Jean-Gabriel Sezgin
email: jean-gabriel@aist.go.jp

Cédric Bosch

Aurore Montouchet

Gilles Perrin

Andrej Atrens

Krzysztof Wolski

Coupled hydrogen and phosphorous induced initiation of internal cracks in a large 18MnNiMo5 component

Keywords: Hydrogen embrittlement, phosphorous segregation, fracture toughness, blistering, 16-20MND5

This study analyzed a crack in a scrapped 18MnNiMo5 80-tonne shell. The root cause of crack formation was the low fracture toughness caused by a combination of a high hydrogen content and phosphorous segregation, caused by an inappropriate heat treatment. Cracking initiated on MnS inclusion clusters enabled by hydrogen desorption at the MnS/matrix interfaces. Crack propagation was due to the internal hydrogen pressure causing the crack tip stress intensity factor to exceed the locally-reduced fracture toughness. This new hydrogen induced cracking (HIC) mechanism can be designated as Hydrogen Induced Localised Damage (HILD). Appropriate process controls can eliminate such defects.

Introduction

The purpose of this paper is to understand the cracking during production in an 80-tonne, hotformed component, made from 18MnNiMo5 steel (herein designated as 18MND5). This component is typical of a component manufactured by forging a large hollow ingot. A typical example is a reactor pressure vessel, of size of 4 m in diameter, 4 m in height, and approximately 200 mm in thickness. Systematic control is required at each production step. The detection of a significant defect by non-destructive inspection requires interruption of production of the component, and scrapping of the component may be required if the defect is confirmed.

For the component analyzed herein, the formation of such a defect typically requires contributions from the following three factors: (i) a heterogeneous microstructure, (ii) the presence of MnS inclusions, and (iii) a sufficient hydrogen concentration.

Firstly, the microstructure of a typical, heavy-forged, low-alloy steel ingot is heterogeneous at several scales: (i) carbon macro-segregation, (ii) ghost lines, and (iii) micro-segregations [START_REF] Naudin | Intergranular fracture stress and phosphorus grain boundary segregation of a M n-Ni-M o steel[END_REF] [START_REF] Bedel | Comparaison des aciers de cuve utilisés dans les Rupther et Krakatoa[END_REF]. The two latter heterogeneities are characterized by increased concentrations of the alloying elements Mn, Mo, Ni, Si, S and especially P, such that the local concentration can be twice the nominal composition.

Secondly, such components contain MnS inclusions, although the concentration of S has been continuously decreased by improved desulfurization techniques, so that the average S content has been decreased from around 20 ppm in 1980 to 2 ppm in the present practice. The elongated MnS inclusions are flat and wrought, sometimes as thick as 15 µm in a forged component. Their distribution is typically heterogeneous. They form flat ellipsoidal clusters, up to 1 mm in size, composed of dozens of individual inclusions. They are preferentially located in, or at the vicinity, of ghost lines.

These three factors impose the following critical requirements on the thermal history of the component during manufacturing. (i) Hydrogen must be sufficiently degassed by maintaining the component above 200 °C for a sufficient duration, which depends on the initial hydrogen content. (ii) Internal stress buildup must be prevented by sufficiently slow cooling rates, but the cooling rates must nevertheless be sufficiently fast to prevent phosphorous segregation at prior austenite grain boundaries. (iii) The bainitic microstructure must be optimized by the heat treatment, to achieve the required mechanical properties of strength, ductility and toughness. The details of these processing requirements depend on the composition of the steel. A deviation from these requirements can lead to defect formation, initially as a small cavity that may grow and compromise structure integrity.

Internal hydrogen can play a critical role, if the hydrogen concentration exceeds the critical value, which is typically below 1 ppm. The precise value depends on the manufacturing details. The mechanisms of hydrogen embrittlement are explained in the literature [START_REF] Coudreuse | Fragilisation par l'hydrogène et corrosion sous contrainte ; Corrosion sous contraintephénoménologie et mécanismes[END_REF] [6] [START_REF] Liu | A review of the influence of hydrogen on the mechanical properties of DP, TRIP and TWIP advanced high strength steels for auto construction[END_REF] [8] [START_REF] Ramamurthy | Stress Corrosion Cracking of High Strength Steels[END_REF] [START_REF] Dunne | Investigation of the effect of electrolytic hydrogen charging of X70 steel : I. The effect of microstructure on hydrogen-induced cold cracking and blistering[END_REF]. Hydrogen influences plastic behavior (e.g. by dislocation displacement [START_REF] Kirchheim | Reducing grain boundary, dislocation line and vacancy formation energies by solute segregation. I. Theoretical background[END_REF] [12] [START_REF] Kirchheim | Revisiting hydrogen embrittlement models and hydrogen-induced homogeneous nucleation of dislocations[END_REF]), can decrease mechanical properties [START_REF] Venezuela | The influence of hydrogen on the mechanical and fracture properties of some advanced high strength steels studied using the linearly increasing stress test[END_REF] [15] [START_REF] Liu | Influence of hydrogen on the mechanical properties of some medium strength Ni-Cr-M o steels[END_REF], can decrease interatomic cohesive forces [START_REF] Troiano | The role of hydrogen and other interstitials in the mechanical behavior of metals[END_REF] [START_REF] Petch | Delayed Fracture of M etals under Static Load[END_REF], can decrease fracture toughness [START_REF] Atsuoka | Effect of hydrogen and prestrain on tensile properties of carbon steel[END_REF], can decrease the threshold stress [START_REF] Petch | The lowering of fracture-stress due to surface adsorption[END_REF] for crack initiation, and can increase the crack growth rate [START_REF] Atsuo | Fatigue crack growth properties of quenched and tempered Cr-M o steel in 0.7 M Pa hydrogen gas[END_REF]. Hydrogen cracks in heavy components can occur during the manufacturing process. These cracks can result from the diffusion of internal hydrogen to cavities at MnS inclusion, where the hydrogen desorbs, pressurizes the cavity sufficiently to cause crack extension when the crack tip stress intensity factor exceeds the locally, hydrogenreduced fracture toughness. The hydrogen concentration is hard to control during manufacturing process, for heavy components exceeding 80 tons [27] [START_REF] Campbell | Complete Casting Handbook : M etal Casting Processes, Techniques and Design[END_REF].

The aims of this paper are (i) to clarify the origin of the internal cracks detected by ultrasonic measurement at the end of the manufacturing process of the 200 mm thick component made of 18MND5 low-alloy steel, (ii) to propose the scenario of crack initiation and propagation, and (iii) to explain the underlying physical mechanisms.

The origin of the internal crack was clarified by undertaking the following. One typical internal crack was machined from the component and subjected to fractographic analysis at EMSE. A series of ultra-high-vacuum heat treatments reproduced the morphology of the internal surfaces of this industrial crack and explained when the cracking occurred. The relevant mechanical properties of the steel were measured by mechanical tests with simultaneous hydrogen charging. Then, a complete scenario of crack initiation and propagation was proposed, together with an explanation of the underlying physical mechanisms. The potential relevance of these findings is related to other industrial situations.

Analysis of the component

Thermo-mechanical history

Table 1 presents the chemical composition of the 18MND5 low-alloy steel used in the production of this component. The component was scrapped at the end of thermo-mechanical treatment because ultrasonic mapping found penny-shaped cracks, 3 to 20 mm in diameter, located at the mid thickness of the component, mainly at three fourth height, where carbon segregation was maximal. Figure 1 presents the actual thermal cycle of this component. The normal thermal cycle consists of two heat treatments: (i) the outgassing heat treatment and (ii) the precaution heat treatment (austenitisation, quenching and tempering). The outgassing heat treatment is aimed to reduce the average content of internal hydrogen. The outgassing maintains the component at a high temperature, in the ferritic phase, in order to enhance hydrogen diffus ivity. No detrimental effects of internal hydrogen are expected after this step. The precaution heat treatment consists of austenitisation, quenching and tempering, and produces the bainitic microstructure to yield the required mechanical properties.

However, the heat treatment was more complex for the component being herein analyzed. Firstly, the standard hydrogen outgassing duration was applied, but it appears that this did not result in a sufficient hydrogen removal. Secondly, after the austenitisation, an unplanned slow cooling occurred in the furnace, instead of the scheduled water quenching. It is expected that this slow cooling produced a ferritic-pearlitic microstructure instead of the required bainitic microstructure produced by the water quenching. After the slow cooling to room temperature, the component was stored for the period of approximately two months. As such a slow cooling was inadequate to produce the required bainitic microstructure, the complete precaution heat treatment was then reapplied to the component. The component was austenitized for the second time, water quenched, tempered and cooled at 30 °C/h to room temperature. The component was machined to the required dimensions. Then, however, ultrasonic measurements revealed the presence of pennyshaped cracks, which led to the decision to stop the production, and scrap the component.

Parts of the component containing ultrasonic crack indications were made available for analysis at EMSE.

Analysis of cracks in the component

Ultrasonic inspection of the material made available to EMSE located a crack within a 30 mm x 30 mm x 40 mm bloc. A thinner sample was extracted from this bloc and analyzed by tomography (Figure 2-a) and by optical microscopy (Figure 2-b). The overall size of the crack was of the order of 20 mm in length on the cross section and its opening was 10 to 20 µm. The position of this crack within the microstructure is presented in Figure 3-a) (SEM image) and Figure 3-b) (EBSD image). The analysis of the crack path with respect to the microstructure usually is extremely helpful [23] [24] in determining the mechanisms of environmentally assisted fracture. However, in this particular case, the second austenitisation must be considered, and for the case the cracks were formed before this austenitisation, the EBSD analysis would be meaningless. A second crack was extracted from the initial bloc, and opened after cooling to the temperature of liquid nitrogen, by a purpose-designed bending device, in order to prevent plastic deformation of the surface and to promote brittle facture of the remaining material. Figure 4 and 5 present the scanning electron microscopy (SEM) fractographic analysis of the penny-shaped crack.

Figure 4a reveals that the crack may have started from the area indicated by a dashed ellipse and propagated in the radial direction. Such cracks usually initiate on microstructural defects [START_REF] Du | Effect of microstructures and inclusions on hydrogen-induced cracking and blistering of A537 steel[END_REF] [26], including inclusions. Figure 4b) provides a magnification of the area within the ellipse of Figure 4a). This micrograph illustrates the smoothed topography of the surface and the presence of small thermal grooves on the surface that indicate that this crack was exposed to a temperature high enough for massive surface diffusion. Furthermore, the smoothing of the whole internal surface of the crack indicates that no propagation occurred after the last heat treatment. In particular, this observation indicates that this crack could not have formed during the one-month period (dedicated to the machining to final dimensions and necessary to perform ultrasonic measurements, see "machining & QC" in Figure1) at room temperature prior the final quality control. Figure 5a) presents an area of mainly intergranular fracture. Figure 5b) presents the same area using backscattered electrons. The contrast at the bottom of the micrograph indicates a series of MnS inclusions forming a large cluster in the immediate vicinity of the intergranular facets. Figure 5c) is a magnification of an area located at the left of Figure 5a). It represents mainly intergranular facets, located on the top of the micrograph, together with some transgranular facets. The smoothing induced by heat treatment may have altered the details of the type of fracture mode and have hidden additional features such as secondary cracks. The mainly intergranular fracture was only present on a limited area of the whole fracture surface, which was mainly transgranular and smoothed, as illustrated in Figure 4a) at low magnification and Figure 4b) at high magnification.

Experimental Methods

UHV heat treatments to reproduce fracture surface smoothing

Experiments were carried out to determine the temperature and duration of the heat treatment, which would transform a mixed brittle intergranular/transgranular fracture surface into the smoothed crack surface, similar to the one extracted from the component. These experiments required (i) a model brittle surface and (ii) heat treatment conditions to reproduce what happened inside the crack within the component, i.e. totally excluding any oxidation in order to allow surface diffusion.

The two model brittle surfaces were obtained by fracturing specimens of 16MND5 steel in liquid nitrogen, i.e. well below the brittle to ductile transition temperature. These resulted in mainly transgranular fracture with some intergranular facets corresponding to the areas of microsegregation. The first specimen had a fully transgranular fracture while the second contained both transgranular and intergranular areas.

These two 1 mm² fracture surfaces were observed by SEM; and then heat treated, and the same areas were again observed in order to quantify the morphological evolution. Heat treatment was done using an ultra-high-vacuum, radiative furnace mounted on an Auger-XPS spectrometer and operated at a vacuum of 10 -9 mbar.

One specimen was subjected to a set of heat treatments aimed at reproducing tempering heat treatments up to the temperature of austenitisation. The specimen was heat treated at 625 °C for 4 h followed by SEM observation, heat treated at 725 °C for 4 h followed by SEM observation, and heat treated at 825 °C for 4 h followed by SEM observation.

Another specimen was heat treated at 950 °C for 3 hours in order to reproduce the austenitisation heat treatment.

Tensile properties and fracture toughness tests

Slow strain rate tensile tests (SSRT) were carried out to characterize the hydrogen influence on the ultimate tensile strength, UTS, maximal elongation and cross section reduction. Two microstructures were tested using material from the component: (i) the bainitic delivery microstructure and (ii) a ferrite-perlite microstructure that was obtained by austenitisation at 900 °C followed by a slow cooling at 6.8 °C.h -1 in order to reproduce the microstructure of the component after the furnace cooling and before the second austenitisation.

The SSRT tests were carried out using axisymmetric smooth tensile specimens with a diameter of 3 mm. They were machined along the radial direction. This direction directly evaluates the influence of planar segregated bands as these are contained in the specimen cross section. The hydrogen charging of the specimens was performed by the application of a cathodic potential on the working electrode using a conventional three-electrode cell with a platinum auxiliary electrode, a saturated calomel electrode (SCE) as a reference electrode, and a 0.3 wt.% NaCl solution. The hydrogen was produced on the specimen outer surface, and a part of the hydrogen entered the steel by adsorption-absorption. Some tests were performed at the free corrosion potential. To ensure penetration of hydrogen and testing repeatability, metal oxidation was inhibited by deoxygenating of the saline solution by bubbling gaseous pure nitrogen into the deoxygenation tank as shown in Figure 6. Slow strain-rate tests (SSRT) at 1.0 x 10 -6 s -1 or 5.0 x 10 -7 s -1 were performed at room temperature using the following four hydrogen charging conditions: 1. 24 h precharging at -1200 mV SCE followed by a SSRT at this cathodic potential. 2. 24 h precharging at -1200 mV SCE followed by a SSRT at the free corrosion potential, E corr .

3. 8 h precharging at -1200 mV SCE followed by a SSRT at the free corrosion potential, E corr . 4. 4 h precharging at -1200 mV SCE followed by a SSRT at the free corrosion potential, E corr .

The fracture surfaces were analyzed resulting from the test for the condition 1. An estimate of the fracture toughness under hydrogen charging conditions was made.

A fracture toughness test was performed using a precracked compact tension (CT) specimen to measure the threshold stress intensity factor for the hydrogen induced cracking, 𝐾 IH , corresponding to the bainitic microstructure. This was required because it was expected that the local fracture toughness at MnS inclusions was decreased by phosphorous segregation. The fracture toughness of the steel is of the order of 200 MPa.m 1/2 and so requires a CT specimen of dimensions greater than 10 cm. In the presence of hydrogen and phosphorous segregation, however, the fracture toughness is significantly decreased, and consequently requires much smaller specimens.

Results

Progressive heat treatment

The progressive heat treatments from 625 °C to 825 °C were performed in three sequential steps, each followed by SEM fractographic observations. Figure 7 shows micrographs of the same area after each heat treatment. Fig. 7a) presents the initial reference condition, and (b,c,d) correspond to observations after 625 °C, 725 °C and 825 °C heat treatments respectively. This figure shows the evolution of the morphology of the intergranular surface with the progressive smoothing of the grainy-texture typical of this type of fracture. After the heat treatment at 825°C the smoothing is significant but less pronounced than that of the industrial crack from the component presented in Figure 5, especially 5c.

Heat treatment at 950 °C

This heat treatment was aimed to reproduce the effect of the austenitisation heat treatment on the second specimen, which had a transgranular fracture surface. Figure 8 presents the fracture surface before (a) and after (b) this heat treatment at 950 °C for 3 hours. Figure 8a shows the pattern of sharp angularities, which are characteristic of this type of brittle fracture. In contrast, after the heat treatment, the surface was smoothed and not angular anymore. Moreover, the comparison of cleavage planes shows that river patterns had vanished following this heat treatment. The tiny thermal grooves might correspond to boundaries between packets of bainite. They can also indicate the emergence of prior austenite grain boundaries due to the heating above the Ac 1 temperature. Figure 8c) and (d) are magnifications of the initially sharp angle marked with a white circle. A comparison between these two micrographs can be quantitatively expressed by a local radius of curvature after heat treatment. This radius represented by a yellow circle on Figure 8d) was 900 nm, measured by an image processing technique. Figure 9 presents another area of the transgranular fracture surface before (a) and after heat treatment (b). The observed smoothing was similar to that presented in Figure 4b. The observation of morphological evolution of the fracture surface as a function of temperature for the relatively short durations of 3 to 4 hours indicated a progressive smoothing of both intergranular and transgranular fracture surfaces. The highest temperature produced smoothing similar to that on the crack extracted from the industrial component. At this temperature the smoothing was so pronounced that microscopic features like the granny morphology on the individual intergranular facet or river lines on individual transgranular facet tended to disappear due to surface diffusion. These observations indicate that the initially brittle surface of the industrial crack was exposed to the austenitisation temperature, and in particular to the second austenitisation heat treatment. Clearly, the cracks were formed before this austenitisation, during the unplanned slow cooling down to the room temperature or during the subsequent 1,5 months of storage.

Hydrogen influence on tensile properties

The results of the slow tensile tests are presented in Figure 10, Table 2 and Table 3. Figure 10a) presents the stress-strain curves for two reference specimens tested in air at room temperature. Both, exhibited ductile behavior: the specimen with the bainitic microstructure and the specimen with the ferrite-pearlitic microstructure. Figure 10b) presents the effect of hydrogen on the bainitic microstructure. The comparison with the air reference specimen indicates a progressive embrittlement as the severity of the charging condition is increased. Testing conditions 3 and 4 produced only a small effect of hydrogen on the mechanical behavior. Loading with charging conditions 1 and 2 significantly changed the stress-strain curves. There was a significant decrease of strength and reduced plasticity, especially for condition 1. The load decrease was compatible with centripetal crack propagation at the UTS across the round section until specimen failure. Figure 10c) presents the results for the ferrite-perlite microstructure. These stress-strain curves lead to similar conclusions. The intensity of HE increased with the testing severity. The presence of hydrogen mainly decreased plasticity. A comparison of Figure 10b) and Figure 10c) indicates that the hydrogen sensitivity of the bainitic microstructure was significantly higher. Figure 11 presents the fracture surface corresponding to the specimen with the bainitic microstructure and tested under condition 1. Figure 11a) presents an overview of the fracture surface. The fracture surface indicates crack initiation at the specimen surface at the bottom of the figure, and crack propagation towards the center of the specimen, followed by final fracture. The reduction of area and almost no necking indicate a macroscopically-brittle fracture of the specimen. The fact that there was necking indicated that the hydrogen induced cracking occurred at or after the ultimate tensile strength. This behavior explains the stress-strain curve from Figure 10b). Detailed observation revealed an abundance of MnS clusters as indicated in Figure 11b). This micrograph also indicates the presence of some ductility, which confirms the assumption of a curve governed by the centripetal crack propagation. This indicates that the brittle fracture occurred first, inducing a stress drop in this strain-controlled SSRT. This effect was followed by a subsequent failure of ductile ligaments. The brittle areas were mostly transgranular but Figure 11c) shows some intergranular regions located nearby MnS. This fact is emphasized in Figure 11d). Table 2 and Table 3 present the parameters extracted from the stress-strain curves, respectively for the bainitic and ferrite-pearlite microstructures: the ultimate tensile stress (UTS), elongation to fracture, 𝜖 𝐴 , and cross-section reduction, 𝜖 𝑁 . For the bainitic microstructure, additionally, hydrogen content measurements by catharometry were carried out on 𝜙 6 mm cylindrical specimens for the different charging conditions. The charging conditions 3 and 4 did not produce significant HE, and the hydrogen content was similar to that of the reference specimen. In contrast, there was significant embrittlement resulting from testing conditions 1 and 2. For the bainitic state, there was a decrease of UTS by 13% in the specimen tested under the condition 1.

The elongation to fracture was decreased by 3.5 times and the section reduced by a factor of 5.4. Under the condition 2, the UTS was reduced by 4%, the elongation to failure by 11% and the section by 44%. For the ferrite-pearlite microstructure, these quantities were reduced by (16%, 2.4 times, 2.9 times) and (4%, 15%, 38%) for conditions 1 and 2 respectively.

Estimation of fracture toughness from the tensile test

Figure 11-a) allows a coarse preliminary estimation of fracture toughness for the bainitic microstructure. The stress intensity factor, for a cylindrically-notched cylindrical tensile specimen, is given by ( 1) and ( 2) [START_REF] Tada | The Stress Analysis of Cracks Handbook[END_REF]. where 𝑟 is the inner radius (at the notch) and 𝑅 is the outer radius of the specimen. Figure 11a) shows two concentric areas. Assuming that the failure occurred in two steps, the peripheral surface crack can be considered as a pre-crack and the inner section as equivalent to the failure surface in a notched toughness test. The peak load, the average outer radius and the average crack length were measured as 4.58 kN, 1.33 mm and 467 µm respectively. The resulting fracture toughness under cathodic charging condition was estimated to be 19 MPa.m 1/2 . This coarsely approximated value can be used to determine the expected order of magnitude of the reduced toughness for the dimensioning of the CT specimen for the measurement of fracture toughness.

Estimation of a local fracture toughness in presence of hydrogen

This experiment targets an evaluation of the local fracture toughness in the presence of hydrogen, in order to quantify the proposed scenario. It is important to underline the word "local" and to recall the complexity of microstructure in such heavy forgings. In fact, the microstructure of the studied materiel is heterogeneous and contains several zones of molybdenum, manganese, carbon and phosphorous over-concentrations in areas named ghost lines and microsegregations. Consequently, typical ASTM standards, like ASTM E399 and E1820 could not be used in order to assess the local value of fracture toughness in the presence of hydrogen, local meaning here : within a millimeter range. Following our first estimation of 19 MPa.m 1/2 , see above, Irwin's model indicates that the plastic zone radius as given by equation ( 3) is 104 µm.

𝒓 𝐘 = 𝟏 𝟐𝛑 ( 𝑲 𝐈𝐇 𝝈 𝐘 ) 𝟐 (3)
Table 4 provides the required specimen size, according to the relations presented in ASTM E647 [START_REF]ASTM E647 -15e1 Standard Test M ethod for M easurement of Fatigue Crack Growth Rates[END_REF]. It is worthy of mentioning that these dimensions satisfy the small-scale yielding (SSY) hypothesis in case of a fracture toughness less than 19 MPa.m 1/2 . For information purposes, the nominal toughness for this steel, typically of the order of 200 MPa.m 1/2 , leads to dimensions of a few dozens of centimeters and extremely high tensile loads. For higher toughness values, a parameter based on the J-integral is more adapted to characterize the elastoplastic material. In the fracture toughness measurement, for repeatability and post processing purposes, the pre-crack was stepwise generated by fatigue as specified in [START_REF] Francois | Essais de mesure de la ténacité -M écanique de la rupture[END_REF] with a sequence of 6 cyclic loading conditions in the 7.96 kN to 2.84 kN range to ensure a progressive decrease of the plastic zone size. The notched CT specimen, of the size given in Table 4, was machined so that crack propagation plane was parallel to the plane of crack formation in the component. Hydrogen was electrochemically charged for 96 h at -1200 mV SCE . Thereafter, hydrogen charging was maintained and progressive deformation was applied to this CT specimen up to the final fracture. The displacement of the testing machine was controlled to target a rate of 2 MPa.m 1/2 .s -1 jointly with continuous load measurement.

Figure 12 presents different aspects of the resulting fracture surface. Figure 12a) presents a global view. At this magnification, a progressive transition from ductile to brittle failure suggests a heterogeneous effect of hydrogen. This progression has been schematized in Figure 12b) which identifies several areas. Near the pre-crack in zone 0, the failure was mainly ductile with some embrittled areas located near MnS inclusions. The fracture path changed to a brittle mode in zones 2-4. These zones are characterized by the continuous reduction of ductility (e.g. ligaments) with increasing crack length. Figure 12c) presents a micrograph focusing on an embrittled area located in the vicinity of the MnS-inclusion in zone 2. Figure 12d) presents a micrograph focused on a representative area of zone 4. This zone was fully brittle transgranular with no trace of ductility. Figure 13a) presents the loading curve for this measurement. The overall shape of the curve indicates load increase up to a maximum of 23 kN followed by the progressive load decrease down to approximately 17 kN, whereupon there was a steep load drop. The critical event happened during the load increase. Figure 13b) presents a magnification of the ascending part of the loading curve. There was a slight load drop during the test, indicating the damage of the material within the elastic region of the curve. The amplitude of the drop was significantly greater than the acquisition noise. In addition, the load recovery led to a load curve in continuity with the initial slope. This damage, that had occurred at 4,3 kN, was linked to the highly heterogeneous microstructure of this steel, and was attributed to micro-fracture of such a heterogeneous area, situated within the ductile region in zone 1. In fact, it is worth emphasising that the 18MnNiMo5 steel is a highly heterogeneous bainitic steel containing ghost lines and microsegregated zones, since it is obtained from a very large ingot. While the nominal toughness of this alloy is of the order of 200 MPa.m 1/2 , individual values well below 100 MPa.m 1/2 are measured. They reflect intrinsic material heterogeneities, namely microsegregations, which are of submillimeter size, therefore much smaller than the CT fracture surface. Under hydrogen charging, much lower values of fracture toughness are expected due to the hydrogen embrittlement phenomenon.

These results, particularly the load drop at 4,3 kN, allow evaluation of fracture toughness based on the CT related expression of stress intensity factor as defined by ( 4) and ( 5) [START_REF] Francois | Essais de mesure de la ténacité -M écanique de la rupture[END_REF].

where 𝑎 is the total crack length, 𝑤 is the maximal crack length to split the CT specimen into two parts, and each of 𝐵 and 𝐵 𝑁 , which were equal in this case, is the specimen thickness.

The fracture toughness corresponding to the peak load of 23 kN was 95 MPa.m 1/2 but this value was invalid since the plain strain conditions were not satisfied at this load for this specimen. This measurement cannot be interpreted.

The fracture toughness evaluated from the load drop at 4.3 kN was 18 MPa.m 1/2 . This measurement provides an estimation of the reduced local fracture toughness since the plain strain conditions are strictly respected for this load. This value of 18 MPa.m 1/2 reflects hydrogen effect on the local fracture toughness of a microsegregated area, and has to be considered as an estimate because significant variations can be anticipated depending on the actual position of a microsegregated areas with respect to the crack tip position. The fracture toughness test showed a heterogeneous effect of hydrogen inducing a progressively brittle fracture mode. The HE heterogeneity can be linked to the hydrogen charging conditions. The hydrogen charging setup used the specimen and both loading jaws as the working electrode. The jaws were protected by a silicone based waterproof coating, which obstructed the front of the specimen and avoided hydrogen charging. Only the back part of the specimen could interact with the counter electrode, so that the specimen was only hydrogen charged from the zones 4-5. The hydrogen charging was performed at a cathodic potential using a deoxygenated saline solution. Such conditions limit the effects of surface and promote adsorption and absorption phenomena.

From the viewpoint of a simplified diffusion problem, the concentration of hydrogen decreases linearly from the charging area (zones 4-5) to the crack tip. Since the electrochemical interactions at the crack tip region were limited, the hydrogen concentration at the crack-tip was assumed to be nearly equal to the as-received material (around 0.72 wt. ppm, see Table 2) whereas the estimated concentration in zone 5 was close to 1.6 wt. ppm, a value based on bulk H content measurements on independent specimen under the same charging conditions.

Hydrogen trapping models suggest an equilibrium driven process, each trap being defined by an activation energy. The site occupancy rate is closely linked to hydrogen activity. The non-zero value of the hydrogen concentration for the as-received reference specimens was measured by catharometry, and was attributed to a significant amount of hydrogen in deep traps (nondiffusible hydrogen). The hydrogen embrittlement induced by the fracture toughness testing conditions was fairly severe, and was attributed to the lowering of cohesion forces at interfaces by the locally-trapped hydrogen.

Discussion

Scenario of crack formation

The starting point of this investigation was the presence of penny-shaped internal cracks, 10-20 mm in diameter, that were detected by ultrasonic measurements on a 18MND5 component after the final step of the thermo-mechanical manufacturing cycle shown in Fig. 1.

One such crack was extracted and analyzed. The internal crack surfaces were smoothed by temperature activated surface diffusion as demonstrated by the series of UHV heat treatments between 625 °C and 950 °C. The morphology obtained after heating at 950 °C was comparable to the morphology of the defect extracted from the industria l component, and indicated that the cracks in the component were formed during the slow cooling following the first austenisitation treatment, and were then smoothed during the second austenitisation. A detailed morphological analysis indicated that the initial internal surface of the cracks was mainly transgranular with some intergranular areas. Intergranular areas containing MnS clusters were associated with crack initiation, while transgranular areas were associated with crack propagation, as discussed below.

Crack initiation

It is proposed that crack initiation was provoked by a combination of the following four effects: (i) local phosphorous segregation facilitating intergranular micro-fracture, (ii) thermal contraction mismatch of the steel and the MnS inclusion during cooling causing separation of the interface between the steel matrix and the MnS inclusion, (iii) the distribution of MnS inclusions in clusters of millimeter size, and (iv) hydrogen desorption at the void at the inte rface between the steel matrix and the MnS inclusion, and the subsequent hydrogen pressure increase Firstly, the local phosphorous segregation is compatible with the heterogeneous microstructure of the 200 mm thick heavy component, where micro-segregations and ghost lines are known to contain an increased amount of phosphorous. The phosphorous is expected to have segregated during the slow cooling, supposed to have been of the order of 10 °C/h within the critical range of temperatures between 550 °C and 350 °C, where thermodynamics and kinetics indicate phosphorous segregation. The resulting phosphorous embrittlement together with increasing internal stresses due to the cooling to room temperature would result in the phosphorous embrittlement of grain boundaries, especially on ghost lines.

Second, the MnS/matrix interface opens during the cooling due to the difference of thermal expansion coefficient, 𝛽, between 12.5 x 10 -6 K -1 for 18MND5 and 16.3 x 10 -6 K -1 for 𝛼-MnS with a NaCl type lattice. Considering a typical MnS inclusion with a radius 𝑟 𝑀𝑛𝑆 of 15 µm, the resulting opening 𝛿 of the metal-MnS interface after a cooling from 620 °C to 20 °C (Δ𝑇 = -600 °C) is 17 nm, calculated using the expression in equation ( 6). The MnS/matrix interface is therefore an important factor contributing to the crack initiation.

𝜹 = ( 𝜷 𝜶𝟏𝟖𝑴𝑵𝑫𝟓 -𝜷 𝜶𝑴𝒏𝑺 ) 𝚫𝑻 𝒓 𝑴𝒏𝑺

Third, the fact that these inclusions are in clusters of a millimetric size induces a significant number of such initiation sites and allows the formation of a hydrogen crack of significant initial size.

The proposed scenario is that (i) for MnS/matrix interfaces located on phosphorous enriched prior austenite grain boundaries, continuous hydrogen desorption to MnS/matrix interfaces at room temperatures results in the progressive stress increase at such interfaces, then (ii) these individual decohesions at MnS/matrix interfaces expand through the region of phosphorous embrittled grain boundaries, and (iii) finally such a crack reaches the external diameter of the ghost lines. The crack initiation is accomplished when such a penny-shaped crack leaves the ghost line and enters the matrix, i.e. the area without MnS clusters, and without such a high hydrogen content as measured in Table 1, and without the phosphorous segregation. The crack initiation is mainly intergranular, and results in a sub-millimeter penny-shaped cavity with the initial thickness given by the range of thermal mismatch between MnS and matrix, i.e. of the order of 20 nm. The kinetics of pressure build-up in such a crack, necessary for crack propagation across the matrix, was quantitatively assessed in our two previous papers [3] [4].

In a search of plausible orders of magnitude for the critical fracture toughness, it is instruct ive to consider the most severe, but still realistic case of such a sub-millimeter initiation site (𝑎 = 0.5 mm). For a penny-shaped pressurized crack, the stress intensity factor is given by the expression in equation ( 7) [START_REF] Urakami | Stress Intensity Factors Handbook[END_REF]. The fracture toughness of "healthy" 18MND5 steel is approximately 200 MPa.m 1/2 , so that the resulting critical pressure required to initiate fracture would be as high as 7 930 MPa.

𝑲 𝑰 = 𝟐𝒑 𝝅 √𝝅𝒂 (7)
The fracture toughness of hydrogen embrittled specimen, as estimated in the present paper by two independent methods is below 20 MPa.m 1/2 , so that the resulting critical pressure required to initiate fracture would drop by a factor of ten, down to 793 MPa. Our previous assessments indicated a maximum pressure of 860 MPa for a hydrogen concentration of 1 wt. ppm [START_REF] Sezgin | M odelling of hydrogen induced pressurization of internal cavities[END_REF]. This indicates that the internal hydrogen pressure can cause the penny shaped cracks at the MnS/matrix interface to grow through the embrittled region.

Crack propagation

Crack propagation is caused by (i) continuous hydrogen desorption to these cracks and the subsequent hydrogen pressure increase, and (ii) the decrease of fracture toughness due to the presence of hydrogen within the phosphorous enriched material at the crack tip, and (iii) discontinuous crack propagation due to these two elements. The crack initiation is considered to have occurred during the cooling to temperatures slightly above room temperature (which caused the hydrogen pressure build-up). Crack propagation is considered to occur at room temperature where hydrogen solubility is minimal, and as a consequence, the resulting hydrogen pressure can reach the highest value. First, hydrogen desorption results in hydrogen pressure build-up within a period of the order of some hours as demonstrated in our recent paper where hydrogen gas behavior was assessed based on recent NIST data and has led to the equilibrium pressure of 860 MPa [START_REF] Sezgin | M odelling of hydrogen induced pressurization of internal cavities[END_REF]. Second, the loss of mechanical properties was expected with the reduction of fracture toughness down to 18 MPa.m 1/2 under hydrogen charging conditions. This represents a local fracture toughness in the presence of hydrogen, K 1H . Finally, the combination of such a high hydrogen pressure and locally low fracture toughness results in crack propagation, given the millimeter size of the initial crack. A penny-shaped crack, 1 mm in diameter, together with material fracture toughness of 18 MPa.m 1/2 leads to a critical pressure of 717.3 MPa according to our previous calculations [START_REF] Sezgin | M odelling of hydrogen induced pressurization of internal cavities[END_REF]. As a consequence, the crack propagates in a brittle manner over a distance corresponding to a characteristic microstructural feature, that was arbitrary taken as a size of a bainitic packet. Such a brittle propagation is expected to be fast. It results in a volume increase and pressure decrease which in turn requires more hydrogen to diffuse again to the crack with a new step of a progressive pressure built up, up to a new, slightly lower critical pressure. Such a mechanism implies a discontinuous crack propagation.

In order to gain an order of magnitude of crack propagation rate, a simplified analysis based on the assumption of a constant crack velocity can be performed as follows. In the case of an industrial penny-shaped crack, starting from a millimeter size initiation site, a centimeter size is obtained during approximately one month, which corresponds to an average crack growth rate of 3.7 x 10 -9 m.s -1 . The differentiation of the stress intensity factor in equation ( 7) provides an expression for the average experimental velocity in equation ( 8) that should be targeted (this value being out of range regarding norms).

𝑲 𝑰 ̇= 𝒑 √𝝅𝒂 𝒂̇ (8)
An internal pressure of 800 MPa leads to a velocity of 9.5 x 10 -5 MPa.m 1/2 .s -1 . An increased effect of HE is then expected in this case of an industrial crack and may lead to an effective value of toughness slightly below 18 MPa.m 1/2 .

A viable path forward to quantify these aspects is a predictive modeling of this industrial problem with a coupled diffusion-mechanical-metallurgical model. The diffusion procedure aims at quantifying the redistribution of hydrogen between the alloy and the hydrogen induced crack. The mechanical procedure targets the kinetic aspects of fracture such as crack length, crack growth rate or failure criterion. Finally, the metallurgical model renders various properties of the alloy such as 𝛼 → 𝛾 transformation and decrease of fracture toughness, induced by hydrogen and phosphorous segregation. This overall scenario of crack propagation can be assessed by a quantitative analysis of crack propagation rate, which is based on above mentioned elements with a strong dependence on fracture toughness, hydrogen solubility and diffusion [START_REF] Sezgin | A model of internal crack extension due to a continuous build-up of hydrogen pressure: application to a pressure vessel component[END_REF].

Possible extension to blister cracking

Premature and unexpected failure of a component can be promoted by hydrogen in solid solution in alloys, which may have various origins. A distinction is made between internal (i.e. initially present) [START_REF] Lynch | Hydrogen embrittlement phenomena and mechanisms[END_REF] [35] and external hydrogen, which requires a contribution from the environment, mainly from electrochemical reactions occurring in operation [START_REF] Kane | Roles of H 2 S in the behavior of engineering alloys: a review of literature and experience[END_REF] [37] [38] [39] .

The overall scenario of embrittlement, together with underlying mechanisms, i.e. (i) segregation of embrittling elements, (ii) hydrogen desorption first to the interfaces with inclusions, and then to the as-initiated crack, and (iii) transgranular crack propagation under the action of hydrogen pressure within the material, itself embrittled by the presence of hydrogen and phosphorous, can be applied to several similar situations, where hydrogen initially comes from external charging, but eventually acts as internal hydrogen in a similar manner to that described in the present paper.

Hydrogen induced cracks (usually shorten to HIC, also known as "blisters" in the Oil and Gas industry [START_REF] Dunne | Investigation of the effect of electrolytic hydrogen charging of X70 steel : I. The effect of microstructure on hydrogen-induced cold cracking and blistering[END_REF]) are a typical example of such cracks, and could be analyzed according to the same scenario of propagation, the only difference being the origin of hydrogen. The subsequent mechanism might be similar to mechanism proposed by [START_REF] Zapffe | Hydrogen embrittlement, internal stress and defects in steel[END_REF] where HIC forms in the presence of H 2 S (external hydrogen source) in sour environment conditions, where the hydrogen fugacity and consequently pressure is known to be high.

Similar cracks would be expected for corrosion in acidified environments, for example in the solutions associated with pits and crevices, as significant hydrogen fugacities have been shown to occur for steels in acid solutions [START_REF] Darken | Behavior of hydrogen in steel during and after immersion in acid[END_REF].

Despite these similarities, additional factors, not directly considered in the present study, could have an impact on cracks caused by hydrogen as raised in [START_REF] Brainin | Shatter cracks and the effect of heat treatment conditions on the elimination of hydrogen from steel[END_REF], [START_REF] Scott | Hydrogen and segregates in flaking[END_REF], [START_REF] Houdremont | Die Entstehungsbedingungen der Flocken im Stahl[END_REF]. In these references, several parameters such as stress resulting from thermal expansion, phase transformation, metallurgical parameters (e.g. grain size, impurities), segregation were investigated bringing some nuance to the analogy between the blisters and the HIC observed in large components.

Hydrogen induced localised damage (HILD)

The approach developed herein can be generalized to other cases of hydrogen induced cracking. The general concept is that there is brittle crack extension for a limited distance in front of a crack tip, if the crack tip stress intensity factor at the crack tip is greater than the locallyembrittled fracture toughness ahead of the crack tip. This process can repeat if the new crack tip region can subsequently become embrittled. This mechanism of hydrogen induced cracking (HIC) can be designated as Hydrogen Induced Localised Damage (HILD). The present work has dealt with embrittlement by hydrogen. The approach is in principle valid, independent of the species causing embrittlement. In this case the mechanism can be designated as Environment Induced Localised Damage (EILD). Furthermore, procedures have been developed herein by which the locally embrittled fracture toughness can be measured.

There are two possible rate-limiting factors on the average crack velocity. One is the distance of brittle crack extension. The brittle crack might extend only through the embrittled region and may be stopped when the crack reaches un-embrittled material. Although the whole steel is hydrogen embrittled, the level of embrittlement is not uniform. In addition, the fast brittle crack propagation is prevented by the slow desorption of hydrogen to put the crack void under high pressure.

It is also worth pointing out, that this mechanism does NOT directly involve any of the current models of hydrogen embrittlement, such as hydrogen enhanced localized plasticity (HELP), hydrogen enhanced decohesion (HEDE) or adsorption induced dislocation emission (AIDE). Instead, the critical event is the brittle fracture through a local region ahead of the crack tip, the region having been embrittled.

Lessons for industrial production

It is worth underlining that the cracks analyzed in this paper formed as a consequence of the combination of (i) a high hydrogen content close to the upper limit of 1 ppm and (ii) the unplanned slow furnace cooling and subsequent long period of storage at room temperature before the second austenitisation. High phosphorous concentrations at former austenite grain boundaries and the presence of large MnS clusters, both mainly within ghost lines, allowed crack initiation by a hydrogen mechanism. Room temperature hydrogen desorption is considered as a main cause of subsequent crack propagation and eventually penny-shaped cracks formation up to 20 mm diameter.

It is however important to underline that the appropriate process control measures allow for the complete elimination of such defects from large forged components.

Conclusions

This study analyzed a crack from the scrapped 18MND5 80-tonne component in order (i) to identify the critical step in crack formation, (ii) to propose a crack initiation and propagation scenario, and (iii) to understand the underlying physical mechanisms. The conclusions are as follows:

1. The smoothed morphology of the internal surface of the crack withdrawn from a 18MND5 shell, together with a series of purpose designed ultra-high vacuum heat treatments on mixed intergranular/transgranular fracture surfaces, indicated that the crack was formed during the unexpected slow cooling in furnace after the first autenitisation heat treatment. 2. The subsequent second austenitisation transformed the brittle fracture surface into a smooth surface strongly affected by surface diffusion. 3. The root cause of crack formation was the low fracture toughness caused by a combination of a high hydrogen content and phosphorous segregation in micro-segregated areas, where phosphorous segregation can be twice as high as in the matrix. 4. Cracks initiated on MnS inclusion clusters of millimeter size, that were located within ghost lines, where there was phosphorous segregation to prior austenite grain boundaries during exceptionally slow cooling. Hydrogen desorption at MnS/matrix interfaces would result in penny-shaped, approximately 1 mm in diameter and thin cracks. 5. Mechanical tests under hydrogen charging conditions demonstrate a significant decrease in fracture toughness, here down to 18 MPa.m 1/2 . This absolute value reflects a local hydrogen effect on a microsegregated area and has to be considered as an order of magnitude because significant variations can be anticipated depending on the actual position of a microsegregated areas with respect to the crack tip position. 6. The fracture toughness value of the order of 18 MPa.m 1/2 together with previously calculated hydrogen pressure, of the order of 860 MPa, due to hydrogen desorption and presence of penny-shaped 1 mm diameter cracks, constitute the condition for transgranular crack propagation to a final dimension of the order of 10 mm in diameter. 7. The cracks in the component were caused by the inappropriate heat treatment (in particular the slow cooling to room temperature) and the presence of an unusually high hydrogen content at temperatures below 200 °C. 8. The proposed hydrogen induced cracking (HIC) mechanism can be designated as Hydrogen Induced Localised Damage (HILD) or more generally Environment Induced Localised Damage (EILD). 9. The appropriate process control measures allow for the complete elimination of such defects from large forged components.

Declaration of Competing Interest

None.

Figure 1

 1 Figure 1 Heat treatment applied to the component. To comply with the microstructure requirements, a second austenitisati on was performed after the unpl anned furnace cooling. Nondestructi ve examination found pennyshaped cracks and the component was scrapped.

Figure 2 Figure 3

 23 Figure 2 Observation of a crack withdrawn from the scraped component (a) by X-ray micro tomography and (b) by optical microscopy on the cross section

Figure 4 S

 4 Figure 4 S EM micrograph of the fracture surface of the crack extracted from the component. a) Possible crack initiation site is indicated by the white dashed ellipse. b) The generalized smoothed crack topography indicates that the crack was exposed to high temperatures after crack formation

Figure 5

 5 Figure 5 Evi dence for large intergranular areas containing MnS inclusions. a) S EM micrograph of a fracture surface of the crack extracted from the component showing a large area of mainly intergranular fracture which was smoothed. b) The same area viewed by backscattered electrons showi ng several MnS inclusions (grey areas at the bottom) close to the intergranular facets. c) Smoothed intergranular facets.

Figure 6

 6 Figure 6 Flow diagram of the de-oxygenation of saline solution for cathodic charging during SSRT.

Figure 7 S

 7 Figure 7 S EM observations of the morphological evolution of intergranul ar facets following progressive heati ng under UHV conditi ons. Four successive steps are represented: (a) initi al reference condi tion, (b) 625 °C for 4 h, (c) 725 °C for 4 h, (d), 825 °C for 4 h.

Figure 8 S

 8 Figure 8 S EM micrographs of a transgranul ar facet before (a,c) and after (b,d) a 950 °C heat treatment under UHV conditi ons for 3 h. a) Transgranular fracture. b) Smoothed facet after exposure to 950 °C. c) Focus on the sharp angularity i ndicated in (a) before heat treatment. d) Focus of the sharp ang ularity indicated in (b) after heat treatment. The heat treatment has introduced a radius of curvature of 900 nm, which is represented on the micrograph by the yellow circle.

Figure 9 S

 9 Figure 9 S EM micrographs of a transgranul ar facet before (a) and after (b) a 950 °C heat treatment in UHV condi tion for 3 h. a) Sharp crack ti p indicated by an ellipse. b) Creation of bri dges linking both si des of the crack tip especially in the thinnest areas

Figure 10

 10 Figure 10 Tensile stress-strain curves for: (a) both bainite and ferrite-perlite microstructures in air at 10 -6 s -1 , (b) bainitic microstructure under different cathodic charging conditions at 5,0.10 -7 s -1 , and (c) ferrite-perlite microstructure under different cathodic charging conditions at 5,0.10 -7 s -1

Figure 11 Fracture

 11 Figure 11 Fracture observation of the bai nitic tensile specimen under the first hydrogen loading condi tion a) global view of the fracture surface, b) presence of elongated MnS inclusions enhancing HE contained in the white elliptic marks c) mixed transgranul ar-intergranular brittle fracture nearby MnS inclusions (dark smooth surfaces) with some ductile areas, and d) magnification of an intergranul ar fracture area, located close to MnS inclusions.

Figure 12 Fracture

 12 Figure 12 Fracture surface obtai ned after a fracture toughness test under hydrogen charging conditi ons at -1200 mV SCE in 30 g L -1 NaCl for 96 h. a) Overview S EM macrograph of the fracture surface. b) Schematic of the different areas on this fracture surface. c) Focus on an elli ptic brittle area surroundi ng several MnS inclusions in the zone 2. d) Fully brittle fracture in zone 4.

Figure 13

 13 Figure 13 Loadi ng curves for the toughness test using a pre-cracked CT s pecimen under hydrogen charging condi tions. a) Overall load-ti me curve. b) Focus on the load drop at around 4.3 kN. This load drop is attri buted to the embrittled areas near MnS inclusion in Zone 1 and 2, and can therefore provi de a quantitati ve value for the fracture toughness in the presence of hydrogen of a microsegregated area.

Table 1 -Composition of the 18MND5 low alloy steel, in wt. %

 1

		C	Mn	Ni	Mo	Cr	Si	Cu	Al	V	Sn	S	P
	18MND5	0.181	1,458	0,724	0,501	0,254	0,173	0,052	0,015	0,005	0,004	0,0002 0,0055

Table 2 Mechanical property values (UTS , elongation to fracture , cross-section reducti on) of speci mens wi th a bainitic microstructure for different charging conditions , and measured average content of hydrogen.

 2

	Condition Precharging Reference --1 -1200 mV SCE , 24 h	E Test --Cathodic -1200 mV SCE	UTS 742 MPa 642 MPa	𝝐 𝐀 21,8 % 6,2 %	𝝐 𝐍 69,3% 12,9 %	𝒄 𝐚𝐯𝐠 (𝑯) 0,72 ppm ±0,07 1,32 ppm ±0,32
	2	-1200 mV SCE ,24 h	E corr	715 MPa	19,3 %	39,1 %	0,91 ppm ±0,08

Table 3 Mechanical property values (UTS , elongation to fracture , cross-section reducti on) of speci mens wi th a ferrite-perlite microstructure for different charging conditions

 3

	Condition Reference	Precharging --	E Test --	UTS 659 MPa	𝝐 𝐍 17 %	𝝐 𝐙 52,2 %
	1	-1200 mV/SCE,24 h	Cathodic -1200 mV/SCE	551 MPa	7 %	18,3 %
	2	-1200 mV/SCE,24 h	E corr	633 MPa	14,5 %	32,5 %
	3	-1200 mV/SCE,8 h	E corr	628 MPa	17,4 %	38,0 %

Table 4 -Dimensions of CT specimen for a hydrogen embrittled steel

 4

		H	B	w	a	w-a	w 1	D	H 1
	Dimension (mm)	15	12,5	25	12,5	12,5	31,25	6,25	8,125

Acknowledgements

This work was done at Mines Saint Etienne (MSE, France) in the frame of Jean-Gabriel SEZGIN's PhD degree between 2014 and 2017. The author acknowledges the colleagues from Mines Saint Etienne, Areva (now: Framatome), the University of Queensland (Australia) and National Institute of Advanced Industrial Science and Technology (AIST, Japan) for their continuous support and kind collaboration on this publication.

Highlights

 An experimental analysis of industrial HIC in heavy components was conducted  SEM imagery shows a smoothed mixt intergranular and transgranular failure surface  In situ SSRT & toughness tests revealed a strong Hydrogen Embrittlement of steel  MnS located near ghostlines (enriched in P) act as millimetric initiation sites

ACCEPTED MANUSCRIPT

 A HIC mechanism accounting for HE, P segregation and H precipitation was proposed


Hamidou Dembélé
email: hamidou.dembele@orange.com

Marie Le Bot
email: marie.lebot@orange.com

Franc ¸ois Gallée
email: francois.gallee@imt-atlantique.fr

Patrice Pajusco
email: patrice.pajusco@imt-atlantique.fr

Multibeam Antenna Modelling and Blockage Impact Mitigation in 5G Millimeter-Wave Bands

Keywords: Multibeam antenna, transmitarray, reflectarray, human blockage, millimeter-wave, 60 GHz, 5G, block error rate

This paper discusses a generic approach proposal for modelling multibeam antenna systems for performance prediction characterized by scan loss and half-power beamwidth. The validation with the existing transmitarrays in the Ka and V bands is carried out. Using the proposed antenna model, the advantage of multibeam antenna to alleviate the impact of human blockage in millimeter-wave bands is highlighted. Communication performance is evaluated in terms of block error rate through the implementation of the 5G physical communication chain in a simulation tool.

I. INTRODUCTION

Millimeter-wave (mmWave) bands constitute real opportunities for 5G enhanced mobile broadband (eMBB) applications because of the huge amount of spectrum available in this region of frequencies. Due to the severe propagation loss of electromagnetic waves in mmWave bands, high-gain antenna systems with directional beams are widely deployed to improve the signal-to-noise ratio (SNR) at the receiver (Rx). However, concentrating all energy towards the main propagation path through directional communication eliminates the diversity offered by the multipath and represents a real risk of sudden interruption of communication in the event of obstacles (human or vehicle) in the beam pointing direction [START_REF] Dogan | Optimization of Antenna Beamwidth under Blockage Impact in Millimeter-Wave Bands[END_REF]. Such a scenario is known as human (or vehicle) blockage phenomenon in mmWaves. To traditionally maintain the communication link on-the-move, the directional beam needs to be mechanically or electronically steered. Mechanical beam steering consists in manually turning the antenna to face the direction of interest. This technique is effective since the antenna gain is maintained and there is a flexibility in the steering range [START_REF] Uchendu | Survey of Beam Steering Techniques Available for Millimeter Wave Applications[END_REF], but it is only amenable to static or very slow changing propagation environments. In this way, electronic steering (or beamforming) antennas are proposed for allowing fast beam pointing. Nonetheless, the feed network and electronic constraints are complex and costly in mmWave bands [START_REF] Smolders | Measurement and calibration challenges of microwave and millimeter-wave phased-arrays[END_REF] - [START_REF] Prabu | Millimeter wave for 5G mobile communication application[END_REF]. In addition, a large number of antennas is required to enhance the system directivity.

To cope with these shortcomings of both electronic and mechanical beam steering, the multibeam antenna (MBA) approach received much attention due to some features such as low profile (small height and width), light weight, cost effective and easy to manufacture [START_REF] Wu | Wide-Angle Beam-Scanning Reflectarray With Mechanical Steering[END_REF]. MBAs generally represent an antenna system capable of simultaneously generating a certain number of concurrent and independent directive beams with a high gain to cover a predefined angular range.

To realize an MBA system, several typical approaches have been employed. One approach is to use multiple feed antennas placed at different positions in front of a reflectarray (RA). Another technique is to illuminate one side of a lens or transmitarray (TA) by the incident waves emitted by the feed antennas, while the other side is used to focus the beams in the desired directions. In contrast to RA and TA techniques based on optic principle, MBA design can also be achieved using beamforming circuits such as Butler matrix that can be integrated with an array of antennas into a single substrate [START_REF] Hong | Multibeam Antenna Technologies for 5G Wireless Communications[END_REF]. Whatever the design choice, the MBA can be considered as a good compromise in the management of user mobility and especially the phenomenon of blockage in the mmWave context. At first sight, a large beam can simply be used when the main path is blocked by an obstacle, in order to capture energy from the other multipath components; but in this case, a large amount of transmitted power is lost, and the cell coverage is ineffective. Designing an MBA system that simultaneously points in well-chosen directions can solve this issue: if a path is blocked, others are maintained with an efficient distribution of the power.

In this paper, a generic model is proposed for the modelling of the MBA system with the following contributions. Indeed, the proposed model is suitable for the purpose of predicting performance in the design and the hardware realization of an MBA. It is more flexible, complete and realistic for mmWave systems than the antenna models suggested in the standards and others (3GPP [START_REF] Gpp | Technical Specification Group Radio Access Network; Study on channel model for frequencies from 0.5 to 100 GHz[END_REF], ITU-R [START_REF] Iut-R | Guidelines for evaluation of radio interface technologies for IMT-Advanced[END_REF], DuFort [START_REF] Dufort | Optimum networks for simultaneous multiple beam antennas[END_REF]) for which the gain reduction for a certain scan range is tolerated. This model is then integrated in a simulation chain of the 5G physical layer, including multipath propagation channels and human blockage. Simulations are run to show the advantage of the MBA.

The rest of the paper is organized as follows. In section II, the proposed model is described with the input parameters as well as the assumptions. Section III is focused on the application of the model and its validation with existing MBA TAs. Section IV is devoted to the link level evaluation in the context of human blockage. Conclusion and future work are given in section V. II. MULTIBEAM ANTENNA MODEL PROPOSAL Typically, MBA systems are made up of a finite number of input ports, each connected to an antenna called feed antenna. The excitation of each port allows to control a single beam pointing at a predefined direction [START_REF] Hong | Multibeam Antenna Technologies for 5G Wireless Communications[END_REF]. Multiple beams pointing in different directions can be transmitted simultaneously from a shared aperture to cover a certain angular range when several ports are excited. In general, the MBA performance is characterized by the scan range which leads to the reduction in antenna gain called scan loss (SL). Other parameters such as aperture efficiency, sidelobe level (SLL), beamwidth can also be used.

A. Study Assumptions and General Expression

In the development of the MBA model, some assumptions were taken into account. First of all, based on the RA and TA structures, we generally consider the MBA system as a two-dimensional planar aperture antenna illustrated in Fig. 1. We assume that the radiation in both perpendicular plans xoy and zox are not dependent, and the waves on the antenna aperture are in phase due to the phase correction performed by the unit-cells (transmitters) of the RAs and TAs. In this case, only the amplitude distribution of the field is considered, from which the radiation pattern at the antenna broadside is defined. Assuming that each of the feed antennas illuminates the center of the radiating surface, the beam patterns in other directions are estimated by the surface projection. However, the model does not depend on the number of unit-cells, but rather on the dimensions of the aperture. And the increase in the level of sidelobes when the scan range becomes wide is not also taken into account. Consequently, the general expression of the antenna gain pattern in the desired pointing direction (θ 0 , φ 0 ) around the antenna broadside is given by,

G (θ, φ) = G 0 SL (θ 0 , φ 0 ) g (θ -θ 0 , φ -φ 0 ) , (1)
where θ and φ represent the angles of elevation and azimuth, respectively. G 0 is the antenna gain at broadside, SL (θ 0 , φ 0 ) defines the scan loss value in the beam direction (θ 0 , φ 0 ) and g (θ -θ 0 , φ -φ 0 ) is the normalized gain radiation pattern.

The scan loss SL is referenced to the antenna broadside.

B. Antenna Gain at broadside

To reduce the complexity of the calculation of the antenna directivity, several approximate expressions are proposed in the literature, among which the Kraus formula given by [START_REF] Tai | An approximate formula for calculating the directivity of an antenna[END_REF],

D 0 ≈ 41253 θ HP BW0 • φ HP BW0 , (2)
is suitable for antenna systems with narrow beams and very low SLLs. D 0 is the antenna broadside directivity from which the gain G 0 is naturally defined by, with η the antenna efficiency.

G 0 = ηD 0 , (3)

C. Normalized Radiation Pattern and Scan Loss

The normalized gain radiation pattern at the antenna broadside is defined by,

g (θ, φ) = |A (θ, φ)| 2 , (4)
where A (θ, φ) represents the two-dimensional field radiation pattern. Using Goudet's formula [START_REF] Eyraud | Théorie et Techniques des Antennes[END_REF] in which we assumed that the radiations in the two perpendicular planes are not dependent and waves on the antenna aperture are in phase, A (θ, φ) can be expressed as,

A (θ, φ) = A (θ) • A (φ) , (5)
with A (θ) and A (φ) the field radiations in elevation and azimuth defined by ( 6) and [START_REF] Gpp | Technical Specification Group Radio Access Network; Study on channel model for frequencies from 0.5 to 100 GHz[END_REF], respectively.

A (θ) = 1 + cos (θ) 2 SF (θ, H, λ). (6)
A (φ) = 1 + cos (φ) 2 SF (φ, W, λ). (7)
SF (θ, H, λ) and SF (φ, W, λ) define the spatial Fourier transform of the electric field in elevation and azimuth, respectively. H and W are the antenna aperture dimensions as illustrated in Fig. 1 and λ the wavelength. Consequently, the normalized radiation pattern in the direction (θ 0 , φ 0 ) is first obtained by replacing θ and φ by θ -θ 0 and φ -φ 0 in the above expressions. And then the aperture size H must be replaced by the equivalent size H equiv = H • cos (θ 0 ) in elevation and the size W by W equiv = W • cos (φ 0 ) in azimuth due to the projected area reduction which is estimated by a cosine reduction of effective aperture surface for the scan directions θ 0 and φ 0 [START_REF] Karki | Beam-Steerable E-Band Lens Antenna for 5G Backhaul Link[END_REF]. In this way, the beamwidth for the scan angles increases from the antenna broadside and is characterized by the half-power beamwidth (HPBW) in elevation by,

θ HP BW (θ 0 ) = θ HP BW0 cos (θ 0 ) , (8)
and in azimuth by,

φ HP BW (φ 0 ) = φ HP BW0 cos (φ 0 ) , (9)
where θ HP BW0 and φ HP BW0 represent the HPBWs at the antenna broadside in elevation and azimuth, respectively. As the beams are scanned in both perpendicular plans independently, the scan loss in the beam direction in elevation θ 0 can be approximated in dB units by,

SL (θ 0 ) = -20log 10 (|cos (θ 0 )|) . ( 10
)
Likewise, if the scan is performed in the azimuth plane, the scan loss is defined by,

SL (θ 0 ) = -20log 10 (|cos (φ 0 )|) . (11)
With the above expression, it can be noted that the SL does not depend on the frequency and dimensions of the antenna unlike HPBWs.

III. APPLICATION AND VALIDATION OF THE MODEL

The purpose of this section is to highlight the application of the proposed method and the performance comparison in terms of beamwidth and scan loss with some existing MBA TAs. To this end, parameters such as frequency, TA dimensions, antenna efficiency and field amplitude distribution on the antenna aperture are required. Using the procedure explained in section II-C, g (θ, φ) can be expressed as,

g (θ, φ) = g (θ) • g (φ) , (12)
with g (θ) and g (φ) the radiation patterns in both planes.

The choice of the distribution really impacts the beamwidth and the SLL. White [START_REF] White | Pattern limitations in multi-beam antennas[END_REF] mentioned that better sidelobe performance can be achieved by MBA systems with cosine or cosine-squared aperture distributions. Nonetheless, any other distribution could be used depending on one's goal. By considering the narrow-beam approximation since the beginning, and adopting the normalized angle coordinate u = πH λ sinθ, radiation characteristics [START_REF] Balanis | Antenna Theory: Analysis and Design[END_REF] in elevation are summarized in Table I for cosine and cosine-squared distributions. Similar expressions are obtained in azimuth by replacing θ by φ, H by W and the angle coordinate u by v = πW λ sinφ. In this work, the choice criterion of a distribution is mainly based on the value of HPBW and eventually the SLL.

In the validation procedure of the proposed approach, several models of existing antennas with different frequency bands were used. But, in this paper, we have only focused on the Ka and V bands when it comes to mmWave communication systems.

A. Performance Comparison with Ka-Band TA

For comparison with Di Palma's electronically reconfigurable TA [START_REF] Palma | Circularly-Polarized Reconfigurable Transmitarray in Ka-Band With Beam Scanning and Polarization Switching Capabilities[END_REF], which is composed of a squared surface of dimension ≈ 100 mm at 29 GHz and a boresight HP BW = 7.5 • , the cosine aperture distribution was assumed. With these parameters, a performance comparison is performed as follows. Fig. 2 shows the scan loss with respect to the directions of the beams. We can note a quite good agreement between the proposed model which can be seen as a prediction model and the measurement results performed by Di Palma et al. for a scan range of ±60 • off broadside. In Fig. 3, the beamwidth comparison is highlighted for the same scan range. We also notice that both beamwidth curves achieve similar performance in the same scan range as for the scan loss.

B. Performance Comparison with V-Band Bifocal TA

In this subsection, the bifocal TA physically realized by Pham et al. [START_REF] Pham | A Wide-Angle Beam-Steering Bifocal Transmitarray Antenna for 5G Channel Sounding at V-band[END_REF] is used in the validation procedure at 60 GHz. The system is composed of a squared surface of dimension 50 mm and a HP BW = 8.2 • is achieved at the broadside. By considering a cosine-squared distribution in addition to system parameters, the scan loss and beamwidth performance versus the beam directions are given in Figs. 2 and3, respectively. Comparing the proposed model with Pham's bifocal TA measurements, we notice similar performance for a scan range of ±35 • around the antenna broadside and a difference of about 1.5 dB at ±45 • in terms of scan loss. Outside the ±45 • angular sector, the discrepancies become significant. This may be justified by the spill-over radiation and hardware component loss that are not taken into account in the proposed prediction model, especially at high frequencies.

IV. LINK LEVEL SIMULATION AT 60 GHZ

In this section, the link level performance evaluation is performed in order to first show the impact of human blockage on 5G quality of service (QoS) in terms of block error rate (BLER). And then, the advantage of the MBA system using the proposed model in the context of blockage is highlighted. In this way, the end-to-end data processing of the 5G physical communication chain shown in Fig. 4 is used. This Tx/Rx chain is mainly composed of key blocks such as cyclic redundancy check (CRC), low-density parity-check (LDPC) channel coding, orthogonal frequency-division multiplexing (OFDM) modulation, clustered delay line (CDL) channel and minimum mean square error (MMSE) equalizer.

A. Human Blockage Modelling

Influence of human movement between the Tx and Rx constitutes a critical aspect of the mmWave propagation. In several experimental studies, measurement results have shown that human body has a significant impact on the signal strength in mmWave bands. The IEEE 802.11ad 60 GHz channel modelling mentions that human shadowing is in the range of 18 to 36 dB [START_REF] Jacob | Human Body Blockage -Guidelines for TGad MAC development[END_REF]. Collonge et al. [START_REF] Collonge | Influence of the human activity on wide-band characteristics of the 60 GHz indoor radio channel[END_REF] have also mentioned an attenuation of 20 dB when the line-of-sight (LOS) path is shadowed by a person at 60 GHz. In fact, in the specification defined by 3GPP, the human body is modelled by a blocking region with an angular space ∆θ centered around θ c in elevation and another ∆φ centered around φ c in azimuth as illustrated in Fig. 5. With this approach, the principle is to attenuate the power of any path whose arrival angles fall in the blocking region, and the attenuation value is given by a knife-edge diffraction formula. With the highest values of blocker parameters, i.e ∆θ = 15 • , ∆φ = 45 • and the distance from the blocker to the Rx r = 2 m [7], an attenuation of 16 dB at 60 GHz is deduced from the 3GPP non-self-blocking model. The non-self-blocking is characterized by the presence of humans or vehicles moving around the user as shown in Fig. 6. Although this model is used for the illustration and simulation purpose in this work, the values of 3GPP human shadowing may be subject to criticism. The difference in the human body attenuation when comparing the 3GPP modelling with other values [START_REF] Jacob | Human Body Blockage -Guidelines for TGad MAC development[END_REF] - [START_REF] Collonge | Influence of the human activity on wide-band characteristics of the 60 GHz indoor radio channel[END_REF], could be justified by the choice of the antenna system in the experiments. In other words, the values are realistic and significant when the directional antenna is used for the human shadowing assessments. In this case, the communication link can be interrupted. However, as explained by Raghavan et al. [START_REF] Raghavan | Modeling and Combating Blockage in Millimeter Wave Systems[END_REF], the 3GPP standard assumes an omnidirectional antenna for the human blockage modelling, which leads to low attenuation values.

B. Simulation Scenarios

Using the 3GPP human blockage model at 60 GHz, three simulation scenarios were considered in this work as follows.

1) Scenario 1: In this first scenario, there is no blocker between the Tx and Rx, and a single communication beam is used and steered towards the main propagation path, i.e the path with the highest power.

2) Scenario 2: The second scenario is exactly scenario 1 with the presence of a human blocker located in the direction of the single beam.

3) Scenario 3: The third scenario can be seen as the solution to scenario 2 with the application of the multibeam aspect. Since blockage occurs in the direction of the first beam, an additional beam is used, and points to the propagation path with the second highest power.

C. Simulation Results

The simulation parameters are summarized in Table II. The 3GPP CDL-C [START_REF] Gpp | Technical Specification Group Radio Access Network; Study on channel model for frequencies from 0.5 to 100 GHz[END_REF] channel model with 24 propagation paths was used. In the scenarios, the proposed MBA is only used at the base station (Tx) with the broadside gain of 20.3 dBi from the parameters of the 60 GHz bifocal TA of Pham et al. [START_REF] Pham | A Wide-Angle Beam-Steering Bifocal Transmitarray Antenna for 5G Channel Sounding at V-band[END_REF]. An omnidirectional antenna with a 0 dBi gain is used at the user equipment (Rx) level for simplicity.

In Fig. 7, a comparison of the performance in terms of BLER versus SNR of the three scenarios is highlighted for a fixed throughput of 1.5 Gbits/s. This data rate is given by the modulation and coding scheme (MCS) 15 [START_REF] Gpp | Technical Specification Group Radio Access Network; NR; Physical layer procedures for data[END_REF], i.e 64QAM modulation with a code rate of 0.65 for a bandwidth of 400 MHz [START_REF] Gpp | Technical Specification Group Radio Access Network; NR; Base Station (BS) radio transmission and reception[END_REF]. Simulation results reveal that to ensure the same 5G eMBB QoS, scenario 2 with human blocking in the single beam direction requires a 16 dB increase in SNR with respect to scenario 1 in which there is no human blockage. However, we only notice a difference of 2.8 dB between scenario 1 and scenario 3. Indeed, in scenario 3, using an additional beam that points towards the second propagation path when the first beam is blocked, almost guarantees the same QoS with the same SNR values and avoids the link interruption between the Tx and the Rx.

With these results, the blockage problem has been analyzed through the SNR variations for a fixed target throughput. Nevertheless, similar performance can be achieved in the three scenarios with the same SNR values as shown in Fig. 8, with a variation of throughputs as follows. Compared to ideal scenario 1, the communication data rates in scenario 2 must be drastically reduced from 1.5 Gbits/s to 227.9 Mbits/s due to the blockage of the single beam direction. This mitigates the advantage of using mmWaves bands. In this scenario, MCS 2 is used, i.e QPSK modulation with a code rate of 0.3. However, in scenario 3 with the use of an additional second beam at the Tx, an improvement in throughput close to scenario 1 is observed, going from 227.9 Mbits/s to 1.15 Gbits/s in the situation of blockage where MCS 2 is changed for MCS 13 (64QAM with a code rate of 0.55). In other words, the blockage induces a loss of about 85% of the throughput, whereas with the MBA this loss is only 23%.

V. CONCLUSION Multibeam antennas (MBAs) are able to generate independent and concurrent narrow beams pointing in different directions. They are low profile, cost effective and amenable for the human blockage resolution to keep the link-on-move in contrast to conventional beam steering mechanisms using multi-antennas systems which present implementation complexity and high hardware costs in mmWave bands. It is in this perspective that a generic approach is proposed in this paper for the MBA modelling. The proposed solution is flexible and can be used in the mmWave channel modelling for simulation purposes instead of using the 8 dBi static 3GPP model defined for mmWave bands up to 100 GHz. Its application shows a quite good agreement with existing MBAs in terms of scan loss and beamwidth. However, the discrepancies become significant for the scan angle from ±45 • off the antenna broadside due the increase in the spill-over radiation that is not considered in the proposed model.

For future work, we will extend the proposed MBA approach to the integration of the spill-over radiation effect from which the increase in the sidelobe level is considerable for high scan angles. In the same way, the impact of the unit-cells of TA or RA could also be investigated in the extension of the model. As we also deal with the 3GPP human blockage model, an improvement can be envisaged in order to make next studies and simulations more amenable.

Fig. 1 .

 1 Fig. 1. Rectangular aperture with dimensions W and H.

Fig. 3 .

 3 Fig. 3. Beamwidth performance at 29 and 60 GHz.

Fig. 4 .

 4 Fig. 4. Communication block diagram of 5G new radio.

Fig. 5 .

 5 Fig. 5. 3GPP Blockage scenario.

Fig. 6 .

 6 Fig. 6. Non-self-blocking illustration.

Fig. 7 .

 7 Fig. 7. Performance under the same throughput.

Fig. 8 .

 8 Fig. 8. Performance with three different throughputs.


O Gonçalves

G Prenat

B Dieny

Radiation Hardened MRAM-Based FPGA

Keywords: Dynamic random access memory (DRAM), field-programmable gate array (FPGA), look-up table, magnetic random access memory (MRAM), magnetic tunnel junction, radiation hardening, scrubbing, soft error

I. INTRODUCTION

O NE of the main issues in current static random access memory (SRAM)-based field-programmable gate arrays (FPGAs) is the radiation sensitivity of SRAM cells in advanced technology nodes. To solve this, one of the solutions is to check and write periodically the configurationmemorytocorrectsoft errors that may have occurred; it is called scrubbing. Then, SRAM FPGAs need an external non-volatile memory to store the configuration data which increases the area of boards. A domain where these constraints are critical is space domain. The use of magnetic random access memory (MRAM) memory cells can be a solution to overcome these specific constraints in space applications. This paper presents an innovative look-up table (LUT) composed of MRAMs and intended for space applications. It combines MRAM and dynamic random access memory (DRAM) memories to make a non-volatile, radiation hardened and dense FPGA.

The paper is organized as follows. Section II describes MRAM memory cells. Section III describes classical architecture of an FPGA and particularly the role of its memory cells. Section IV describes the innovative FPGA architecture. Section V shows some results of the silicon demonstrator. Then, we conclude on the use of MRAM cells for FPGAs.

II. MRAM DESCRIPTION

The basic element of MRAM is the magnetic tunnel junction (MTJ, Fig. 1), a nanostructure basically composed of two ferromagnetic (FM) layers separated by an insulating layer. The magnetization of one layer is fixed while that of the other layer (storage layer, SL) can be switched between two stable positions depending on the data to be written. The digital information stored in the MTJ is represented by its resistance level, which changes according to the relative magnetization of the two layers. Reading the memory state consists in measuring this resistance, typically by comparing the current through the MTJ to a reference current. Writing the information consists in switching the magnetization orientation of the SL, using a pulse of magnetic field or spin polarized current through the junction [START_REF] Slonczewski | Currents and torques in metallic magnetic multilayers[END_REF]. The demonstrator has been designed with thermally assisted switching (TAS) MRAMs [START_REF] Prejbeanu | Thermally assisted switching in exchange-biased storage layer magnetic tunnel junctions[END_REF]. In this technology, the MTJ is temporarily heated during write by a pulse of current flowing through the MTJ. This heating releases the SL magnetization, thus enabling writing at a lower field. This approach significantly lowers the write power consumption in comparison to t h et o g g l eM R A M. The TAS MRAM has been developed by Spintec and Crocus Technology. It is very suited for space applications thanks to its thermal stability and radiation hardness.

III. FPGA DESCRIPTION

A. FPGA

An FPGA is a programmable digital component. It can implement any digital function if it is complex enough. It is composed of a matrix of LUTs. An LUT is a simple circuit that can compute any simple digital function. For example, Fig. 2 shows a 3-input LUT. It is composed of a multiplexer (MUX) and memory cells connected to it. The 3 inputs , ,a n d select the memory cell whose content is routed to the output S. In the example, if the input code is "001", the memory cell at the address "001" is selected and the output S is in the 0 logic state. If the input code is "111", then the output is in the 1 logic 1 state. To implement a function, the truth table has to be directly written in the memory cells at the correct addresses.

To make a complex function, such as an adder or a multiplier, several LUTs are connected together with programmable interconnects. A programmable interconnect is simply made of a transistor connected to two lines, which is controlled by a memory cell connected to its gate, and that can be activated or not whether the two lines has to be connected or not.

In an FPGA, between 50% and 80% of silicon area is constituted by the interconnection network [START_REF] Kuonandj | MeasuringtheGapbetweenFPGAsandASICs[END_REF] and it contributes to dynamic power consumption between 60% and 80% [START_REF] Lin | A low-power field-programmable gate array routing fabric[END_REF]. That is why the choice of the configuration memory technology is important. Indeed, SRAM cells are composed of at least 5 transistors, so most of the surface of an FPGA is composed of memory cells.

B. Radiation Induced Errors in FPGAs

Circuits used in space systems are influenced by particles such as high energy protons, electrons, gamma, and X-rays. These particles, when they go through a circuit, hit the silicon and generate electron-hole pairs. When these pairs are under an electric field, at the drain of a deactivated transistor, the electrons and holes are separated and a current is generated. The pulse of current, in the order of hundred of picoseconds, can change the state of a transistor and create an error in the circuit andintheendamalfunctionatthe system level, that can be dangerous according to the circuit's functionality. The errors do not damage physically the circuit and disappear quickly, so they are called soft errors.

Memory cells are particularly sensitive to soft errors because a particle can change their content and hence lead to permanent e r r o ro ft h em e m o r yu n l e s si ti scorrected. This is a problem, particularly in the case of FPGAs because if the content of ac o n figuration memory is changed, the functionality of the FPGA is modified and can cause an error at the system level. To avoid this, SRAM-based FPGA can use a technique called scrubbing which consists in reading and checking the configuration memory periodically to detect and correct configuration errors that may have occurred. This technique avoids the accumulation of errors in the circuit.

IV. NEW ARCHITECTURE

The LUT we propose combines MRAM and DRAM memories (Fig. 3). The configuration memories connected to the multiplexer are made of a capacitance that stores the configuration data and a selection transistor. It is the active part of the circuit as it is computing when the FPGA is in operation. The data are refreshed periodically thanks to an MRAM block in which each cell is composed of two MTJs operating in differential mode. The configuration is stored permanently since MRAM are non-volatile and reliable because MTJs are immune to radiation. That is why this part is called reference part. Moreover, the power supply of the MRAM block can be switched off between two refreshing phases to decrease power consumption. Thanks to the refreshing phase, every possible error in the configuration memories is corrected. The refresh phase replaces the scrubbing technique in SRAM-based FPGA so it does not add power consumption to the circuit compared to the same circuit implementedinanSRAM-basedFPGA.

A pending patent [START_REF]Programmable Logic Device With Memory Refresh Based on Single Event Upset Occurrence to Maintain Soft Error Immunity by T[END_REF] describes an FPGA based on DRAM memory cell refreshed by an external ROM memory component. In our case, the use of MRAM allows a local implementation of the non-volatile memory. The power consumption is reduced as it is a local memory and there is the possibility to switch off LUT and interconnection circuits that are not used. So we can optimize the power consumption of the circuit.

V. D EMONSTRATOR

The demonstrator was fabricated using an industrial hybrid CMOS/magnetic manufacturing process developed by Tow-erJazz and Crocus Technology. The circuit implemented is a 2-input LUT as illustrated in Fig. 3. It comprises 4 configuration memory cells. Each memory cell is composed of a capacitance made of the parasitic capacitance of a transistor. The corresponding MRAM memory cell is composed of two MTJs in opposite states. In this way, the reading phase is differential and more reliable. Figs. 4 and5 show, respectively, a picture of the test chip and the layout of the circuit with its different parts. Its 2 . Most of the area is composed of the current generator and the reading circuit. However, their surface can be neglected for bigger LUT because they are shared between several MTJs.

An XOR function was first implemented. The corresponding truth table was first written in the MTJs. The capacitances were then refreshed and eventually the LUT was tested by applying all its input codes. The tests showed that these different steps are successfully achieved. First, the programming phase (Fig. 6) has been tested. The XOR function is written in the MRAM block. Each MRAM is selected sequentially (signals "sel0" to "sel3"). For each MTJs selected, the "Heat" signal is set to "0" to heat them. Then, it is set to "1" to stop the heating. The current to produce the magnetic field "Ifield" is applied at the end of the heating period for about 15 ns corresponding to the duration of the cooling phase. Then, the refreshing phase is tested (Fig. 7). Each MRAM cell and each capacitance is selected sequentially (signals "sel0" to "sel3" and "s0" to "s3" successively). Each time, the latch is updated to read the MRAM selected. The "EN" signal is activated to enable refreshing the capacitance with the data read, while the "AZ" signal is used to convert the convert the magnetic state of the MTJs into electrical information.

ThenextstepistotestthattheLUTisimplementingthecorrect function. The test vectors of this test and the result are presented in Fig. 8. All the inputs ("e0", "e1") codes are displayed. We see that the function XOR is successfully executed.

Fig. 9 shows a dynamic reconfiguration of the LUT. It means that the function implemented by the LUT is changed when the circuit is in operation. The old function was "0010" and is reconfigured into the XOR function. All the possible functions have been implemented during the tests, meaning that 16 functions were tested successfully. It shows that the MRAM were written correctly. To test the non-volatility, the power has been shut down and then powered on again. After the refresh phase, the XOR function was still implemented demonstrating that the configuration memory is non-volatile. Fig. 10 shows a parametric test in which the pulse width and magnetic fields (by varying the power supply of the magnetic field driver) are changed. Four zones can be identified.Inzone1,themagnetic field and MTJ temperature are too low to correctly write the MTJ. In zone 2, the temperature is high enough but the field is too low. In zone 3, the field is large enough but the temperature istoolow.Inzone4,thefield and temperature are high enough to write correctly the MTJs. It shows that the test is successful when the pulse width is longer than 40 ns and the supply voltage 3 Compared to Flash memories used in most of FPGA for space applications, MRAMs have a lower power consumption, are faster, and have a higher endurance [START_REF] Tehrani | Magnetoresistive random access memory using magnetic tunnel junctions[END_REF]. In addition, MTJs have a high tolerance to total ionizing dose radiation [START_REF] Conraux | Effectsof swift heavy ion bombardment on magnetic tunnel junction functional properties[END_REF] compared t oF l a s hm e m o r i e s [START_REF] Troxel | SEAKR Engineering, Inc. Seminar Day, Military and Aerospace Programmable Logic Devices (MAPLD)[END_REF] .I nc o m p a r i s o nt oS R A MF P G A ,t h e scrubbing period is faster and can be applied more frequently. In addition, the circuit benefits from the radiation immunity of MTJs to correct every configuration soft error. The MRAMs are placed locally in the FPGA, whereas SRAM FPGAs need an external radiation hardened non-volatile memory. Hence, this new MRAM FPGA architecture is more reliable and has lower power consumption than Flash and SRAM-based FPGAs.

Fig. 1 .

 1 Fig. 1. TAS writing scheme.

Fig. 2 . 3 -

 23 Fig. 2. 3-inputs look-up table example implementing an AND logic function.

Fig. 3 .

 3 Fig. 3. Schematic of the proposed LUT architecture.

Fig. 4 .

 4 Fig. 4. Picture of the silicon demonstrator.

Fig. 5 .

 5 Fig. 5. Layout of the demonstrator.

Fig. 6 .

 6 Fig. 6. Programming phase.

Fig. 7 .

 7 Fig. 7. Refreshing phase.

Fig. 8 .

 8 Fig. 8. LUT in operation.

Fig. 9 .

 9 Fig. 9. Dynamic reconfiguration.

Fig. 10 .

 10 Fig. 10. Parametric passed/failed as a function of the supply voltage of the magnetic field driver and the heating pulse duration.

  VI. CONCLUSION An MRAM-based FPGA has been designed. It combines MRAM and DRAM to make a high density, low power, and radiation hardened FPGA. A 2-input LUT demonstrator has been designed with the Tower 130 nm technology for the CMOS layers and Crocus TAS MRAM technology for the magnetic layer. The test of the demonstrator has validated the concept of configuration circuit presented in this paper.

ACKNOWLEDGMENT

The authors would like to thank TowerJazz/Crocus Technology for having given access to their technology in order to build this demonstrator and for their assistance in the tests.

This work was supported by ERC Adv grant HYMAGINE 246942.