruction from US scan).

Figure 13 :

 13 Figure 13: Metallographic observation of cracking in X70 SwS steel (L plan).

Figure 14 :

 14 Figure 14: Metallographic observation of cracking in X65 SwS steel (L plan).

Figure 15 :

 15 Figure 15: Correlation chart representing absolute energy as a function of signals duration for a test conducted in the EFC 16 solution at pH 3.5 under 1bar H 2 S with an applied stress of 90% YS on a X65 SS specimen. HIC signals were detected during the tensile test performed on the X65 SS steel. This was not the case during the test on C110 SS steel. Hence, from AE results, HIC cracks are expected in the X65 SS specimen.

Figure 16 :

 16 Figure 16: Fracture surface of X65 SS steel tested in the EFC 16 solution at pH 3.5 under 1bar H 2 S with an applied stress of 90% YS.

Figure 17 :

 17 Figure 17: Profile of the fracture surface of X65 SS steel tested in the EFC 16 solution at pH 3.5 under 1bar H 2 S with an applied stress of 90% YS.(red line on figure16).

Figure 18 :Figure 19 :

 1819 Figure 18: Fracture surface of C110 steel tested in the EFC 16 solution at pH 3.5 under 1bar H 2 S with an applied stress of 90% YS.

Figure 20 :

 20 Figure 20: Detail of X65 SS fracture.

Figure 21 :

 21 Figure 21: Detail of the brittle area of C110 fracture surface.

Figure

  Figure 22: Detail of the ductile area of C110 fracture surface.

Table 1 :

 1 Chemical composition (wt %) of steels.

	Steel	C	Mn	Si	P	S	Cr	Ni	Mo	Cu	Nb

Table 2 :

 2 Mechanical tensile properties of studied steels.

	Steel	YS(MPa)	UTS (MPa)	El%
	X65 SwS	523	649	24
	X65 SS	529	571	47
	X70 SwS	548	615	28
	C110 SS	798	887	20
	(YS: Yield strength, UTS: Ultimate Tensile Strength, El: Elongation)

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Acknowledgement

Financial support from Total and IFP is gratefully acknowledged. The authors express their appreciation for assistance with the 3D observations of fracture surfaces to N. Mary from INSA-Lyon.