Ecole Sigle

Doctorale

Renée El

Melhem Bât

Pascal Blaise

R Gourdon

M Stéphane

ÉLECTROTECHNIQUE E E A Électronique

M Gérard Scorletti
email: gerard.scorletti@ec-lyon.fr

M Philippe
email: philippe.normand@univ-lyon1.fr

Mme Sylvie

Jean-Yves Buffière
email: jean-yves.buffiere@insa-lyon.fr

Stéphanie Cauvin

M Jocelyn Bonjour

Véronique Guichard

J Y Toussaint

M Christian Montes

Baba Dada & Ma

Teju

M Arjan Mol

M Herman Terryn

Mme Marie

Kurty, Zhiheng, Philippe Yu Shi

Quentin

Ying Chloé

Ho

Liza

Amandine Nicolas Kathleen

Jonathan Lucas

Advanced High Strength Steels (AHSS) are increasingly used as fabrication material for vehicle Body In White (BIW), owing to their superior properties and ability to reduce carbon footprint. However, its susceptibility to hydrogen embrittlement (HE) restricts the use of AHSS. The present study aims to understand the H influence on four commercialgrade AHSS steels, two Dual Phase (DP), one Complex Phase (CP), and one Twinning Induced Plasticity (TWIP) steel. Results show high HE susceptibility for DP and TWIP steel compared to CP steel. The superior HE resistance in CP steel was attributed to a more homogeneous microstructure, smaller yet stronger trap density, and lower H concentration.

In DP steels, a high density of weak traps and high H uptake increased HE susceptibility.

During charging, H preferentially adsorbed along the grain boundaries and interfaces for all steels along with grain interior in TWIP steels. Dislocations and grain boundaries were the main trap sites for all steels, along with cementite particles in CP steels and AlN particles and austenitic grain interior in TWIP steels. For all steels under stress, hydrogen desorption increased up to yield point due to lattice expansion and dislocation movement, whereas decreased in the plastic region due to defect generation. For CP steel, strongly trapped hydrogen desorbed at UTS whereas in TWIP steel, generation of deformation twinning released hydrogen. The study of the galvanized layer showed that at higher cathodic overpotential, the Zn layer behaved as a barrier layer protecting the steel, while at a lower potential, it increased the HE susceptibility due to Zn layer dissolution. Overall, CP steel was the most resistant steel to HE, followed by TWIP and DP steels.

General Introduction and Thesis Outline 1.1 Introduction

Automobiles are one of the significant contributors to air pollution emitting a myriad of pollutants such as particulate matters (PMs), Volatile Organic Compounds (VOCs), Cox, NOx, SOx, and greenhouse gases. These pollutants not only impose a severe threat to human health and the environment but are also the principal cause of global warming and climate change. According to a study conducted by the International Council on Clean Transportation (ICCT), George Washington University and the University of Colorado Boulder, in 2015, the vehicular emissions were responsible for about 385,000 premature deaths worldwide of which 70% of the cases occurred in four largest vehicle markets in the world, i.e., China, India, the European Union, and the US. Consequently, global legislators are issuing stringent and aggressive targets for the coming years in order to curb pollution caused by vehicular emissions [1].

In an attempt to reduce the fuel emissions, automotive industries have adopted a strategy to reduce the overall weight of the manufactured vehicle. The idea behind the strategy is that the weight reduction will lead to higher fuel efficiency or lower fuel consumption, which in turn would improve vehicular emission. A 10% reduction in vehicular weight is estimated to reduce the fuel emission by 4.9-5.5 % [2], [3]. However, weight reduction should not compromise on the safety performance of the vehicle, as it is a strong factor, which influences the consumer purchasing decision. Thus, automotive manufacturers are always on a lookout for newer materials and better engineering capabilities to address these requirements. They realize that in order to develop safe automobiles with reduced weight, particular emphasis must be given to the material selection and design. Presently, weight reduction can be achieved either by using low-density materials (such as Al, Mg, or composites) or the use of Advanced High Strength Steels (AHSS). The low-density materials are expensive, restricting their use to luxury class vehicles. AHSS, as the name suggests, are steels with high tensile strength and reduce automobile weight by allowing the use of thinner sheet thickness as compared to conventional steels.

Moreover, the additional advantage of the use of AHSS over low-density alloys is its better recyclability and lower greenhouse gases (GHG) emission during production. With the above factors considered along with the vehicular emission, AHSS proves to be an quantity of hydrogen, in terms of few ppm, is sufficient to exhibit its deleterious effect on the mechanical properties of the material in question. AHSS are particularly prone to HE due to their high strength and complex microstructures. Thus in order to fully explore the applications of AHSS in automotive components, the issue of HE needs to be correctly understood and addressed.

1.2 Objectives and Layout of the Thesis HE has been widely studied in the past and still is an active research topic, given its importance in maintaining material integrity and the zest of steel industries to reinvent their steels in order to cater to the diverse demands. Despite the vast amount of studies present in the available literature, which makes the study difficult or rather interesting is that the precise role of hydrogen in the embrittlement process is still unclear. The difficulty arises due to the small size of hydrogen, the numerous parameters involved, and the fact that the embrittlement issue is caused differently in different material systems. The main goal of this study is to develop a deeper understanding of the nature of the interaction of hydrogen with different AHSS and factors which lead to their embrittlement. The present study adds to the present knowledge about the embrittlement process and aide in improving the resistance of AHSS to HE.

With the above objectives in mind, the project commenced as a collaboration between INSA Lyon, France, Indian Institute of Technology Bombay, India, and TATA Steel in the Netherlands. Tata Steel Europe is a steelmaking company that has its operations in the United Kingdom and the Netherlands. Evidently, with the growing AHSS market, the study of the detrimental influence of hydrogen in steels is an essential subject for TATA Steel.

As a result, they were kind enough to provide every possible help for the successful completion This chapter aims to summarize the information present in the available literature related to the embrittlement of advanced high strength steels. This chapter starts with the review of different types of AHSS steels with particular focus on dual phase, complex phase, and twinning induced plasticity steels. Subsequently, the chapter discusses the different mechanisms of hydrogen embrittlement, the factors responsible for HE, and specific case studies related to hydrogen embrittlement issues. Finally, the chapter highlights the H entry mechanism in iron, the nature of the H interaction, and the different techniques to study them.

Chapter 3: Study of Diffusible Behavior of Hydrogen in First Generation

AHSS

This chapter studies the interaction of hydrogen with AHSS using the electrochemical permeation cell. The permeation test is a useful tool to investigate the diffusivity of hydrogen in the material, the activation energy of traps, the trap density, and the sub-surface concentration of hydrogen in the specimen. The study allows us to compare the behavior of hydrogen in different materials and predict their sensitivity to HE. This chapter aims to understand the effect of microstructure on the diffusion behavior of hydrogen in AHSS.

Chapter 4: Study of H Trapping in AHSS using TDS and Ag Decoration Technique

In this chapter, using Thermal Desorption Spectroscopy (TDS) and Ag decoration tests, the H distribution in the material was identified. In TDS, the applied thermal energy enables us to identify the location of hydrogen in the bulk of the material, the energy of the Based on the results, a model is proposed to explain the Zn layer's role in the change in the interaction of hydrogen with steel.

Chapter 7: Summary and Perspectives

This chapter is dedicated to summarize the significant findings of the entire manuscript and highlights some potential future works.

Introduction

The transport sector is one of the most energy-consuming industries, reported contributing to over half of the world's oil consumption and around 30% of the total energy consumption [1]. Thus, to produce more fuel-efficient cars, auto industries have adapted vehicle weight reduction as a simple strategy to improve the fuel efficiency. A 10% reduction of vehicle weight is estimated to improve fuel efficiency by 4.9-5.5% [1], [2].

The materials used for Body In White (BIW) fabrication offer the highest possibility to reduce the vehicle weight. Advanced High Strength Steels (AHSS), as the name suggests, are developed to allow higher strength compared to the conventional steels, to address the issue of weight reduction. AHSS are the fastest growing materials in the automotive industry as a direct result of performance flexibility, as well as the benefits which include low cost, mass reduction capabilities, safety attributes, reduced greenhouse gas emission, and superior recyclability [3].

Generally, AHSS are multiphase steels with strength higher than 600 MPa. The multiphase microstructure results from the careful chemical composition selection and the precise heat treatment, giving the desired strength, ductility, toughness, and fatigue properties [4]. In general, AHSS have strength higher than 600 MPa. However, high strength and the complex microstructure of AHSS, also makes them susceptible to Hydrogen Embrittlement (HE), which is one of the main constraints in their use. In the auto components, steels can absorb H during the steelmaking process, production process, and under service conditions [5].

Since the discovery of HE or delayed fracture in steel [6], there is a consensus that the diffusible hydrogen causes HE. The hydrogen diffuses towards the stress field and detrimentally changes the fracture mechanism of steels [7]. Since H can cause embrittlement in low concentrations, sometimes even less than one ppm, it is essential to minimize the detrimental effect of hydrogen by either reducing the H adsorption or by making the absorbed H benign [7]. It could be possible to reduce HE by changing the surface structure of steel with by microstructural control or by the addition of coating(s). Another method could be to make the absorbed H immobile by trapping them in specific microstructural sites. In either case, the

Advanced High Strength Steels (AHSS) in Automotive Industry

Requirement for lightweight, high strength steels led to the development of High Strength Low Alloy (HSLA) steels. However, a tradeoff between their strength and ductility limits the use of HSLA [8]. To overcome such limitations, AHSS were developed in the late '70s and early '80s, with the development of low carbon steels, followed by dual phase and subsequently with TRIP steels [1], [9]. AHSS are flat steel products, and the principal difference between the conventional high strength steels and AHSS lies in their microstructure [1]. Conventional high strength steels usually have ferritic microstructure, whereas AHSS have complex structures [9]. AHSS are multiphase steels and may have a combination of polygonal ferrite, bainitic ferrite, austenite, retained austenite, martensite, and possibly other phases to obtain a wide range of properties [1], [10]. The strengthening mechanism for AHSS includes solid solution strengthening, precipitation strengthening, grain refinement, phase transformation, mechanical twinning, and cold deformation [9], [11].

Consequently, AHSS have become the materials of choice for sill reinforcements, A-pillars, Bpillars, side-impact beams, waistline reinforcements, bumpers, roof bows, and seats [1], [4], [10].

Different generations of AHSS

Based on the steel's microstructure, AHSS can be divided into three generations [12], as shown in the banana diagram in Figure 2.1.

© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon

The first generation of AHSS includes Dual Phase (DP) steels, Martensitic Steels, Complex Steels, and transformation induced plasticity steels (TRIP Steels). They have a ferrite matrix in combination with other phases such as martensite, bainite, and retained austenite imparting better mechanical properties than conventional high strength steels.

However, the banana diagram clearly shows that ductility suffered when strength increased, prompting the development of second-generation AHSS steels [4].

The second-generation AHSS have austenitic matrix. A high Mn, Al, C alloy content makes the otherwise unstable austenite to be stable at room temperature [13], [14].

Twinning Induced Plasticity (TWIP) steels, Lightweight steels with Induced Plasticity (L-IP) steels, Shear band Induced Plasticity (SIP) steels are some of the examples of secondgeneration AHSS [2], [9], [14], [15]. These steels offer superior strength and ductility, but the high alloying cost and susceptibility to delayed cracking restricts their commercial application [4], [13].

To address these problems, the third-generation AHSS are being developed, which offer high strength and increased ductility at moderate alloying concentrations [16], [17]. The third-generation AHSS include Quenching and Partitioning (Q&P) steels and medium Mn steels [4], [9]. The microstructure of third-generation steels are predominantly bainitic ferrite with martensite as the second phase [18]. These steels are considered as the future opportunities in automotive and other structural applications, as they address the strength and ductility issues concerning the first generation AHSS and are much more economical than the second-generation AHSS [2]. The next section briefly discusses some of the first and second-generation AHSS and their properties.

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© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon IF is interstitial -free steels, CMn is Carbon Manganese steels, HSLA is high strength low alloy steels, DP is dual phase steels, TRIP is transformation induced plasticity steels, CP is complex phase steels, TWIP is twinning induced plasticity steels, Q&P is quenching and partitioning steels. Adapted from Venezuela et al. [9] 2.1.1 Examples of first and second-generation AHSS steels

Dual Phase Steels

.2(a) presents the SEM image of DP steel, characterized by a soft ferrite matrix embedded with hard martensite islands, offering high strength (up to 1200 MPa) and good ductility (10% or higher) [19]. The ferrite is generally distributed continuously in the microstructure but becomes discontinuous when the volume fraction of martensite exceeds 30% [2], [20]. Though the term 'dual phase' suggests that the microstructure is composed of two phases, DP steels may contain some amount of bainite and retained austenite, depending on the processing conditions [2], [21]. DP steels use low C-Mn-Si as the elemental chemical composition [2], [22]. Carbon is usually used in the range of 0.06-0.15 wt.% and acts as an austenitic stabilizer and also strengthens martensite. Mn is a ferrite solution strengthener and retards ferrite formation and is usually added between 1.5-2.5 wt.% to stabilize austenite, while Si promotes ferrite formation. Other elements, such as Cr and Mo, are used to retard pearlite and bainite formation [2].

© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon strength of DP steel is proportional to the amount of martensite [20], [23]- [25] up to a certain extent beyond which the strength decreases with further increase in martensitic volume fraction due to decrease in the carbon content of martensite [26], [27]. DP steels exhibit a low yield strength ratio of about 0.5 and a high work-hardening rate [22], [28]. DP steels exhibit continuous yielding due to the presence of internal stresses from austenitic transformation to martensite, allowing dislocations to remain mobile [1], [29]. During the transformation from austenite to martensite, the volume expansion results in local plastic deformation of ferrite, which leads to high dislocation density and residual stresses in the ferrite phase near the martensite [1], [19]. The different mechanical properties of the constituent phases and the heterogeneity in the ferrite phase result in ductile fracture under tensile stress [30], [31]. The DP steel can fail either due to martensitic cracking or void nucleation caused by decohesion of the ferrite-martensite interface [32], [33]. The martensite morphology also influences the fracture mechanism. In the case of coarse martensite, the due to martensite cracking at low strain results in void formation followed by interfacial decohesion at ferrite/martensite interface at higher strain. However, in the case of fine martensite, the majority of the voids formed are due to the decohesion of the ferrite-martensite interface [34].

BIW uses DP steels as underbody cross members, bumpers, B-pillars, and door beam [14], [18]. The improved formability and ductility of specific grades of DP enable them to replace thicker and conventional high strength steels [35]. One limitation of DP steel, however, is its low plastic strain ratio value (rm), making them a weak candidate for applications that require high drawability. They usually exhibit poor hole expansion ratio values. Ti addition can strengthen ferrite and reduces the hardness difference between the two phases.

© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon The microstructure of CP steel is made up of ferrite with a large volume fraction of hard phases such as martensite and bainite and usually does not contain retained austenite [18], [22], [36], [37]. CP steels have a C-Mn-Al type alloying system similar to that of TRIP steels with additional alloying elements such as Nb, Ti, and V to form fine precipitates for precipitation strengthening effect [2], [22]. Titanium or niobium addition helps in second phase particle precipitation and retards recrystallization, helping grain refinement [2], [22], [38], [39]. Using the same chemical composition, but by adjusting the cooling rate, the volume © 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon fraction of different phases can be controlled, offering a wide range of mechanical properties.

Complex Phase Steels

The cooling rate controls the formation of ferrite and bainite/tempered martensite, where a high cooling rate favors martensitic transformation and suppresses ferrite formation [2].

Figure 2.2(d) presents the stress-strain curve of a CP steel. In comparison to DP steels, CP steels have much higher yield strength (YS) for the same strength level. The strengthening mechanisms in CP steels include precipitation hardening, grain refinement, and solid solution strengthening [2]. The high strength is because of the high dislocation density of bainitic ferrite (> 10 12 .cm -2 ) and the fine dispersion of the second particles and carbonitrides or carbides [22]. The ductility of CP steel decreases drastically from 13% to 3% when the strength increases from 800 MPa to 1200 MPa [39]. CP steels with a bainitic microstructure exhibits superior formability due to a relatively small difference in the hardness of bainite and martensite [40]. Due to the low hardness ratio, CP grades exhibit high local formability, but the global formability is low due to high dislocation density pileups [2].

Pathak et al. [41] investigated the fracture behavior of CP800 steels and observed that the TiN particles are the primary location for void nucleation, and the martensite-bainite interface is the secondary void nucleation region. Scott et al. [42] studied the damage evolution in CP steels and confirmed that void nucleation occurs at TiN particles due to the fracture or decohesion of the particle-matrix interface. They further added that the void growth occurs in the loading direction, resulting in ellipsoid shape at moderate to high level of strain. During deformation, CP steels exhibit continuous yielding behavior leading to high energy absorption, high residual deformation capacity, and excellent hole expansion [2]. With high yield strength, good local formability, better bendability, edge ductility, structural parts with dominating compressive load use CP steels, such as cross members or longitudinal members, which are

commonly produced by roll forming and bending operations [18], [37].

TWIP Steels

Twinning Induced Plasticity (TWIP) steels are the second generation AHSS with austenite matrix with high potential for automotive forming applications [13], [43]. steel is gaining interest in the automotive industry because of its high strength and ductility © 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon [13], [43], as shown in Figure 2.2(d). TWIP steels have a high work-hardening rate due to the generation of deformation-induced twins [44]. The high strength results from the 'dynamic hall petch effect' as the twins generated during straining act as grain boundaries, refining grain size, and impeding the dislocation movement during deformation resulting in high strain hardening [14].

The twin formation is closely related to the stacking fault energy (SFE) SFE is the energy required to change the stacking sequence of atoms during deformation, and in TWIP steel, SFE is affected by chemical composition, temperature, segregation of alloying elements at stacking faults [39]. If SFE is < 20 mJ.m -2 , strain-induced transformation is likely to take place, and if it exceeds 50 mJ.m -2 , the formation of twins is suppressed [14]. Thus, by precisely controlling the alloy composition, the SFE is maintained between 20-50 mJ.m -2 , so that twins can be induced during straining, leading to enhanced mechanical properties [2], [9],

[45]- [48]. In TWIP steels, C-Mn-Al are the major alloying elements. High Mn concentration is required (15-30 wt%) to stabilize austenite at room temperature, but with an increased amount of C (> 0.6 wt%), less amount of Mn (around 12%) is required [14]. Al increases the SFE of austenite and prevents the austenite to martensite transformation during deformation [49], [50]. Silicon addition helps in solid solution strengthening but reduces the SFE [51]. Al is reported to (a) increase the SFE (b) minimize susceptibility to delayed fracture (c) increase the YS (d) reduction in density of the steel (e) suppress the formation of martensite and (f) minimize HE issues [2], [13], [14], [43], [51]- [55] The strain hardening behavior of TWIP steel is due to the generation of dislocation slip, deformation twinning, and martensitic transformation. A minimum deformation is required to initiate deformation twins [48]. With an increase in strain, the density of twins increases, providing more obstacles, increasing strength, and work-hardening rate, till the twinning volume fraction reaches a saturation level [56]. On further straining, unstable austenite may transform to martensite [57]. If the austenite is stable, secondary twinning may nucleate between the interspaces of the primary twins [56], increasing the material strength.

© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon

Due to high strength and ductility, structural frame and safety parts with complex shape such a bumpers and B-pillars, can use TWIP steel for fabrication [58]. The potential application of TWIP steel is not likely to be limited to the automotive industry and may extend to other transportation areas, like pipe production, shipbuilding, and special applications, which may require nonmagnetic panels [14]. However, the application of TWIP steels have not been extensively explored due to its susceptibility to delayed fracture and the low YS for given tensile strength [54], [56].

Hydrogen Embrittlement issues in AHSS

Advanced High Strength Steels offer excellent mechanical properties, but unfortunately, are prone to hydrogen embrittlement [29]. In the automotive industry, steel sheets undergo forming operations, heat treatment during pretreatment, and painting after assembly. Hydrogen can enter before, during, or after these stages from the electrolytic coating, phosphate, or primer application, the presence of water vapor during spot welding or corrosion in service [59]. The failure due to H can be instantaneous upon attack or delayed, making failure prediction difficult, increasing the liability, maintenance, and repair cost [60].

In the next section, some of the existing theories related to HE are reviewed, along with some specific case studies for AHSS steels.

Hydrogen Embrittlement in steels

Usually, HE susceptibility increases with the increase in H concentration in the steels up to a threshold concentration, beyond which there is little change in the degree of embrittlement [61], [62]. HE can influence the mechanical properties by subcritical crack growth leading to fracture at crack tip stress intensity (KI) lower than fracture toughness (KIC) [61]. Though a vast number of studies are available in the literature, there is no single mechanism that can thoroughly explain every observed embrittlement behavior [62], [63].

The complexity of this subject can be highlighted by the fact that hydrogen influences the degradation behavior differently for different metal/alloy systems [64]. As a result, several mechanisms that may explain the embrittlement behavior, singularly, or in combination, are proposed [62].

Mechanisms of Hydrogen Embrittlement

Hydrogen Enhanced Dislocation Emission (HEDE)

Pfeil, in 1926 proposed that hydrogen negatively influenced the cohesion across cubic planes and grain boundaries [66], [67]. Later, Troiani explained the atomic bond weakening by hydrogen in terms of charge transfer and proposed the HEDE mechanism [68].

Oriani and later, other researchers further developed these mechanisms in their study [69], [START_REF] Mcmahon | Hydrogen-induced intergranular fracture of steels[END_REF]. Decohesion theory became widely accepted in the 1960s and early 1970s but presently faces stiff competition with newer theories based on plasticity [67].

HEDE theory states the decrease of metals' atomic bond strength due to H segregation at grain boundary or other interfaces [START_REF] Katzarov | Hydrogen embrittlement II. Analysis of hydrogen-enhanced decohesion across (111) planes in α -Fe[END_REF]. The decrease in the cohesive interatomic force is due to the donation of 1s electron of hydrogen to the unfilled 3d orbital of iron. The charge transfer leads to the weakening of the interatomic bonds resulting in atomic decohesion in preference to slip [START_REF] Lynch | Hydrogen embrittlement phenomena and mechanisms[END_REF]. The atomic decohesion in the presence of hydrogen can occur at (a) crack tips, (b) ahead of cracks where dislocation shielding effect results in maximum tensile stress, (c) position of maximum hydrostatic stress and (d) particle-matrix interface ahead of cracks [START_REF] Katzarov | Hydrogen embrittlement II. Analysis of hydrogen-enhanced decohesion across (111) planes in α -Fe[END_REF]. The basic idea of HEDE is that H damage occurs in the fracture process zone when the stress to open the crack tip (lowered due to the presence of hydrogen) exceeds the maximum atomic local cohesion strength [START_REF] Gangloff | Hydrogen Assisted Cracking of High Strength Alloys[END_REF].

© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon

It is challenging to propose direct experimental evidence for HEDE as there is no experimental technique for directly observing the events occurring at the atomic scale on the crack tips in bulk material [START_REF] Lynch | Hydrogen embrittlement phenomena and mechanisms[END_REF]. Also, the dissolved H concentration in the bulk material is orders of magnitude lower than the H concentration required to accumulate locally near the crack tip [START_REF] Gangloff | Hydrogen Assisted Cracking of High Strength Alloys[END_REF]. The most direct evidence for HEDE is the 'easier field evaporation' of surface atoms, which is observed by the field ion microscopy with hydrogen as the imaging gas [START_REF] Lynch | Hydrogen embrittlement phenomena and mechanisms[END_REF].

Sometimes, a featureless fracture observed by scanning electron microscopy (SEM) at high magnification is considered as the proof of decohesion in the presence of hydrogen. However, the resolution of SEM may not be high enough to resolve the shallow dimples associated with localized plasticity during crack growth [START_REF] Gangloff | Hydrogen Assisted Cracking of High Strength Alloys[END_REF]. Nevertheless, there is indirect evidence for HEDE, which includes (a) quantum mechanical calculations supporting the reduction of interatomic bond strength due to hydrogen and (b) high concentration of hydrogen observed in the fracture zone like grain boundary and precipitate/matrix interface [9].

Hydrogen Enhanced Localized Plasticity (HELP)

In 1972, Beachem proposed that hydrogen facilitates dislocation motion, resulting in localized plastic deformation with macroscopic brittle fracture characteristics [START_REF] Beachem | A new model for hydrogen-assisted cracking (hydrogen "embrittlement[END_REF]. It was in contradiction with the popular belief at that time that H restricts dislocation movement to cause brittle fracture. He said, "instead of hydrogen locking dislocations in place, it unlocks them and allows them to multiply or move at reduced stresses." It was a significant development in the understanding of HE as previously, the ductile features observed on the fracture surface were considered as the after-effects of the embrittlement process and was not relevant to the HE mechanism. Since then, much research is focussed in this direction [62], [63], [START_REF] Birnbaum | Hydrogen-enhanced localized plasticity--a mechanism for hydrogenrelated fracture[END_REF], [START_REF] Birnbaum | Hydrogen-enhanced localized plasticity-a mechanism for hydrogenrelated fracture[END_REF]. In the late 1980s, Birnbaum et al. proposed the HELP mechanism, where the hydrogen in the solution, ahead of a crack tip, and not the one adsorbed at a crack tip, is responsible for a highly localized plasticity behavior [START_REF] Birnbaum | Hydrogen-enhanced localized plasticity-a mechanism for hydrogenrelated fracture[END_REF], [START_REF] Dear | Mechanisms of hydrogen embrittlement in steels: discussion[END_REF].

The fundamental concept of the HELP mechanism is that the solute hydrogen shields dislocations from elastic interactions with the physical obstacles present in the material [9]. High-energy elastic interactions result in a reduction of slip movement by hindering the dislocation motion. Lowering of elastic energy in the presence of hydrogen facilitates dislocation mobility and plastic deformation. According to the HELP mechanism,

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© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon solute hydrogen atoms diffuse and concentrate around dislocations and move along with the dislocations [67], [START_REF] Lynch | Hydrogen embrittlement phenomena and mechanisms[END_REF], [START_REF] Birnbaum | Hydrogen-enhanced localized plasticity-a mechanism for hydrogenrelated fracture[END_REF].

Unlike HEDE, there are considerable evidences to support the HELP mechanism.

'Illinois group' led by Birnbaum, Robertson, and Sofronis provided direct evidence to show the role of hydrogen in enhancing dislocation movement [62]. They conducted deformation experiments in an environmental transmission electron microscopy (TEM), where they observed the dislocation behavior in both the presence and absence of hydrogen. First, they observed the dislocation velocity change with the introduction of the hydrogen environment on a specimen with a constant rate of displacement of the specimen stage. In the second type of experiment, hydrogen was introduced in the stressed specimen with stationary dislocations to observe the effect of hydrogen. In the first case, the dislocation velocity increased on introducing hydrogen gas, whereas in the second case, the stationary dislocations started moving.

Further, on removal of the gaseous H environment, the dislocation behavior returned to initial conditions. They found this effect to be independent of the material but depended on the material purity and hydrogen gas pressure. It not only validated the HELP mechanism but also showed the reversible nature of HE.

Adsorption Induced Dislocation Emission (AIDE)

AIDE mechanism is a variation of hydrogen assisted cracking where the presence of H stimulates the dislocation movement. Clum first reported the role of adsorbed hydrogen on the dislocation motion using field ion microscopy but did not correlate the observation with the embrittlement mechanism [START_REF] Lynch | Hydrogen embrittlement phenomena and mechanisms[END_REF], [START_REF] Clum | The role of hydrogen in dislocation generation in iron alloys[END_REF]. Lynch proposed and later developed the AIDE mechanism in his papers [67], [START_REF] Lynch | Hydrogen embrittlement phenomena and mechanisms[END_REF], [START_REF] Lynch | Environmentally assisted cracking: Overview of evidence for an adsorption-induced localised-slip process[END_REF]. In the AIDE model, 'dislocation emission' includes the dislocation nucleation and subsequent dislocation movement away from the crack tip.

Lynch [START_REF] Lynch | Hydrogen embrittlement phenomena and mechanisms[END_REF] emphasized that adsorbed H facilitates dislocation nucleation, and once nucleated, the dislocations readily moves away from the crack tip under the applied stress.

In the AIDE mechanism, crack growth under increasing stresses occur due to the combined effect of nucleation and growth of microvoids ahead of crack tips and dislocation emission from the crack tips. Voids nucleate at second phase particles, slip band intersections,

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© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon or other sites in the plastic zone ahead of the crack tips [START_REF] Lynch | Hydrogen embrittlement phenomena and mechanisms[END_REF]. The linking of voids aides the crack advancement and maintains a sharp crack tip by interacting with intense slip bands, from crack tip dislocation emission [START_REF] Gangloff | Hydrogen Assisted Cracking of High Strength Alloys[END_REF]. The intergranular cracking in the AIDE mechanisms indicate the preferential H adsorption along the line of intersection between the grain boundary plane and the crack front, and perhaps the high density of precipitates that form preferentially along the grain boundaries [START_REF] Gangloff | Hydrogen Assisted Cracking of High Strength Alloys[END_REF]. Lynch mentions that the AIDE mechanism is more complicated than the HEDE or HELP mechanisms [67].

Fractography studies mainly contribute to the evidence of the AIDE mechanism.

The presence of a high density of tiny dimples in the intergranular facet surface in several alloys and polycrystalline microstructures is observed [START_REF] Gangloff | Hydrogen Assisted Cracking of High Strength Alloys[END_REF]. These features are similar to that caused by Liquid Metal embrittlement (LME), where cracking is restricted to a crack surface, reportedly consistent with the AIDE mechanism [START_REF] Gangloff | Hydrogen Assisted Cracking of High Strength Alloys[END_REF].

Though the above mechanisms have been widely used in the literature to explain the embrittlement process, the mechanisms are debatable due to their weaknesses in supporting evidences. For example, in the experiment carried out to support the HELP mechanism [62], the surface issues might have been exacerbated due to the high H fugacity.

High H fugacity, produced by H2 dissociation in the electron beam, is capable of reducing surface oxides, which could affect the dislocation motion [START_REF] Gangloff | Hydrogen Assisted Cracking of High Strength Alloys[END_REF]. Gangloff presents a comprehensive overview of the limitations of each of the mechanisms proposed above [START_REF] Gangloff | Hydrogen Assisted Cracking of High Strength Alloys[END_REF]. To summarize, the effect of hydrogen on the embrittlement yet unresolved, evidenced by newer models regularly replacing the older ones [62]. However, each model provides a valuable insight for the advancement of the understanding of the hydrogen embrittlement phenomenon.

Factors affecting HE

The principal factors, which can cause HE issues, are the material itself, the environment, and the associated stress [START_REF] Nazarov | Scanning Kelvin Probe for detection of the hydrogen induced by atmospheric corrosion of ultra-high strength steel[END_REF]. Some factors may predominate the others. For example, the strength influences the susceptibility of a material to HE [65], [67], [START_REF] Zhang | Effect of the cathodic current density on the sub-surface concentration of hydrogen in X80 pipeline steels under cathodic protection[END_REF], [START_REF] Park | Effect of microstructure on the hydrogen trapping efficiency and hydrogen induced cracking of linepipe steel[END_REF].

However, HE susceptibility of most materials increases with an increase in strength provided that the microstructures are similar [START_REF] Lynch | Hydrogen embrittlement phenomena and mechanisms[END_REF]. Usually, the effect of microstructure overrides the © 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon effect of strength. Grain size, amount of retained austenite, and the nature of martensite (twinned martensite or lath martensite) also play a decisive role in determining the susceptibility to HE [67], [START_REF] Lynch | Hydrogen embrittlement phenomena and mechanisms[END_REF].

Microstructure

The HE susceptibility of steel is closely related to the microstructure [START_REF] Luppo | The influence of microstructure on the trapping and diffusion of hydrogen in a low carbon steel[END_REF]. A systematic study on the effect of carbon and heat treatment on hydrogen diffusivity and trapping capacity of different microstructures was performed [START_REF] Chan | Effect of carbon content on hydrogen occlusivity and embrittlement of ferrite-pearlite steels[END_REF]- [START_REF] Chan | Effect of retained austenite on the hydrogen content and effective diffusivity of martensitic structure[END_REF]. In the case of ferrite/pearlite steels, the maximum amount of H trapped was at 0.69 wt% of carbon, which corresponds to the composition close to eutectoid composition with maximum ferrite/pearlite interface (Figure 2.3(a)). The ferrite/pearlite interface behaved as hydrogen traps, which resulted in the high hydrogen uptake, with the ferrite/cementite interface in pearlite is a weak trap as compared to the ferrite/pearlite interface.

[84]- [START_REF] Chan | Effect of retained austenite on the hydrogen content and effective diffusivity of martensitic structure[END_REF] In another study, in the case of martensitic microstructure, both the high and low carbon steels exhibit higher hydrogen trapping as compared to steels with medium carbon concentration [START_REF] Chan | Hydrogen trapping ability of steels with different microstructures[END_REF]. The martensite morphology plays an essential role in the hydrogen trapping behavior (Figure 2.3(b)). In the study, the as-quenched martensitic structure has weak HE resistance, whereas tempered martensite or tempered bainite offers better HE resistance [START_REF] Chan | Hydrogen trapping ability of steels with different microstructures[END_REF]. A linear relationship between the hydrogen content and retained austenite was proposed, with the hydrogen concentration per unit volume of retained austenite was independent of the carbon concentration [START_REF] Chan | Effect of retained austenite on the hydrogen content and effective diffusivity of martensitic structure[END_REF]. The hydrogen present in the steel is more likely to be trapped by the martensite-austenite interface than in retained austenite itself.

© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon Hydrogen trapping of the martensitic structure as a function of carbon content [START_REF] Chan | Hydrogen trapping ability of steels with different microstructures[END_REF] Tempered martensites are considered as one of the best microstructure for steels to resist hydrogen assisted fatigue crack propagation [START_REF] Tau | Hydrogen enhanced fatigue crack propagation of bainitic and tempered martensitic steels[END_REF]. The detrimental effect of hydrogen charging on elongation decreases with increasing tempering temperature [20]. Tempering tends to soften high carbon martensite, reducing H sensitivity. A sufficiently tempered martensite can significantly reduce the embrittlement issue in DP steels, but it will severely affect the material strength. Uncharged martensitic specimens showed ductile fractures, whereas, the H charged specimen's fracture surface show a mixed type due to intergranular and interlath cracking. On tempering, the fracture surface showed a combination of ductile and intergranular cracking.

Bainitic steels also offer excellent resistance to embrittlement in conditions that cause HE and corrosion fatigue [START_REF] Escobar | Thermal desorption spectroscopy study of the interaction between hydrogen and different microstructural constituents in lab cast Fe-C alloys[END_REF]. In high strength steels, lower bainite reduces the susceptibility to slow strain rate hydrogen embrittlement as compared to quenched and tempered martensite, normalized ferrite/pearlite, spherodized microstructure, and tempered martensite [START_REF] Escobar | Thermal desorption spectroscopy study of the interaction between hydrogen and different microstructural constituents in lab cast Fe-C alloys[END_REF]. It is due to the homogeneous, refined, and stress-free microstructure, though other factors such as carbide type and other microphases should be considered as well [START_REF] Tau | Hydrogen enhanced fatigue crack propagation of bainitic and tempered martensitic steels[END_REF].

Tau et al. studied three different bainitic structures and compared them with tempered martensitic structures of similar yield strength [START_REF] Tau | Hydrogen enhanced fatigue crack propagation of bainitic and tempered martensitic steels[END_REF]. They observed that the carbide particles provide a large area for the hydrogen trapping. With the increase in the cementite particle size, the total interfacial area decreased the hydrogen content. They concluded that the The role of cementite is, however, debatable in the literature. Some researchers observed slow H diffusion in pearlite due to the effective H trapping of cementite. Though the exact nature of trapping was unclear, it was speculated either due to the stress field induced by cementite precipitation, boundary layers, or by cementite itself. [START_REF] Escobar | Thermal desorption spectroscopy study of the interaction between hydrogen and different microstructural constituents in lab cast Fe-C alloys[END_REF], [START_REF] Hanada | Profiling of hydrogen accumulation in a tempered martensite microstructure by means of tritium autoradiography[END_REF]. Others reported that the H diffusivity decreases with the increase in the ferrite/cementite interface in pearlite, and pro eutectoid ferrite acts as the main H diffusion [START_REF] Hagi | Diffusion Coefficients of Hydrogen in Carbon Steels between 278 and 318 K and Hydrogen Trapping Effect of Interface between Cementite and Ferrite[END_REF], [START_REF] Hong | The interaction of hydrogen and the cementite-ferrite interface in carbon steel[END_REF]. Some researchers failed to find any meaningful relationship between hydrogen diffusion and cementite [START_REF] Akiyama | Hydrogen Visualization in Steels Using Ag Decoration Method[END_REF], [START_REF] Wei | Quantitative analysis on hydrogen trapping of TiC particles in steel[END_REF]

Environment

One of the primary H sources for HE of steels is the cleaning solutions and the application of protective coatings. Heat treatment in hydrogen bearing environment and application of cathodic protection with impressed current or sacrificial anodes can also lead to hydrogen uptake [START_REF] Lynch | Hydrogen embrittlement phenomena and mechanisms[END_REF]. Corrosion of sacrificial coatings during service can cause HE issues, especially if the coatings are porous [START_REF] Lynch | Hydrogen embrittlement phenomena and mechanisms[END_REF]. Environmental pH also has a marked influence on the hydrogen uptake capacity with acidic conditions favoring hydrogen pickup [START_REF] Zhang | Effect of the cathodic current density on the sub-surface concentration of hydrogen in X80 pipeline steels under cathodic protection[END_REF].

Temperature also plays an essential role in the embrittlement behavior. For most steels, the HE susceptibility is maximum at 0 °C, with the susceptibility gradually decreasing with increasing temperature [START_REF] Louthan | Hydrogen Embrittlement of Metals: A Primer for the Failure Analyst[END_REF].

Stress

Applied stress has a profound influence on the hydrogen behavior within the material. The H concentration increases with the applied stress due to the lattice expansion and the increase in microstructural defects, which trap more hydrogen [START_REF] Kim | Electrochemical hydrogen permeation measurement through high-strength steel under uniaxial tensile stress in plastic range[END_REF], [START_REF] Kim | Effect of tensile stress in elastic and plastic range on hydrogen permeation of high-strength steel in sour environment[END_REF]. The presence of stress concentrators such as notches, surface scratches, and machining marks also increases the sensitivity to HE [9]. The paper of Wang et al. [START_REF] Wang | Effect of hydrogen and stress concentration on the notch tensile strength of AISI 4135 steel[END_REF] (and references therein) describes that increasing stress concentration in the material makes it more susceptible to HE.

Manufacturing processes, which create residual stresses in the material such as cold drawing or excessive grinding, are also detrimental from the perspective of HE [START_REF] Walton | The Influence of Residual Stresses on the Susceptibility to Hydrogen Embrittlement in Hardened Steel Components Subjected to Rolling Contact Conditions[END_REF].

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Alloying Elements

Alloying addition is a common practice in steel design to achieve desired microstructures and properties. However, the addition of alloying elements also influence the HE tendency of steels. The addition of carbon (C) in steels increases the strength of the steel.

However, the phase hardening due to C addition also increases the HE susceptibility. An increase of 1 wt% C in steel can increase H solubility by 70 times and reduce the H diffusion coefficient by two orders of magnitude [START_REF] Hadam | Absorption of hydrogen in tensile strained iron and high-carbon steel studied by electrochemical permeation and desorption techniques[END_REF]. Mo decreases the H diffusion coefficient by increasing H solubility, as electrons from hydrogen fill-up the unfilled Mo d orbitals [START_REF] Ningshen | Diffusion behaviour of hydrogen in nitrogen containing austenitic alloys[END_REF].

Mn, which is prone to form MnS, also increases the solubility of hydrogen, whereas Al is known to decrease the solubility of H due to the lowering of dislocation density [53]. The effect of Cr on the solubility of hydrogen is not straight forward and depends on its concentration in the alloy. When the Cr concentration increases over 20 wt.% there is a decrease in the hydrogen embrittlement resistance of austenitic steels [START_REF] Martin | Effect of alloying elements on hydrogen environment embrittlement of AISI type 304 austenitic stainless steel[END_REF]. Alloying with Cr affects both the crystal and electron structure of the material, which generates a competing effect. Martin et al. [START_REF] Martin | Effect of alloying elements on hydrogen environment embrittlement of AISI type 304 austenitic stainless steel[END_REF] The addition of alloying elements such as V or Ti can lead to the formation of VC or TiC, which acts as strong binding sites for hydrogen, thus pinning down hydrogen in the lattice [START_REF] Escobar | Thermal desorption spectroscopy study of the interaction of hydrogen with TiC precipitates[END_REF], [START_REF] Nagao | Visualization of hydrogen transport in high strength steels affected by stress fields and hydrogen trapping[END_REF]. The nature of the carbide precipitates also plays an essential role in the interaction of the precipitates with H. For example, when the precipitates are small and incoherent, the particles effectively pin down or trap the hydrogen at precipitate/matrix interface, and the amount of energy required to release hydrogen from such locations is very high [START_REF] Pressouyre | An example of the effect of hydrogen trapping on hydrogen embrittlement[END_REF]. The activation energy to release hydrogen from the incoherent TiC particles can be about 88-95 kJ.mol -1 [START_REF] Pressouyre | An example of the effect of hydrogen trapping on hydrogen embrittlement[END_REF], [START_REF] Lee | crystals structures by the desorption thermal analysis technique[END_REF]. However, for coherent particles, the trapping is not very effective, with H easily diffusing out of the traps.

Case Studies related to HE in AHSS

This section discusses some case studies from the available literature related to the H influence in AHSS. Two first-generation steels, DP and CP, and one second-generation steel, TWIP, will be discussed.

© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon

Dual Phase Steels

Davies [20], [START_REF] Davies | Influence of martensite content on the hydrogen embrittlement of dual-phase steels[END_REF] conducted studies to determine the susceptibility of dual phase steel to hydrogen embrittlement with changes in the martensitic content. With an increase in the martensitic content, the ferrite grain size decreases, and the martensitic regions tend to percolate together to form a semi-continuous network. Hydrogen charging did not significantly affect the yield strength or the tensile strength but had a noticeable impact on the ductility. He observed a linear relationship between the material strength and the martensite content. The relationship between the embrittlement susceptibility and the martensite content was, however, not linear. DP steels with up to 10% martensite was not significantly susceptible to embrittlement, and the embrittlement susceptibility increased with the increase in martensite content from 10% to 30% and beyond 30% martensite when the martensite formed a network, the susceptibility to embrittlement remained constant.

Davies [20] in another study examined the fracture surface of DP steels with and without hydrogen charging. In the absence of hydrogen, the DP steels failed by ductile shear with local necking, and the fracture surface had ductile fractures with occasional inclusions in the dimples. After hydrogen charging, the steels failed in a brittle manner with coarser fracture and transgranular cleavage facets. The fractures were at right angles to the tensile axis with no necking. He concluded from the above results that the DP steels were susceptible to HE due to the presence of high strength martensitic islands [20], [START_REF] Davies | Influence of martensite content on the hydrogen embrittlement of dual-phase steels[END_REF]. Fracture in dual phase steels involves cleavage of both ferrite and martensite phases. The cleavage cracks initiate in the high strength martensite or at the ferrite-martensite interface and propagates through ferrite. When the martensite content is low, the crack propagation through soft ferrite is difficult, which retards its susceptibility to embrittlement. With an increase in the martensite content, the ferrite grain size decreases. As the grain boundary node of the ferrite is anchored by martensite, with decrease in the ferrite grain size the distance between two martensite islands reduces. This leads to easier propagation of cracks through the martensite particles. Thus, the ability of the crack to propagate in the dual phase steel depended on the strength, size and martensite distribution. Ronevich et al. [55] observed the fracture surface change of DP steels with increased H concentration in the specimen (from as-received condition to 26 ppm). The asreceived specimen failed perpendicular to the tensile axis and showed significant necking; however, with an increase in the H concentration, the necking gradually decreased. At 26 ppm of H, the specimen failed with negligible necking and brittle tortuous surface. Fractography revealed an increased number of voids with an increase in the H concentration. They suggested that the hydrogen weakened the cohesive force between martensite and ferrite, facilitating interface separation at low strain levels-the cracks generated at the martensite, which later propagated into the ferrite matrix.

Complex Phase Steels

Not much literature on the effect of hydrogen on complex phase steels is available as compared to that of other AHSS. However, the beneficial effects of bainite is extensively reported in the literature. [START_REF] Tau | Hydrogen enhanced fatigue crack propagation of bainitic and tempered martensitic steels[END_REF], [START_REF] Escobar | Thermal desorption spectroscopy study of the interaction between hydrogen and different microstructural constituents in lab cast Fe-C alloys[END_REF], [START_REF] Rehrl | Mechanical properties and fracture behavior of hydrogen charged AHSS/UHSS grades at high-and low strain rate tests[END_REF]. One of the reasons for the improved resistance of CP steels to HE is the lesser difference in the hardness between bainite and martensite as compared to the ferrite-martensite matrix in dual phase steels. Rehrl et al. [START_REF] Rehrl | Mechanical properties and fracture behavior of hydrogen charged AHSS/UHSS grades at high-and low strain rate tests[END_REF] observed a significant difference in the fracture surface of CP1200 steels between the H charged and uncharged conditions. The hydrogen charged specimens revealed quasi-cleavage and whereas the uncharged specimens showed ductile features. They reported that quasi-cleavage fracture initiated at the edge of tensile specimens and around the inclusions, which became ductile with increasing distance from the edge of the specimens to the center of the specimens.

TWIP steels

Chun et al. [52] investigated the effect of Al content in the hydrogen embrittlement susceptibility of the TWIP steels. In their study, they prepared TWIP steels with Al content ranging from 0 to 3.5 wt% and further cold rolled some specimens to induce defects before charging them with hydrogen. For the uncharged specimens, they observed an increase in the yield strength and decrease in the tensile strength and the elongation with increasing Al content of steel. They attributed the increase in the yield strength to the solid solution hardening and the decrease in tensile strength to the low strain hardenability, due to higher stacking fault energy (SFE) by Al addition. For cold rolled uncharged specimens, however, the YS decreased, but the elongation increased with an increase in the Al addition.

The dynamic hall petch strengthening due to twins is weaker in Al added steels, leading to higher elongation and lower YS with an increasing amount of Al in the steel. In H charged specimens, HE susceptibility decreased with an increase in Al content, due to the decrease in the dislocation density. With the increase in Al content, the residual stress reduces and strong <111> and <100> textures develop, which decreases the driving force for the H diffusion and arrests the cracks at specific grain boundaries, leading to enhanced embrittlement resistance.

© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon

Ronevich et al. [54] reported limited hydrogen sensitivity of TWIP steels, both with and without Al addition. They charged the TWIP steel (0.6C-18Mn-1.5Al) with hydrogen by applying cathodic current density in acidic solution. H had a negligible effect on yield strength, tensile strength, and elongation of the tensile specimen. However, in their test, the specimen thickness was too high for the charging period (1.2 mm thickness and the maximum charging period is 6 hours). As a result, the hydrogen remained at the surface and did not affect the bulk specimen.

Park et al. [START_REF] Park | The advantage of grain refinement in the hydrogen embrittlement of Fe-18Mn-0.6C twinning-induced plasticity steel[END_REF] investigated the effect of grain size on the embrittlement behavior of Fe-18Mn-0.6C TWIP steels. Increasing grain size decreased HE resistance. For smaller grain size (< 20 µm), the specimens showed ductile failure at all regions. However, with larger grain size, the tensile specimen edges exhibit a quasi-cleavage fracture. They attributed the higher HE tendency of large grain specimens to the more active mechanical twins in a coarsegrained structure. They suggested that the crack is generated at the twin boundaries, twintwin junction, and twin grain boundary junctions and propagate along the twins and the grain boundaries, making the alloy more susceptible to HE.

Hydrogen entry in steels

The presence of free interstitial sites and the small size of H enables it to diffuse freely through the metals. Hydrogen can enter the metal either through a gaseous environment or by electrochemical methods. The entry of hydrogen through the electrochemical process is quite useful as it can sustain the high equivalent pressure of hydrogen gas at the metal surface [START_REF] Perng | A brief review note on mechanisms of hydrogen entry into metals[END_REF]. In the case of steels for automotive applications, electrochemical corrosion results in H2 evolution, introducing H in the steels. The dissolved hydrogen concentration determines the efficiency of hydrogen entry from the environment, at and just beneath the metal surface, where it forms a solid solution with the metal. Ideally, hydrogen is distributed randomly in the metal lattice in the absence of any inhomogeneity but is distributed preferentially in case of microstructural inhomogeneity. This distribution of hydrogen determines its role in the HE mechanism.

© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon

In order to understand the influence of hydrogen on the AHSS, it is essential to know about (a) H evolution and adsorption on steel surface, (b) absorption into the metal, (c) transportation through the lattice and (d) interaction with different microstructural features.

2𝑀𝑀𝐻𝐻 𝑎𝑎𝑎𝑎𝑎𝑎 ↔ 2𝑀𝑀𝐻𝐻 𝑎𝑎𝑎𝑎𝑎𝑎 ( 6
)
Where M is the metal surface, MHads is the hydrogen adsorbed on the surface, MHabs is the hydrogen absorbed in the metal. The absorbed hydrogen can either be released from the surface electrochemically according to reaction (3) and ( 4), chemically by reaction (5) or absorbed by the material as per reaction (6).

Hydrogen absorption in steels

Hydrogen evolution and absorption into the metal are two competing reactions.

In the case of hydrogen embrittlement, one is concerned with the absorbed hydrogen. Usually, the material absorbs only a small fraction of hydrogen reduced at the surface. Sometimes poisons such as sulfides, cyanides, and arsenic ions are used to retard the recombination reaction and promote hydrogen absorption in hydrogen research. The factors that dictate the above reactions are the surface coverage by hydrogen, the surface potential, the distance between two adsorbed hydrogen, temperature, and cathodic current density. The literature proposes two models for the entry of hydrogen into a metal.

Bockris and his co-workers developed the first model, where they proposed hydrogen absorption reaction and the hydrogen recombination processes as the competing reactions [START_REF] Bockris | Recent Developments in the Study of Hydrogen Overpotential[END_REF], [START_REF] Flitt | Hydrogen/metal interactions with special reference to electrochemical approaches[END_REF]. In other words, the nature of hydrogen on the metal surface is identical for both the recombination and absorption reactions. In this model, the fraction of hydrogen that enters the metal is proportional to the coverage of the metal by adsorbed hydrogen. The other model states that the nature of hydrogen, which undergoes absorption and recombination, are not identical and hence are not competitive and occur independently to each other [START_REF] Pickering | A Review of the Mechanism and Kinetics of Electrochemical Hydrogen Entry and Degradation of Metallic Systems[END_REF].

© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon

Hydrogen diffusion behavior in lattice

The hydrogen entry efficiency is characterized by the concentration of hydrogen just beneath the metal surface [START_REF] Zakroczymski | Adaptation of the electrochemical permeation technique for studying entry, transport and trapping of hydrogen in metals[END_REF], [START_REF] Moro | Hydrogen embrittlement susceptibility of a high strength steel X80[END_REF]. Hydrogen in solid solution can be distributed either randomly among the atoms in a lattice or can be attached to microstructural defects.

The freely moving hydrogen in the lattice obeys Fick's law of diffusion and is termed as diffusible hydrogen. On the contrary, the hydrogen associated with microstructural defects

show reduced mobility and is tightly bound or trapped. Darken and Smith first observed the H diffusivity in steels impeded by the lattice imperfections in a cold-rolled steel and was later confirmed by several researchers [START_REF] Asahi | Hydrogen Trapping Behavior in Vanadium-added Steel[END_REF]- [START_REF] Darken | Behavior of Hydrogen in Steel During and After Immersion in Acid[END_REF]. Hydrogen diffusion is a concentration driven process where hydrogen jumps between interstitial lattice sites [START_REF] Liu | Reversible hydrogen trapping in a 3.5NiCrMoV medium strength steel[END_REF]. The rate of hopping of hydrogen from one interstitial site to the other is influenced by defects (or commonly known as traps) in the microstructural sites such as grain boundaries, vacancies, and interfaces [START_REF] Liu | Reversible hydrogen trapping in a 3.5NiCrMoV medium strength steel[END_REF].

As traps provide a favorable environment for the hydrogen atom, there is a reduction in energy when a hydrogen atom leaves the lattice position to enter in a trap [7].

Lee et al. [START_REF] Lee | crystals structures by the desorption thermal analysis technique[END_REF], [START_REF] Lee | The trapping and transport phenomena of hydrogen in nickel[END_REF] suggested that the trapping behavior can be well understood by characterizing the trap activation energy of hydrogen (Ea), the trap binding energy (Eb), the density of traps (N), the and saddle point energy (Es). A trap with high activation energy will behave as an irreversible trap, but if the saddle point energy is much higher than the activation energy of hydrogen diffusion through the regular lattice, both the trapping and detrapping process is complicated. For example, incoherent TiC particles are known to be a strong-traps, but as H must overcome an energy barrier (saddle point energy) of 60-90 kJ.mol - 1 to enter the traps, the hydrogen trapping at the carbon vacancies in the TiC particles only occurs at a relatively high temperature [START_REF] Winzer | Hydrogen diffusion and trapping in Ti-modified advanced high strength steels[END_REF]. On the other hand, if the saddle point energy, for example of grain boundaries, is lower than the activation energy of the lattice diffusion, the trapping and detrapping of hydrogen should be easy. weak trapping site [START_REF] Lee | crystals structures by the desorption thermal analysis technique[END_REF]. Sn is the normal lattice site, and ST is the trap site. EaD is the activation energy for lattice diffusion, Ea is the trap activation energy for the escape of hydrogen, Eb is the trap binding energy, and Es is the saddle point energy.

Trapping determines the hydrogen distribution in the microstructure, and the hydrogen embrittlement process is closely related to the hydrogen-trap interaction [START_REF] Park | Effect of microstructure on the hydrogen trapping efficiency and hydrogen induced cracking of linepipe steel[END_REF].

Based in order of increasing trap activation energy (Ea) [START_REF] Lynch | Hydrogen embrittlement phenomena and mechanisms[END_REF], the potency of trapping effect can be ranked such as (a) some solute atoms, (b) free surfaces and sites between the first few atomic layers beneath the surface, (c) mono-vacancies and vacancy clusters (which are present well above thermal equilibrium values as hydrogen reduces vacancy forming energy), (d) dislocation cores and strain fields, (e) grain boundaries, (f) precipitation/matrix interface and strain fields around precipitates, (g) inclusion/matrix interface and (h) voids and internal cracks. Some of the trapping sites that can be present in steels are schematically summarized in Figure 2.7 [61]. Traps not only reduce the H mobility in steels but also increase the H uptake when the H concentration that enters the steel is high [7]. As the steel is in physiochemical equilibrium with a fixed external chemical potential of hydrogen, it will absorb hydrogen up to the solubility limit of the lattice, while an additional amount of hydrogen will occupy the traps, resulting in increased H uptake [START_REF] Krom | Hydrogen trapping models in steel[END_REF]. Equilibrium establishes when the chemical potentials of the hydrogen in lattice and traps and the external chemical potential scale [START_REF] Pundt | HYDROGEN IN METALS: Microstructural Aspects[END_REF] Reversible traps are particularly dangerous in terms of HE susceptibility. As the H trapped in such sites can be detrapped relatively quickly, they act as hydrogen sources, supplying hydrogen in the lattice, which can subsequently diffuse towards the stress field, increasing the hydrogen concentration and detrimentally altering the fracture properties [7]. Grain boundaries, dislocations, coherent interface between precipitate-matrix interfaces are some of the reversible traps present in the material. The susceptibility of steel to HE increases with an increase in the coherent boundaries in the microstructure, supporting the above suggestions [52].

In literature, there is a disagreement on the influence of irreversible traps in the HE susceptibility. Some published articles report that the presence of irreversible traps leads to increased susceptibility to HE in steels [START_REF] Michler | Hydrogen environment embrittlement of an ODS RAF steel -Role of irreversible hydrogen trap sites[END_REF], [START_REF] Huang | Effect of microstructure and inclusions on hydrogen induced cracking susceptibility and hydrogen trapping efficiency of X120 pipeline steel[END_REF], [START_REF] Nakatani | Fatigue crack growth acceleration caused by irreversible hydrogen desorption in high-strength steel and its mechanical condition[END_REF]. However, others believe that the presence of irreversible traps decreases the susceptibility to HE as they reduce the amount of diffusible hydrogen present in the lattice [7], [START_REF] Hirth | Effects of hydrogen on the properties of iron and steel[END_REF]- [START_REF] Tsay | Hydrogen embrittlement susceptibility and permeability of two ultra-high strength steels[END_REF]. Lynch [START_REF] Lynch | Hydrogen embrittlement phenomena and mechanisms[END_REF] believes that microstructure with strong traps that are easily fractured are particularly susceptible to HE as they not only behave as the hydrogen concentration sites but also crack initiation sites.

However, numerous moderate to strong traps, which do not initiate cracks, reduce the quantity of diffusible hydrogen, and increase HE resistance [START_REF] Lynch | Hydrogen embrittlement phenomena and mechanisms[END_REF]. Bhadeshia [7] explains that though the presence of strong-trap increases the H concentration, but will not contribute to HE. Instead, he suggests to use the strong traps to reduce the mobile hydrogen present in the material as a material design route to improve hydrogen compatibility. H transport by dislocations is essential in moving hydrogen from the grain volume to the grain boundaries, causing intergranular fracture [START_REF] Lynch | Hydrogen embrittlement phenomena and mechanisms[END_REF]. The enhanced H transport by dislocations lead to greater H mobility within the metal than the typical diffusion process [9].

Hydrogen is attracted to the dislocation stress field, especially the triaxial stress present in the edge dislocation, and the dislocation core traps H along the length of the dislocation line [9]. Hydrogen can also enhance the dislocation motion in two possible ways. It can lower the dislocation core energy to reduce the Peierls energy, initiating the dislocation motion [9], [62].

H can also shield the dislocation from the lattice defect's elastic interactions, reducing dislocation motion resistance [START_REF] Birnbaum | Hydrogen-enhanced localized plasticity--a mechanism for hydrogenrelated fracture[END_REF]. Dislocations can result in multiple trap sites with different trapping energies. The binding energy around the edge dislocation is of the range between 28-38 kJ.mol -1, whereas the trapping energy at the dislocation core can be as high as 58-67 kJ.mol -1, which leads to strong hydrogen trapping [START_REF] Frappart | Hydrogen trapping in martensitic steel investigated using electrochemical permeation and thermal desorption spectroscopy[END_REF]- [START_REF] Taketomi | Atomistic study of hydrogen distribution and diffusion around a {112}<111> edge dislocation in alpha iron[END_REF].

Grain boundary:

Grain boundary has a significant influence on the interaction of hydrogen with metals [START_REF] Mcmahon | Hydrogen-induced intergranular fracture of steels[END_REF]. The role of grain boundaries on the H diffusion behavior is somewhat divided.

Though there are ample evidences to show that grain boundary retards the H diffusion due to the trapping effect [START_REF] Aucouturier | Grain Boundary Segregations and Hydrogen Embrittlement[END_REF]- [START_REF] Koyama | In situ observations of silverdecoration evolution under hydrogen permeation: Effects of grain boundary misorientation on hydrogen flux in pure iron[END_REF], some researchers argue that the grain boundaries provide a fast diffusion path to the hydrogen [START_REF] Lynch | Hydrogen embrittlement phenomena and mechanisms[END_REF], [START_REF] Harris | Grain boundary diffusion of hydrogen in nickel[END_REF], [START_REF] Gavriljuk | Comments to the paper "in situ observations of silver-decoration evolution under hydrogen permeation: Effects of grain boundary misorientation on hydrogen flux in pure iron[END_REF]. Lynch [START_REF] Lynch | Hydrogen embrittlement phenomena and mechanisms[END_REF] proposed that the H diffusion of hydrogen may be faster in the grain boundaries in comparison with lattice in some instances, but when the trap density becomes higher at the boundaries, the traps retard the H diffusion.

Carbides or second phase precipitates:

Asahi et al. [START_REF] Asahi | Hydrogen Trapping Behavior in Vanadium-added Steel[END_REF] measured the activation energy for vanadium carbide (VC) to be around 33-35 kJ.mol -1 . They observed that the hydrogen trapping was rapid, but the detrapping was small. One probable reason for such observation could be that the saddle point energy of such traps could be lower as compared to the lattice diffusion activation energy with © 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon large binding energy. Ohnuma et al. [START_REF] Ohnuma | Direct observation of hydrogen trapped by NbC in steel using small-angle neutron scattering[END_REF] studied the trapping behavior of NbC and found them to be effective hydrogen traps. Many researchers have also studied the trapping behavior of TiC, and all concluded that TiC is effective traps with H trapped firmly at the precipitate/matrix interface [START_REF] Escobar | Thermal desorption spectroscopy study of the interaction of hydrogen with TiC precipitates[END_REF], [START_REF] Pressouyre | An example of the effect of hydrogen trapping on hydrogen embrittlement[END_REF], [START_REF] Lee | crystals structures by the desorption thermal analysis technique[END_REF]. Lee et al. [START_REF] Lee | crystals structures by the desorption thermal analysis technique[END_REF] estimated the activation energy of TiC traps to be around 87 kJ.mol -1 .

2.5 Evaluation of Hydrogen Detection Techniques

Permeation Test

As hydrogen diffusion has a direct impact on the transport of hydrogen to the critical areas, hydrogen permeation is an essential parameter in the overall embrittlement process. The hydrogen permeation technique is a convenient and accurate method to determine hydrogen flux in a metal, with time [START_REF] Zakroczymski | Adaptation of the electrochemical permeation technique for studying entry, transport and trapping of hydrogen in metals[END_REF]. In this electrochemical arrangement, hydrogen is first adsorbed on to a metallic membrane surface, followed by absorption and diffusion, and oxidation at the exit side. This technique, developed by Devanathan and

Starchuski [START_REF] Devanathan | The Adsorption and Diffusion of Electrolytic Hydrogen in Palladium[END_REF], is based on certain boundary conditions such as:

(a) The permeation of hydrogen should be diffusion controlled and not be affected by other factors such as surface effects, and (b) The concentration of hydrogen atoms at the input side should be constant during the test.

The concentration of hydrogen at the exit side should be zero, i.e., the hydrogen that reaches the exit surface of the specimen after diffusion should be immediately oxidized.

The permeation technique assumes the immediate establishment of the boundary conditions, and the permeation current density measured is only diffusion controlled. It enables the determination of H diffusion coefficient (D) in the membrane and, subsequently, the subsurface hydrogen concentration (C0). The techniques applied to calculate different permeation parameters are discussed later in the section 2.5.1.3.

However, in typical situations, most of these assumptions are not fulfilled due to the slow surface kinetic processes related to hydrogen entry [START_REF] Turnbull | Perspectives on hydrogen uptake, diffusion and trapping[END_REF], [START_REF] Chaudhari | Boundary conditions during the electropermeation of hydrogen through pure iron[END_REF], and hydrogen © 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon trapping during the charging of hydrogen-free specimens as shown in Figure 2.8. It leads to a wide scattering in published diffusion data [START_REF] Haq | Effect of microstructure and composition on hydrogen permeation in X70 pipeline steels[END_REF]. Several models attempt to explain the discrepancies between the experimental and theoretical curves [69], [START_REF] Mcnabb | A new analysis of the diffusion of hydrogen in iron and ferritic steels[END_REF], [START_REF] Kumnick | Hydrogen transport through annealed and deformed armco iron[END_REF]. However, in these models, it was necessary to make certain assumptions and include some new variables [START_REF] Zakroczymski | Adaptation of the electrochemical permeation technique for studying entry, transport and trapping of hydrogen in metals[END_REF]. A modified permeation technique involving partial H charging and decay in the specimen, to minimize the effect to surface changes and traps, is developed to make the process indeed diffusion-controlled [START_REF] Zakroczymski | Adaptation of the electrochemical permeation technique for studying entry, transport and trapping of hydrogen in metals[END_REF]. does not influence the kinetic process of diffusion in the specimen [START_REF] Svoboda | Determination of trapping parameters and the chemical diffusion coefficient from hydrogen permeation experiments[END_REF].

The role of Pd coating has been the subject of discussion in the past with permeation tests performed: (a) without any Pd layer [START_REF] Su-Il Pyun | The permeation of hydrogen through the passivating films on iron and nickel[END_REF], [START_REF] Szklarska-Smialowska | Effect of EDTA on the Cathodic Reduction of Oxide Films on Iron in Sodium Hydroxide Solution[END_REF] (b). with the Pd layer at the detection side [START_REF] Chaudhari | Boundary conditions during the electropermeation of hydrogen through pure iron[END_REF] (c). Pd layer at the charging side [START_REF] Su-Il Pyun | The permeation of hydrogen through the passivating films on iron and nickel[END_REF] and (d). Pd layer at both sides [START_REF] Casanova | The influence of an oxide layer on hydrogen permeation through steel[END_REF],

[151]. With the Pd layer at the charging side, the steady-state current decreased as the H evolution occurred in the Pd layer instead of the iron layer which possibly retarded the reduction kinetics and subsequent H coverage on the surface, changing the absorption efficiency of hydrogen into the material [START_REF] Zakroczymski | Adaptation of the electrochemical permeation technique for studying entry, transport and trapping of hydrogen in metals[END_REF], [149][150]. However, it is important to make sure that the Pd layer is stable during the testing process; otherwise, it could influence the obtained permeation current.

Measurement of Hydrogen Diffusivity by Permeation Technique

The H diffusivity and solubility of hydrogen in steels depend on the phases present, as shown in Figure 2.10. Austenite has higher H solubility for but has a lower diffusivity. On the other hand, ferrite has higher diffusivity but lower solubility of hydrogen. The presence of defects can decrease the diffusivity of hydrogen, as discussed above. Diffusivity is the movement of hydrogen from one lattice point to the other. In an anisotropic medium, Fick's first law of diffusion can represent the diffusion flux (J) as:

𝐽𝐽 = -𝐷𝐷 𝜕𝜕𝜕𝜕 𝜕𝜕𝜕𝜕 Equation 2.1
Where 'D' is the diffusion coefficient in m².s -1 . and austenite [7] and (b) the Fe-H phase diagram showing the solubility of hydrogen in different phases [START_REF] Zinkevich | Thermodynamics of Fe-Sm, Fe-H, and H-Sm systems and its application to the hydrogen-disproportionation-desorption-recombination (HDDR) process for the system Fe17Sm2-H2[END_REF].

© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon

The 'D' can be experimentally determined by:

(a) The time lag method has been used in the literature to calculate the apparent diffusion coefficient of hydrogen in the specimens [START_REF] Zhang | Effect of the cathodic current density on the sub-surface concentration of hydrogen in X80 pipeline steels under cathodic protection[END_REF], [START_REF] Luppo | The influence of microstructure on the trapping and diffusion of hydrogen in a low carbon steel[END_REF], [START_REF] Svoboda | Determination of trapping parameters and the chemical diffusion coefficient from hydrogen permeation experiments[END_REF], [START_REF] Addach | Study of the electrochemical permeation of hydrogen in iron[END_REF].

𝐷𝐷 = 𝐿𝐿 2 6𝑡𝑡 0.63 Equation 2.2
Where L is the specimen thickness, the time lag t0.63 is the elapsed charging time when the normalized permeation current density is 0.63 (s), and D is the diffusion coefficient (m 2 .s -1 ).

(b)

The breakthrough time method is the time required for hydrogen to reach the output side.

Breakthrough time is defined by the intersection point of the time axis with the tangent drawn at the inflection point of the permeation curve. [START_REF] Luppo | The influence of microstructure on the trapping and diffusion of hydrogen in a low carbon steel[END_REF], [START_REF] Grabke | Absorption and Diffusion of hydrogen in Steels[END_REF], [START_REF] Frappart | Hydrogen trapping in martensitic steel investigated using electrochemical permeation and thermal desorption spectroscopy[END_REF], [START_REF] Garcia | Influence of microstructure in the hydrogen permeation in martensitic-ferritic stainless steel[END_REF].

𝐷𝐷 = 𝐿𝐿 2 15𝑡𝑡 𝑏𝑏 Equation 2.3 (a) (b) © 2020 -Dwaipayan MALLICK -Lab.

MATEIS, INSA Lyon

Where L is the specimen thickness (m), tb is the breakthrough time (s), and D is the diffusion coefficient (m 2 .s -1 ). (c) In The Permeation Transient Method, the experimentally observed permeation curves are fitted (both the rise and fall transients) assuming that the boundary conditions for the permeation measurements are satisfied [START_REF] Grabke | Absorption and Diffusion of hydrogen in Steels[END_REF], [START_REF] Garcia | Influence of microstructure in the hydrogen permeation in martensitic-ferritic stainless steel[END_REF] Rise Transient:

𝑖𝑖 𝑛𝑛 = 2𝐿𝐿 √𝜋𝜋𝜋𝜋𝑡𝑡 ∑ exp (- (2𝑛𝑛+1) 2 𝐿𝐿 2 4𝜋𝜋𝑡𝑡 ) ∞ 𝑛𝑛=0 Equation 2.4
Fall Transient:

𝑖𝑖 𝑛𝑛 = 1 - 2𝐿𝐿 √𝜋𝜋𝜋𝜋𝑡𝑡 ∑ exp �- (2𝑛𝑛+1) 2 𝐿𝐿 2 4𝜋𝜋𝑡𝑡 � ∞ 𝑛𝑛=0 Equation 2.5
Where L is the specimen thickness (m), t is time (s), D is the diffusion coefficient (m 2 .s -1 ), and in is the normalized current density.

© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon

H diffusion coefficient in pure metals at ambient temperatures depend on the crystal structure with hydrogen diffusion coefficient, generally 4 to 5 times higher for bodycentered cubic (bcc) structure as compared to face-centered cubic (fcc) structure and hexagonal close-packed (hcp) structure (with an exception for Pd (fcc) and Co(hcp) ) [START_REF] Lynch | Hydrogen embrittlement phenomena and mechanisms[END_REF], [START_REF] Volkl | 5 -Hydrogen Diffusion in Metals[END_REF]. In alloys with complex microstructure, the D decreases with an increase in the number of traps.

Trapping is not expected to affect the steady-state current density but tends to slow down the diffusivity of hydrogen [START_REF] Liu | Reversible hydrogen trapping in a 3.5NiCrMoV medium strength steel[END_REF]. Hydrogen diffuses through the lattice in response to gradients in (a) hydrogen concentration, (b) temperature, and (c) hydrostatic stress fields [START_REF] Lynch | Hydrogen embrittlement phenomena and mechanisms[END_REF]. The apparent diffusion coefficient is the representation of the combined effect of the lattice diffusion coefficient and the trapped hydrogen [START_REF] Luppo | The influence of microstructure on the trapping and diffusion of hydrogen in a low carbon steel[END_REF]. Using successive transients can help in rapid data accumulation and to obtain statistically improved analysis compared to the single transient [START_REF] Early | Hydrogen diffusion in palladium by galvanostatic charging[END_REF].

Hydrogen permeation technique has been used to determine diffusion coefficients of several metals and alloys, particularly steels. Zacroczymski [START_REF] Zakroczymski | Adaptation of the electrochemical permeation technique for studying entry, transport and trapping of hydrogen in metals[END_REF] performed permeation measurements on Armco iron and reported the diffusion coefficient of hydrogen in ferrite to be 7.3 X 10 -9 m².s -1 . Koyama et al. [43] reported the diffusion coefficient of hydrogen in TWIP steel (austenite) to be 1.29 X 10 -13 m².s -1 . Several multiphase steels have been investigated in the literature to determine the diffusivity of hydrogen in complex microstructures with values reported to lie between that of ferrite and austenite steels [START_REF] Luppo | The influence of microstructure on the trapping and diffusion of hydrogen in a low carbon steel[END_REF], [START_REF] Ichitani | Visualization of hydrogen diffusion in steels by high sensitivity hydrogen microprint technique[END_REF], [START_REF] Luu | The influence of microstructure on hydrogen transport in carbon steels[END_REF].

Measurement of Sub Surface Concentration of Hydrogen

As mentioned previously, hydrogen entry efficiency is characterized by the H concentration just beneath the metal surface. The diffusion coefficient can be used to calculate the H concentration at the subsurface of the entry side. The hydrogen concentration at the charging side is evaluated from the diffusion coefficient and steady-state current density as [START_REF] Zhang | Effect of the cathodic current density on the sub-surface concentration of hydrogen in X80 pipeline steels under cathodic protection[END_REF], [START_REF] Liu | Reversible hydrogen trapping in a 3.5NiCrMoV medium strength steel[END_REF], [START_REF] Zhao | Hydrogen permeation and embrittlement susceptibility of X80 welded joint under high-pressure coal gas environment[END_REF]:

© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon 𝐶𝐶 0 = 𝑖𝑖 𝑠𝑠𝑠𝑠 * 𝐿𝐿 𝐹𝐹 * 𝜋𝜋 Equation 2.6
Where L is the specimen thickness, D is the H diffusion coefficient, iss is the steadystate current density, 𝐶𝐶 0 is the H concentration at the charging subsurface, and F is the Faraday constant.

Trap Density Measurement

Measuring the difference in the area of the 'fall transient' between the experimental curve and the corresponding theoretical fit, the H concentration released from the reversible traps can be calculated. With an assumption that each trap can hold only one hydrogen atom, the trap density in a specimen is determined using Equation 2.7 [START_REF] Liu | Reversible hydrogen trapping in a 3.5NiCrMoV medium strength steel[END_REF].

𝑁𝑁 𝑇𝑇 = 2𝐴𝐴 * 6.24 * 10 18 𝐿𝐿 Equation 2.7
Where, 𝑁𝑁 𝑇𝑇 is the trap density of the material, and 'A' is the difference in the area between the theoretical and the experimental curve.

Thermal Desorption Spectroscopy (TDS) Measurement

Hydrogen permeation experiments can provide information related to the measured hydrogen flux at the discharge side of a material but do not provide direct information on the traps and their activity in it [START_REF] Svoboda | Determination of trapping parameters and the chemical diffusion coefficient from hydrogen permeation experiments[END_REF]. Thermal Desorption Spectroscopy (TDS) can help gain important information related to the activation energy of the traps present in the material as well as quantify the amount of hydrogen trapped in reversible and irreversible specimens. TDS was initially developed as a technique to study the adsorption and desorption kinetics of adsorbed layer, its binding states, and thermal evaluation of the adsorbed layer and soon gained prominence in the steel research [START_REF] Escobar | Influence of experimental parameters on thermal desorption spectroscopy measurements during evaluation of hydrogen trapping[END_REF].

TDS is a non-isothermal technique where a specimen pre-charged with hydrogen is exposed to linearly increasing temperature, and a mass spectrometer measuring the desorbed hydrogen [START_REF] Escobar | Influence of experimental parameters on thermal desorption spectroscopy measurements during evaluation of hydrogen trapping[END_REF]. The TDS data obtained contains a plot of the mass spectrometer Nevertheless, the distinction of different trapping sites using TDS is not as straightforward due to (a) overlapping of temperature regions from different kinds of traps sites. For example, the activation energy for the grain boundaries, dislocations, and microvoids often overlap; [START_REF] Louthan | Hydrogen Embrittlement of Metals: A Primer for the Failure Analyst[END_REF], [START_REF] Svoboda | Determination of trapping parameters and the chemical diffusion coefficient from hydrogen permeation experiments[END_REF], [START_REF] Darken | Behavior of Hydrogen in Steel During and After Immersion in Acid[END_REF], [START_REF] Casanova | The influence of an oxide layer on hydrogen permeation through steel[END_REF], and (b) heating the specimens during TDS can affect on the microstructure and its trapping capabilities [START_REF] Escobar | Influence of experimental parameters on thermal desorption spectroscopy measurements during evaluation of hydrogen trapping[END_REF]. The specimen microstructure is assumed to be stable during the TDS analysis. It could be valid for stable sites such as grain boundaries, impurity atoms, or second phase precipitates [START_REF] Nagumo | Nature of hydrogen trapping sites in steels induced by plastic deformation[END_REF]. However, defects that can agglomerate or annihilate at relatively low temperatures, the heating has a significant effect and must be taken into account [START_REF] Nagumo | Nature of hydrogen trapping sites in steels induced by plastic deformation[END_REF]. TDS data does not take into consideration the specimen geometry, which becomes particularly important when the trap activation energy is similar to that of the energy required for lattice diffusion. In such cases, a diffusion-controlled hydrogen evolution peak is observed. Thus, in TDS measurements, the nature of the peak will be determined by the overall rate-controlling step [START_REF] Lee | crystals structures by the desorption thermal analysis technique[END_REF].

Calculation of Activation Energy (Ea) of Traps

With an assumption that the hydrogen escaping from a trap site is a thermally activated process, the hydrogen evolution rate from a trapping site was proposed by Kissinger as [START_REF] Kissinger | Reaction Kinetics in Differential Thermal Analysis[END_REF] 𝑎𝑎𝑋𝑋

𝑇𝑇 𝑎𝑎𝑡𝑡 = 𝐴𝐴(1 -𝑋𝑋 𝑇𝑇 ) exp � -𝐸𝐸 𝑎𝑎 𝑅𝑅𝑇𝑇 � Equation 2.8
Where XT is the fraction of the hydrogen released from the traps at a given time 't,' A is the reaction constant, Ea is the activation energy of the traps, R is the gas constant, and T is the absolute temperature. The parameter A is athermal part of the frequency factor © 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon in the equation and represents the probability per unit time of successful desorption of hydrogen from a specific trap [START_REF] Escobar | Thermal desorption spectroscopy study of the interaction of hydrogen with TiC precipitates[END_REF]. The exponential term represents the probability of hydrogen escaping from the trap sites to a normal lattice site.

The above equation was simplified by Lee et al. [START_REF] Lee | crystals structures by the desorption thermal analysis technique[END_REF], [START_REF] Lee | The trapping and transport phenomena of hydrogen in nickel[END_REF] to calculate the activation energy of the traps present in the material as

∂�ln� ϕ T c 2 �� ∂� 1 T c � = - E a R Equation 2.9
Where, ϕ is the heating rate, and Tc is the peak temperature.

The above equation is widely used to determine the activation energy of the traps present in materials [START_REF] Escobar | Thermal desorption spectroscopy study of the interaction of hydrogen with TiC precipitates[END_REF], [START_REF] Asahi | Hydrogen Trapping Behavior in Vanadium-added Steel[END_REF]. The trap activation energy Ea can be calculated from the slopes of ln(ϕ /Tc²) versus 1/Tc. However, if the activation energy of lattice diffusion is comparable to that of the trap, diffusion-controlled hydrogen evolution will be measured.

Thus, in TDS, the character of the hydrogen evolution is determined by the rate-controlling step of the overall reaction [START_REF] Lee | crystals structures by the desorption thermal analysis technique[END_REF].

Calculation of Diffusion Coefficient using TDS

In literature, there are two techniques proposed to determine the diffusion coefficient of the specimen using TDS measurements. The first, proposed by Lee et al. [START_REF] Lee | The trapping and transport phenomena of hydrogen in nickel[END_REF],

determines the diffusion coefficient of the specimen using the following equation:

ln �1 - 𝑄𝑄 𝑡𝑡 𝑄𝑄 ∞ � = ln 8 𝜋𝜋 2 -𝐷𝐷𝜋𝜋 2 𝑡𝑡 𝐿𝐿 2 ⁄ Equation 2.10
Where, 𝑄𝑄 𝑡𝑡 = the amount of hydrogen absorbed in time 't'; 𝑄𝑄 ∞ = the total amount of hydrogen in a fully saturated specimen; D is the diffusion coefficient of the specimen, and 'L' is the thickness of the specimen.

Thus, for a fixed thickness and charging temperature, the slope of the plot of ln{1-𝑄𝑄 𝑡𝑡 /𝑄𝑄 ∞ )} as the function of time, the diffusion coefficient can be determined. In another © 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon method [START_REF] Iacoviello | A thermal outgassing method (t.o.m.) To measure the hydrogen diffusion coefficients in austenitic, austeno-ferritic and ferritic-perlitic steels[END_REF], the specimens are isothermally heated to determine the H diffusivity in the specimen. The time required to desorb half of the total hydrogen present in the specimens (𝑡𝑡 1 2 ⁄ ) is measured and the diffusion coefficient is calculated using the following equation:

𝐷𝐷 = 0.0492 𝐿𝐿² 𝑡𝑡 1 2 ⁄ Equation 2.11
Where the terms have their usual meaning.

Choo et al. [START_REF] Choo | Thermal analysis of trapped hydrogen in pure iron[END_REF] calculated the activation energies of different types of traps in Fe.

They observed three distinct peaks at 385, 488, and 578 K on heating the specimen with a heating rate of 156 K/h. In order to relate the traps with the data available from TDS measurements, they divided the subsequent tests into three different parts. In the first stage, they took specimens with different grain sizes and kept the degree of dislocations and microvoids constant by controlling the extent of cold work. They observed that with an increase in grain boundary area, the area under the peak at 385 K increased while the other peaks remained almost constant. Since the area under the curve represents the concentration of hydrogen, they concluded that the peak at 385 K corresponded to the hydrogen released from the grain boundaries. Similarly, they tried to vary the dislocation density keeping other factors constant. They found that the peak at 488 K corresponded to the hydrogen released from the dislocations, and the peak at 578 K corresponded to the release of hydrogen from microvoids. The activation energy calculated for grain boundaries, dislocations, and microvoids was 17.2, 26.8, and 35.2 kJ.mol -1 , respectively.

Escobar et al. [START_REF] Escobar | Influence of experimental parameters on thermal desorption spectroscopy measurements during evaluation of hydrogen trapping[END_REF] studied the influence of various experimental parameters on H trapping using the TDS method. They observed that the influence of electrolyte on the lowtemperature peak is less significant but showed an additional high-temperature peak with arsenic in the solution. The peak height, however, changed with electrolyte, as the amount of hydrogen absorbed in the specimen is known to be dependent on the charging environment.

They observed that the peak temperature shifted towards higher temperature values on increasing the specimen thickness for low energy traps, whereas the high energy trap's peak temperature was independent of the thickness. It is consistent with the observations of Lee et al. [START_REF] Lee | The trapping and transport phenomena of hydrogen in nickel[END_REF] that when the trap activation energy and the lattice diffusion energy are comparable,

© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon

the peak observed is the combined effect of the two processes mentioned above. Also, the polished specimen absorbed more hydrogen as compared to the ground specimens as polished specimens lead to higher charging current density (ground surface have higher surface area due to surface roughness as compared to the polished surface) in a galvanostatically controlled cathodic charging condition.

Hydrogen Microprinting/Ag Decoration Technique

Hydrogen diffusion coefficient, trapping parameters, and hydrogen content in irreversible and reversible traps are extensively studied using the hydrogen permeation cell and TDS techniques. However, the tests do not provide much information on the H diffusion path and the desorption sites. There is a need for visual understanding and identifying the trapping sites in steels to compliment the above studies. Various methods have been developed

for the visual examination of hydrogen on the surface of the steels such as tritium autoradiography, hydrogen microprint technique (HMT), Ag decoration technique, Scanning Kelvin Probe and secondary ion mass spectroscopy (SIMS) [START_REF] Nazarov | Scanning Kelvin Probe for detection of the hydrogen induced by atmospheric corrosion of ultra-high strength steel[END_REF], [START_REF] Akiyama | Hydrogen Visualization in Steels Using Ag Decoration Method[END_REF], [START_REF] Takai | Visualization of the hydrogen desorption process from ferrite, pearlite, and graphite by secondary ion mass spectrometry[END_REF].

HMT was earlier developed Ovejero-Gracia [START_REF] Pérez | Direct observation of hydrogen evolution in the electron microscope scale[END_REF] and later improved by Ichitani et al. [START_REF] Ichitani | Visualization of hydrogen diffusion in steels by high sensitivity hydrogen microprint technique[END_REF]. Hydrogen microprinting technique (HMT) and Ag decoration are similar techniques, which take advantage of the reduction effect of hydrogen atom released from the specimen surface, and hydrogen exit sites are visualized, as H reduces the Ag particles. In Ag decoration test, the reduction of Ag + by the absorbed hydrogen (𝐻𝐻 𝑎𝑎𝑎𝑎𝑎𝑎 ) on the metal surface, leads to Ag particle precipitates [START_REF] Akiyama | Hydrogen Visualization in Steels Using Ag Decoration Method[END_REF]. The H distribution, thus, can be visualized by Ag particle distribution in the specimen.

𝐻𝐻 𝑎𝑎𝑎𝑎𝑎𝑎 + 𝐴𝐴𝐴𝐴 + → 𝐻𝐻 + + 𝐴𝐴𝐴𝐴 ↓ (7)
The HMT or Ag decoration technique (here the terms would be used interchangeably) is investigated in different steels such as low carbon steels [START_REF] Luppo | The influence of microstructure on the trapping and diffusion of hydrogen in a low carbon steel[END_REF], mild steel [START_REF] Luu | The influence of microstructure on hydrogen transport in carbon steels[END_REF], ferritic steels [START_REF] Koyama | In situ observations of silverdecoration evolution under hydrogen permeation: Effects of grain boundary misorientation on hydrogen flux in pure iron[END_REF], hypo and hypereutectoid steels [START_REF] Matsuda | Visualization of hydrogen diffusion path by a high sensitivity hydrogen microprint technique[END_REF], austenitic steels [START_REF] Akiyama | Hydrogen Visualization in Steels Using Ag Decoration Method[END_REF], Ni and Pd [START_REF] Yao | The use of silver decoration technique in the study of hydrogen transport in metallic materials[END_REF], duplex stainless steels [START_REF] Yalçì | Application of the hydrogen microprint and the microautoradiography techniques to a duplex stainless steel[END_REF], micro-alloyed [START_REF] Mohtadi-Bonab | A comparative study of hydrogen induced cracking behavior in API 5L X60 and X70 pipeline steels[END_REF] and high strength steels [START_REF] Nagao | Visualization of hydrogen transport in high strength steels affected by stress fields and hydrogen trapping[END_REF].

Besides, the Ag decoration test has been used to study the hydrogen distribution around stress fields [START_REF] Nagao | Visualization of hydrogen transport in high strength steels affected by stress fields and hydrogen trapping[END_REF], [START_REF] Ichitani | Visualization of hydrogen diffusion in steels by high sensitivity hydrogen microprint technique[END_REF] and to perform hydrogen mapping across a crevice [START_REF] Sundararajan | Hydrogen Mapping Across Crevices[END_REF], [START_REF] Sundararajan | Hydrogen mapping across a crevice: Effect of applied potential[END_REF] Ag decoration test has also been used for identifying the diffusion paths by Koyama et al. [START_REF] Koyama | In situ observations of silverdecoration evolution under hydrogen permeation: Effects of grain boundary misorientation on hydrogen flux in pure iron[END_REF] and Akiyama et al. [START_REF] Akiyama | Hydrogen Visualization in Steels Using Ag Decoration Method[END_REF].

Ichitani et al. [START_REF] Ichitani | Visualization of hydrogen diffusion in steels by high sensitivity hydrogen microprint technique[END_REF] investigated the diffusion behavior of low carbon steel and hypo eutectoid steel using the microprinting technique. They used a setup similar to the permeation cell with Ag depositing on the detection side. They observed a uniform distribution of Ag particles in ferrite in low carbon steels with the size of the Ag particles increasing with the charging time. They did not observe any preferential diffusion along the ferrite grain boundaries. In hypereutectoid steels, the Ag particles initially localized along the ferrite/carbide interface of the pearlite and then subsequently appeared in the pro eutectoid ferrite and the ferrite in pearlite. They concluded that the pro eutectoid ferrite and ferrite in pearlite were the main diffusion paths, whereas attributed the diffusion of hydrogen through the ferrite/carbide interface to the trapping effect.

Akiyama et al. [START_REF] Akiyama | Hydrogen Visualization in Steels Using Ag Decoration Method[END_REF] investigated the hydrogen distribution behavior of SUS304 austenitic steel and martensitic and spherodized SUS440 steels. Figure 2.12 shows the SEM and AFM observations of the Ag decoration test for SUS304 steel. The Ag particles appear as white dots, are selectively deposited on the slip lines. The Ag grains did not cover the grain boundaries completely but were deposited selectively in the regions where the slip lines intercepted the grain boundaries. The selective precipitation Ag particles reveal that there was no Ag deposition in the annealed twins (Figure 2.12(b)). Since Ag deposition correlates to the H distribution, the H in austenite were believed to be selectively distributed along the slip lines in a strained specimen.

The limitation of the HMT test and the Ag decoration test is its efficiency. The efficiency of the test is quite low, which means there is no one to one correspondence between the deposited silver and the escaping hydrogen atoms [START_REF] Yao | The use of silver decoration technique in the study of hydrogen transport in metallic materials[END_REF]. As a result, the Ag decoration test fails to provide quantitative information about the hydrogen flux [START_REF] Nagao | Visualization of hydrogen transport in high strength steels affected by stress fields and hydrogen trapping[END_REF]. As a result, there are high chances of Ag particles not being deposited when the hydrogen flux is low, as pointed out by Koyama et al. [START_REF] Koyama | In situ observations of silverdecoration evolution under hydrogen permeation: Effects of grain boundary misorientation on hydrogen flux in pure iron[END_REF]. As a result, the Ag decoration test may not give an accurate overview in case of low hydrogen concentration or case of extremely low diffusivity [START_REF] Akiyama | Hydrogen Visualization in Steels Using Ag Decoration Method[END_REF]. Allen et al. [START_REF] Allen | Microstructural evaluation of hydrogen embrittlement and successive recovery in advanced high strength steel[END_REF] studied the effect of crystallographic orientation on hydrogen diffusion in high strength steels using microprinting technique and EBSD characterization.

© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon

Based on the Ag particles deposited, they observed that the hydrogen flux rate and desorption rate had a strong dependence on the crystallographic orientation. In the steels investigated, they showed that <001> and <101> orientations exhibit high levels of early hydrogen flux and decay rates during hydrogen charging and discharging. <111> orientation, on the other hand, had lower hydrogen flux but retained hydrogen for a more extended period. They explained that in the bulk though the hydrogen atom is free to move in any direction (due to cubic structure), the path is determined by the energy associated with the orientation. Since <111> has higher surface energy as compared to the other orientations, the diffusion of hydrogen is slower as compared to other orientations.

Summary and Conclusions

The study highlights the importance of understanding the effect of hydrogen on AHSS. The AHSS offers excellent mechanical properties advantageous to automobile industries, but HE issues restrict their use. The susceptibility to HE is affected by several environmental, material, and mechanical parameters. Two first-generation AHSS (DP and CP) and a second-generation TWIP steel were discussed in greater detail. The susceptibility of DP steels was attributed to its microstructure (ferrite matrix with martensitic islands),

where the ferrite phase is inhomogeneous with higher lattice defects near the phase boundaries.

The hydrogen trapped in these sites promote the crack propagation along the ferrite/martensite interface. Not much information is available on the HE susceptibility of CP steels. However, CP steels show relatively improved HE resistance (compared to DP steels) due to the lower difference in the mechanical strength of the two hard phases (bainite and martensite) present. The HE susceptibility of TWIP steel is not entirely clear as the published literature report contradictory findings. However, Al addition enhances the HE resistance as it reduces the dislocation density. Coarser grained microstructure is more susceptible to HE as compared to fine-grained structures attributed to their more active mechanical twins. Though the detrimental influence of hydrogen on steels have a subject of study for decades, the exact nature of hydrogen once in the metal is still under investigation. The challenges associated with this research is due to the small size of hydrogen and its ability to cause damage at minimal concentration levels. In order to highlight the significance of the elusive nature of hydrogen in embrittlement behavior, when asked to Prof. David Dye (Imperial College, London) that, if given a chance, how does he plan to use sixty million dollars in this embrittlement research [START_REF] Dear | Mechanisms of hydrogen embrittlement in steels: discussion[END_REF]. His answer, "I think we want to know where the hydrogen is."

© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon

Introduction

Reducing vehicle weight is a key challenge for the automobile industry to increase fuel efficiency and reduce vehicular emission without compromising on safety. Advanced High Strength Steels (AHSS), which offer superior strength, therefore are gaining attention as the body in white material (BIW), with its use increasing from 7% in 2010 to 21% in 2017 (and projected to increase up to 40% by 2025) [1]. However, its susceptibility to hydrogen embrittlement (HE) continues to pose an important bottleneck for its utilization [2]- [8].

Hydrogen present in steels, present even less than 1 ppm, tends to diffuse towards triaxially stressed regions, increasing its concentration locally, causing premature failures. Such failure can occur instantaneously or delayed, making failure prediction difficult, increasing liability and repair costs [9]- [11]. Since H diffusion has a direct impact on H transport to stressed areas, studying H permeation behavior in steels is an important step in understanding the overall embrittlement behavior [12]- [15].

Electrochemical permeation study using Devanathan and Stachurski (DS) cell is an effective way to understand H permeation behavior with good sensitivity by continuously recording the diffusive flow of hydrogen with time [3], [12], [16]. The study of permeation behavior is particularly important in the purview of microstructure, as H uptake and its diffusivity strongly depend on H solubility, segregation, and trapping tendency of different microstructures. For example, at ambient temperature, the lattice diffusion coefficient of H in pure iron is reported to be around 1.3 x 10 -9 to 7.5 x 10 -9 m².s -1 whereas it is 4.47 x 10 -11 m².s - 1 for fully martensitic and 5.19 x 10 -11 m².s -1 for fully bainitic steels [17]- [20]. These microstructural features, such as grain interior, phase interfaces, dislocations, grain boundaries, precipitates, etc., can trap H and reduce the diffusion coefficient by up to two to three orders of magnitude, depending on their activation energy (Ea) [17]- [19]. The activation energy of traps dictates whether, under stressed conditions, they will behave as sources or sinks for diffusible hydrogen, in turn, increasing or decreasing HE susceptibility [21]- [24].

Traps also exhibit a strong influence on H solubility, with an equilibrium existing between the chemical potentials of lattice and the traps [25]. For example, studies show that DP steels have higher retention for H than bainitic ferrite steel of similar strength and have lower threshold values for HE [26]. In the case of martensitic structure, high carbon steels are

65

© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon reported to trap larger quantities of H compared to low carbon steel, increasing H susceptibility [27]. However, tempering reduces the sensitivity to hydrogen, with HE resistance increasing with increasing tempering temperature [28]. Bainitic steels also offer good resistance to embrittlement due to the homogeneous, refined, and stress-free microstructure, though other factors such as carbide type and other micro phases also needs to be considered as well [28], [29].

Needless to emphasize, AHSS being multiphase steels with varying amounts of ferrite, martensite, bainite, or austenite exhibit a complex interaction with H influencing its diffusivity, solubility, and distribution within the material [30]- [32]. Thus, especially in the context of AHSS, there is a need to understand the influence of microstructure on the permeation behavior as material microstructure plays an important role in H distribution [33].

This study aims to investigate the influence of microstructure of three first-generation AHSS on H permeation behavior, specifically in terms of diffusion coefficient, the H concentration in the lattice and the traps, and trapping capabilities of the investigated steels. The lattice diffusion coefficient of H was experimentally determined and compared among the investigated specimens. The influence of material microstructure on H solubility, segregation between lattice and trap sites, and its distribution across the material is studied in detail.

Finally, the difference in the trapping behavior among the different steels were discussed, and an average Ea was calculated for each material. In summary, the H permeation behavior in three steels were studied extensively to understand the role of their microstructure.

Experimental

For the present study, TATA Steel Netherlands provided two DP steels having a tensile strength of about 800 MPa (DP1000) and 1000 MPa (DP1000) and a CP steel having a tensile strength of about 1000 MPa (CP1000) in the form of 30 cm x 30 cm sheet of different thicknesses. The DP steels had an average thickness of 0.98 mm, while CP steel had 1.60 mm average thickness. The chemical compositions of materials are as shown in Table 3.1.

© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon

© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon

electrochemically reduce H + at the cathodic/charging side and the other to subsequently oxidize H at the detection/anodic side, were used. A fraction of the H + reduced at the cathode surface of the specimen is absorbed as H into the specimen, diffuses across the specimen, and finally oxidized at the anode surface (Pd coated) of the specimen. In this study, the H + was reduced galvanostatically with varying current densities and the oxidized potentiostatically at +0.2 V w.r.t. RE The potential represented an overpotential of +250 mV ensuring complete ionization of H at the detection side as per Hads → H + + e -. Heated water circulated in the double jacket to maintain a constant temperature of 35°C at each compartment.

Figure 3.1: Schematic of the permeation cell used for this study. H + is reduced as H at the left compartment and oxidized at the right compartment

Experimental Procedure: The electrochemical hydrogen permeation experiment was carried out according to the ISO 17081 specifications [3], [33], [34]. Specimens with an exposed surface area of 7 cm² were clamped between the two compartments. The anodic compartment was filled with 0. ) was added slowly at the charging side and the cathodic surface was charged galvanostatically. Ar gas flowed through both the electrolytes throughout the experiment for deaerating the electrolytes. In the absence of oxygen, H oxidation is the only electrochemical reaction at the anodic surface and provides a direct measure of the H atoms permeating through the specimen.

The galvanostatic charging at the cathodic surface involved the following three steps. As an example, Figure 3.2(a) represents the permeation current density as obtained due to the charging schedule at the cathodic surface for CP1000. Similar testing was conducted for other specimens.

Step I (Cathodic Pre-charging): When the background current density at the detection side was below 0.3 µA.cm -², a cathodic current density of -10 mA.cm -² was applied on the charging surface until an observable steady-state current permeation current density

(Figure 3.2(b)).
Step II (Subsequent Partial Transient): After step I, the charging current density was raised to -20 mA.cm -². Once the permeation current density reached a new steady-state, the charging current density was lowered back to -10 mA.cm -². The applied current density was held for 5000 s for each step. The steps ware repeated to verify the permeation curves' reproducibility (as shown in Figure 3.2(c)).

Step III (Partial Transient Loop): After completing step II, the charging current density was raised from -10 mA.cm -² to -50 mA.cm -² and after that lowered to -10 mA.cm -² (thus creating a loop), with a step size of -10 mA.cm -² ( The above charging schedule reduces the surface modification influence in the permeating current density occurring at the charging surface, which is one of the key challenges in applying ISO 17081. Additionally, as a cathodic potential is always associated © 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon with the charging surface, it minimizes the possibility of corrosion issues, particularly during fall transients [12], [18], [35], [36]. It is reported that surface modifications (dissolution/formation of oxides/hydrides/reaction products) occurring during H charging can significantly facilitate/block hydrogen entry in the material during the test and thus influence the observed permeation behavior [37]- [41]. According to Turnbull, during the early stages of the test, there will always be a partial surface reaction, and its effect could be minimized by constant charging [34]. A long term cathodic pre-charging, as carried out in step I, will, therefore, stabilize the specimen surface and minimize the influence of surface modification on permeation curves [42]- [44].

To verify pre-charging's effectiveness in reducing surface modification's influence on permeation current density, the steady-state permeation current density was plotted against the square root of applied cathodic current densities for pre-charged specimens as shown in Figure 3.3. The linear relationship between the two parameters validate that the permeation current density is indeed controlled by the impressed cathodic current density and not influenced by surface modifications and is attributed to the activation of cathodic surface facilitating hydrogen absorption [34], [45]. Zakroczymski et. al. [35] compared this relation in non-pre-charged and pre-charged specimens and observed a sigmoidal curve in case of nonpre-charged specimens, whereas followed a linear relationship in the case of pre-charged specimens. Step II followed step I, where the partial rise and fall transients' reproducibility were analyzed, as shown in filling up all the irreversible traps leading to the interaction of hydrogen with only the reversible traps and lattice sites after a long pre-charging; (b) the trapping effect from the reversible traps to be similar for both transients; (c) the surface effects for both the transients to be similar [37], [39], [46]. The reproducibility of the curves also indicate that there was no permanent damaged caused to the material (also confirmed by observing the specimen under an optical microscope) due to high current density of the cathodic charging, a concern which has been raised by some authors in the past [36], [47].

© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon

© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon

Results

Microstructure

Figure 3.5 presents the scanning electron images (SEM) of the three alloys investigated in this study. The phase fractions and the average grain sizes were determined using EBSD and presented in Table 3.2, along with the mechanical properties. Figure 3.5 (a and c) presents the surface microstructure of DP800 and DP1000 specimen, respectively and reveal the martensitic islands embedded in ferrite [48]. The high martensite phase fraction in the alloys make the ferrite discontinuous, an observation similar to the works of Davies et al.

[49], [50]. Between the two DP steels, DP1000 contains a higher martensite phase fraction than DP800, which is responsible for the higher strength and reduced ductility of DP1000.

With higher martensitic phase fraction and similar size of the martensitic islands in DP1000, the average ferritic grain size in DP800 was significantly larger than DP1000. revealing martensite as the predominant phase followed by ferrite and bainite. CP steels contain significantly higher quantities of bainite than DP steels. Further, the average size of the ferrite grains and the martensitic islands is similar in CP1000 but is smaller than DP steels.

The phase fraction of retained austenite in all the investigated steels was below 0.25%.

Hydrogen Diffusion Coefficient

H diffusion in metals is a critical parameter to understand their sensitivity to HE as it is the diffusible H that migrates to crack tip, causing embrittlement. The hydrogen diffusion coefficient was calculated by fitting the experimentally obtained partial permeation transient data to Equation 1 & 2 assuming that during H charging, the H concentration at the charging side attains a new value just after the cathodic current density is varied [3], [17], [21].

𝑖𝑖 𝑛𝑛 = 2𝐿𝐿 √𝜋𝜋𝜋𝜋𝑡𝑡 ∑ 𝑒𝑒𝑒𝑒𝑒𝑒 �- (2𝑛𝑛+1) 2 𝐿𝐿 2 4𝜋𝜋𝑡𝑡 � ∞ 𝑛𝑛=0
Equation 1 (Rise Transient)

𝑖𝑖 𝑛𝑛 = 1 - 2𝐿𝐿 √𝜋𝜋𝜋𝜋𝑡𝑡 ∑ 𝑒𝑒𝑒𝑒𝑒𝑒 �- (2𝑛𝑛+1) 2 𝐿𝐿 2 4𝜋𝜋𝑡𝑡 � ∞ 𝑛𝑛=0 Equation 2(Fall Transient)
Where, in is the normalized current density (the initial and final current densities for the partial rise and partial fall transients were considered to be 0), D (m².s -1 ) is the H diffusion coefficient, L is the specimen thickness (m), and t is the time (s) for each partial transient.

The H diffusion coefficient so-measured corresponds to the apparent diffusion coefficient (Dapp) and not to the lattice diffusion (DL) as traps present in the alloy impede the hydrogen movement [18]. The experimental data fitted to the above equations for the various initial and final charging test conditions for DP800 is presented in Figure 3.6 for illustration.

The goodness of fit of the curves suggests that the hydrogen diffusion is not affected by the slow surface processes [17], [18], [39], [44]. A similar fitting procedure was carried out for the other steels as well.

The diffusion coefficient values obtained in this work for the tested steel specimens are well within the range of diffusion coefficients reported in the literature [36], [37], [51]. In case of DP800 the diffusion coefficient values obtained is in the range of 0.83 to 1.68 x 10 -10 m 2 .s -1 , 0.77 to 1.60 x 10 -10 m 2 .s -1 for DP1000 and 1.07 to 3.16 x 10 -10 m ² .s -1 for CP1000, as presented in Figure 3.7.

© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon cathodic overpotential is consistent with the literature [17], [36], [39], [44]. The above increase in Dapp results from the incremental filling of traps and the consequent reduction of the trapping capacity of microstructural features with a step-wise increase in cathodic overpotential/charging current density. Another factor that might contribute to the rising Dapp value with the applied current density is the increase in the H concentration gradient across the membrane facilitating the H diffusion across the membrane.

© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon

Similarly, the Dapp value decreases with the reduction in the charging current density. With the reduction in applied current density, H is released from both the lattice sites and the trap sites, maintaining equilibrium [33]. It results in the increasing influence of trapping sites impeding the H movement and consequently decreasing the Dapp value. With the evident decreasing influence of traps with an increase in charging current density/cathodic overpotential, the lattice H diffusion coefficient (DL) in the investigated steels is considered from the data obtained at the highest overpotential performed in the study [3], [21], [39]. The DL values obtained for DP800, DP1000, and CP1000 are 1.65 x 10 -10 m ² .s - 1 , 1.58 x 10 -10 m ² .s -1 , and 3.07 x 10 -10 m ² .s -1 respectively The so-obtained DL values are assumed to be a close approximation of the actual lattice H diffusion coefficient in the investigated steels.

Data shows that CP1000 has the highest DL value, while both the DP steels show nearly the same DL values. As expected, the DL is much higher than the Dapp for all observed cases.

Zakroczymski et al. [18] made similar observations where the diffusion coefficient of H in non pre-charged pure iron was an order of magnitude smaller than the lattice diffusion coefficient.

Hydrogen Concentration

The sub-surface and lattice hydrogen concentrations are other important parameters contributing to HE of steels with limited information in the published literature.

Using the steady-state current density and the calculated diffusion coefficient values, both the total subsurface H concentration (CT) and the lattice H concentration(CL) in mol.m -3 , are determined. Under equilibrium conditions, CT and CL are related to the steady-state current density (iss) via the corresponding diffusion coefficients using Equations 3 and 4, respectively [12], [52].

𝐶𝐶 𝑇𝑇 = 𝑖𝑖 𝑠𝑠𝑠𝑠 * 𝐿𝐿 𝐹𝐹 * 𝜋𝜋 𝑎𝑎𝑎𝑎𝑎𝑎 Equation 3 𝐶𝐶 𝐿𝐿 = 𝑖𝑖 𝑠𝑠𝑠𝑠 * 𝐿𝐿 𝐹𝐹 * 𝜋𝜋 𝐿𝐿 Equation 4
Where F is the Faraday's constant.

The difference between the CT and CL corresponds to the H present in the traps (CTrapped). The above three H concentration values are plotted against the charging current density in Figure 3.8, and the corresponding data is presented in Table 3.3. In general, all three H concentration values increase with an increase in the applied current density. The data suggest that CT and CTrapped for the steels vary in the order DP1000>DP800>CP1000, while a cross-over, from the above trend, occurs for DP800 at higher charging current density with respect to DP1000 for CL, indicating that DP800 has the highest value at higher charging current densities. The subsurface H concentration also has an inverse relationship with the DL. The steel with lower DL has higher sub-surface concentrations (Figures 3.7 and 3.8). It is understandable since the H that enters the steel due to cathodic charging diffuses faster into the steel as it has a high DL value [53].

Further, it is observed that CTrapped saturates beyond 20 mA.cm -2 because of the finite trap sites. Such a distinct saturation is not noted for CT (Figure 3.8.a). It appears that the increase in the H fugacity with applied current density enables higher solubility of hydrogen in the steel, although at high hydrogen fugacity, hydrogen recombination may be favored. Further, with an increase in the charging current density, the available free sites for H reduction are expected to decrease, due to the increased surface coverage by the adsorbed H atoms, limiting the H adsorption. As shown in Figure 3.8, when the charging current density is increased from -20 mA.cm -2 to -30 mA.cm -2 for DP800, the increase in total hydrogen concentration is 0.55 mol.m -3 , whereas at higher charging current densities when its raised from -40 mA.cm -2 to -50 mA.cm -2 , the change in hydrogen concentration is only 0.19 mol.m -3 .

© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon

Hydrogen Trapping

Reversible H traps also affects the HE tendency of steels, as these traps act as reservoirs for hydrogen atoms and release them into the lattice over time. The experimental permeation fall transient curve is sensitive to the hydrogen in the reversible traps, with the area under the curve representing H desorbed from the lattice as well as reversible trap sites (CT*). The detrapping kinetics of hydrogen from the reversible traps is calculated by subtracting a decay curve corresponding to DL (Equation 2) from the experimentally obtained current decay curve [3], [18], [37], [39]. Where NT is the number of hydrogen trapping sites per unit volume (sites.cm -3 ),

A is the difference in the area under the two curves (s), L is the thickness of the specimen (cm), in is the steady-state current, and 6.24 x 10 12 is the number of hydrogen atoms in 1µA.s.cm -2 .

The calculated NT summarized in Table 3.3 indicates that with increasing charging current density, the reversible trap sites of all the three steels decrease. Further, the calculated reversible trap density varies in the order DP800 > DP1000 > CP1000. The trap density emptied during fall transients increases with decreasing charging current densities.

For example, in DP800, lowering charging current density from -50 mA.cm -² to -40 mA.cm - ² emptied only 12.5% of the total calculated traps, whereas, with lowering the charging current density from -20 mA.cm -² to -10 mA.cm -², 46.2% of the total calculated traps were emptied.

All the steels investigated demonstrated similar behavior. In DP steels, the calculated trap density was higher in DP800 than DP1000. It is possibly related to the energy associated with the traps. If the trap strength of DP1000 is stronger than DP800, it is understandable that fewer traps would be emptied with a reduction in charging current density.

Trap Activation Energy

The binding energy of hydrogen with traps play a decisive role in determining the diffusion/trapping of H in the material. Following the works of Oriani [33] and later modified by Dong [54], the activation energy of the traps can be calculated from the obtained permeation data [55] using Equation 6:

𝐸𝐸 𝑎𝑎 = 𝑅𝑅𝑅𝑅{𝑙𝑙𝑙𝑙 � 𝜋𝜋 𝐿𝐿 𝜋𝜋 𝑎𝑎𝑎𝑎𝑎𝑎 -1� -𝑙𝑙𝑙𝑙 𝑁𝑁 𝑇𝑇 𝑁𝑁 𝐿𝐿 } Equation 6
Where NT is the number of hydrogen trap density (sites.cm -3 ), NL is the density of interstitial sites in the steel, Ea is the activation energy for the traps, R is the gas constant and T is the absolute temperature.

Based on similar works available in the literature, the interstitial site density (NL) is assumed 7.52 x 10 22 sites. cm -3 [13], [56]. The average Ea calculated for DP800 is 34.05 kJ.mol -1 , 37.61 kJ.mol -1 for DP1000 and 38.94 kJ.mol -1 for CP1000. The activation energy measured from the permeation technique is usually higher than the values obtained by TDS measurements [57]. In TDS, when the thermal energy is applied, all the hydrogen present in the material effuses out. However, in the permeation method, the decrease in the cathodic overpotential may not be sufficient to release all the hydrogen present in the traps. As trap density is inversely proportional to the activation energy, the Ea calculated from permeation tests present a higher value. Though the activation energy measured from the permeation technique may not be accurate; nevertheless, it can provide a reasonable idea about the nature of traps.

Discussion

The present study suggests that the H diffusion coefficient in CP steel was higher compared to the DP steels. The H diffusion path influences the observed diffusion coefficient values of these steels. As shown in Figure 3.5(b,d, and f), DP steels have a banded structure across its cross-section, which was not present in CP steel. The banded structure indicates a longer mean diffusible path for hydrogen in DP steels than in CP steel. In DP steels, the grain boundaries are the preferred diffusion pathway as the matrix is predominantly ferrite [58].

Koyama et al. [59], using the Ag decoration technique, reported that grain boundaries were the preferred diffusion pathways for hydrogen in ferrite. Haq et al. [36] reported that effective hydrogen diffusivity in specimens with banded structure perpendicular to the direction of hydrogen diffusion tends to be an order of magnitude lower than specimens with bands parallel to the direction of hydrogen diffusion. Olden et al. [60] reported that the greater the diffusion path and the higher the tortuosity of hydrogen, the lower is the diffusion coefficient of hydrogen. Zakromyzski discussed that the flux of hydrogen is directed along the layers of fibers of inclusions [18]. A schematic showing the influence of lamellar structure on the actual hydrogen diffusion path is presented in The calculated hydrogen diffusion path (specimen thickness)

The actual hydrogen diffusion path Among the three steels investigated, the CT, CL, and the CTrapped (Figure 3.8) were significantly lower for CP1000 compared to both DP1000 and DP800. Looking at the carbon content in all the steels, as presented in Table 3.1, the C content in CP steel is relatively less than DP steels. Carbon concentration in steels is known to influence its HE susceptibility and, consequently, the hydrogen concentration in steel and its diffusion coefficient. Hadam et al.

observed that 1 wt.% increase of C in steel increases the solubility of hydrogen in steel by 70 times and could reduce the diffusion coefficient by two orders of magnitude [61].

In the absence of the knowledge of the charging schedule, an inference could be drawn that the hard phases, particularly martensite in CP1000 steel is softer as compared to the DP1000. This is due to the higher phase fraction of hard phases in CP steel with lower C concentration as compared to DP1000. Softer martensite contains lesser dislocation density and less stressed ferrite regions around the ferrite/martensite interface [28]. As dislocation density present in the martensite is considered a reversible trap site, CP1000 has a lesser © 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon number of reversible trap sites compared to DP steels, even though CP1000 has a higher grain boundary area due to its refined microstructure. It is consistent with the trap density observed in Table 3.3, explaining the small quantity of trapped H in CP steels.

However, though the trap density is smaller in CP1000, under similar charging conditions, a higher fraction of H is trapped in CP steels as compared to the DP steels. From figure 3.8, if we calculate the fraction of the total absorbed hydrogen present in traps, it is evident that a higher fraction of H is trapped in CP1000, as shown in Figure 3.11. In CP steels, there is a significant portion of bainite, which is composed of cementite precipitates. These cementite particles are known to be strong hydrogen traps and are the major trapping sites in the bainitic phase [28]. The size and morphology of the carbide particles also play an essential role in determining the permeability of hydrogen. Ramunni et. al. [62], observed finer carbides as more efficient in hydrogen trap sites. The higher trap activation energy of CP steel also supports the present observation (Table 3.3).

Chun et al. [63] reported that the Ea associated with dislocations and grain boundaries to be around 35 kJ.mol -1 . Luppo et al. [64], in their study of low carbon martensitic steel, found that Ea for lath interfaces to be 46 kJ.mol -1 . Wong et al. [65] found the energy required to release hydrogen from dislocations to be around 21 kJ.mol -1 . Based on these previous results, the hydrogen released from the reversible traps may be present in the grain boundaries, dislocations, and martensite lath interfaces in DP and CP steels along with cementite particles present in CP steels. The low hydrogen uptake, the low trap density, and the presence of stronger traps in CP steels make it more resistant to HE as compared to DP steels [9], [28], [43], [56], [66]- [68]. The ratio between the change in the hydrogen concentration at the discharging side and at the charging side is 0.36 for DP800, 0.33 for DP1000, and 0.30 for CP1000 steels (Table 3.

3). This helps to get an insight into the possible hydrogen distribution within the metal.

Hadam et al. [61] suggested that H distribution within the material is uneven along the material cross-section, with most of the H trapped at the sub-surface on the charging side of the specimen. Zakroczymski [44] showed that most of the hydrogen desorbed through the charging side of the specimen. After the permeation current reached steady state, Zakroczymski interrupted the cathodic charging, and an anodic potential equal to the discharging side was applied to the charging side, which allowed the measurement of desorbed hydrogen from both surfaces. The hydrogen detected from the discharging side accounted for only 12.4% of the total hydrogen released. In the present test conditions, there is the possible release of hydrogen from both sides (charging and discharging). As a result, the total hydrogen desorbed from the discharge side is around one-third of the total change in the hydrogen concentration at the charging side. It must be remembered that in this case, there is no anodic potential applied at the charging side after steady-state, but the only driving force of desorption of hydrogen from the charging side is the lowering of cathodic current density.

This interpretation is however, just speculative and needs further validation as the H concentration profile is extremely difficult to be verified experimentally [START_REF] Mcmahon | Hydrogen-induced intergranular fracture of steels[END_REF]. However, if true, such observations can provide better insight to interpret the trapping effect of different microstructural structures in the material.

Conclusions

Permeation characteristics of DP and CP steels were investigated in the study The Dapp for all specimens increased with an increase in charging current density due to the reducing influence of traps and increased H concentration gradient. Moreover, DL in CP steel (3.16 x 10 -10 m ² .s -1 ) was observed to be twice compared to the DP steels (1.65 x 10 - 10 m 2 .s -1 for DP800 and 1.58 x 10 -10 m 2 .s -1 for DP1000). The smaller diffusion coefficient of hydrogen in DP steels is attributed to the higher trap density and lamellar microstructure along the cross-section, perpendicular to the H diffusion path.

CP steel has low H concentration at the charging sub-surface compared to DP steels. A higher DL resulted in easier H diffusion, decreasing H concentration at the subsurface. The relatively lower carbon concentration in CP steels is a possible reason for the low H sub-surface concentration. Though DP steels have higher H pick up and trap density, a higher fraction of the absorbed hydrogen is trapped in CP steel, indicating better resistance of the steel to HE. This is due to the presence of cementite particles in CP steels which are known to be s strong trap with high trap activation energy. CP steel had the least trap density compared to DP steels but have higher trap strength. The Ea calculated for DP800 is 34.05 kJ.mol -1 , 37.61 kJ.mol -1 for DP1000 and 38.94 kJ.mol -1 for CP1000.

The CL of DP1000 was observed to be higher than DP800 at low charging current density, but at higher charging current density, CL is higher for DP800 due to easier saturation of traps in DP800 leading of H enrichment in the lattice.

For all steels, H concentration effusing from the discharge side is almost a third from that of the charging side, 0.36 for DP800, 0.33 for DP1000, and 0.30 for CP1000 steels indicating the non-uniform distribution of H across the cross-section, with higher H concentration near the charging surface.

It is possible to suggest that CP1000 steel is likely to be more resistant to HE than the two DP, namely DP 800 and DP 1000 steels due to the fact that the CP steels have higher hydrogen diffusion coefficient, lower H pickup, smaller but stronger trap density.

Introduction

In auto industries, Advanced High Strength Steels (AHSS) are often used to fabricate Body In White (BIW), to increase fuel efficiency and reduce carbon footprint. AHSS offer superior mechanical properties compared to conventional high strength steels but suffer from increased hydrogen embrittlement (HE) susceptibility [1]- [3]. The material's microstructural features or traps attract the absorbed H, as a lowering of energy is associated with H transfer from lattice to these defect sites [4]. Once trapped, if the activation energy (Ea) of the traps are high, H remains trapped, until such trap sites are fully saturated, or an adequate amount of energy is applied [5]. Conversely, if Ea is low, trapped H can be easily detrapped, diffuse towards the triaxial region of stress and cause detrimental influence on material's fracture behavior [6]- [9]. Thus, the trap's characteristics determine its role as H source or sink under stress, and plays a significant role in determining material's susceptibility to HE.

Thermal Desorption Spectroscopy (TDS) is a useful technique to determine the H concentration in a material, distributed in strong or weak traps, and their activation energy.

Initially developed to study the dangling bond termination by hydrogen on semiconductor surfaces, TDS is a non-isothermal technique popularly adapted to study the hydrogen trapping in steels [10]. It can also be used to determine the hydrogen diffusion coefficient in steels and is particularly useful for steels with a low H diffusion coefficient, which otherwise would take a significantly longer time using hydrogen permeation technique [11].

Hydrogen pre-charged specimens are exposed to linearly increasing temperature in TDS, while desorbing H is measured using a mass spectrometer [12]. Hydrogen releases from the trap sites on absorbing sufficient thermal energy to overcome the barrier energy [13].

TDS has been used to identify the Ea of microstructural features such as dislocations [14]- [16], grain boundaries [16], [17], crystal lattice [15], [18], [19], dislocation cores [15],

precipitates [20], [21], lath boundaries [22], [23] and other possible trapping sites [8], [16], [24]. The knowledge of the activation energy of different trapping sites is beneficial for the understanding and development of materials with improved HE resistance. For example, the presence of well-distributed strong-traps reduces the HE susceptibility, as it reduces the

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volume of diffusible hydrogen, particularly if a finite amount of hydrogen is available [25].

One particular advantage of TDS is its ability to detect hydrogen up to sub-ppm (parts per million) range, which is useful given that sometimes only a small concentration of hydrogen in the material can induce failure due to HE [26].

TDS, though a useful technique to characterize the hydrogen present in the bulk of the metal, does not provide direct information about the distribution of hydrogen on the metal surface. Ag decoration, on the other hand, provides information about the distribution of adsorbed hydrogen on the specimen surface through visual observation [27]- [29]. The Ag decoration technique assumes that the adsorbed H is active and can reduce Ag + in the solution leading to precipitation of metallic Ag on specimen surface according to Equation 4.1 [27]:

𝐻𝐻 𝑎𝑎𝑎𝑎𝑎𝑎 + 𝐴𝐴𝐴𝐴 + → 𝐻𝐻 + + 𝐴𝐴𝐴𝐴 ↓ Equation 4.1
The reduced Ag particles precipitate on the specimen surface at sites where adsorbed hydrogen is present. Such sites are identified by visual observation under optical microscopy or scanning electron microscopy (SEM). [28]. One limitation of the Ag decoration test is the lack of one to one correspondence between the deposited Ag atoms and the adsorbed H atoms on the surface. As a result, the test results can at best be treated qualitatively and not quantitatively [29]- [31].

Table 4.1: Chemical Composition of the steels in wt% used for TDS and Ag decoration study

In the previous chapter, the microstructure of the two DP steels and CP steel have been discussed in detail (Chapter 3). DP steel consists of ferrite and martensite phases, with a higher phase fraction of martensite for DP1000 as compared to DP800. CP1000 microstructure consists of ferrite, bainite, and martensite. TWIP steel is completely austenitic, as determined by EBSD analysis. For TDS measurements, 15mm X 20 mm coupons from each specimen were ground up to 600 grit SiC, followed by washing twice with acetone and once with alcohol in an ultrasonic bath and then air-dried. For the Ag decoration test, 15mm X 20 mm coupons from each specimen were ground up to 4000 grit SiC and polished up to 1 µm diamond paste followed by similar cleaning and drying procedure. The polished surfaces were etched with a 2% Nital solution before exposing them to the Ag decoration solution.

Electrochemical Charging

For TDS analysis, the specimens were charged electrochemically in 3 wt.% NaCl solution with 5 g.L -1 NH4SCN as hydrogen recombination poison (deactivate recombination

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reaction by irreversible adsorption on the metal surface at active centers [12]) at a charging current density of -1 mA.cm -2 . The charging time required to saturate the specimens with hydrogen was determined using the H diffusion coefficient in the specimen. From our previous study (Chapter 3) the diffusion coefficient of hydrogen in DP800, DP1000 and CP1000 were 1.68 X 10 -10 , 1.60 X 10 -10 and 3.16 X 10 -10 m².s -1 respectively. The diffusion coefficient of TWIP steel at room temperature was assumed from the published literature to be 1.29 X 10 - 13 m².s -1 [32]. Lee et al. [7] suggested a relationship between the specimen H concentration and hydrogen diffusion coefficient as per Equation 4.2:

𝑄𝑄 𝑡𝑡 𝑄𝑄 ∞ = 1 - 8 𝜋𝜋 2 exp �- 𝜋𝜋 2 𝜋𝜋𝑡𝑡 𝐿𝐿 2 � Equation 4.2
Where, 𝑄𝑄 𝑡𝑡 is the amount of hydrogen absorbed in time 't', 𝑄𝑄 ∞ is the total amount of hydrogen in a fully saturated specimen, 𝐷𝐷 is the diffusion coefficient of hydrogen in the specimen, and '𝐿𝐿' is the thickness of the specimen.

To saturate the specimens with H, the charging period for DP and CP steel specimens was determined using Equation 4.3, to be below 1 hour whereas, for TWIP steel, the charging period estimated to be up to 120 days. A similar charging period for DP steels proposed in the published literature, where charging time was experimentally determined using a hot extraction technique [11], [33]. Thus, to adequately charge the specimens with H, DP and CP steels were charged for 3 hours while the TWIP steel specimens were charged for 7 days, estimating that during this period, H could penetrate to a depth of 300 µm using Equation 4.3 [34] 𝑙𝑙 = (𝐷𝐷. 𝑡𝑡) -1/2

Equation 4.3

Where 𝑙𝑙 is the distance up to which the hydrogen can diffuse in time 't', 𝐷𝐷 is the H diffusion coefficient, and 𝑡𝑡 is the time.

The H charged specimens were washed with alcohol and quickly transferred in liquid N2 to arrest H desorption. Immediately before the TDS analysis, the specimens were taken out from liquid N2, cleaned with acetone in an ultrasonic bath (the time required for cleaning was sufficient to bring the specimens to ambient temperature), weighed, and inserted in the TDS instrument.

Thermal Desorption Spectrometer

The TDS instrument consisted of a specimen introduction chamber or loading chamber, a programmable furnace, and an analysis chamber connected to a Q-mass controller for measuring the desorbed H. The TDS instrument and its schematic of the same is as shown in Figure 4.1. The introduction chamber was brought to atmospheric pressure by a flow of pure nitrogen gas. Once at atmospheric pressure, the chamber door was opened to place the specimen in the chamber's specimen holder. After placing the specimen, the chamber door was locked, after stopping the gas flow and was vacuumed using a pump to a pressure of 10-3 Pa.

Subsequently, the chamber was vacuumed to 9.99 X 10 -5 Pa using a turbo-molecular pump.

The interval between loading the specimen and vacuuming was around 30 minutes. After reaching the vacuum, the gate valve separating the introduction chamber and the furnace was opened, transferring the specimen in the furnace using a transfer rod. After placing the specimen, the transfer rod was retracted, and the gate valve was closed. Before each measurement, the instrument was calibrated using small doses of hydrogen. Once the specimen was in the furnace, it was switched on with different temperature ramp rates (100, 200, 400, and 800 K.h -1 ) from room temperature to a maximum temperature of 1023 K.

The linear heating rate during the specimen analysis was maintained using a proportional integral derivative (PID) controller. The PID controller was monitored through a thermocouple, located outside the furnace, and positioned symmetrically halfway through the length of the tube. The transparency of the cylindrical quartz tube helped to consistently

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position the specimen at the same place, helping reproducibility. The partial pressure of the desorbed H was then measured using a quadrupole mass spectrometer (QMS) present in the analysis chamber.

Ag decoration Test

Using the Ag decoration test, the H distribution on the specimen surface was observed, in two ways. For the first method, the H charged specimens were immersed in an aerated 4.3 mM K[Ag(CN)2] (Ag decoration solution) solution at OCP for 10 minutes at room temperature. The specimens were then cleaned with acetone in an ultrasonic bath and airdried. The Ag deposition on the surface was observed using Jeol JSM-7800F PRIME SEM.

In the second method, the specimen was placed between two electrochemical compartments to behave as the common working electrode. The setup was similar to a hydrogen permeation cell setup (Chapter 3). The etched specimen surface faced the detection side of the permeation cell. At the charging side, the surface was charged electrochemically in 3 wt.% NaCl + 5 g.L -1 NH4SCN solution at a charging current density of -1 mA.cm -², whereas the solution at the detection side consisted of 4.3 mM K[Ag(CN)2] solution at OCP. The idea was to observe the H distribution at the detection side after H permeated through the specimen. The specimen was then removed, washed in an ultrasonic bath with acetone, airdried, and observed under SEM. The second method was not carried out for TWIP steel, due to the small H diffusion coefficient in TWIP steel as compared to the other steels.

Results and Discussion

Hydrogen concentration

Figure 4.2 presents the TDS spectra of the four steel specimens for the heating rate of 800 K.h -1 . The obtained TDS curves are divided into two distinct regions: 'Region 1', from room temperature to 600 K, and 'region 2', from temperature beyond 600 K. In 'region 1', H is desorbed from weak-traps and can quickly diffuse out of the specimen when kept at room temperature for an extended period whereas, in 'region 2', H is released from strongtraps and require a higher amount of energy for desorption [3], [10], [36]. The desorption peak

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© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon temperature (Tc) observed for DP800, DP1000, and CP1000 steel specimens were close to each other (461 K, 477 K, and 482 K respectively) but with different peak intensity. For TWIP steel, Tc was higher than other investigated steels at 556 K. The difference in peak intensity indicates a different concentration of H desorbed, with the highest amount for TWIP steel, followed by DP1000, DP800 and CP1000. Higher Tc in TWIP steel is associated with higher trap strength and slow H diffusion in austenitic microstructure [6], [7]. As a result, H from weak-traps in TWIP steel continues to desorb beyond 600 K, and the total H desorbed from the weak trap sites is considered to be the inflection point between the two regions (shown in The measure of H released from each region provides insight on H concentration in strong and weak traps. Figure 4.3 shows the total H desorbed and its distribution in strong and weak traps for the investigated steels. Comparing the two DP specimens, the total H concentration is higher in DP1000 than DP800. However, H concentration is almost identical in the strong-traps, leading to higher H concentration in weak-traps, suggesting increasing HE susceptibility in DP steels with strength. Between CP1000 and DP1000, under similar charging conditions, H pickup in CP steel is smaller than DP steel. However, in CP1000, a higher hydrogen concentration is present in strong-traps compared to DP1000. A similar

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tendency was observed and reported using the hydrogen permeation technique in the previous chapter (Chapter 3). In TWIP steel, the total H concentration was higher than the other steels due to longer charging period and high H solubility in austenite [6], [7]. A common observation for all investigated steels was that most of the H was present in weak-traps.

Although the trapped H quantity in a specimen depends on both the strength and microstructure, the microstructure has more influence on H distribution. [37]. . With increasing heating rate, Tc shifts towards higher temperatures with an increasing peak intensity, as with an increase in heating rate, the detection of H diffusing out of the specimen is at a higher temperature. The observed curves are the convoluted spectra of desorbed H from different traps present in the material. Thus, for calculating the Ea of each trap site, the TDS spectra need to be deconvoluted and fitted with different sub-peaks to identify Tc for each sub-peak. Consequently, the curves were deconvoluted using the peak fitting module in OriginPro 2016, assuming that the sub-peaks followed the Gaussian shape [15], [38]. similarly deconvoluted for all steel specimens, at all heating rates, with the results summarized in Table 4.2. DP and CP steel specimens were deconvoluted using two sub-peaks in 'region 1'.

Determination of Activation Energy (Ea) of traps

400
However, in region 2, DP steels exhibit a single peak at the heating rate of 800 K.h -1 (shown in the inset of Where ϕ is the heating rate in K.h -1 , Tc is the peak temperature in K, Ea is the activation energy of the traps in kJ.mol -1, and R is the universal gas constant in J.(K.mol) -1 .

The Ea of different traps in DP800 and DP1000 are below 20 kJ.mol -1 . The low Ea suggests the presence of weak-traps, which is consistent with the earlier observation that most H trapped in DP steels are present in weak-traps. For CP1000, the Ea of the weak-traps are also under 20 kJ.mol -1 . Additionally, trap sites with Ea in the range of 31 to 45 kJ.mol -1 are also present in CP steels. These traps possibly correspond to the cementite particles of bainite present in the CP steel. Galindo et al. observed that the energy required to remove hydrogen from cementite is close to 30 kJ.mol -1 [40]. or AgCN byproducts' precipitation, which could be distinguished by their particle morphology [28]. The Ag particles precipitated from reduction reaction are spherical, whereas the AgCN precipitates form needle-shaped crystals. In the case of DP1000 and CP1000, the spherical Ag particles are deposited preferentially along the ferrite-martensite interface. Additionally, in CP1000, some Ag particles are observed along the cementite particles present in the bainite (marked as B in Figure 4.8(e)). In the case of TWIP steel (

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Discussions

The total H concentration in TWIP steel (3.65 wppm) was significantly higher compared to the other steels attributed to its austenitic structure and longer charging period.

In 'region 1', the Tc for TWIP steel specimen was at 556 K, which also is at a much higher temperature compared to the other steel specimens, due to the presence of traps with higher activation energy, and the influence of activation energy required for lattice diffusion. In TWIP steel, the H diffusion coefficient is low, and the activation energy required for the lattice diffusion is comparable to the energy required for hydrogen desorption from weak-traps.

According to Lee et al., if the trap activation energy and the activation energy for lattice diffusion are comparable, the TDS curves are influenced by the bulk diffusion effect, shifting

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© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon the peak temperature (Tc) to a higher temperature value [6]. In the case of Ni (with f.c.c. structure similar to austenite), the peak temperature increased by 100 K, when the specimen thickness changed from 0.35 mm to 0.91 mm.

Between DP and CP steels, the H concentration in CP steel was much lower than DP steel, consistent with the observation in the previous chapter (Chapter 3). Under similar charging conditions, H concentration in CP steels is smaller because of lower trap density and comparatively lower C content in its chemical composition. However, though the trap density is less in CP steels, its trap strength is higher (as observed from activation energies of traps), resulting in a higher fraction of trapped H present in strong traps compared to DP steels. The higher trap strength in CP steel is due to the presence of cementite particles that have higher Ea for trapping H than ferrite or martensite [42]. It is an important consideration to design materials with higher HE resistance, as hydrogen present in strong-traps has a lesser influence on the embrittlement phenomenon [25].

Comparing the microstructure of DP steels, DP1000 contains a higher phase fraction of martensite than DP800. Further, due to the smaller martensitic grain size in DP1000, it has a larger volume of ferrite martensite interfaces (Chapter 3). Both martensite and ferrite-martensite interfaces are known to be potential trapping sites [8]. Additionally, the small size and higher volume of martensite, results in a higher density of strained ferrite around the interface, leading to higher dislocation density which serves as potential trap sites as well [43]. Thus, under similar charging conditions and specimen geometry, with the increase in martensitic content in dual phase steel, the trap density increase leading to higher hydrogen trapping. However, results indicate that there was no considerable change in the trapped H concentration in strong-traps. The above suggests that the martensite, ferritemartensite interface, and the ferrite near the ferrite-martensite interface, behave as weak-traps with low Ea. The Ea measured for both DP steels in the present study indicates the same, with Ea calculated to be below 20 kJ.mol -1 . The observed activation energy is associated with dislocations, grain boundaries, ferrite-martensite interface, and martensite, which are the major trapping sites in DP steels. However, the wide scatter and the overlapping nature of the data available in literature makes it difficult to associate the obtained activation energies

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© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon to a particular trap site. To highlight the challenge associated with one to one correspondence of traps sites with activation energy, Table 4.3 summarizes some available data.

Similar to DP steels, CP steels also exhibit traps with Ea below 20 kJ.mol -1 . These traps, similar to DP steels, are attributed to dislocations, grain boundaries, phase interfaces, and martensite. Also, it could be associated with the activation energy required for lattice diffusion of H in ferrite, reported as 8 kJ.mol -1 [44]. Besides, CP steels contain traps with Ea in the range of 30-45 kJ.mol -1 . The presence of different traps with similar activation energy values (31-45 kJ.mol -1 ) could be due to the different cementite particle sizes present in the material. The particle morphology affects the lattice strain around the particle, influencing the trapping energy [20], [39], [42], [45]. Though the particle size was not measured, TDS being a thermally destructive process alters the microstructure, or specifically the trapping sites, during H desorption measurement [10]. As a result, the carbide morphology during TDS analysis is modified, resulting in a series of activation energies, as observed in this case. As discussed previously, the peak obtained in 'region 1' of the TDS spectra for DP and CP steels, is fitted with two sub-peaks using Gaussian function where each sub-peak corresponds to a particular activation energy value. When the area under each sub-peak was measured with time, indicating the change in H concentration of weak-traps, it was dependent on the activation energy. For example, in DP1000, the area under the sub-peak with activation energy of 12.95 kJ.mol -1 reduced faster as compared to the sub-peak with the activation energy of 18.4 kJ.mol -1 . Figure 4.9 presents the change in the area of sub-peaks for the two steel specimens. Considering that the hydrogen desorption follows an exponential behavior, a time constant (τ) for the release of H can be associated with each trap site. A time constant is the period for an exponentially decaying quantity to be reduced to 36.8% of its original value. For DP1000, the time constant of hydrogen desorption was 2.74 and 4.78 hours for traps with activation energy of 12.95 and 18.4 kJ.mol -1 , respectively. Similarly, for CP1000, the time constant was 3.29 and 3.24 hours for traps with activation energy of 15.9 and 17.52 kJ.mol -1 , respectively. The τ value was proportional to the trap activation energy, which means that with an increase in trap activation energy, the desorption rate of H from the traps will reduce. [29], [52]. With this consideration, the above results indicate that the ferrite-martensite interfaces are the preferred hydrogen reduction sites for both DP and CP steels. Ferrite-martensite interface is known to be susceptible to HE due to the inherent difference in the mechanical properties of the two phases [53]- [55]. In the case of TWIP steel, the presence of Ag particles at both grain boundary and the grain interior indicates that both act as H trapping sites.

According to Akiyama et al. [28], the Ag deposition on the specimen surface is improved if the H concentration at the surface is high or if the diffusion of hydrogen is fast through a particular site to supply hydrogen for Ag reduction. At the discharge/detection side, after 2 hours of charging, Ag particles were deposited at the ferrite matrix and along the ferrite-martensite interface, indicating them to be the main diffusion path for H in DP and CP steel specimens. Ferrite is known to be the principal diffusion path for hydrogen in steels with ferrite matrix [56], [57]. The opinion on the role of grain boundaries as H diffusion path is not clear in the literature. Some researchers believe that grain boundaries act as trapping sites [27], [29], [58], whereas others consider grain boundaries as fast diffusion paths for hydrogen [37], [59], [60]. Since observations were ex-situ in the study, it is difficult to comment on initial positions where the Ag particles precipitated. However, Koyama et al. [29] observed in their in-situ study that the Ag particles initially precipitated at the grain interior, followed by precipitation along the grain boundary, in a pure iron specimen. For CP steels, the Ag particles around the cementite phase confirmed the trapping effect of cementite [30], [56], [57].

However, the H concentration around cementite was lesser than the ferrite or ferritemartensite interface, indicated by the smaller quantity of Ag particles precipitated. The low H concentration around cementite is due to the small H trapping capacity of cementite.

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Moreover, the H concentration around cementite is low as most of the H diffuse through ferrite or ferrite-martensite interface.

Conclusions

The objective of this study was to examine the H trapping and H adsorption behavior of DP, CP, and TWIP steel specimens using TDS and Ag decoration techniques. Among all investigated steel specimens, the H concentration was maximum in TWIP steel specimens due to the longer H charging period and austenitic structure. Under similar charging conditions between DP and CP steels, DP absorbs more hydrogen than CP steel due to its high carbon content and higher trap density. Among DP steels, with an increase in strength, the H concentration in the weak-traps increases while H concentration in strongtraps remains unchanged, increasing its HE susceptibility. The activation energy of traps observed in DP steels were low (< 20 kJ.mol -1 ), attributed to dislocations, grain boundaries, ferrite-martensite interface, or even the energy required for lattice diffusion of H. CP steels had similar activation energies for weak-traps, with some strong-traps (cementite) between 30-45 kJ.mol -1 . The higher activation energy of traps in CP steels compared to DP steels resulted in slow desorption of H. In TWIP steel, the activation energy of traps was higher than the other steel specimens. The main trapping sites in TWIP steel specimen were the austenite grain, the grain boundaries, and the AlN precipitates. In DP and CP steels, H was preferentially reduced along the ferrite-martensite interface whereas, in TWIP steel, H reduced at both grain interior and grain boundaries. Ferrite and ferrite-martensite interface were the H diffusion path in DP and CP steels whereas, for TWIP steel, the diffusion path could not be ascertained.

One notable result of this study was that the total hydrogen concentration is not the best parameter for comparing material sensitivity to HE but is one of many including, trap activation energy, H distribution in the trap, and its effusion rate from materials.

Acknowledgments

For this study, I would like to thank the ERASMUS+ scholarship for supporting my stay at Tohoku University, Japan. I would also like to express my heartfelt gratitude to

Introduction

Auto industries are increasingly replacing conventional high strength steels with Advanced High Strength Steels (AHSS) for fabricating Body In White (BIW) components, for environmental and economic reasons [1]- [7]. However, due to its high strength and complex microstructure, AHSS are susceptible to hydrogen embrittlement (HE), which restricts the use of AHSS to its full capacity [8]- [10]. Hydrogen could be introduced in the steels by any surface electrochemical reactions that occur during steelmaking, processing operations, or during service [6].

Hydrogen influences the mechanical strength by reducing the cohesive force between the metal atoms (HEDE), facilitating dislocation movement and formation of localized plastically deformed regions (HELP), aiding dislocations nucleation near the surface (AIDE), either singularly or in combination of more than one process [11]- [15]. Slow strain rates tensile tests are often employed to characterize HE by comparing the mechanical properties of specimens, with and without H [16], [17]. Results indicate that hydrogen assisted failure generally occurs in the plastic region, causing ductility loss and appearance of brittle fracture regions [6], [16]. Though the ductility loss is commonly acknowledged, the available literature presents mixed views on hydrogen's impact on tensile and yield strength. While some studies suggest a decrease in tensile s and yield strength, others report no significant influence of hydrogen [4], [10], [17], [18]. In some rare cases, HE occurs in the elastic region of the tensile curve, and in these cases, HE susceptibility is supposed to be very high [19].

Nevertheless, though there are several divergent opinions on the mechanisms and effects of HE, most of the researchers agree that HE is caused by diffusible hydrogen in the material. Hydrogen can either be present as trapped or as mobile hydrogen, which can diffuse towards the region of triaxial stress under stress, causing embrittlement [20]- [22]. Thus, the study of hydrogen transport during tensile tests can provide valuable information about the material's HE susceptibility. Though much data is available in the literature associated with hydrogen desorption during deformation, a little study is available related to the in-situ study of hydrogen desorption during deformation [23].

© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon

This study aims to analyze the hydrogen desorption behavior when mechanical energy is applied. The hydrogen charged AHSS specimens were subjected to tensile tests in vacuum equipped with quadrupole mass spectrometer (QMS) to detect desorbed hydrogen in situ at room temperature. The H desorption behavior was observed to evaluate the steels' susceptibility to hydrogen embrittlement. The test analysis is divided into three parts. The first part compares the tensile curves of uncharged and H charged specimens. The second part deals with the fracture surface analysis, and finally, the desorption spectra obtained during the tensile test measurements are studied.

Experimental Procedure

Materials

For the present study, TATA Steel Netherlands provided two DP steel sheets having a tensile strength of about 800 MPa (DP1000) and 1000 MPa (DP1000), a CP steel sheet having a tensile strength of about 1000 MPa (CP1000) and a Twinning Induced Plasticity (TWIP) steel sheet having a strength of about 1000 MPa in the form of 30 cm X 30 cm sheet of different thicknesses. The DP steels had an average thickness of 0.98 mm while CP and TWIP steel had 1.60 mm average thickness. The chemical compositions of materials are as shown in Table 4.1.

Table 5.1: Chemical Composition of the steels in wt% used in the tensile test study

The four steel specimen's microstructures were discussed in detail in the previous chapters (Chapter 3 and 4). The microstructure of DP steel contains ferrite and martensite with increasing martensitic phase fraction with an increase in strength. The CP steel had a more complex microstructure with ferrite, bainite, and martensite. The TWIP steel had a complete austenitic microstructure. The tensile specimens were machined along the rolling direction for the tests, according to the geometry shown in Figure 5

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for the tensile specimen of DP steels was 0.98 mm, and for CP1000 and TWIP steel was reduced to 1 mm.

Electrochemical Charging

Hydrogen was introduced in the tensile specimens by electrochemical charging with hydrogen. H diffusion coefficient in the specimen and its thickness, as discussed previously, determined the charging time (Chapter 4). In the case of DP and CP steels, a charging period of 3 hours was sufficient to saturate the specimens with hydrogen. For TWIP steel, it was difficult to saturate the specimen with H within a reasonable period due to the austenitic structure. So a 7-day charging schedule was followed, expecting H to penetrate to a depth of 300 µm (Chapter 4). The specimens were charged in 3% NaCl + 5 g.L -1 of NH4SCN aqueous solution with an applied current density of -1 mA.cm -2 . A platinum (Pt) wire was used as the counter electrode. After charging, the specimens were cleaned with acetone and ethanol in an ultrasonic bath and placed the tensile test chamber.

Tensile Test

Hydrogen desorption during the tensile test was measured using a tensile test machine installed in a vacuum chamber. The instrument used a dry rotary and a turbomolecular pump for vacuuming. A quadrupole mass spectrometer (QMS) was connected in the setup to detect the desorbed H. The equipment used for the study is as shown in Figure 5.2. The aim was to observe hydrogen desorption spectra during the applied strain.

The chamber was brought to the atmospheric pressure by a flow of pure nitrogen gas into the chamber. Once at atmospheric pressure, the nitrogen flow was stopped, and the specimen was placed in the tensile test machine. The tensile specimen was fixed in the tensile test machine using a set of jigs and fastened tightly with screws. The chamber was then closed and the vacuumed using the dry and turbomolecular pumps.

After 30 minutes of vacuuming, the tensile test measurement started with a strain rate of 0.03 min -1 . In this interval of 30 minutes, a certain amount of hydrogen desorbed from the specimen. Due to the setup limitation, it was difficult to reach a satisfactory vacuum level, but it was necessary to start the tensile test to prevent further hydrogen loss. Once fractured, the specimens were carefully removed for SEM observation. The transverse surface of the fractured specimens were etched in 2% Nital to observe microstructural influence on the surface cracks.

Results and Discussion

Hydrogen concentration

Reference coupons of 15 mm X 20 mm were charged under the same charging conditions as the tensile specimens, and the H concentration was measured using thermal desorption spectroscopy (TDS) according to the procedure described previously (Chapter 4).

Figure 5.3 presents the results with total hydrogen concentration in DP800, DP1000, CP1000, and TWIP steel were 1.20, 1.77, 1.08, and 3.65 wppm, respectively. The high hydrogen content in TWIP steel is attributed to the longer charging period and the higher H solubility in the austenitic phase. In the case of DP steels, the DP1000 contained a higher amount of hydrogen as compared to DP800 due to the presence of higher trap density (Chapter 3). The hydrogen content in CP1000 was lower than DP1000 due to its microstructural difference discussed previously (Chapter 3 and Chapter 4).

Mechanical Properties

Figure 5.4 (a to d) presents the engineering stress-strain curves of the uncharged and hydrogen charged specimens, with the mechanical properties summarized in Table 5.1.

The shift in linearity in the elastic region is due to the machine slackness, caused by relative motion between the tensile specimen and the fixture block at low strain. researchers have observed no change in the deformation behavior whereas, some have reported a significant loss in Y.S. after H charging [3], [6]. The ductility loss was normalized in terms of the 'Hydrogen Embrittlement Index' (HEI) using Equation 5.1. HEI can vary between 0 and 100%, where 0% indicates no ductility loss and, 100% indicates complete ductility loss [24], [25]. The HEI calculated using Equation 5.1 is presented in Table 5.1.

𝐻𝐻𝐸𝐸𝐻𝐻 = 𝐸𝐸𝐸𝐸 𝑢𝑢𝑛𝑛𝑐𝑐ℎ𝑎𝑎𝑎𝑎𝑎𝑎𝑎𝑎𝑎𝑎 -𝐸𝐸𝐸𝐸 𝑐𝑐ℎ𝑎𝑎𝑎𝑎𝑎𝑎𝑎𝑎𝑎𝑎 𝐸𝐸𝐸𝐸 𝑢𝑢𝑛𝑛𝑐𝑐ℎ𝑎𝑎𝑎𝑎𝑎𝑎𝑎𝑎𝑎𝑎 𝑋𝑋 100 % Equation 5.5
Where, 𝐸𝐸𝑙𝑙 𝑢𝑢𝑛𝑛𝑢𝑢ℎ𝑎𝑎𝑎𝑎𝑎𝑎𝑎𝑎𝑎𝑎 and 𝐸𝐸𝑙𝑙 𝑢𝑢ℎ𝑎𝑎𝑎𝑎𝑎𝑎𝑎𝑎𝑎𝑎 are the elongation of the uncharged charged specimen, respectively. Compared to CP steel, The DP and TWIP steels show significant ductility loss.

CP steel is expected to have better HE resistance compared to other investigated steels, especially DP steels, due to the presence of bainite in the microstructure [3], [26]- [28]. Bainite is associated with stress-free, homogeneous, and refined structure, with mechanical strength between ferrite and martensite, reducing the strength difference between the two phases [26].

Additionally, the cementite particles act as strong-traps and decrease the diffusible H concentration [29]. The argument supports the earlier works wherein CP steels, the amount

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© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon of hydrogen present in strong traps are much higher compared to DP steels (Chapter 4).

However, the observed 2% HEI could not be conclusively associated with HE, as it could be within the range of experimental error.

Between DP800 and DP1000, both steels exhibit similar HEI. A higher HEI was expected for DP1000 due to the higher tensile strength and martensitic content, as cracks usually initiate at the ferrite-martensite interface, or in the martensite, and propagates through ferrite [18], [30], [31]. HE behavior of TWIP steel is inconsistent in the available literature. While Some studies suggest the susceptibility of TWIP steel to HE, others report TWIP to be resistant to HE [6], [32]- [35]. HE in TWIP steel is suspected from hydrogen assisted grain boundary decohesion, hydrogen promoted deformation twinning or martensite transformation [36], [37]. However, due to the high Al content, the martensitic transformation can be ruled out [38].

The HEI values in the present study are lower than those reported in the available literature due to the experimental conditions of the study [10], [24], [29], [31], [36], [37], [39].

Usually, during the tensile test, the HEI values reported in the literature are under in-situ hydrogen charging conditions, resulting in higher hydrogen concentration, leading to higher ductility loss. In this study, the specimens were charged with H ex-situ, followed by the tensile test. Another probable reason was the time interval between H charging and the tensile test for creating a vacuum, which resulted in H desorption from the specimen. At atmospheric conditions, 3 hours after H charging, about 61% of diffusible hydrogen from DP1000, and 65% from CP1000 was desorbed (Chapter 4). Additionally, H desorption is faster in a vacuum leading to the desorption of a higher H concentration from the specimens [40]. center to the surface, depleting hydrogen concentration from the specimen center. Thus, the mixed-mode of failure is observed only near the specimen surface for the charged specimen.

Fractographic Analysis

In the absence of hydrogen to influence the stress field (lowering crack propagation resistance), the quasi cleavage fracture areas change into ductile fracture with increasing distance from the specimen surface towards the center. The observation was in agreement with the low HEI value observed from the stress-strain curve. and due to H accumulation during plastic deformation [43], [44]. In the necking region, due to the applied strain, new defects are generated, trapping H. If the hydrogen concentration reaches a critical value, it will cause embrittlement with cracks perpendicular to the tensile axis. As the charging in the present study was ex-situ, these crack were fewer in number.

Among DP1000 and CP1000 charged specimens, CP1000 had a higher density of such cracks compared to DP1000. Since H desorption is faster in DP steels than CP steel, and CP steels have stronger traps (Chapter 4), under deformation, the H concentration in CP steel around the necking region in CP steel could have been higher.

In Figure 5.7(b), there are some longitudinal cracks along the tensile axis in CP steel. However, these cracks are the consequence of tensile specimen failure, rather than the cause. The decrease in the work hardening rate triggers plastic flow localization, which contains numerous possible sites for crack nucleation [45]. As the stress in the thickness direction is tensile, with a significant stress value near the specimen half-thickness, it tends to open the cracks parallel to the sheet plane in 'Mode 1'. As the plastic flow localization zone becomes increasingly damaged, the two half specimens separate, and the longitudinal crack opens up [45]. With similar cracks present in an uncharged specimen, it could be concluded that hydrogen promotes, and does not initiate, crack formation and propagation. Koyama et al. [30] suggested that the hydrogen's influence on cracking of martensite or ferrite/martensite interface is due to the reduction of cohesive energy, i.e., the HEDE mechanism. Furthermore, the subsequent crack propagation in ferrite is explained by HELP.

Martensite cracking and ferrite martensite decohesion is also observed in the CP1000 (Figure 5.8 (e and f)). The faceted features around the inclusion show brittle failure, resulting from stress concentration around the inclusions. Figure 5.9(e and f) shows the fracture surface of the uncharged and charged specimen around the edge, respectively. In the case of an uncharged specimen surface, the edge shows a ductile failure with regions of shear tearing and ductile dimples. In charged specimen, brittle fracture is observed with increasing ductility with increasing distance from the specimen surface. In the presence of hydrogen, the cohesion between twin boundaries of freshly formed mechanical twins, during tensile deformation, is low, and fracture can propagate through the twin boundaries causing brittle failure by HEDE [17]. A steady increase in the hydrogen desorption with an increase in the stress was observed in the elastic region, reaching a maximum at the point corresponding to the yield point. Beyond the yield point, the hydrogen desorption decreased with the increase in the plastic strain. The results are consistent with the literature, where the hydrogen desorption increases in the elastic stage, promoted around the yield point and then decreased with further increase in the plastic strain [16]. In the case of a charged CP1000 specimen, an additional peak could be observed in the plastic region, corresponding to the necking region in the stressstrain curve. In the elastic region, the H diffusion coefficient increases with tensile stress due to the lattice expansion [46]. In addition to the lattice expansion, the following factors influence the hydrogen diffusion: (a) the temperature rise caused by the heat generation during desorption, (b) oxide film fracture on the specimen surface, (c) uneven specimen surface caused due to deformation, (d) hydrogen transport due to moving dislocations (e) increase in the hydrogen trapping sites due to plastic deformation and (f) increase in hydrogen absorption capacity due to decrease in chemical potential [16], [23]. Many researchers observed H desorption commences only near the onset of yielding and not in the initial elastic region, unlike the present observation [23], [47]- [49].

Hydrogen desorption during tensile test

Matsumoto et al. [16] suggested the difference is in the initial dislocation density of the investigated steel specimen. If the dislocation density is initially high, small plastic deformation can occur even under an elastic region. Since the present study involves AHSS, they have a high initial dislocation density due to their complex microstructures [50], [51].

Hydrogen is desorption in the elastic region by the moving edge dislocations, due to high initial dislocation density. When the stress is near the yield point, the H desorption becomes prominent due to the dislocation multiplication by mechanisms such as the 'Frank-Read mechanism' [52]. Beyond the yield point, the hydrogen desorption decreased for DP800, DP1000, and CP1000 specimens due to the generation of defects in the plastic region, which behave as hydrogen traps. In the available literature, the decrease in the hydrogen diffusion coefficient of hydrogen due to plastic deformation is well documented [46], [53]- [55].

The lattice defects created by plastic deformation, trap the diffusible hydrogen, reducing the amount of hydrogen available for desorption. The formation of vacancy clusters during plastic deformation is promoted in the presence of a hydrogen-containing environment, as hydrogen stabilizes the vacancies formed during plastic deformation [56], [57]. Besides, after the yield point, the distance traveled by the dislocations is shorter, as multiple slip systems are activated, and screw dislocations also start to move [16]. The binding energy of hydrogen with screw dislocations is usually smaller than edge dislocations resulting in a lower H concentration of hydrogen (Cottrell atmosphere of hydrogen) around the screw dislocations [23]. Furthermore, the dislocations intersect each other, leading to a tangled cell structure, reducing the hydrogen concentration reaching the specimen surface, resulting in lower H desorption [16], [23]. Thus, vacancy generation, screw dislocation mobility, and multiple slip planes activation, all contribute to the reduction of hydrogen desorption during plastic deformation. The sharp peak observed for DP1000 and CP1000 specimens at fracture, was

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© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon due to the sudden desorption of accumulated hydrogen, on the abrupt exposure of the new surfaces [23].

The small peak near the necking region in the CP1000 steel is due to H released from the inherent strong-traps present in the material (Chapter 4). Near this region, the H present in the strong-traps gain sufficient energy and are released into the lattice, increasing the diffusible H concentration. It agrees well with the earlier observations where crack perpendicular to the tensile axis, near the necking region, was observed in CP1000 specimens.

In the case of TWIP steel, the hydrogen desorption did not reduce immediately after the onset of plastic deformation, as shown in Figure 5.11(a). The hydrogen desorption continued, during the early stages of the plastic deformation, before the H desorption decreased. For further investigation, the TWIP steel was charged longer (14 days), and tested.

The 14 days H charged specimen revealed a small peak near the yield point, followed by a higher peak in the plastic region, following which hydrogen desorption from the steel specimen progressively decreased until fracture (Figure 5.11(b)). The hydrogen desorption spectra for the TWIP was similar to other steels up to the yield point. However, beyond the yield point, the hydrogen desorption in TWIP steel increases, unlike the other investigated steels. The observed peak in the plastic region, in the 14 days hydrogen charged TWIP steel, could be due to the hydrogen released due to the © 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon formation of deformation twins. Koyama et al. [58] observed that deformation twin boundaries act as crack initiation sites as well as crack propagation paths during the tensile test. The deformation twin boundaries act as potential hydrogen trapping sites, influencing the HE behavior of the TWIP steels [24], [58]. During the initial period of plastic deformation, hydrogen trapped in the twins is released into the lattice [59] and transported to the surface by the moving dislocations facilitating H desorption. At the same time, the newly generated defects trap the diffusible hydrogen. Thus, there exists a competing effect, and the quantity of desorbed H increases, until the amount of hydrogen released from the twins, is more than the amount of hydrogen trapped. The H desorption rate decreases when the quantity of hydrogen being trapped exceeds the quantity of hydrogen released. The H desorption continues to decrease until fracture. Though the above argument appears to be plausible, further investigation is required to support the above reasoning.

Conclusion

The study focuses on the influence of hydrogen on the mechanical properties of DP, CP, and TWIP steel, as well as the desorption behavior of hydrogen under stress. For this study, the tensile specimens were electrochemically charged with H and tested under vacuum. The salient findings of the study are summarized as follows:

• There was an observable ductility loss in all tested specimens in the presence of H, but the deformation behavior remained unaffected.

•

The susceptibility to HE was TWIP > DP > CP steels. The high HE resistance of the CP steels is attributed to the bainite microstructure, low H uptake, and presence of strongtraps, which reduce the diffusible H concentration in the lattice.

•

In DP and CP steels, cracks resulted from martensite cracking and decohesion at the ferrite-martensite interface, initiated due to the strain incompatibility between the two phases.

• During the tensile test, the H desorption increased with increasing strain up to the yield point due to lattice expansion and H transport by an edge dislocation.

• Beyond the yield point, the hydrogen desorption markedly decreased for DP800, DP1000, and CP1000 due to the generation of defects in the plastic region, which trapped the diffusible hydrogen.

• A sharp increase in H desorption was observed near fracture due to the sudden desorption hydrogen on abrupt exposure of the new surfaces.

•

In CP1000, an H desorption peak is observed near the necking region, indicating the H desorption from the strong-traps.

•

In TWIP steel, a desorption peak is observed in the plastic region, caused by the release of hydrogen due to the formation of deformation twins.

Influence of Hot Dip Galvanized Coatings on the Hydrogen Permeation

Behavior of DP steels

Resumé

The investigates the hydrogen permeation characteristics of dual phase steel with and without the galvanized layer. Dual phase (DP) steels are first-generation advanced high strength steels (AHSS), known to be susceptible to hydrogen embrittlement (HE). The electrochemical investigations were carried out in a modified Devanathan-Starchuski cell in deaerated 0.1 M K2SO4, 0.02M H2SO4 solution at two charging current densities (-1.5 and -10 mA.cm-2). The diffusion coefficient of H in Zn was found to be four orders of magnitude smaller than that in Fe. The Zn layer appeared to inhibit the hydrogen entry and permeated a small volume of hydrogen when charged at high current density. On the other hand, the Zn layer enhanced the hydrogen entry into the bulk of the Fe substrate when charged at low charging current density. The increase in the volume of absorbed hydrogen was attributed to the Zn dissolution that occurred during the electrochemical test. The Zn layer was observed to retard the effusion of absorbed hydrogen from the bulk of the Fe substrate, increasing the concentration of hydrogen in the material. Finally, a model is proposed explaining the role of the Zn layer in increased hydrogen absorption by Fe substrate under low charging current density.

Introduction

Advanced High Strength Steels (AHSS) offer a lighter, sturdier, safer, and more sustainable alternative to the conventional steels for the fabrication of Body In White (BIW) in the auto industry [1], [2]. Dual Phase (DP) steels are the first-generation AHSS designed specifically for BIW and are widely used in vehicle components like underbody cross members, bumpers, B-pillars and door beam [3]- [6]. DP steels use low C-Mn-Si composition and have a ferrite and martensite microstructure [7], [8]. Though DP steels provide superior mechanical attributes, the use is limited by their susceptibility to hydrogen embrittlement (HE) [9]- [16].

It is common in the auto industry to add a Zn layer on steel, to protect the substrate from corrosion, in the form of cathodic and barrier protection [17]. Zn coating is applied by either electrodeposition or by hot-dip galvanization [18], [19]. Zn electrodeposition provides uniform coating thickness and a good surface but is expensive. Galvanization has poor paintability and surface finish, but is cost-effective and preferred for the inner parts of BIW. When a sacrificial coating corrodes (cathodic protection), H is generated on the surface and is absorbed by the exposed Fe substrate. The absorbed hydrogen can cause hydrogen embrittlement (HE) if present in sufficient quantities. Though the practice of hot-dip galvanization is common in the auto industry, there is limited knowledge of its effect on the HE susceptibility [20].

In the published literature, the role of Zn coating on the embrittlement susceptibility of the underlying steels substrate remains ambiguous. Several authors have reported that the electrodeposited Zn layer effectively inhibits the H absorption in steels providing high HE resistance [21]- [25]. Boiadjieva et al. [21] reported that coatings as thin as 0.5 µm, successfully block the hydrogen entry in the underlying substrate. According to the references, the high H overvoltage on Zn and the low H diffusion coefficient through the electro-galvanized layer is responsible for the improved HE resistance. However, contrary to electro-galvanized coatings, hot-dip galvanized coatings are reported to increase the substrate's HE susceptibility [17], [26]- [30]. Roberge et al. [27] suggested that the microstructural defects present in the galvanized layer can lead to increased hydrogen 150 © 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon susceptibility. However, the role of the galvanized Zn layer on increased susceptibility of the steel substrate to HE is not thoroughly investigated.

The study aims to investigate the influence of the galvanized layer on the H uptake in DP steels. The H diffusion coefficient in the substrate and the galvanized layer were compared using the electrochemical permeation technique. The influence of the galvanized layer on H absorption in the substrate was investigated in detail. Finally, the role of cathodic current density on the H uptake, in the presence of Zn coating, was discussed.

Experimental Procedure

Materials

Dual Phase steels, with (DP1000 Zn) and without a galvanized layer (DP1000), having a tensile strength of about 1000 MPa, were used in this study. TATA Steel Netherlands provided the 30 cm X 30 cm steel sheets with different thicknesses. The chemical composition of DP1000 was 0.14 wt.% C, 2.14 wt.% Mn, 0.05 wt.% Si and 0.57 wt.% Cr whereas the chemical composition of DP1000 Zn was 0.15 wt.% C, 2.31 wt.% Mn, 0.1 wt.% Si and 0.56 wt.% Cr. The small chemical composition difference was assumed not to have significantly affected the measured electrochemical properties. DP1000 sheet was 0.98 mm thick, and the DP1000 Zn sheet was 1.24 mm thick.

Specimen Preparation

For the permeation tests, 35 mm X 35 mm coupons were cut and ground to 600 grit SiC. Both specimen surfaces were ground in DP1000, whereas, only one specimen surface was ground in DP1000 Zn, to remove the Zn layer (detection surface). Subsequently, the specimens were cleaned twice with acetone and once with alcohol in an ultrasonic bath and air-dried. A Pd layer was deposited on the detection surface, according to the procedure in For the present study, the hydrogen reduction was carried out galvanostatically with two different current densities (-1.5 mA.cm -2 and -10 mA.cm -2 ), and the oxidation was carried out potentiostatically at +0.2 V w.r.t. saturated Ag/AgCl electrode. A graphite auxiliary electrode (AE) was used at the detection side and a Pt mesh as the AE at the charging side. All experiments were performed at 35 °C, maintained using a water jacket.

The electrochemical hydrogen permeation experiment was carried out according to the ISO 17081 specifications [31], [32]. Specimens with the exposed surface area of 7 cm² were clamped between the two compartments. The anodic compartment was filled with 0.2 M KOH solution (pH = 13; Conductivity = 45.3 mS.cm -1 ) and anodic potential (+0.2 V w.r.t. R.E.) was applied. The results were measured using Gamry Framework data acquisition software v6.32. Once the measured current density stabilized below 0.3 µA.cm -2 , deaerated 0.1 M K2SO4, 0.02M H2SO4 solution (pH = 2.08; Conductivity = 26.5mS.cm -1 ) was added slowly at the charging side and the cathodic surface was charged galvanostatically. Ar gas flowed through both the electrolytes throughout the experiment for deaerating the electrolytes. In the absence of oxygen, H oxidation is the only electrochemical reaction at the anodic surface and provides a direct measure of the H atoms permeating through the specimen. [20], [24], [33]. The galvanized layer was measured to be 6-10 µm thick, consistent with the thickness reported by other authors [30], [30], [34], [35]. indicates the sacrificial behavior of the Zn layer. Zn has a lower exchange current density for hydrogen evolution than Fe, accelerating H.E.R, and reducing the dissolved H concentration at the Zn coated specimen surface [20], [37]. The H mobility through a coating depends on its diffusion coefficient, and the coating's ability to adsorb and absorb hydrogen [38]. As shown later, using the permeation test, the sub-surface H concentration in DP1000 Zn is much lower than DP1000, at higher current density.

Results and Discussion

Potentiodynamic Polarization Test

Hydrogen Permeation Test

For the H permeation test, two different charging regimes were followed (a) -10 mA.cm -2 and (b) -1.5 mA.cm -2 . The former would be referred to as the 'high charging current density' and the latter, as the 'low charging current density.' Figure 6. In the case of high charging current density (Figure 3(a)), the permeation current density of DP1000 is significantly higher than DP1000 Zn. The steady-state permeation current density of DP1000 was 7.79 ± 0.57 µA.cm -2 whereas, for DP1000 Zn, the steady-state current density was 0.67 ± 0.17 µA.cm -2 . The H diffusion coefficient in DP1000 was calculated according to the Equation 6.1 [39], [40]

𝑖𝑖 𝑛𝑛 = 2𝐿𝐿 √𝜋𝜋𝜋𝜋𝑡𝑡 ∑ exp �- (2𝑛𝑛+1) 2 𝐿𝐿 2 4𝜋𝜋𝑡𝑡 � ∞ 𝑛𝑛=0 Equation 6.19
Where, in is the normalized current density, D (m 2 .s -1 ) is the diffusion coefficient, L is the specimen thickness (m), and t is the time (s). Equation 6.1 is not useful for H diffusion coefficient calculation in DP1000 Zn, as the DP1000 Zn represents a bimetallic layer. H diffusion coefficient in DP1000 Zn was calculated using a model proposed by Song et al. (Equation 6.2) [41]. The model assumes that during the H charging, the proportion of the H concentration across the interface is equal to the ratio of their corresponding equilibrium solubility. Zn layer thickness was considered as 8 µm and Fe substrate thickness as 1.1 mm. According to Cao et al. [23], the K for a Zn plated specimen on Fe is 2.6.

𝑡𝑡 𝐸𝐸𝑎𝑎𝑎𝑎 = 𝐸𝐸 1 2 2𝜋𝜋 1 + 𝐸𝐸 2 2 2𝜋𝜋 2 + 𝐾𝐾𝐸𝐸 1 𝐸𝐸 2 𝜋𝜋 2 Equation 6.20
Where tlag is the lag time [26], [42]- [44],.D1 and D2 are the H diffusion coefficients in Zn and Fe respectively, l1 and l2 are the thickness of Zn, and Fe layers and K is the ratio of the concentration of hydrogen across the interface.

The calculated H diffusion coefficient in DP1000 and DP1000 Zn was 8.4 X 10 -11 m².s -1 and 4.59 X 10 -15 m 2 .s -1 , respectively. The permeation current density behavior in DP1000 and DP1000 Zn at 'high charging current density' is explained in terms of hydrogen interaction with Fe and Zn. The adsorbed H is either absorbed into bulk material or leaves the surface as hydrogen gas, influenced by the H diffusion coefficient, and the hydrogen

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© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon absorption rate constant of bulk metal [45]. As a result, hydrogen permeation is weaker in systems with a high concentration of active centers for H reduction, and have weak bonds between these centers and hydrogen atoms [46]. Zn has a weak H adsorption capability than Fe due to its h.c.p structure [21]. Additionally, the H diffusion coefficient in Zn is four orders of magnitude slower than Fe. As a result of these factors, the permeation current density in DP1000 Zn is low. The H bubbles partially cover the surface, increasing the actual current density under galvanostatic conditions due to the reduced exposed area, increasing the cathodic overpotential. As these bubbles leave the surface, the surface is exposed, shifting the current and the potential to their original values. Since the HER overpotential is higher in DP1000

Zn specimen, the potential fluctuation is more prominent in DP1000 Zn. The explanation also supports the observed spherical particles at the specimen surface.

© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon

The potential evolution trend for DP1000 is similar for both the charging condition. However, the potential-time curve for DP1000 Zn is different at low charging current (Figure 6.5(b)). At low charging current density, the potential remains steady for a period during charging. This potential corresponds to the Zn potential for the applied current density (as verified from the polarization curves). With further H charging, the potential shifts to a more noble value and stabilizes at a potential corresponding to that of DP1000, when charged at the same cathodic current density, clearly showing that the Zn layer dissolution during H charging and exposing the bare underlying substrate. Roberge et al. [27] studied the effect of the galvanized layer on HE of steel in the deaerated cement environment at OCP and reported a similar evolution of the potential-time curve. They suggested that the metallurgical defects of the galvanized layer facilitated the hydrogen entry, and these defects would be the preferred hydrogen evolution and absorption site. They further proposed that such microstructural features would act as stress concentrators if there is any applied stress.

The effect is accelerated if the oxygen content is low, which is similar to the present case.

Zhang et al. [29], [30] also reported increased hydrogen uptake in the case of defects in the Zn coating of galvanized steel. Results indicate that during cathodic charging at lower overpotential, the galvanized Zn layer initially dissolves along the grain boundary, exposing the bare substrate, which results in the increased hydrogen entry. From the above observations, a model is proposed to explain the observed trends for DP1000 Zn specimen, when charged at low cathodic current density. However, before proposing the model, it was considered interesting to digress a bit to determine if the Zn coating trapped H during electrochemical charging or not.

GDOES Measurements

The DP1000 and DP1000 Zn specimens were electrochemically charged in 3 wt% NaCl + 5 g.L -1 NH4SCN solution for seven days, at an applied current density of -1 mA.cm -2 .

After the electrochemical charging of the specimens with hydrogen, the profile of hydrogen to the distance from the surface was measured using glow discharge optical emission spectroscopy (GDOES), as shown in Figure 6.7. For DP1000, the H concentration decreased with increasing distance from the surface. In DP1000 Zn, the H concentration was small in the Zn layer; however, across the Zn-Fe interface, there was an H build-up, increasing the H concentration in the Fe substrate. specimen. The potential at the Zn layer was initially stable, followed by a rapid change and then finally stabilized to a value similar to that of the bare substrate. With a higher Zn dissolution rate, the exposed area of the steel increases, leading to a further increase in the cathodic current density, as shown in Figure 6.9(b). At this point, the cathodic potential remains unaffected, and only the cathodic and anodic reactions are affected to maintain the galvanostatic condition. This process, in some form, is like an autocatalytic process. Initially, the H concentration reduced at the Fe surface was not significant, as the exposed area was small, and H reduction on the Zn surface was dominant.

However, with more Fe surface exposed to the electrolyte, the H reduction reaction in Fe becomes dominant. Due to higher reduction kinetics of H in iron and higher cathodic overvoltage, Zn dissolves rapidly. Since the actual cathodic current density is higher than the applied cathodic current density, the permeation current is quite high. The reaction continues until the Zn is dissolved to an extent, that the Zn dissolution reaction is no longer sufficient to compensate for the high actual cathodic current density (Figure 6.9(c)). , which corresponds to the maximum of the permeation curve and the stable region of the potential curve. At the point II, the potential rises rapidly whereas, the permeation current starts to decrease as the Zn is no longer sufficient to compensate for the high cathodic reaction, the potential of the system shifts to a noble value, to reduce the cathodic overpotential of H reduction on the Fe surface. It reduces the H concentration at the surface resulting in lower hydrogen entry into the material, which leads to a decrease in the permeation current.

At point III, the potential reaches a stable value, and the permeation current density reaches a steady-state (Figure 6.9(d)). These values correspond to the values of the bare metal with the complete dissolution of the Zn layer.

Conclusion

The effect of the galvanized layer on the HE susceptibility of the underlying DP The results in this work are yet to be published in a scientific journal but presented in several international conferences held in India, Poland, Czech Republic, and China, allowing to gain immense exposure to current research in this field. This chapter summarizes the salient results.

The Hydrogen diffusion coefficient

The quantitative and qualitative nature of the hydrogen diffusion behavior was investigated in DP and CP steels using hydrogen permeation studies. Precharged specimens were subjected to partial rise and fall transients instead of the complete rise and fall transients to ensure maximum stability of the charging surface during cathodic charging and minimum corrosion-related issues during discharging. The diffusion coefficients were calculated from experimentally obtained curves using equations based on Fick's second law and modified by Laplace transformation. The range of diffusion coefficient values obtained for DP800, DP1000 and CP1000 were 1.02 to 1.68 X 10 -10 m 2 .s -1 , 0.87 to 1.60 X 10 -10 m 2 .s -1 and 1.79 to 3.16 X10 - 10 m 2 .s -1 respectively. The H diffusion coefficient in DP steels was attributed to the banded structure in the specimen cross-section perpendicular to the direction of the diffusion of hydrogen, which increased the effective diffusion pathway. The diffusion coefficient values observed were well within the range presented in the literature and was an order lower than that of pure iron reported in the literature. It was due to the H interaction with different traps in the material, which effectively retards the diffusion of hydrogen in the material.

© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon

With the lowering of the charging current density, the apparent diffusion coefficient, as well as the lattice and total hydrogen concentration decreased. As the lattice hydrogen concentration is in dynamic equilibrium with trap sites, it leads to the release of hydrogen from reversible trap sites. As more and more reversible trap sites emptied, it slows down the diffusion of hydrogen in the steel. As a result, the apparent diffusion coefficient of H decreases in reducing the charging current density. Another possible reason could be that as we lower the charging current density, the H concentration gradient across the specimen decreased, leading to a decrease in the apparent diffusion coefficient.

Hydrogen Concentration in AHSS

On increasing the cathodic current density or the cathodic overpotential, the trapping capacity of the traps decreased. The trapped hydrogen concentration tends to saturate at higher current density value. However, it tends to saturate much faster than the total hydrogen concentration. It means that when the charging current density increases, the increase in the trapped hydrogen concentration is small, and the diffusion coefficient of hydrogen increases, indicating the decreasing influence of traps on the apparent diffusion coefficient.

The total hydrogen concentration among DP and CP steels was DP1000 > DP800 > CP1000. However, though the lattice hydrogen concentration values are higher for DP1000 than DP800 at lower charging current densities, at higher charging current densities, the concentration of hydrogen in the lattice of DP800 is higher than DP1000. Martensite, due to its high dislocation density, acts as a hydrogen trap and as a result, DP1000 has higher trap density as compared to DP800, as more and more hydrogen enters the specimens at higher current densities, the traps of DP800 saturates quickly as compared to DP1000 leading to enrichment of hydrogen in the lattice of DP800. On the other hand, since DP1000 has higher trap density, hydrogen continues to fill up the trap sites leading to comparatively lower hydrogen concentration in the lattice sites. This observation was particularly interesting in terms of embrittlement phenomenon, as the crack initiation in dual phase steels occurs due to the cracking of martensite and propagating through the ferrite-martensite interface. The © 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon martensite or the martensite/ferrite interface thus can act as a trapping site where hydrogen can be reversibly or irreversibly trapped depending upon the activation energy.

The H distribution in the bulk material during the permeation experiment was evaluated as the ratio between the change in the hydrogen concentration at the discharging side and the charging side for DP and CP steels. The ratio obtained for the different steels were 0.30 for DP800, 0.34 for DP1000, and 0.28 for CP1000. It is consistent with the observation of other researchers that the distribution of hydrogen along the cross-section of the material is nonuniform, and a higher concentration of trapped hydrogen near the charging surface. Though the observation is just speculative, further information in this context could provide valuable information about the distribution of hydrogen in the material during hydrogen charging.

From the permeation experiment, if we compare the total reversible traps emptied during the fall transient, we observe that the value is highest for DP800, followed by DP1000 and CP1000. If all the reversible traps emptied during the permeation fall transient, then it should indicate that DP800 has the highest number of reversible traps, which logically should not be the case. One possible explanation is that during fall transient, not all reversible traps are emptied but only the ones with very low binding energies. Thus DP1000 and CP1000

possibly have hydrogen stored in traps with higher binding energy, and conditions during the fall transient are not sufficient enough to release hydrogen from such sites. Alternatively, as explained above, a fraction of the hydrogen released from the reversible traps effuse out through the charging side, thus remain undetected. As a result, the value of the total trap density calculated is different than the actual trap density present in the material.

Diffusible and Total Hydrogen Concentration

The traps present in the material attract hydrogen as the lowering of energy is associated with the H transfer from t lattice sites to these defects. On the application of stress, these traps can behave as source or sinks for diffusible hydrogen based on their activation energy (Ea).

Under the present investigated conditions, the amount of hydrogen trapped by weak trapping sites in DP800, DP1000, CP1000, and TWIP steel were 1.08, 1.65, 0.85, and 2.63 wppm, respectively. On the other hand, the amount of hydrogen trapped by strong trapping sites in DP800, DP1000, CP1000, and TWIP steel were 0.12, 0.12, 0.23, and 1.02 wppm, respectively. In DP steels, this indicates that with an increase in strength, the amount of diffusible hydrogen in the material increases, making it more susceptible to HE, whereas, with an increase in strength, the strongly trapped hydrogen is not significantly affected. The conclusion drawn from the observation is that in the investigated steels, most of the trapped H was present in weak-traps, and microstructure influences the HE susceptibility more than the strength levels.

Comparing DP1000 and CP1000, the total hydrogen present in CP1000 is lower than DP1000 but contains a higher amount of strongly trapped hydrogen, indicating less numerous but stronger traps in CP1000 as compared to DP1000. This behavior was due to the presence of cementite particles in CP steels, which act as strong-traps, whereas the banded cross-section of the DP steels increases the number of traps that hydrogen encounters while diffusing through the bulk material.

Activation energy of the trap sites

The activation energy (Ea) values of the hydrogen traps present in the material can provide valuable insight about the susceptibility of a material to HE. In this study, the activation energy of the investigated steels were determined both by permeation studies and thermal desorption spectroscopy. it difficult to do a one to one correspondence between the activation energy and the trap sites present in the material. In the case of CP steels, the traps with activation energy between 30-45 kJ.mol -1 are associated with the cementite particles present in the bainite. In the case of TWIP steel, the activation energy of 51.75 kJ.mol -1 is associated with the release of hydrogen from austenite.

Using the permeation technique, the calculated activation energy for DP800 was 33.28 kJ.mol -1 , 37.82 kJ.mol -1 for DP1000, and 40.33 kJ.mol -1 for CP1000. The discrepancy in the activation energy calculated from both methods is due to the low sensitivity and resolution of the permeation technique to determine the different traps present in the material. Though the values obtained from both methods, do not coincide, but the trend that they follow is the same, i.e., among DP and CP steels, CP steel has traps with higher activation energy.

Furthermore, in DP steels, DP1000, on with higher strength, has higher activation energy for traps compared to DP800.

The rate of H desorption from DP steel was significantly higher than CP steels at atmospheric conditions. For DP1000, the time constant of hydrogen effusion for traps with activation energy of 12.95 and 18.4 kJ.mol -1 was 2.74 and 4.78 hours, respectively, whereas, for CP1000, the time constant for traps with activation energy of 15.9 and 17.52 kJ.mol -1 was 3.29 and 3.24 hours respectively.

Role of microstructure on HE

In DP steels, the microstructure is primarily ferrite and martensite, whereas, for CP1000, the microstructure consists of martensite, bainite, and ferrite. If we consider the average grain size of each phase, CP1000 has the smallest average grain size, followed by DP1000 and DP800. Since grain boundaries are considered to be the sites of hydrogen trapping, one would expect the diffusivity of H in CP1000 to be the least, which was in contradiction to the observation made by the permeation studies. The dual phase steels have a banded orientation of microstructure perpendicular to the direction of the movement of hydrogen. In the case of CP steel, the distribution of microstructure along the cross-section is more homogeneously distributed. Thus, due to the orientation of microstructure along the © 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon cross-section, the mean path for hydrogen to completely diffuse through the specimen is much higher than complex phase steel, greater the diffusion path, and higher the tortuosity of hydrogen, lower is the diffusion coefficient of hydrogen. As a result, though DP steels have a higher fraction of ferrite, the diffusion coefficient of dual phase steels are lower as compared to complex phase steel. DP1000 had higher martensitic content and grain boundary area as compared to DP800, increasing the amount of potential trap sites. However, from the TDS measurements, it was evident that the above factors only influenced the hydrogen content in the reversible trap sites, and the hydrogen content in strong trap sites were relatively unaffected. In CP steels, the amount of hydrogen trapped in strong traps are much higher than DP steels. It is due to the cementite particles (in bainite) present in the CP steels. In the case of TWIP steel, the overall hydrogen content was higher than DP and CP steels. It is due to the longer charging time of TWIP steel as compared to the other investigated steels.

Additionally, the higher hydrogen content is attributed to the austenitic microstructure, which is known to have higher solubility of hydrogen as compared to DP and CP steels. Thus, it could be concluded that the distribution of hydrogen in weak and strongtraps were influenced more by the microstructure of the material as compared to the strength of the material. Also, in the case of TWIP steel, the energy required for bulk diffusion was comparable to the trap activation energy, which influenced the TDS measurements.

From the Ag decoration test, it was observed that the ferrite-martensite interfaces are the preferred hydrogen reduction sites for both DP and CP steels. The spherical Ag particles were preferentially deposited along the ferrite-martensite interface. Additionally, in CP1000, few Ag particles were observed along the cementite particles present in the bainite.

Ferrite-martensite interface is known to be susceptible to HE due to the inherent difference in the mechanical properties of the two phases. In the case of TWIP steel, both the grain interior and the grain boundary act as the reduction site for hydrogen. It was also observed that the resistance offered to hydrogen motion is least for ferrite, followed by ferritemartensite interface and cementite and most by martensite.

Influence of Hydrogen on the Mechanical Properties

The study was carried out to understand the changes in the mechanical properties of the investigated steels under the influence of hydrogen. In the presence of hydrogen, a distinct loss of ductility was observed; however, no significant changes in the yield strength and tensile strength were observed. The ductility loss was calculated to be 10.46 % for DP800, 10.71 % for DP1000, 2.05 % for CP1000, and 13 % for TWIP steel. The low ductility loss in CP steel was reasoned to be due to the stress-free homogeneous microstructure associated with bainite. The above observation is in good agreement with the available literature according to which the bainite phase has a strong resistance towards HE. The ductility loss in TWIP steel is possibly due to the reduction of grain boundary cohesion in the presence of hydrogen and/or hydrogen promoted deformation twinning.

The lower ductility loss compared to the literature could be attributed to the absence of hydrogen charging during the tensile test and the time interval between the charging and testing, which could lead to significant effusion of hydrogen. In DP steels, the HE occurred only near the surface of the specimen as the hydrogen diffused towards the surface from the center, depleting the hydrogen from the center, resulting in ductile fracture at the center and brittle fracture near the surface. In hydrogen charged DP and CP steels, cracks perpendicular to the tensile axis were observed, which is considered to be due to the presence of hydrogen. Such cracks were only observed in the vicinity of the necking region.

In addition, in CP steels, some cracks parallel to the tensile axis were observed. However, these cracks are believed to be the consequence of tensile specimen failure rather than the cause. In TWIP steel, the fracture surface contained multiple fish eye features with a central inclusion (AlN) surrounded by brittle faceted features. Both the hydrogen charged and uncharged specimens exhibited ductile failure. In hydrogen charged TWIP steel, the failure near the surface showed brittle behavior

Influence of tensile stress on hydrogen desorption

In hydrogen charged specimens, distinct hydrogen desorption peaks could be observed during the tensile test, which were absent in the uncharged specimens. In all the specimens, a steady increase in the hydrogen desorption with an increase in the stress value was observed in the elastic region, which reached a maximum at the point corresponding to the yield point. This was attributed to the lattice expansion due to the application of tensile stress and the increased hydrogen transport due to the moving edge dislocations.

In the case of DP and CP steels, the hydrogen desorption progressively decreased with an increase in the plastic strain due to the generation of defects in the plastic region, which behave as hydrogen traps for the diffusible hydrogen. In addition, after yield point, the distance traveled by the dislocations is shorter as multiple slip systems are activated, and screw dislocations also start to move, which have smaller binding energy between hydrogen and consequently carry a lesser quantity of hydrogen.

In the investigated steels, a sharp peak was observed at fracture due to the sudden desorption of hydrogen, that had accumulated at the surface, on the sudden exposure of the new surface. In addition, in CP steel, a small peak was observed near the necking region due to the inherent strong traps present in the material. In the case of TWIP steel, beyond the yield point, a stronger peak was observed, possibly due to the release of hydrogen during the formation of deformation twins. Under high charging current density, the permeation current density in Zn coated specimens were observed to be significantly low as compared to the bare substrate owing to the low diffusion coefficient of hydrogen in Zn and its weak interaction and adsorption capability of H. Thus at high charging current densities, the presence of Zn layer appears to inhibit the entry of hydrogen in the specimen. At lower cathodic charging current density, the absorption of hydrogen in bulk DP steel was increased in the presence of the Zn layer. This was attributed to the role of the Zn layer as a sacrificial anode. The dissolution of the Zn layer led to an increased reduction of hydrogen on the Fe substrate, increasing the volume of hydrogen absorbed by the bulk of the Fe substrate. It was also noticed that the Zn-Fe interface could potentially trap hydrogen and restrict the hydrogen from effusing out of the substrate, which is already present in the material, increasing the susceptibility to HE significantly. Thus in the case of galvanized coating, its influence on the HE susceptibility of the substrate should be an essential consideration.

Perspectives

In accordance with the work carried out in the present thesis, some key areas were identified, with possibilities to be explored further, which in turn would actively enhance the current understanding of the subject. Some of the areas where future works could be carried out are as follows:

a. Hydrogen Permeation experiments on TWIP steels would provide several critical pieces of information about the diffusion coefficient and hydrogen trapping behavior in TWIP steels. However, as the diffusion coefficient of hydrogen in austenitic steels are quite low, modifications in the experimental setup, probably in terms of temperature or material thickness, could be done to allow measurements at a reasonable period.

b. It was speculated that a significant fraction of hydrogen was desorbed from the specimen surface at the charging side of the permeation experiment, which Les aciers AHSS sont largement utilisés pour les caisses des véhicules, en raison de leurs bonnes propriétés mécaniques et de leurs capacités à réduire l'empreinte carbone. Toutefois, leur utilisation est limitée par leur sensibilité à la fragilisation par l'hydrogène (HE). La présente étude vise à comprendre l'influence de l'hydrogène sur quatre aciers AHSS : deux biphasés (DP), un phase complexe (CP) et un à plasticité induite (TWIP). Les résultats montrent une forte susceptibilité à l'HE pour les aciers DP et TWIP par rapport aux aciers CP. Le comportement de CP est attribuée à une microstructure plus homogène, une densité de piège plus petite (mais à énergie élevé) et une concentration en H plus faible. Dans les aciers DP, la forte densité de pièges à faible énergie et la forte absorption de H augmentent la susceptibilité à l'HE. Les dislocations et les joints de grains sont les principaux sites de piégeage pour tous ces aciers, ainsi que la cémentite dans les aciers CP et les particules AlN et l'austénite pour les aciers TWIP. Sous chargement mécanique, la désorption de l'hydrogène s'accélère avec l'expansion du réseau cristallin et les mouvements des dislocations (jusqu'à la limite d'élasticité), alors qu'elle diminue en raison de la génération de défauts dans la domaine plastique. Pour l'acier CP, l'hydrogène piégé fortement désorbe à l'UTS alors que dans l'acier TWIP, la génération de défauts libère l'hydrogène. Pour l'acier DP galvanisés, la couche de Zn se comporte une couche barrière à l'hydrogène sous polarisation fortement cathodiques, tandis qu'à potentiels cathodique plus faible, elle favorise la perméation de l'hydrogène en raison de sa dissolution.

MOTS-CLÉS

© 2020 -

 2020 Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon of the project and the manuscript. Some thermal and mechanical experiments (Chapter 4 and Chapter 5) were performed at IMR, Tohoku University, under the supervision of Prof. Eiji Akiyama and Asst. Prof. Tomohiko Hojo. The study in Japan was made possible due to the grant provided by Erasmus Mundus. Their support and guidance were invaluable for the completion of the manuscript. The entire manuscript is divided into seven chapters (including the Introduction chapter). The remaining chapters and their content are in the following order: Chapter 2: Bibliography

© 2020 -Chapter 5 :Chapter 6 :

 202056 Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon traps, and the relative amount of hydrogen present in different traps. Ag decoration test, on the other hand, provides information on the location of hydrogen adsorbed at the surface of the material. Influence of Tensile Stress on Hydrogen Diffusion Behavior in AHSS Steels This study aims to analyze the desorption behavior of hydrogen when mechanical energy is applied. The hydrogen charged AHSS specimens are subjected to tensile tests with the tensile test instrument equipped with quadrupole mass spectrometer (QMS) to detect the desorbed hydrogen in situ at room temperature. Thus, the observed H desorption behavior during the tensile test aids in evaluating the sensitivity of the mechanical properties of the steels in the presence of hydrogen. The analysis of the above investigation was performed in three parts. The first part compares the tensile test results of hydrogen charged and uncharged specimens of the investigated AHSS. The second part deals with the fracture surface analysis, and the final part deals with the desorption spectra obtained during the tensile test measurements. Influence of Hot Dip Galvanized Coatings on the Hydrogen Permeation Behavior of DP Steels Hot-dip galvanized coating is commonly applied on steel surfaces to provide corrosion resistance to Body In White (BIW) components in the automobile industry. The chapter investigates the influence of the galvanized coating on the substrate in terms of H absorption. Permeation tests are carried out in Zn coated and uncoated steel substrate at different charging current density to observe the effect of the Zn layer on hydrogen pickup.

Figure 2 . 1 :

 21 Figure 2.1: Comparison of strength and ductility of different kinds of AHSS and conventional steels.

Figure 2 .

 2 Figure 2.2(d) presents the stress-strain curve of a dual phase steel. The tensile

Figure 2 . 2 :

 22 Figure 2.2: SEM images of (a) DP steel, (b) CP steel. (c) Band contrast image of TWIP steel and (d) Stress-strain curve of DP, CP, and TWIP steels.

Figure 2 .

 2 Figure 2.2(b) shows the microstructure of CP steels. Complex Phase (CP) steels are another first-generation AHSS characterized by high tensile strength of over 800 MPa.

Figure 2 .

 2 2(c) presents the band contrast image of TWIP steel obtained during EBSD analysis. TWIP

Figure 2 . 3 :

 23 Figure 2.3: (a) Effect of carbon content on the hydrogen uptake of ferrite pearlite steels [84] (b)

  Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon bainitic structure might dominate the hydrogen transport, and the fineness of the carbides distributed in the bainitic structure determined the hydrogen diffusivity and content.

Figure 2 . 4 :

 24 Figure 2.4: Damage evolution curves plotted against local plastic strain with and without hydrogen charging: (a) damage area fraction, (b) the number of damage incidents per area, and (c) average crack size [19].

Figure 2 .

 2 5 presents a schematic of the Volmer, Tafel, and Heyrovsky reactions in acidic media. © 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon

Figure 2 . 5 :

 25 Figure 2.5: Elementary Hydrogen Evolution Reaction (HER) reactions. Adapted from [111]

Figure 2 . 6 (

 26 a and b). presents a schematic of the hydrogen energy levels around a strong and weak-trap. However, though it might present a wholesome picture, it is quite challenging to obtain such precise data for all the parameters involved. Pressouyre et al.[START_REF] Pressouyre | An example of the effect of hydrogen trapping on hydrogen embrittlement[END_REF],[START_REF] Pressouyre | Trap theory of Hydrogen embrittlement[END_REF] came up with a much-simplified classification of the traps based on the activation energy of the traps. Though there is a continuum of traps present, they suggested that the traps with activation energy more than © 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon 77 kJ.mol -1 (0.79 eV) are irreversible and can trap hydrogen firmly till the traps are saturated[9]. At elevated temperature or beyond the saturation point of the traps, the hydrogen can diffuse into the lattice[START_REF] Michler | Hydrogen environment embrittlement of an ODS RAF steel -Role of irreversible hydrogen trap sites[END_REF]. Reversible traps or weak traps are traps with activation energy lower than 60 kJ.mol -1 (0.62 eV) where capture and release of hydrogen is relatively more straightforward.

Figure 2 . 6 :

 26 Figure 2.6: Energy Level of hydrogen around trapping sites for (a) a strong trapping site and (b) a

  Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon are equal. It leads to an increase in the apparent H solubility as compared to the lattice solubility.

Figure 2 . 7 :

 27 Figure 2.7: Schematic illustration of traps sites for hydrogen in (a) atomic scale and (b) microscopic

36 ©

 36 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon 2.4.4 Examples of different traps present in a material 2.4.4.1 Dislocations :

Figure 2 . 8 :Figure 2 . 9

 2829 Figure 2.8: Schematic presentation of the ideal (a) and real (b and c) permeation transients. (a) Ideal transients; C0 = constant, fully diffusion control; (b) first build-up and complete decay; (c) partial decay and build-up; C0 = constant, trapping and detrapping not revealed.[START_REF] Zakroczymski | Adaptation of the electrochemical permeation technique for studying entry, transport and trapping of hydrogen in metals[END_REF]

Figure 2 . 9 :

 29 Figure 2.9: Schematic of the permeation cell used for the present study

Figure 2 .

 2 Figure 2.10 : (a) The variation of diffusivity with temperature in ferrite (with and without traps)

Figure 2 .

 2 11 presents a schematic of the permeation curve, indicating the time lag and breakthrough time.

Figure 2 . 11 :

 211 Figure 2.11: A schematic of the permeation curve indication the time and the breakthrough time.

© 2020 -

 2020 Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon intensity with temperature. The spectrum contains peaks from desorbed hydrogen at different temperatures. The peak temperature gives an intuitive idea of the trap binding energy. An adequately designed TDS experiment can offer information about the H diffusion coefficient in the specimen. It is mainly useful when applied to metals or alloys with a low H diffusion coefficient as it provides faster results as compared to other techniques such as hydrogen permeation test[START_REF] Escobar | Evaluation of hydrogen trapping in high strength steels by thermal desorption spectroscopy[END_REF].

Figure 2 . 12 :

 212 Figure 2.12 : (a) and (c) SEM image of the Ag decorated surface of the deformed SUS304 specimen; (b) and (d) AFM image of the Ag decorated surface of the SUS304 specimen; GB shows grain boundary, TW twin and SL slip lines. The specimens were strained and charged galvanostatically (1 A/m²) in 0.1 M NaOH solution for two hours, followed by an Ag decoration test.

51 ©

 51 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyonnot only depend on the presence of grain boundary but also the degree of misfit between the grain boundary, which directs the amount of local enrichment of hydrogen.

  Subsequently, mechanisms proposed to describe HE in steels were discussed, among which the three most cited mechanism are, (a) HEDE which proposes a decrease in the cohesive bond strength between the metal atoms in presence hydrogen (b) HELP which proposes enhanced dislocation mobility under the influence of hydrogen and (c) AIDE which considers localized plasticity near the vicinity of the metal surface at the regions of stress concentrations.Finally, techniques popularly used to determine the diffusible and trapped hydrogen in the material were discussed. The hydrogen permeation technique, where hydrogen is reduced to its atomic form at one surface of the material and oxidized at the other, provides information about the diffusion coefficient, the hydrogen concentration, and the trap density of the material. The TDS technique measures the total concentration of hydrogen in reversible or irreversible trap sites and their respective activation energy. The Ag decoration test provides information about the distribution of reduced hydrogen at the metal surface.

Table 3 . 1 :Figure 3 . 1 .

 3131 Figure 3.1. Each compartment was equipped with an auxiliary electrode (AE: graphite at the detection side and a Pt mesh at the charging side) and a saturated Ag/AgCl reference electrode (RE). Two independent Gamry Interface 1000 potentiostats, one to

Figure 3 . 2 :

 32 Figure 3.2: (a) The complete data of three-step permeation experiment involving pre-charging, successive partial transient and partial transient loop at different charging current densities for CP1000. Expanded plots of (b) the pre-charging step; (c) the successive partial transient; (d) partial transient loop for CP1000 Each point in the figure (d) corresponds to the applied current density at the charging surface during the test (P1: -10 mA.cm -²; P2: -20 mA.cm -²; P3: -30 mA.cm -²; P4: -40 mA.cm -²; P5: -50 mA.cm -²; P6: -40 mA.cm -²; P7: -30 mA.cm -²; P8: -20 mA.cm -²; P9: -10 mA.cm -²)

5 Figure 3 . 3 :

 533 Figure 3.3: The variation of steady-state permeation current density with the square root of charging current density. A linear relationship indicates the permeation current density is controlled by the impressed cathodic current density.

Figure 3 . 4 (

 34 (a) and (b)). Both curves overlap each other validating their reproducibility and increasing the reliability of the data obtained from the subsequent steps. Step I and Step II are the surface conditioning steps, whereas step III would be used to provide the permeation data. The subsequent partial transients, Figure 3.4(a) and (b) were in a good argument with each other. The attributed the good reproducibility to (a) pre-charging

Figure 3 . 4 :

 34 Figure 3.4: The normalized curve for two successive (a) rise transients for CP1000 steel when the charging current density was raised from -10 mA/cm² to -20 mA/cm² and (b) fall transients for CP1000 steel when the charging current density was lowered from -20 mA/cm² to -10 mA/cm².

Figure 3 .

 3 Figure 3.5(e) presents the surface microstructure of the CP1000 specimen,

Figure 3 .

 3 Figure 3.5 (b, d, and f) presents the cross-sectional images along the rolling direction for all the steel sheets under study. Notably, both the DP800 and DP1000 have banded structure aligned along the rolling direction, whereas CP1000 is free from it.

Figure 3 . 5 :

 35 Figure 3.5: SEM images of the surfaces of (a) DP800 steel (c) DP1000 steel and (e) CP1000 steel and cross-sections of (b) DP800 (d) DP1000 and (f) CP1000 steel sheets. The orange, blue and green arrows represent ferrite, martensite, and bainite respectively

Figure 3 . 6 :

 36 Figure 3.6: Partial Permeation Transients for DP800 (a) Rise Transient from -10 to -20m.cm -2 and Fall transient from -20 to -10mA.cm -2 (b) Rise Transient from -20 to -30mA.cm -2 and Fall transient from -30 to -20mA.cm -2 (c) Rise Transient from -30 to -40mA.cm -2 and Fall transient from -40 to -30mA.cm -2 (d) Rise Transient from -40 to -50mA.cm -2 and Fall transient from -50 to -40mA.cm -2 .

Figure 3 .

 3 Figure 3.7 presents the variation of Dapp with the charging current density for all three steel specimens. The increase in Dapp with increasing charging current density or

Figure 3 . 7 :

 37 Figure 3.7: The evolution of the apparent diffusion coefficient of hydrogen in different steels investigated in this study with the change in cathodic current density during the rise transient. The values indicate the mean diffusion coefficient value obtained for each transient.

Figure 3 . 8 :

 38 Figure 3.8: The evolution of (a) The total hydrogen concentration at the sub-surface of the charging side (b) The lattice hydrogen concentration at the sub-surface of the charging side and (c) The trapped hydrogen concentration at the sub-surface of the charging side with change in the cathodic current density.

Figure 3 .

 3 Figure 3.9 shows the schematic of the fall transients observed for DP1000 steel at different charging current densities and the corresponding calculated current, due to oxidation of lattice H, using DL values. Assuming only one kind of trap site present in the material, and each trap site can hold only one H atom, the reversible trap density is calculated using Equation 5 [37].

Fall 2 Figure 3 . 9 : 84 ©

 23984 Figure 3.9: The schematic normalized current density for the decay transient for each fall transient and the theoretical curve calculated using DL = 1.58 x 10 -10 m².s -1 for DP1000 steel.

Figure 3 .

 3 10, wherein DP steels (Figure 3.10a), due to 86 © 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon the banded structure the actual H diffusion path is significantly longer than the calculated hydrogen diffusion path as compared to CP steels (Figure 3.10b).

Figure 3 . 10 :

 310 Figure 3.10: Schematic of the influence of microstructural cross-section on the actual hydrogen diffusion path of (a) DP steels and (b) CP steels.

Figure 3 . 11 :

 311 Figure 3.11: CTrapped/CT values as calculated from the partial permeation rise transient.

Figure 3 .

 3 Figure3.8 suggests that CL is higher for DP1000 at low charging current density, but at higher charging current density, CL is higher for DP800 (Figure3.8b). Both DP steels have ferrite and martensite, with higher martensite fraction in DP1000. Martensite, due to its high dislocation density, acts as a hydrogen trap, and resultantly DP1000 has higher trap density than DP800. Thus the traps in DP800 are quickly saturated compared to traps in DP1000, leading to enrichment of hydrogen in DP800 lattice. This observation is particularly interesting in terms of embrittlement phenomenon as Koyama et al.[69] observed that the crack initiation in DP steels occurred through martensite cracking and propagated through the ferrite/martensite interface. The martensite or the martensite/ferrite interface thus can act as strong or weak H trapping sites depending on the activation energy.

  using the electrochemical hydrogen permeation technique to understand the microstructural influence. The influence of charging surface modification during permeation experiments was successfully mitigated by pre-charging the specimens before the permeation experiment. The salient findings of the work are summarized below: 90 © 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon

99 ©

 99 The present study aims to understand trap interaction with H in four commercial grades AHSS. TDS measurements determined hydrogen concentration present in the material, their distribution in different types of traps, and their activation energy. Hydrogen effusion rates of different steels at ambient temperature was also determined. The Ag decoration test complimented TDS tests, providing a visual indication of the hydrogen adsorption or release sites on the specimen surface. 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon 4.2 Experimental Procedure Materials For the present study, TATA Steel Netherlands provided two DP steels having a tensile strength of about 800 MPa (DP1000) and 1000 MPa (DP1000), a CP steel having a tensile strength of about 1000 MPa (CP1000) and a Twinning Induced Plasticity steel having a strength of about 1000 MPa (TWIP) in the form of 30 cm X 30 cm sheet of different thicknesses. The DP steels had an average thickness of 0.98 mm while CP and TWIP steel had 1.60 mm average thickness. The chemical compositions of materials are as shown inTable 4.1.

Figure 4 . 1 :

 41 Figure 4.1: (a) The TDS instrument used in the study and (b) the schematic diagram of the TDS instrument. Adapted from[35]

Figure 4

 4 Figure 4.2).

Figure 4 . 2 :

 42 Figure 4.2: The TDS spectra of DP800, DP1000, CP1000 and TWIP steel after electrochemical charging in 3 wt % NaCl + 5 g.L -1 NH4SCN solution. Heating rate 800 K.h -1 .

Figure 4 . 3 :

 43 Figure 4.3: The total hydrogen concentration, concentration of hydrogen in weak-traps, and the concentration of hydrogen in strong-traps for DP800, DP1000, CP1000, and TWIP steel after electrochemical charging in 3 wt % NaCl + 5 g/L NH4SCN solution.

Figure 4 .

 4 Figure 4.4 presents the TDS curves of the test specimens at different heating rates of 100, 200, 400, and 800 K.h -1 . With increasing heating rate, Tc shifts towards higher

Figure 4 . 4 :

 44 Figure 4.4 : TDS results for four different heating rates for (a) DP800; (b) DP1000; (c) CP1000 and (d) TWIP steel. The heating rates applied are 100 K.h -1 , 200 K.h -1 , 400 K.h -1 and 800 K.h -1 .

Figure 4 .

 4 Figure 4.5 presents the experimentally obtained and deconvoluted curves for DP800, CP1000, and TWIP steel specimens at the heating rate of 800 K.h -1 . All curves were

Figure 4 .

 4 4(a)), which disappears at lower heating rates. Region 2 of the CP107© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon steels were deconvoluted using multiple sub-peaks for a satisfactory fit. The entire TDS spectra of TWIP steel is deconvoluted using four sub-peaks. According to Wei et al.[39], the number of sub-peaks required to deconvolute the TDS spectra is an arbitrary choice and subjected to the best fit. The trap's activation energy (Ea) is calculated using Equation4.4, with the Ea determined from the slope of the curve plotted with the natural logarithm of (ϕ / Tc²) against 1/T.

Figure 4 .

 4 4(d) presents the plot with the activation energy of different traps in the TWIP steel as an example and summarized results for all specimens in Table4.2.

Figure 4 . 5 :

 45 Figure 4.5: The curve fitting for TDS measurements carried out at 800 K.h -1 for (a) DP800; (b) CP1000 and (c) TWIP steels. (d) The activation energy for the thermal analysis peak in TWIP steel.

Figure 4 . 109 ©

 4109 6(a) shows the band contrast image of the TWIP steel and the AlN particles observed under SEM (Figure 4.6(b)) and verified by EDS analysis over a fractured TWIP specimen. Ea of 51.75 kJ.mol -1 corresponds to the hydrogen released from the austenite. Cota et al. determined the Ea of retained austenite to be about 50 kJ.mol -1 [41]. Escobar et al. observed peaks at 790 K in the TDS spectra of TRIP steel and associated it to the hydrogen released from retained austenite at a heating rate of 400 K.h -1 [11]. At the same heating rate, in the present study, a peak was observed at 771 K. 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon

Figure 4 . 6 : 110 ©

 46110 Figure 4.6: (a) Band contrast image of the TWIP steel and (b) the AlN particles present in the material as observed under SEM of a fractured surface of TWIP steel.

Figure 4 . 7 : 112 ©

 47112 Figure 4.7: TDS results for (a) DP1000 and (b) CP1000 exposed to room temperature for various time intervals before carrying out the TDS measurements. The heating rate for the measurement was 800 K/h.

Figure 4 .

 4 8(a), (e), and (h) present the backscattered electron images (BSE) under SEM for Ag decorated DP1000, CP1000, and TWIP steels at the charging surface. Energy-dispersive X-ray analysis confirmed the white particles present on the specimen surface as Ag particles. The detected Ag particles could be either H reduced Ag +

Figure 4 .

 4 8(h): the dotted line highlights a grain), the Ag particles are present at both the grain boundary and the grain interior.

Figure 4 .

 4 Figure 4.8(b), (c), and (d) show the Ag decoration results at the detection side of DP1000 specimens after 1, 2, and 4 hours of H charging. The specimens were charged in an electrochemical cell as described in the experimental section and the discharge side of the cell containing the Ag decoration solution at OCP. During the first two hours, Ag particles reduced in the ferrite grains and along the grain boundary. After 4 hours, the Ag particles are seen in the martensite as well.Figure 4.8(f) and (g) presents the surface at the detection side

Figure 4 .

 4 8(f) and (g) presents the surface at the detection side of CP1000 after 2 and 4 hours of charging, respectively. The Ag particles were initially observed in the ferrite grain and subsequently in the martensite. Additionally, after 2 hours of charging, Ag particles were also detected beside the carbide particles.

Figure 4 . 8 :

 48 Figure 4.8: The Ag decorated specimens as observed under SEM for (a) DP1000 charging side; (b) discharging side 1h; (c) discharging side 2 h; (d) discharging side 4 h; (e) CP1000 charging side; (f) CP1000 discharging side 2 h; (g) CP1000 discharge side 4 h; (h) TWIP steel charging side. F is ferrite, M is Martensite and B is Bainite

Figure 4 . 9 :

 49 Figure 4.9: The detrapping of hydrogen from the two weak-traps present in (a) DP1000 and (b) CP1000 with time.

Figure 5 . 1 :

 51 Figure 5.1: Geometry of the tensile specimen used in the tensile test study.

Figure 5 . 2 :

 52 Figure 5.2: Image of the vacuum tensile instrument used in this experiment. (a) The entire setup and (b) the enlarged image of the opened vacuum chamber with a tensile specimen placed between the jigs.

Figure 5 . 3 :

 53 Figure 5.3: The hydrogen content in different AHSS specimens after electrochemical charging.

Figure 5 . 4 :

 54 Figure 5.4: The stress-strain curve of uncharged and charged (a) DP800 (b) DP1000 (c) CP1000 and (d) TWIP specimens

Figure 5 . 5 (

 55 Figure 5.5(a and b) present the mid-thickness, and Figure 5.5(c and d) present the edge of the fracture surface of the uncharged and H charged DP800 steel. The uncharged specimen showed a cup and cone type of failure with fracture surface covered with dimples formed by the void formation and their subsequent coalesce, indicating ductile failure as the predominant failure mechanism [41]. For the charged specimen, the fracture surface at the

Figure 5 . 5 :

 55 Figure 5.5: Fracture surface of uncharged DP800 specimen at the (a) mid thickness region and (c) edge of the specimen. Fracture surface of charged DP800 specimen at the (b) mid thickness region and (d) edge of the specimen

Figure 5 . 6 :

 56 Figure 5.6: Fracture surface of the hydrogen charged CP1000 specimen (a) at the edge of the specimen and (b) a particle embedded in the fracture surface

Figure 5 .

 5 Figure 5.7 presents the transverse section of the charged and fractured DP1000 and CP1000 specimens.Figure 5.7(a) presents the shear bands around the necking region of

Figure 5 .

 5 7(a) presents the shear bands around the necking region of the DP1000 specimen. DP800 and CP1000 specimens exhibited similar features. Shear bands, formed at around 45° to the tensile axis in both charged and uncharged steel specimens, 135 © 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon indicate the direction of maximum shear stress during plastic deformation. Besides, some cracks perpendicular to the tensile axis were observed in DP and CP steels, implying HE, as indicated by arrows in Figure 5.7(a). These cracks appear in the vicinity of the necking region

Figure 5 . 7 :Figure 5 . 8 .

 5758 Figure 5.7: The SEM image shows the transverse section of charged DP and CP steels near the necking region. The surface colored in yellow (inset picture) indicates the observed surface. (a) The transverse section of the hydrogen charged DP1000 tensile specimen. Similar features were observed in DP800 and CP1000 as well. (b) Transverse section of CP1000 showing cracks parallel to the tensile axis observed only in the CP1000 specimens

Figure 5 . 8 :

 58 Figure 5.8: (a), (b), (c), (d) Micrograph of the transverse section of the hydrogen charged DP1000 specimen after the tensile test. (e), (f) ) Micrograph of the transverse section of the hydrogen charged CP1000 specimen after the tensile test

Figure 5 .

 5 Figure 5.9(a) and Figure 5.9(b) present the fracture surface of the uncharged and charged TWIP steel, respectively. The charged specimen surface was flatter than the uncharged specimen surface. Both surfaces exhibited ductile failure at the mid-thickness region of the fracture surface (Figure 5.9(c)). The similar fracture surface at the center for both specimens was expected, as, during H charging, hydrogen could penetrate only to a depth of 300 µm. The ductile dimples observed on the surface of TWIP steel was smaller than DP and CP steel. Both specimens contained multiple fish eye features, with a central inclusion surrounded by faceted features.Figure 5.9(d) is a magnified image of one of such fish eye

Figure 5 .[ 4 ]. 138 ©

 54138 9(d) is a magnified image of one of such fish eye features. EDS analysis reveals the inclusion to be AlN, formed during the steelmaking process 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon

Figure 5 . 9 :

 59 Figure 5.9: (a) Fracture surface of uncharged TWIP steel, (b) fracture surface of charged TWIP steel, (c) fracture surface of the charged TWIP steel at the mid-center region, (d) fish eye observed in the fracture surface of TWIP steel, (e) fracture surface of uncharged TWIP steel at the edge of the surface, (f) fracture surface of uncharged TWIP steel at the edge of the surface

Figure 5 .

 5 Figure 5.10 (a and b) presents the hydrogen desorption spectra and the corresponding stress-time curve for DP800 steel in uncharged and charged condition, respectively. The experimentally obtained hydrogen desorption curves were in terms of the vacuum chamber pressure. The desorption spectra were extracted by subtracting the baseline using the 'Peak Analyzer' function in OriginPro 2016 and presented in the form of hydrogen pressure. The Uncharged DP800 specimen exhibited no desorption peak in the spectra, similar to other the uncharged steel specimens investigated. In the case of charged DP800, DP1000, and CP1000 steel specimens (Figure 5.10(b,c, and d)), peaks observed in the desorption spectra indicate H desorption during tensile stress.

Figure 5 . 10 :

 510 Figure 5.10: Correlation between the hydrogen desorption spectra and stress-time curve after background correction during tensile testing for (a) Uncharged DP800 (b) Charged DP800 (c) Charged DP1000and (d) Charged CP1000 specimens.

  Factors (a) to (d) promote H 141 © 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyondesorption, whereas factors (e) and (f) retard hydrogen desorption. Among the different factors mentioned, the enhanced hydrogen desorption in the elastic region is predominantly caused by the edge dislocations, which can move under lower stresses[16].

Figure 5 . 11 :

 511 Figure 5.11: Correlation between the hydrogen desorption spectra and stress-time curve after background correction during tensile testing for (a) TWIP steel charged for seven days and (b) TWIP steel charged for 14 days

Chapter 3 .

 3 The Pd layer minimizes surface passivation, and minimize background current, during the electrochemical test.151 © 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA LyonElectrochemical TestFor the hydrogen permeation test, a modified Devanathan and Starchuski cell (hydrogen permeation cell) was used, with the working mechanism described in Chapter 3.

6. 3 . 1

 31 Characterization of the Zn layer

Figure 6 .

 6 Figure 6.1 presents the scanning electron microscopy (SEM) images of the Zn layer. The Zn layer (Figure 6.1(a)) completely covers the underlying steel substrate. The Zn layer cross-section appears to be compact except the occasional presence of a small number of pores (indicated by arrows in Figure 6.1(b)). The presence of pores or defects, is possibly due to the rapid solidification of the Zn coating, as previously reported in the published literature

Figure 6 .

 6 Figure 6.1(c) presents the Energy Dispersive X-Ray Spectroscopy (EDS) of the Zn layer cross-section, with a pure Zn layer over Fe substrate. Zn-Fe intermetallics, usually formed during galvanization, was absent, attributed to the Al added in the molten Zn bath[19],[20]. The Al present at the Zn-Fe interface forms the Al-Fe intermetallic layer and prevents Fe diffusion into the Zn layer[36].

Figure 6 . 1 :

 61 Figure 6.1: The SEM image of (a) the Zn layer, (b) the cross-section of the Zn layer, and (c) the EDS analysis of the cross-section of the Zn coating. Arrows indicate the pores present in the cross-section (b).

Figure 6 .

 6 Figure 6.2 presents the polarization curve of the DP1000 and DP1000 Zn specimens in deaerated 0.1 M K2SO4, 0.02M H2SO4 solution. The corrosion potential of DP1000 and DP1000 Zn was -0.55 mV and -0.97 mV w.r.t. saturated Ag/AgCl electrode whereas, the corrosion current density was 0.47 mA.cm -2 and 1 mA.cm -2 respectively. The higher corrosion potential and corrosion current density

Figure 6 . 2 :

 62 Figure 6.2: Potentiodynamic polarization curves of DP1000 and DP1000 Zn in deaerated 0.1 M K2SO4, 0.02M H2SO4 solution.

155 ©

 155 3(a) and (b) present the permeation curves of the investigated specimens at high and low charging current densities, 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon

Figure 6 .

 6 Figure 6.3(b) presents the permeation curve of DP1000 and DP1000 Zn under low charging current density. The permeation current density of DP1000 Zn was higher than DP1000 (2.21 ±0.41 µA.cm-2). The permeation current density of DP1000 Zn was initially high and subsequently decreased to the value similar to that of the steady-state current density of DP1000. If the permeation current density behavior of the DP1000 Zn is compared between high and low charging current densities, the Zn layer appears to inhibit H entry at high charging current density and facilitates the H entry at a low charging current density.

Figure 6 . 3 : 158 ©

 63158 Figure 6.3: The permeation current density for DP1000 and DP1000 Zn specimens for charging current density at (a) -10 mA.cm -2 and (b) -1.5 mA.cm -2 .

Figure 6 . 4 :

 64 Figure 6.4: The SEM image of the DP1000 Zn specimen after the permeation test (a) when charged with high cathodic current density (b) Enlarged image of the same specimen, (c) when charged with low cathodic current density (d) XRD of DP1000 Zn specimen after the permeation test with cathodic charging current density of -1.5 mA/cm² compared with DP1000 specimen.

Figure 6 .

 6 Figure 6.5 presents the potential-time curve of the investigated specimens during the galvanostatic charging. In Figure 6.5 (a), the DP1000 Zn specimen show potential fluctuations when charged under high charging current density. These fluctuations result from the hydrogen bubble forming and leaving the cathodic surface during cathodic charging.

Figure 6 . 5 :

 65 Figure 6.5: Potential time curve during galvanostatic charging of the charging surface for DP1000 and DP1000 Zn for (a) charging current of -10 mA/cm² and (b) charging current of -1.5 mA/cm²

Figure 6 . 6 : 161 ©

 66161 Figure 6.6: (a) The combined permeation and potential evolution curve for DP1000 Zn charged at low charging current density. The permeation tests were interrupted at the points I, II and III to observe the polarized surface under SEM; (b) the SEM image of the surface at the point I; (c) the SEM image of the surface at point II and (d) the SEM image of the surface at point III.

Figure 6 .

 6 6(c). shows that the Zn layer at point II is partially damaged. EDS measurements at point C detect 83 at% Zn and 4.8 at.% Fe whereas, at point D, 20.6 at.% Zn, 61 at.% Fe, and 4.1% Al. The Zn layer, which initially dissolved along the grain boundaries, exposed the bare substrate resulting in the further Zn layer dissolution. At point III, the Zn layer has wholly dissolved, exposing the bare substrate (Figure 6.6(d)).

Figure 6 . 7 :

 67 Figure 6.7: The variation in the intensity of H with depth for DP1000 and DP1000 Zn as measured by GDOES.

Figure 6 .[ 47 ] 163 ©

 647163 Figure 6.8 presents the macroscopic surface image of the DP1000 Zn specimen, before and after H charging. The Zn surface damage after H charging was visible with the bare eye. After H charging, the literature reports the presence of Zn hydrides on the surface[47] The slow H diffusion coefficient in Zn and the formation of the hydride in the Zn layer probably generates intense tensile stress and causes localized rupture and blistering. The Zn layer was partly damaged as compared to the permeation test, due to the milder charging solution. The remaining Zn layer behaved as a barrier layer to desorbing hydrogen and prevented the desorption of hydrogen from the specimen[23].

Figure 6 . 8 :

 68 Figure 6.8: Macroscopic image of the surface of the Zn layer (a) before and (b) after electrochemical charging in 3 wt% NaCl + 5 g NH4SCN solution.

Figure 6 .

 6 Figure 6.9(a) presents the schematic of the specimen at the point I, where the potential at the charging surface is -1.05 VAg/AgCl, similar to the potential observed in the polarization curve for the corresponding polarization current of -1.5 mA.cm -2 . Due to the charging condition's proximity to OCP, the Zn dissolves along the grain boundaries because of the residual anodic reactions. Boiadjieva et al.[21] have also reported crack formation in the grain boundaries due to the H impact leading to the localized penetration of the electrolyte within the Zn layer to reach the Fe surface. As a result, in certain localized regions, the steel substrate is exposed, resulting in H reduction at the steel surface. According to the mixed potential theory:

165 ©

 165 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon

Figure 6 . 9 :

 69 Figure 6.9: Mechanism of the dissolution of the Zn layer when electrochemically charged with hydrogen at low charging current density.

7. 1 . 8

 18 Influence of Zn coating on the HE susceptibility of the substratePermeation tests were carried out in order to understand the influence of the galvanized coating on the HE tendency of DP steels. The diffusion coefficient of hydrogen in Zn was calculated to be 4.59 X 10 -15 m².s -1 , which was four orders of magnitude lower than the steel substrate. It was observed that the influence of the galvanized layer on HE of the substrate had a strong dependence on the charging current density. The galvanized layer inhibited the entry of hydrogen when charged under high charging current density, whereas © 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon it appeared that under low charging current density, the Zn layer increased the susceptibility of HE for the underlying steel substrate.

  Figure 1.1: Graph indicating the growth of AHSS in the automobile industry [1]. .................................................. 3

1

 1 Figure 1.1: Graph indicating the growth of AHSS in the automobile industry [1]. .................................................. 3

Figure 2 . 7 :

 27 Figure 2.7: Schematic illustration of traps sites for hydrogen in (a) atomic scale and (b) microscopic scale [126] .............................................................................................................................................................................................. 35

Figure 2 . 8 : 38 Figure 2 . 9 : 39 Figure 2 . 10 : 41 Figure 2 . 11 : 42 Figure 2 . 12 : 50 Figure 3 . 1 : 67 Figure 3 . 2 : 69 Figure 3 . 3 :

 283829392104121142212503167326933 Figure 2.8: Schematic presentation of the ideal (a) and real (b and c) permeation transients. (a) Ideal transients; C0 = constant, fully diffusion control; (b) first build-up and complete decay; (c) partial decay and build-up; C0 = constant, trapping and detrapping not revealed. [115]................................................................................................... 38 Figure 2.9: Schematic of the permeation cell used for the present study...................................................................... 39Figure 2.10 : (a) The variation of diffusivity with temperature in ferrite (with and without traps) and austenite [7] and (b) the Fe-H phase diagram showing the solubility of hydrogen in different phases [152]. ......................... 41 Figure 2.11: A schematic of the permeation curve indication the time and the breakthrough time. ........................ 42 Figure 2.12 : (a) and (c) SEM image of the Ag decorated surface of the deformed SUS304 specimen; (b) and (d) AFM image of the Ag decorated surface of the SUS304 specimen; GB shows grain boundary, TW twin and SL slip lines. The specimens were strained and charged galvanostatically (1 A/m²) in 0.1 M NaOH solution for two hours, followed by an Ag decoration test. ......................................................................................................................... 50 Figure 3.1: Schematic of the permeation cell used for this study. H + is reduced as H at the left compartment and oxidized at the right compartment .................................................................................................................................... 67 Figure 3.2: (a) The complete data of three-step permeation experiment involving pre-charging, successive partial transient and partial transient loop at different charging current densities for CP1000. Expanded plots of (b) the pre-charging step; (c) the successive partial transient; (d) partial transient loop for CP1000 Each point in the figure (d) corresponds to the applied current density at the charging surface during the test (P1: -10 mA.cm -²; P2: -20 mA.cm -²; P3: -30 mA.cm -²; P4: -40 mA.cm -²; P5: -50 mA.cm -²; P6: -40 mA.cm -²; P7: -30 mA.cm -²; P8: -20 mA.cm - ²; P9: -10 mA.cm -²) ............................................................................................................................................................ 69Figure 3.3: The variation of steady-state permeation current density with the square root of charging current density. A linear relationship indicates the permeation current density is controlled by the impressed cathodic current density. .................................................................................................................................................................................. 71

Figure 3 . 4 : 72 Figure 3 . 5 : 74 Figure 3 . 6 : 77 Figure 3 . 7 : 78 Figure 3 . 8 : 82 Figure 3 . 9 : 83 Figure 3 . 10 : 86 Figure 3 . 11 :

 34723574367737783882398331086311 Figure 3.4: The normalized curve for two successive (a) rise transients for CP1000 steel when the charging current density was raised from -10 mA/cm² to -20 mA/cm² and (b) fall transients for CP1000 steel when the charging current density was lowered from -20 mA/cm² to -10 mA/cm². .................................................................................. 72 Figure 3.5: SEM images of the surfaces of (a) DP800 steel (c) DP1000 steel and (e) CP1000 steel and crosssections of (b) DP800 (d) DP1000 and (f) CP1000 steel sheets. The orange, blue and green arrows represent ferrite, martensite, and bainite respectively ................................................................................................................................... 74 Figure 3.6: Partial Permeation Transients for DP800 (a) Rise Transient from -10 to -20m.cm -2 and Fall transient from -20 to -10mA.cm -2 (b) Rise Transient from -20 to -30mA.cm -2 and Fall transient from -30 to -20mA.cm -2 (c) Rise Transient from -30 to -40mA.cm -2 and Fall transient from -40 to -30mA.cm -2 (d) Rise Transient from -40 to -50mA.cm -2 and Fall transient from -50 to -40mA.cm -2 . .......................................................................... 77 Figure 3.7: The evolution of the apparent diffusion coefficient of hydrogen in different steels investigated in this study with the change in cathodic current density during the rise transient. The values indicate the mean diffusion coefficient value obtained for each transient. .................................................................................................................... 78 Figure 3.8: The evolution of (a) The total hydrogen concentration at the sub-surface of the charging side (b) The lattice hydrogen concentration at the sub-surface of the charging side and (c) The trapped hydrogen concentration at the sub-surface of the charging side with change in the cathodic current density. ........................................................ 82 Figure 3.9: The schematic normalized current density for the decay transient for each fall transient and the theoretical curve calculated using DL = 1.58 x 10 -10 m².s -1 for DP1000 steel. ............................................................. 83 Figure 3.10: Schematic of the influence of microstructural cross-section on the actual hydrogen diffusion path of (a) DP steels and (b) CP steels. ......................................................................................................................................... 86 Figure 3.11: CTrapped/CT values as calculated from the partial permeation rise transient. ............................ 88

Figure 4 . 1 :

 41 Figure 4.1: (a) The TDS instrument used in the study and (b) the schematic diagram of the TDS instrument. Adapted from [35] .......................................................................................................................................................... 102

Figure 4 . 2 :

 42 Figure 4.2: The TDS spectra of DP800, DP1000, CP1000 and TWIP steel after electrochemical charging in 3 wt % NaCl + 5 g.L -1 NH4SCN solution. Heating rate 800 K.h -1 . ............................................................................ 104

Figure 4 . 3 :

 43 Figure 4.3: The total hydrogen concentration, concentration of hydrogen in weak-traps, and the concentration of hydrogen in strong-traps for DP800, DP1000, CP1000, and TWIP steel after electrochemical charging in 3 wt % NaCl + 5 g/L NH4SCN solution. ................................................................................................................................ 105

Figure 4 . 4 :

 44 Figure 4.4 : TDS results for four different heating rates for (a) DP800; (b) DP1000; (c) CP1000 and (d) TWIP steel. The heating rates applied are 100 K.h -1 , 200 K.h -1 , 400 K.h -1 and 800 K.h -1 . .................................................. 106

Figure 4 . 5 :

 45 Figure 4.5: The curve fitting for TDS measurements carried out at 800 K.h -1 for (a) DP800; (b) CP1000 and (c) TWIP steels. (d) The activation energy for the thermal analysis peak in TWIP steel. ........................................... 108

Figure 4 . 6 :

 46 Figure 4.6: (a) Band contrast image of the TWIP steel and (b) the AlN particles present in the material as observed under SEM of a fractured surface of TWIP steel. ....................................................................................................... 109

Figure 4 . 7 :

 47 Figure 4.7: TDS results for (a) DP1000 and (b) CP1000 exposed to room temperature for various time intervals before carrying out the TDS measurements. The heating rate for the measurement was 800 K/h. ....................... 111

Figure 4 . 8 :

 48 Figure 4.8: The Ag decorated specimens as observed under SEM for (a) DP1000 charging side; (b) discharging side 1 h; (c) discharging side 2 h; (d) discharging side 4 h; (e) CP1000 charging side; (f) CP1000 discharging side 2 h; (g) CP1000 discharge side 4 h; (h) TWIP steel charging side. F is ferrite, M is Martensite and B is Bainite ........................................................................................................................................................................................... 114

Figure 4 . 9 :

 49 Figure 4.9: The detrapping of hydrogen from the two weak-traps present in (a) DP1000 and (b) CP1000 with time. ................................................................................................................................................................................... 117

Figure 5 . 1 : 127 Figure 5 . 2 : 128 Figure 5 . 3 :

 511275212853 Figure 5.1: Geometry of the tensile specimen used in the tensile test study. ............................................................. 127Figure 5.2: Image of the vacuum tensile instrument used in this experiment. (a) The entire setup and (b) the enlarged image of the opened vacuum chamber with a tensile specimen placed between the jigs. ............................................ 128Figure 5.3: The hydrogen content in different AHSS specimens after electrochemical charging. ........................ 129

Figure 5 . 4 :

 54 Figure 5.4: The stress-strain curve of uncharged and charged (a) DP800 (b) DP1000 (c) CP1000 and (d) TWIP specimens ........................................................................................................................................................................... 130

Figure 5 . 6 : 134 Figure 5 . 7 : 136 Figure 5 . 8 :

 561345713658 Figure 5.6: Fracture surface of the hydrogen charged CP1000 specimen (a) at the edge of the specimen and (b) a particle embedded in the fracture surface ....................................................................................................................... 134Figure 5.7: The SEM image shows the transverse section of charged DP and CP steels near the necking region. The surface colored in yellow (inset picture) indicates the observed surface. (a) The transverse section of the hydrogen charged DP1000 tensile specimen. Similar features were observed in DP800 and CP1000 as well. (b) Transverse section of CP1000 showing cracks parallel to the tensile axis observed only in the CP1000 specimens ................. 136 Figure 5.8: (a), (b), (c), (d) Micrograph of the transverse section of the hydrogen charged DP1000 specimen after the tensile test. (e), (f) ) Micrograph of the transverse section of the hydrogen charged CP1000 specimen after the tensile test .......................................................................................................................................................................... 137

Figure 5 . 11 : 142 Figure 6 . 1 : 153 Figure 6 . 2 : 154 Figure 6 . 3 : 156 Figure 6 . 4 : 158 Figure 6 . 5 : 159 Figure 6 . 6 : 160 Figure 6 . 7 : 162 Figure 6 . 8 : 163 Figure 6 . 9 :

 511142611536215463156641586515966160671626816369 Figure 5.11: Correlation between the hydrogen desorption spectra and stress-time curve after background correction during tensile testing for (a) TWIP steel charged for seven days and (b) TWIP steel charged for 14 days ......... 142Figure 6.1: The SEM image of (a) the Zn layer, (b) the cross-section of the Zn layer, and (c) the EDS analysis of the cross-section of the Zn coating. Arrows indicate the pores present in the cross-section (b). ............................... 153 Figure 6.2: Potentiodynamic polarization curves of DP1000 and DP1000 Zn in deaerated 0.1 M K2SO4, 0.02M H2SO4 solution. ................................................................................................................................................................ 154Figure 6.3: The permeation current density for DP1000 and DP1000 Zn specimens for charging current density at (a) -10 mA.cm -2 and (b) -1.5 mA.cm -2 . ...................................................................................................................... 156 Figure 6.4: The SEM image of the DP1000 Zn specimen after the permeation test (a) when charged with high cathodic current density (b) Enlarged image of the same specimen, (c) when charged with low cathodic current density (d) XRD of DP1000 Zn specimen after the permeation test with cathodic charging current density of -1.5 mA/cm² compared with DP1000 specimen. ................................................................................................................................. 158 Figure 6.5: Potential time curve during galvanostatic charging of the charging surface for DP1000 and DP1000 Zn for (a) charging current of -10 mA/cm² and (b) charging current of -1.5 mA/cm² ........................................ 159 Figure 6.6: (a) The combined permeation and potential evolution curve for DP1000 Zn charged at low charging current density. The permeation tests were interrupted at the points I, II and III to observe the polarized surface under SEM; (b) the SEM image of the surface at the point I; (c) the SEM image of the surface at point II and (d) the SEM image of the surface at point III. ................................................................................................................... 160 Figure 6.7: The variation in the intensity of H with depth for DP1000 and DP1000 Zn as measured by GDOES. ........................................................................................................................................................................................... 162Figure 6.8: Macroscopic image of the surface of the Zn layer (a) before and (b) after electrochemical charging in 3 wt% NaCl + 5 g NH4SCN solution. ............................................................................................................................ 163Figure 6.9: Mechanism of the dissolution of the Zn layer when electrochemically charged with hydrogen at low charging current density. ................................................................................................................................................. 165

  FOLIO ADMINISTRATIF THESE DE L'UNIVERSITE DE LYON OPEREE AU SEIN DE L'INSA LYON NOM : MALLICK DATE de SOUTENANCE : le 13 Juillet 2020 Prénoms : Dwaipayan TITRE : Hydrogen Behavior in First and Second Generation of Advanced High Strength Steels NATURE : Doctorat Numéro d'ordre : 2020LYSEI052 Ecole doctorale : ED 34 Matériaux de Lyon Spécialité : Matériaux RESUME :

  : AHSS, DP Steel, CP steel, TWIP steel, HE, Diffusion, H Trapping, H Permeation, TDS, Ag Decoration Test, SSRT Laboratoire (s) de recherche : Matériaux, Ingénierie et Science (MATEIS) Directeur de thèse: Bernard NORMAND Président de jury : Composition du jury :

Table 3 . 2 :

 32 Summary of microstructural (phase fraction, average grain size) and mechanical (tensile strength and % Elongation) data of the steels.

Table 3 . 3 :

 33 The trap density, total amount of hydrogen released, and the activation energy of the specimens obtained from the fall transient of the permeation curve.

	H Charging		Nt (10 15 site.cm -3 )		CT* (mol.m -3 )			CT*/ CT X 100 %	Ea (kJ.mol -1 )
	(mA.cm -2 )	DP800 DP1000 CP1000 DP800 DP1000 CP1000 DP800 DP1000 CP1000 DP800 DP1000 CP1000
	-50 to -40	7.76	1.59	0.86	0.073	0.032	0.017	40.00	33.44	32.66
	-40 to -30	8.62	3.72	1.48	0.084	0.044	0.029	37.77	32.18	28.52
										34.05	37.62	38.94
	-30 to -20	16.80	10.04	2.21	0.136	0.067	0.043	37.25	32.40	28.21
	-20 to -10	28.50	24.10	23.2	0.217	0.127	0.212	32.45	36.99	31.22

Table 4 . 2 :

 42 The calculated peak temperature and activation energy values for investigated steel specimens as obtained from the peak fitting data. charging, after 9 hours, the total H concentration in both specimens was comparable (0.28 wppm for DP1000 and 0.31 wppm for CP1000). At this point, though the total H concentration in both steels are similar, H distribution in strong and weak traps are different, which might lead to different HE resistance behavior. Most of the trapped H is present in weak-traps for DP1000 and strong-traps for CP1000. Under similar H concentration and stress condition, DP steel is more susceptible to HE than CP steels, since it is easier for H to enter into the lattice sites from weak-traps than from strong-traps. It shows that the total H concentration is not the ideal parameter for comparing material sensitivity to HE. It would also influence the post-treatment schedule for increasing HE resistivity. For example, CP steels are more HE resistant than DP steel, however, as DP steels have higher diffusible H concentration, baking the steel specimens for a sufficient amount of time can reduce its risk to HE.

		Tc 1 (K)	Tc 2(K)	Tc 3(K)	Tc 4(K)	Tc 5(K)	Tc 6(K)
				DP800
	800 K.h -1	424	476	957
	400 K.h -1	353	420
	200 K.h -1	335	391
	Ea (kJ.mol -1 )	10.60	17.67
				DP1000
	800 K.h -1	437	494	939
	400 K.h -1	408	443
	200 K.h -1	351	392
	Ea (kJ.mol -1 )	12.95	18.4
				CP1000
	800 K.h -1	450	500	672	779	908	1061
	400 K.h -1	404	445	621	722	781	904
	200 K.h -1	371	401	577		748	881
	Ea (kJ.mol -1 )	17.52	15.90	36.69	44.38	31.11	35.31
				TWIP
	800 K.h -1	556	670	726	831
	400 K.h -1	518	593	632	771
	200 K.h -1	471	539	600	731
	100 K.h -1	433	469		677
	Ea (kJ.mol -1 )	25.06	18.03	26.64	51.75
	4.3.3 Hydrogen desorption in DP and CP steels
	Figure 4.7 presents the TDS spectra of DP1000 and CP1000 at 800 K.h -1 for
	specimens kept at room temperature after H charging for different time intervals. Results
	show that more than half of the total trapped H diffused out during the first 3 hours (Table
	4.3.). Hydrogen desorption was faster in DP1000 than CP1000 with H concentration reducing
	from 1.65 and 0.85 wppm immediately after charging to 0.2 and 0.11 wppm respectively after
	9 hours. However, H concentration in strong trapping sites did not change significantly with

time. Results show that though the total H in DP1000 steel was higher than CP1000 steel 111 © 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon immediately after

Table 4 . 3 :

 43 The total, weakly trapped, and strongly trapped hydrogen evolution in DP1000 and CP1000 with an increase in desorption time before TDS measurements.

	Desorption	Total H (wppm)	Weakly Trapped H (wppm)	Strongly Trapped H (wppm)
	Time
			DP1000
	0 h	1.77	1.65	0.12
	3 h	0.77	0.64	0.13
	6 h	0.45	0.32	0.13
	9 h	0.28	0.20	0.08
			CP1000
	0 h	1.08	0.85	0.23
	3 h	0.52	0.28	0.23
	6 h	0.37	0.15	0.21
	9 h	0.31	0.11	0.20
	4.3.4 Ag decoration test
	Ag decoration tests revealed H distribution on the specimen surface,
	complementing TDS results.

Table 4 .

 4 is released from the trap sites. Lower the activation energy of traps faster will be the process, and resultantly faster will be the H detrapping and effusion. Thus, H has a faster effusion rate in DP steel than in CP steel. The hydrogen concentration present in strong-traps, however, does not change significantly with time.

4: Published data for the activation energy of different trapping sites in steels. Several methods have been used in different publications to evaluate the activation energy of traps such as TDS, hydrogen permeation, local equilibrium model, and numerical simulations. It provides an insight into the widely scattered and overlapping data for a particular trapping site making it difficult to attribute obtained activation energy to a particular trap site. 117 © 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon equilibrium, hydrogen

  .1. The specimen thickness

		C	Mn	Si	Al	Ni	Cr	Ti	Fe
	DP800	0.14	1.73	0.26	0.07	0.01	0.56	0.02	Bal.
	DP1000	0.14	2.14	0.05	0.04	0.02	0.58	0.03	Bal.
	CP1000	0.11	2.12	0.05	0.05	0.02	0.56	0.03	Bal.
	TWIP	0.7	14.48	-	2.3	-	-	-	Bal.

Table 5 . 2 :

 52 The mechanical properties of the different uncharged and charged steel specimens investigated in the study.

			Yield Strength (MPa)	Tensile Strength (MPa)	HEI (%)
	DP800	Uncharged	774	865
		Charged	777	857	10.46
	DP1000	Uncharged	996	1073
		Charged	992	1059	10.71
	CP1000	Uncharged	998	1023
		Charged	995	1021	2.05
	TWIP	Uncharged	630	1094
		Charged 7 days	639	1089	13.00

  steel substrate was investigated in this study using electrochemical techniques. The H diffusion coefficient in the Zn layer was four orders of magnitude lower than in Fe. At lower cathodic charging current density, the H absorption in bulk DP1000 increased in the presence of the Zn layer. It was attributed to the role of the Zn layer as a sacrificial anode. The Zn layer dissolution led to an increased H reduction on the Fe substrate, increasing the H concentration absorbed by the bulk. Once hydrogen enters into the bulk material, the Zn layer acts as a7.1 SummaryThe project investigated the H influence on the material property of some commercial first and second-generation Advanced High Strength Steels (AHSS). Prof. Bernard Normand Assoc. Prof. Nicolas Mary guided the project from INSA Lyon, and Prof.V.S. Raja from the Indian Institute of Technology, Bombay (IIT Bombay). TATA Steel Netherlands provided the specimens in the form of steel sheets. The experiments to study the thermal and mechanical properties were performed in IMR Laboratory, Tohoku University, Japan, under the supervision of Asst. Prof. Tomohiko Hojo and Prof. Eiji Akiyama.

  The activation energy of traps for TWIP steel was calculated only by thermal desorption spectroscopy due to charging time limitations. The activation energy (Ea) of different traps present in DP800 were calculated to be 10.60 and 17.67 kJ.mol -1 , 12.95 and 18.4 kJ.mol -1 for DP1000, 17.52, 15.90, 36.69, 44.38, 31.11, 35.31 kJ.mol -1 for CP1000 and 25.06, 18.03, 26.64, 51.75 kJ.mol -1 for TWIP steels. The traps with Ea below 20 kJ.mol -1 could be associated with martensitic lath interface, dislocations, grain boundaries, or even lattice diffusion. The wide scatter and the overlapping nature of the

activation energy associated with a particular kind of trap in the published literature, makes © 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon

:

  Graph indicating the growth of AHSS in the automobile industry[1]. .................................................. 3 Figure 2.1: Comparison of strength and ductility of different kinds of AHSS and conventional steels. IF is interstitial -free steels, CMn is Carbon Manganese steels, HSLA is high strength low alloy steels, DP is dual phase steels, TRIP is transformation induced plasticity steels, CP is complex phase steels, TWIP is twinning induced plasticity steels, Q&P is quenching and partitioning steels. Adapted from Venezuela et al. [9] ................................ 12Figure 2.2: SEM images of (a) DP steel, (b) CP steel. (c) Band contrast image of TWIP steel and (d) Stressstrain curve of DP, CP, and TWIP steels. ...................................................................................................................... 14 Figure 2.3: (a) Effect of carbon content on the hydrogen uptake of ferrite pearlite steels [84] (b) Hydrogen trapping of the martensitic structure as a function of carbon content [85] ................................................................................... 23Figure 2.4: Damage evolution curves plotted against local plastic strain with and without hydrogen charging: (a) damage area fraction, (b) the number of damage incidents per area, and (c) average crack size [19]. ..................... 28Figure 2.5: Elementary Hydrogen Evolution Reaction (HER) reactions. Adapted from [111] ............................ 32Figure 2.6: Energy Level of hydrogen around trapping sites for (a) a strong trapping site and (b) a weak trapping site[START_REF] Lee | crystals structures by the desorption thermal analysis technique[END_REF]. Sn is the normal lattice site, and ST is the trap site. EaD is the activation energy for lattice diffusion, Ea is the trap activation energy for the escape of hydrogen, Eb is the trap binding energy, and Es is the saddle point energy.

.............................................................................................................................................................................................. 34

  ....................................................................................................................................................................... 130Figure 5.5: Fracture surface of uncharged DP800 specimen at the (a) mid thickness region and (c) edge of the specimen. Fracture surface of charged DP800 specimen at the (b) mid thickness region and (d) edge of the specimen ........................................................................................................................................................................................... 133

  ...................................................................................................................................................................... 137 Figure 5.9: (a) Fracture surface of uncharged TWIP steel, (b) fracture surface of charged TWIP steel, (c) fracture surface of the charged TWIP steel at the mid-center region, (d) fish eye observed in the fracture surface of TWIP steel, (e) fracture surface of uncharged TWIP steel at the edge of the surface, (f) fracture surface of uncharged TWIP steel at the edge of the surface .......................................................................................................................................... 138 Figure 5.10: Correlation between the hydrogen desorption spectra and stress-time curve after background correction during tensile testing for (a) Uncharged DP800 (b) Charged DP800 (c) Charged DP1000and (d) Charged CP1000 specimens. .......................................................................................................................................................... 140

Table 3 . 1 :

 31 Chemical composition of the steels in wt% used in hydrogen permeation study ......................................... 66

Table 3 . 2 :

 32 Summary of microstructural (phase fraction, average grain size) and mechanical (tensile strength and % Elongation) data of the steels. ............................................................................................................................................... 75

Table 3 . 3 :

 33 The trap density, total amount of hydrogen released, and the activation energy of the specimens obtained from the fall transient of the permeation curve. .......................................................................................... 85

Table 4 . 1 :

 41 Chemical Composition of the steels in wt% used for TDS and Ag decoration study................................. 99

Table 4 . 2 :

 42 The calculated peak temperature and activation energy values for investigated steel specimens as obtained from the peak fitting data. ...................................................................................................................................................110

Table 4 . 3 :

 43 The total, weakly trapped, and strongly trapped hydrogen evolution in DP1000 and CP1000 with an increase in desorption time before TDS measurements. ...................................................................................................112

Table 4 . 4 :

 44 Published data for the activation energy of different trapping sites in steels. Several methods have been used in different publications to evaluate the activation energy of traps such as TDS, hydrogen permeation, local equilibrium model, and numerical simulations. It provides an insight into the widely scattered and overlapping data for a particular trapping site making it difficult to attribute obtained activation energy to a particular trap site. ..[START_REF] Moro | Hydrogen embrittlement susceptibility of a high strength steel X80[END_REF]

Table 5 . 1 :

 51 Chemical Composition of the steels in wt% used in the tensile test study ...................................................126

Table 5 . 2 :

 52 The mechanical properties of the different uncharged and charged steel specimens investigated in the study.

...............................................................................................................................................................................................131

Cette thèse est accessible à l'adresse : http://theses.insa-lyon.fr/publication/2020LYSEI052/these.pdf © [D. Mallick], [2020], INSA Lyon, tous droits réservés

© 2020 -Dwaipayan MALLICK -Lab. MATEIS, INSA Lyon

Acknowledgments

For this study, I would like to thank the ERASMUS+ scholarship for supporting the stay at Tohoku University, Japan. I would also like to express my heartfelt gratitude to Prof. Eiji Akiyama and Asst. Prof. Tomohiko Hojo, who helped me unconditionally during my visit to carry out the experiments and interpret the results.

heartfelt gratitude to Prof. Raja, for his in-depth gratitude to Akiyama Sensei and Hojo Sensei from IMR, Tohoku University, Japan, for their support and guidance during my visit to Japan. The work carried out in Japan was instrumental for this thesis, and I would like to thank the ERASMUS MUNDUS fellowship for making it happen.

Trapping site

Ea (kJ.mol -1 ) Reference Trapping site Ea (kJ.mol -1 ) Reference Dislocations (Austenite) 26-35 [16], [46] Grain Boundary (Austenite) 19-35 [16], [47] Grains Volume (Austenite) 49-58 [48] Lattice diffusion 15-25 [15], [18], [18] The slower hydrogen effusion from CP steels than DP steels (Figure 4.7) is attributed to trap strength [17], [51]. There exists an equilibrium between the lattice hydrogen and trapped hydrogen in the material. After H charging, when the specimens are left at room temperature, H diffuses out. As the hydrogen in lattice decreases, to maintain