Investigations on Hydrogen Embrittlement of Advanced High

Investigations on Hydrogen Embrittlement of Advanced High

Strength Steels for Automotive Applications

Von der Fakultät für Georessourcen und Materialtechnik

der Rheinisch-Westfälischen Technischen Hochschule Aachen

zur Erlangung des akademischen Grades eines

Doktors der Ingenieurwissenschaften

Genehmigte Dissertation

Vorgelegt von M.Sc.

Qi Gao

Aus Nei Mongol, China

Berichter: Univ.-Prof. Dr.-Ing. Wolfgang Bleck

Univ.-Prof. Dr.-Ing.habil. Brita Daniela Zander

Tag der mündlichen Prüfung: 06. März 2017

Diese Dissertation ist auf den Internetseiten der Hochschulbibliothek online verfügbar

Abstract

The utilization of Advanced High Strength Steels (AHSS) in automotive applications is motivated by a

push to reduce the weight of cars in order to minimize CO2 emissions. The potential risk of Hydrogen

Embrittlement (HE) limits the widespread application of AHSS, making it one of the hottest topics in

materials engineering and the automobile industry.

The current work focused on the hydrogen-induced cracking of 1180 MPa steels for car body

constructions. Various mechanical tests and forming tests to investigate internal hydrogen assisted

cracking as well as hydrogen environment assisted cracking were conducted, concerning the automobile

production processes and conditions during the life of a car. Tests of component-similar samples were

carried out to compare laboratory results with those from the situation close to reality. The

susceptibilities of various steels to HE were characterized and compared. The test methods were

evaluated in the terms of differentiation and simplicity. Additionally, the material- and process-

immanent factors on HE-susceptibility were elucidated and discussed.

The results obtained in this work reveal that the immersion test with U-bend samples and the

incremental loading test are appropriate to evaluate the HE-susceptibility for the material release

process. Factors such as the presence of retained austenite, the high content of carbon equivalent and

the coarse grain size increase the HE-susceptibility of AHSS. However, the addition of carbide-

immanent elements Ti and Nb, and the alloying elements related to pitting corrosion resistance like Cr,

Mo, N reduces the material HE-susceptibility. In the production processes the welding point and

punched edges are sensitive to HE. The strain state during the forming process and the corrosive

surroundings during the life of a car are also significant in the safe application of AHSS in car bodies.

On the basis of these results, by controlling the factors influencing on HE-susceptibility, roughly

divided into the three areas of material, hydrogen and stress, significant steps toward the safe usage of

AHSS in component parts have been made.

Kurzfassung

Die Verwendung von höchstfesten Stählen in Automobilen hat eine hohe Bedeutung bei der

Verringerung des Fahrzeuggewichts und damit bei der Reduzierung der CO2-Emissionen. Der Einsatz

dieser Stähle ist jedoch zum Teil noch limitiert. Ein Grund dafür ist die bei diesen Stählen potentiell

auftretende Wasserstoffversprödung, die daher in den Materialwissenschaften und in der

Automobil-Industrie eine bedeutende Rolle spielt, um diese Stähle industriell einsetzen zu können.

Der Schwerpunkt der vorliegenden Arbeit liegt auf der Untersuchung des wasserstoffinduzierten

Sprödbruchverhaltens von Stählen mit einer Zugfestigkeit von 1180 MPa, für den Einsatz in

Karosseriebauteilen. Die Stähle wurden in verschiedenen mechanischen und umformtechnischen

Versuchen, die die Verhältnisse bzw. die Beanspruchungen der Stähle bei der Produktion sowie der

Verwendung im Fahrzeug widerspiegeln, untersucht und bewertet. Nachfolgend erfolgte eine

Validierung anhand von bauteilähnlichen Versuchen aus dem Karosseriebau zum Abgleich der

Ergebnisse zwischen Laborversuchen und Realität. Gleichzeitig erfolgte auf Basis dieser Bewertungen

ein Vergleich der verschiedenen Prüfmethoden im Hinblick auf die Eignung als Standard für zukünftige

Untersuchungen zur Wasserstoffversprödung. Zusätzlich wurden die wichtigsten Einflüsse der

Materialien sowie der Verarbeitungsprozesse auf die wasserstoffinduzierte Rissbildung von

höchstfesten Stählen für Karosserieanwendung bewertet.

Die Ergebnisse zeigen, dass zwei Methoden sehr geeignet sind, um die Wasserstoffversprödung von

höchstfesten Stählen für Karosserieanwendung zu bewerten: Versuche mit U-Biege-Proben, die in

korrosive Medien getaucht werden und Stufenversuche, bei denen eine Last an zugprobenähnliche

Geometrien in korrosiver Umgebung stufenförmig gesteigert wird. Faktoren wie die Anwesenheit von

Restaustenit, ein hohes Kohlenstoffäquivalent und eine große Korngröße steigern die

Wasserstoffversprödungsneigung. Entgegengesetzt wirkt die Zugabe von Carbid-Bildnern wie Ti

und Nb sowie von Cr, Mo und N, die die Anfälligkeit gegen Lochkorrosion vermindern. Im

Herstellungsprozess haben Schweißpunkte sowie hohe Spannungen an der Kante, wie sie beim

Stanzen auftreten, einen negative Einfluss auf die Neigung zur Wasserstoffversprödung. Weitere

untersuchte und bewertet Einflüsse sind die Dehnungszustände im Umformprozessen sowie

verschiedene korrosive Einflüsse. Die Arbeit bietet damit einen wissenschaftlich fundierten Beitrag zur

Bewertung der Wasserstoffversprödung von höchstfesten Stählen für den sicheren Einsatz von

höchstfesten Stählen in Karosseriebauteilen.

VII

Contents

Abstract ................................................................................................................................................. III

Kurzfassung ........................................................................................................................................... V

Abbreviations and symbols ................................................................................................................... XI

1 Introduction and Objective.............................................................................................................. 1

2 Theoretical Fundamentals ............................................................................................................... 3

2.1 Developments and applications of HSS and AHSS ................................................................ 3

2.2 Hydrogen sources in automotive production processes .......................................................... 7

2.3 Hydrogen absorption and transport in steels ........................................................................... 9

2.3.1 Hydrogen absorption ....................................................................................................... 9

2.3.2 Hydrogen transport ....................................................................................................... 11

2.4 Hydrogen-induced cracking in steels .................................................................................... 16

2.5 Mechanism of hydrogen-induced cracking ........................................................................... 17

2.6 Important factors influencing hydrogen-induced cracking ................................................... 20

3 Materials and Experimental Methods ........................................................................................... 24

3.1 Investigated materials ........................................................................................................... 24

3.2 Material characterization....................................................................................................... 24

3.3 Hydrogen charging and analysis ........................................................................................... 27

3.3.1 Galvano-static charging and electro-galvanizing .......................................................... 27

3.3.2 Thermal desorption spectroscopy measurement ........................................................... 28

3.4 Mechanical tests to investigate HE behavior ........................................................................ 29

3.4.1 Slow Strain Rate Test (SSRT) ...................................................................................... 29

3.4.2 Incremental Loading Test (ILT).................................................................................... 29

3.4.3 Constant Loading Test (CLT) ....................................................................................... 30

3.4.4 U-bending test ............................................................................................................... 31

3.4.5 Four-point-bending test ................................................................................................. 33

3.5 Forming tests ......................................................................................................................... 34

3.5.1 Determination of the Forming Limit Curve (FLC) ....................................................... 34

VIII

3.5.2 Determination of the Bending Limit Curve (BLC) ....................................................... 36

3.6 Investigation of the car component-similar samples ............................................................. 37

4 Characterization of As-delivered Materials .................................................................................. 40

4.1 Chemical composition........................................................................................................... 40

4.2 Microstructure characterization ............................................................................................ 40

4.3 Mechanical properties ........................................................................................................... 47

5 Characterization of Pre-charged Materials.................................................................................... 49

5.1 TDS analysis of pre-charged samples ................................................................................... 49

5.2 GDOES analysis of electro-galvanized samples ................................................................... 51

6 Effects of Hydrogen on Mechanical Properties ............................................................................ 53

6.1 Results from Slow Strain Rate Test (SSRT) ......................................................................... 53

6.2 Results of Incremental Loading Test (ILT) .......................................................................... 62

6.2.1 Hydrogen-induced cracking on holed samples ............................................................. 62

6.2.2 Hydrogen-induced cracking on welded samples ........................................................... 65

6.3 Results of Constant Loading Test (CLT) .............................................................................. 69

6.4 Results of U-bending tests .................................................................................................... 71

6.4.1 U-bending on pre-charged samples ............................................................................... 71

6.4.2 U-bend samples in salt solution .................................................................................... 77

6.4.3 U-bend samples in other corrosion mediums ................................................................ 82

6.4.4 Characterization of different edges ............................................................................... 85

6.5 Results of four-point-bending test ......................................................................................... 89

7 Effects of Hydrogen on Forming Behavior ................................................................................... 92

7.1 Forming Limit Curve (FLC) ................................................................................................. 92

7.2 Bending Limit Curve (BLC) ................................................................................................. 96

8 Behaviors of Component-similar Samples.................................................................................... 98

8.1 Results of 90° V-bend samples ............................................................................................. 98

8.2 Results of cross-die samples ................................................................................................. 99

9 Discussion ................................................................................................................................... 102

9.1 Evaluation of testing methods to investigate HE of AHSS ................................................. 102

9.2 Evaluation of susceptibility to HE in different materials .................................................... 109

IX

9.3 Influence of process-immanent parameters ........................................................................ 116

9.3.1 Influence of welding ................................................................................................... 116

9.3.2 Influence of the edge condition ................................................................................... 118

9.3.3 Influence of strain state ............................................................................................... 120

9.3.4 Influence of corrosive surroundings............................................................................ 121

9.4 Validation of mechanical tests with component-similar tests ............................................. 122

9.5 Consequences for using AHSS for car body applications .................................................. 123

10 Conclusion .............................................................................................................................. 127

Reference ............................................................................................................................................ 130

Curriculum Vitae ................................................................................................................................ 139

X

XI

Abbreviations and symbols

AHSS Advanced High Strength Steels

bcc Body-centered cubic

BH Bake-Hardenable

BLC Bending Limit Curve

CE Carbon Equivalent

CLT Constant Loading Test

CP Complex Phase

DF Delayed Fracture

DP Dual Phase

EBSD Electron Back Scatter Diffraction

eBLC enhanced Bending Limit Curve

EDX Energy-Dispersive X-Ray

EI Embrittlement Index

EG Electro-galvanized

FE Finite Element

fcc Face-centered cubic

FLC Forming Limit Curve

FPZ Fracture Process Zone

GDOES Glow Discharge Optical Emission Spectrometry

GI Hot-dip galvanized

HAR Hydrogen Absorption Reaction

HAT Hydrogen Adsorption Theory

HAZ Heat Affected Zone

HE Hydrogen Embrittlement

HEAC Hydrogen Environment Assisted Cracking

HEDE Hydrogen Enhanced De-cohesion

HELP Hydrogen Enhanced Localized Plasticity

HER Hydrogen Evolution Reaction

HFT Hydrides Formation Theory

HIC Hydrogen-Induced Cracking

HISCC Hydrogen-Induced Stress Corrosion Cracking

HSLA High-Strength, Low-Alloy

HSS High Strength Steel

IHAE Internal Hydrogen Assisted Embrittlement

ILT Incremental Loading Test

HPT Hydrogen Pressure Theory

IQ Image Quality

IR-trap Irreversible trapping site

IZ Inside Zone

KAM Kernel Average Misorientation

LME Liquid Metal Embrittlement

LOM Light Optical Microscope

LS Localized Slip

MS Martensitic

MS-W Martensitic (warm rolled)

PREN Pitting Resistance Equivalent Number

QP Quenching and Partitioning

R-trap Reversible trapping site

RD Rolling Direction

SCC Stress Corrosion Cracking

SEM Scanning Electron Microscopy

SFE Stacking Fault Energy

SSRT Slow Strain Rate Test

SZ Surface Zone

TBF TRIP-aided Bainitic-Ferrite

TD Transverse Direction

TDS Thermal Desorption Spectroscopy

TRIP Transformation Induced Plasticity

TWIP Twinning Induced Plasticity

WZ Welded Zone

XII

Au Uniform elongation

Af Total elongation

c Equilibrium concentration of hydrogen under stress

C Hydrogen concentration

c0 Constant of hydrogen concentration

cH Hydrogen solubility in iron substrate

D Hydrogen diffusion coefficient in crystallographic lattice of metallic solid

D0 Pre-exponential factor of hydrogen diffusion coefficient

E´, ∆Ex Barrier energy to activate the hydrogen from trapping sites

Ea Activation energy for hydrogen diffusion in perfect crystal lattice

∆Etrap Activation energy for hydrogen interaction with trapping sites

F Force

Fmax Maximal force until fracture

n Number of cracks

H2 Molecule hydrogen

Hab Absorbed hydrogen atoms

Had Adsorbed hydrogen atoms

∆H Partial molar heat

∆HV Hardness difference

JD Diffusion flux density

JH Diffusion flux density of hydrogen

ltot Total length of cracks

L Distance between the mounted rollers in bending test

PH Hydrogen partial pressure

Q Activation energy of hydrogen

R Universal gas constant, 3.814×10-3

R Bending radius

Ra Surface roughness

Rm Tensile strength

Rp0.2 Yield strength

Rp0.2/ Rm Yield ratio

t Thickness of steel sheet

tc Time to crack initiation

tf Incubation time to delayed fracture

T Temperature

Mole volume of hydrogen

y Displacement of lower points in four-point-bending test

ymax Maximum allowed bending grade

α Pre-coefficient of hydrogen solubility in iron substrate

α Ferrite phase

α´ Alpha martensite phase

γ, γA retained austenite

ε Epsilon martensite phase

σ Applied stress

σf Applied stress to fracture

Hydrostatic stress

σlocal Local stress

σmarco Marco-stress

σmax Maximum allowed stress

φv Effective strain

φ1 Major deformation degree

φ2 Minor deformation degree

1

1 Introduction and Objective

The utilization of Advanced High Strength Steels (AHSS) in the automobile industry has brought great

advances in increased passenger safety and weight reduction of the car body construction to meet the

requirements of fuel efficiency. In recent years, the steel grades of AHSS have typically been developed

to get a better balance between strength and ductility. These materials are required to produce the

components in body-in-white structures with minimal forming, such as the center tunnel and crash

boxes [WOR14]. As reported by Voestalpine, the newly developed third generation of AHSS with

excellent ductility and no loss of strength are a cost-efficient option in comparison to Twinning Induced

Plasticity (TWIP) Steel and press hardening steels [HEB14].

To make further progress in the automotive lightweight construction, the phenomenon of Hydrogen

Embrittlement (HE) in the materials needs to be paid attention to. HE in metals is defined as a loss in

ductility leading to brittle fracture due to the introduction and diffusion of hydrogen into the crystal

structure of the metal [KAE03]. It has been shown in previous work that with increasing tensile strength

the hydrogen influence on the material properties is more pronounced [DEP11]. Hydrogen uptake can

take place in steel-making or the automobile manufacture processes because of the high mobility of the

hydrogen atom in the steel crystal structure. In the service life of an automobile, Hydrogen-Induced

Cracking (HIC) may occur along with environment-assisted corrosion. Accordingly, a good

understanding of the susceptibility to HE of steels, especially AHSS with tensile strength over

1000 MPa, is necessary to ensure their safe application in automobiles.

The influence of hydrogen on the mechanical properties of steels has been reported in several scientific

works [MOL14] [THI11a] [DEP11] [GUO14] [ZHO14]. The previous investigations usually focused

on the mechanism of HIC and the mechanical properties of special steel grades via different testing

methods. Few studies compared the various testing methods to evaluate the HE of AHSS. There is still

a limited understanding of how to interpret the research results and apply them to industrial applications.

Hydrogen embrittlement of AHSS is still a big challenge for the automobile industry since it can be

influenced by many factors and its mechanisms are quite complex.

The motivation of the present work is to provide a guideline to the industrial automobile production and

in-service conditions to prevent HE from occurring in the vehicle. As illustrated in Figure 1.1, this

thesis focuses on two main aspects in AHSS for car body construction. On the one hand, suitable

methods to evaluate the HE are identified. On the other hand, the related influences of materials and

production processes are investigated. Combining the test methods and related influences, the hydrogen

embrittlement of AHSS in automobile industry can be better understood, leading to recommendations

for the material selection and car manufacture.

2

For the aims of this work, various steel grades with tensile strengths of 1180 MPa were selected and

characterized; the 980 MPa steels were investigated as a reference. The phenomena of Internal

Hydrogen Assisted Cracking (IHAC) and Hydrogen Environment Assisted Cracking (HEAC) of the

materials are both studied. To investigate IHAC, the materials were pre-charged to modify their internal

hydrogen content. The hydrogen content was determined by Thermal Desorption Spectroscopy (TDS)

measurement. Corrosive mediums were applied to investigate the HEAC. Different mechanical tests,

such as Slow Strain Rate Test (SSRT), Incremental Loading Test (ILT), Constant Loading Test (CLT),

U-bending test and four-point-bend testing were carried out to investigate the mechanical properties

and delayed fracture behavior as a function of hydrogen content or stress distribution. The influence of

hydrogen on the formability of the materials was quantified by determining the Forming Limit Curve

(FLC) and Bending Limit Curve (BLC). Additionally, the component-similar parts were tested to

validate the HE of the components similar to typical automobile parts. Fractography using a Scanning

Electron Microscopy (SEM) helped to understand the damage mechanisms of different materials in

varying test conditions.

Figure 1.1: Structure of the present work, to investigate the HE of AHSS for automotive applications.

3

2 Theoretical Fundamentals

2.1 Developments and applications of HSS and AHSS

In the last decades, many efforts have been made to meet the requirements of lightweight constructions

in the field of steels development. The European legislation requirement specifies a target in the

reduction of CO2 emissions. How to make the car body lighter is therefore one of the biggest challenges

in the automobile industry. The purposes of high vehicle safety combined with thin-walled car

components are realized by the application of steels with high strength and elongation.

In 1990s, the High-Strength, Low-Alloy (HSLA) steels became the commonly-used materials in the

automobile industry, and its application achieved approx. 50% in mass. Since then, High Strength Steels

(HSS) and Advanced High Strength Steels (AHSS) have been replacing mild steels due to the higher

deformation energy absorption in crash situation [MA11]. The HSS, e.g., Bake-Hardenable (BH) steels

were used in many car components like body closures [BIL13]. In the 3rd International Conference on

Steels in Cars and Trucks, the FutureSteelVehicle programs showed that with the intensive use of new

HSS combining optimized steel design, a weight-reduction of 25%-35% in body-in-white structure

could be reached [CLA11]. In recent decades, AHSS was developed and replaced parts of HSS, thus

the application was significantly growing [WOR14]. Thanks to an excellent balance of tensile strength

(over 600 MPa) and formability, currently AHSS are developed and applied in various components in

body-in-white constructions, such as crash boxes, front end structures, A and B pillars, roof rails and

sill reinforcements [WOR14], as shown in Figure 2.1. As reported by TATA steel in 2014, the

application of new developed HSS/AHSS instead of HSLA led to a 23% weight reduction of the

longitudinal parts and reduction of 42.6 kg CO2 emissions, considering the product lifecycle [LAH14].

Figure 2.1: Application of steel and aluminum in the body-in-white construction [LAH14].

The combination of strength and elongation of steels are distinguished by the crystallographic phases,

alloying contents, deformation mechanisms, and etc. As seen in global strength-formability diagram in


Figure 2.2, it illustrates the range of properties available from today’s AHSS grades. AHSS is

developed from the first generation to the third generation, towards better performances with more

excellent formability, lightweight potential and higher strength level [BIL13]. However, material

strength improves at the expense of its ductility.

Figure 2.2: Global strength-formability diagram, illustrating the range of properties available from today’s

AHSS grades [BIL13].

In the 1st generation, AHSS consisted of typical ferritic multiphase steels, which possess primarily

ferrite-based microstructures [MAT09]. Figure 2.3 demonstrates the four typical microstructures in the

1st generation of AHSS: Dual-Phase (DP), Complex Phase (CP), Transformation Induced plasticity

(TRIP) and Martensitic (MS) steels.

Figure 2.3: Typical microstructures of various steels in the first generation of AHSS [LOI14].

Dual-Phase (DP) steel plays a significant role, which nowadays is a standard material for car body

constructions. In DP steel, the martensite-islands are distributed in the soft ferritic matrix. The strength

level is largely dependent on the volume fraction and hardness of martensitic phase [OLI07]. This type

of microstructure enables a tensile strength in the range of 500 to 1400 MPa [BIL13] [KUZ08]. DP

steels have the advantages of low elastic limit, continuous yielding, low yield ratio (Rp0.2/Rm around 0.5)

and high initial strain hardening rate. The differing hardness of the constituent phases leads to stress

localization at inter-phase boundaries during forming. As a consequence, DP steels display limited

2.1 Developments and applications of HSS and AHSS 5

drawability and bendability. The additional of Ti can induce precipitation strengthening in ferrite

reducing the differences in hardness [KUZ08].

In the production process of CP steel, the cooling rate is controlled to promote the formation of the

bainitic phase. The microstructure of CP steel consists of very fine ferrite with a high volume fraction

of martensite and bainite. The less diffusion-controlled bainite transformation leads to a smaller

hardness difference between bainite and ferrite and allows a higher yield point and better formability in

CP steel compared to DP steel. The yield ratio of CP steel can reach 0.7 by controlling the average

temperature of the cold rolling process [BHA05]. The existence of soft ferrite exhibits better strain

hardening properties than that in full bainitic steel. The grain refinement of ferrite is associated with the

additional micro-alloying elements like Nb, Ti and V, which promote precipitation strengthening.

In TRIP steel, the transformation induced plasticity is related to the retained austenite (γA, 7-15%) in

addition to ferrite (50-55%), bainite (30%-35%) and possible martensite (1-5%). During deformation,

phase transformation from austenite to martensite occurs, providing high strength behavior. Compared

to traditional DP steel, TRIP steels display high n-values up to the limit of uniform elongation [MOR02]

and bake hardening [YAK02]. The excellent formability and drawability of TRIP steel enables forming

into complex components in car bodies. By adjusting alloying elements such as Si and Mn for solution

hardening, the available tensile strength level of low alloyed TRIP steel can currently reach 1200 MPa

[BIL13].

Martensitic steel consists of 100% martensitic structure, which mostly consists of martensite with lath

or needle-like morphology. The martensitic structure can be achieved by quenching after annealing (e.g.,

water quenching) to reach the martensitic transformation temperature. As generally known, martensite,

in which C is concentrated, is the hardest phase in steel. Thus martensitic steels display the highest

strength compared to other steel groups in the 1st generation of AHSS. The uniform tensile strength can

be as high as 1600 MPa. The disadvantage of the uniformly hard phase is the lowest formability [MA11].

In order to improve the formability/ductility, a tempering process is industrially applied. In some cases,

the temper-modified martensitic steel is also named as tempered martensitic steel.

In the group of 1st generation of AHSS, the formability is greatly reduced due to the introduction of a

secondary hard phase to achieve a higher strength level. Thus, the available geometries are quite limited.

The development of second generation of AHSS enables excellent formability. In the austenitic high

Mn steels the strength is achieved through a plastic forming process. The Transformation Induced

Plasticity (TRIP) steels and Twinning Induced Plasticity (TWIP) steels are involved in the 2nd

generation of AHSS. The high Mn content (15-30%) stabilizes austenite at room temperature, achieving

the change from body centered cubic (bcc) matrix to face centered cubic (fcc) matrix in the steel

[BLE14].

An essential concept for understanding high Mn steels is the Stacking Fault Energy (SFE). SFE

describes the interaction between the extended dislocations on the closely spaced slip planes [GUO12]


[VIT06] in the fcc crystal. SFE is the most crucial nucleation parameter related to the dominant

deformation mechanism of austenitic Mn steels [CUR10]. The deformation mechanisms in high Mn

steels are discussed based on SFE [LEE12] [CUR10] [COO11]. Table 2.1 summarizes the dominant

deformation mechanism according to different ranges of SFE. Various calculations and research

revealed that SFE depends on the alloying elements such as C and Mn [SCO06]. The additional of Al

in austenitic steel greatly raises the SFE [DUM08] [KIM11].

Table 2.1: SFE and the dominant deformation mechanism in high-Mn steels [GUO12].

SFE [mJ·m-2] Dominant deformation mechanism

<12 austenite→α´martensite or austenite→ε martensite→ α´martensite

12 – 18 formation of ε-martensite 18 – 45 twinning mechanism

>45 dislocation gliding

Owing to the fcc matrix and deformation mechanisms, the 2nd generation of AHSS display excellent

formability and drawability combined with high strength. Typical TWIP steel can achieve over 50%

total elongation at a strength level of around 1000 MPa [BIL14a]. Application in the automobile

industry is, however, greatly limited by its tendency for Delayed Fracture (DF), difficulties in

weldability and the high cost of alloying elements.

The research and development of the 3rd generation AHSS focused on the improvement of formability

of steel without significant sacrifice of strength. The new version multiphase steels with more balanced

formability are economically attractive [SUG08] [HEB08] [HEB14]. In order to provide a maximum

of strain hardening, elongation and thus deep drawability, the introducing of austenite into the

microstructure is the key step [HEB14]. The 3rd generation of AHSS have a medium alloying content

of maximum 10 vol.% of Mn, Si and Cr combined in the steelmaking process of 3rd generation AHSS,

quenching is interrupted by subsequent partitioning, where the C and Mn partition between bcc and fcc

phases during the long diffusion treatment [BIL14b] [BLE14]. The addition of Mn, Si and Cr avoids

iron-carbides precipitation. The redistribution and enrichment of C in austenite effectively increases the

strength of steel due to solid solution strengthening [HEB14]. The existence of retained austenite, as

the key component in this kind of AHSS, leads to phase transformation via the TRIP mechanism into

martensite. The multiphase microstructure can achieve the targeted mechanical properties such as high

tensile strength (uniform tensile strength more than 1000 MPa [BLE14]) and combined better

formability than 1st AHSS; on the other hand, the medium alloying content makes the production cost

is lower than 2nd AHSS.

Figure 2.4 shows a comparison of the characteristic mechanical properties, volume fractions of retained

austenite and additional Mn contents in three typical grades of AHSS for automotive applications. The

Index of Rm·Af is referred to a rough indicator for the overall mechanical properties [BLE14]. The 3rd

2.2 Hydrogen sources in automotive production processes 7

generation of AHSS is still under development. In this classification of AHSS with medium Mn-content,

the Quenching and Partitioning (QP) steel and TRIP-aided Bainitic-Ferrite (TBF) are nowadays

commercially produced by steel makers.

Figure 2.4: Summary of three generations of AHSS for automotive applications [BLE14].

2.2 Hydrogen sources in automotive production processes

Hydrogen atom has small radius of 0.32 Å (1 Å=10-10 m), and its high mobility makes it especially

active in the metal lattice. The hydrogen atoms accumulate and interact with the matrix, especially when

defects in steel exist [PUN05]. Figure 2.5 schematically illustrates the hydrogen accumulation in the

lattice defects in the steel [PUN05]. Thus, the hydrogen transfer from an ambient medium into the solid

steels involves several necessary steps, including the uptake of hydrogen on the medium/steel interface,

the adsorption of the hydrogen atoms, the trapping and diffusivity of hydrogen in steels. The kinetics

of each step is influenced by various conditions such as the ambient medium, surface quality,

temperature, partial pressing and so on. The details of hydrogen behavior in steels will be outlined in

the following sections.

a) hydrogen solubility in the matrix

b) trapping sites on the surface

c) trapping sites on the subsurface

d) cylindrically shaped region

e) edge dislocations

f) vacancies

Figure 2.5: Schematic presentation of defects in metal and accumulation of hydrogen atoms in the low-concentration range [PUN05].


Concerning automotive production and application, hydrogen uptake during steelmaking and car

production processes is considered.

Steel manufacturing processes

In steelmaking processes, hydrogen uptake starts during melting. In metallurgical processes Ca(OH)2

is formed in a highly humid atmosphere. It can decompose into CaO in the slag and water vapor. During

the reaction between iron and water atomic hydrogen enters into the melt [NAB10]. During the followed

procedures of annealing, rolling and cooling, various oxides and hydroxides are formed on the sheet

surface. In the procedure of pickling prior to cold rolling or galvanizing, the sheet surfaces are treated

with solution of sulphuric acid or hydrochloric acid and hydrogen uptake occurs [MAA08].

Before galvanization, the steel sheets are cleaned in an alkaline electrolyte. The steel coils are connected

as a cathode in direct current cycles in order to remove the oil, emulsion and impurities on the band

surfaces. During cathodic cleaning, hydrogen is partially absorbed into the steel [SCH01]. In the process

of hot-dip galvanizing, a mixed nitrogen and hydrogen gas atmosphere is used. Before going to Zn-bath,

an annealing process was carried out in order to achieve aimed microstructure and reduce the iron-oxide

production on the surfaces. Therefore hydrogen uptake takes place [SCH01]. In comparison, hydrogen

is absorbed on the Fe/Zn interface during electro-galvanizing. And the following galvanizing provides

the Zn-layer which plays a role of barrier to prevent hydrogen from effusion [HIL04] [THI11b].

Car body manufacturing processes

The different manufacturing processes in the body-in-white plant also lead to the risk of hydrogen entry.

Figure 2.6 indicates the trend in of hydrogen content and stress levels in variable manufacturing steps.

In car production processes, the hydrogen uptake is measured from joining onwards. During welding

the hydrogen content in car components is increased. The operation of pre-treatment before painting is

carried out in order to optimize the adherence of coating. In the pre-treatment, the car body is firstly

cleaned and degreased. After phosphating, the hydrogen concentration in car body is increased. The

process of cathodic electrophoresis for the painting also causes hydrogen uptake. However, during the

bake operation to finish the E-coating the absorbed hydrogen can effuse from the car body, resulting in

a reduction in hydrogen content [LOI14] [COR14].

2.3 Hydrogen absorption and transport in steels 9

Figure 2.6: Schematic hydrogen and stress levels during steel making, vehicle production and in-service

processes.

Service life of vehicle

During the life of the vehicle, hydrogen uptake occurs during the corrosion process. The painting from

E-coating and Zn-coating on the substrate steel protects the car body from corrosion. But in some special

locations, especially when there are defects on the coating, the direct connection between steel and

surroundings allows the entry of hydrogen. Many environmental factors such as temperature and

humidity affect the corrosion behavior of the car. Acid corrosion takes place when acidic solution or

surroundings exist. In contrast, when oxygen concentration is sufficient or in the condition of alkaline

electrolyte, oxygen corrosion is preferred. The hydrogen entry reactions are described in the following

section.

In general, hydrogen uptake occurs not only from various sources during the steelmaking process

[RAT07], but also from the body-in-white shop during the vehicle manufacturing. Nowadays most

sources of hydrogen are well recognized and several methods are applied to reduce the hydrogen entry.

For instance, the addition of pickling inhibitors effectively reduces the solution of iron and formation

of hydrogen in steels by pre-sorption on the steel surface [MAA08] [SCH01]. However, the introduction

of hydrogen from manufacturing procedure and service life of vehicle is still unavoidable and the

hydrogen concentration cannot be reduced by process control. Therefore, further treatment is essential

to achieve a lower hydrogen concentration.

2.3 Hydrogen absorption and transport in steels

2.3.1 Hydrogen absorption

Atmosphere

In the atmosphere hydrogen is present as a molecule in the gas phase. The thermodynamic equilibrium

on the interphase H2/steel surface is achieved. At room temperature and atmospheric pressure the


tendency to hydrogen absorption is low, because the thermodynamic driving force to dissociate the

molecule hydrogen into the atomic hydrogen is not sufficient. Only in the condition of high purity of

hydrogen gas in partial pressure over 100 bar molecular hydrogen adsorb from gas phase into the solid

metallic phase [SPA00]. The hydrogen adsorption under high partial pressure involves three procedures,

which are illustrated in Figure 2.7 [GEH94] [REH13]. After the deposition of the molecule hydrogen

on the steel surfaces the adsorbed hydrogen molecule is dissociated into hydrogen atoms. The chemical

absorption of hydrogen atoms ( ) may take place.

Figure 2.7: Illustration of the reactions in hydrogen absorption from atmosphere [REH13].

Corrosion medium

When the steel is exposed to corrosive mediums, cathodic corrosion plays a crucial role for hydrogen

absorption. For blank steels, the reduction reaction takes place on the surfaces:

(2-1)

A zinc coating can protect the bare steel from cathodic corrosion, due to the fact that zinc has more

negative corrosion potential than iron, which is more favorable to be dissolved as an anode. However,

when there is a defect in the Zn-coating or the anode is dissolved, the hydrogen attacks the base material

directly. The hydrogen activity deeply depends upon the presence of catalysts or re-combination

inhibitors. Once the hydrogen activity at the interphase between steel substrate and corrosion medium

is achieved, hydrogen adsorption occurs. Regarding the corrosive medium, Hydrogen Evolution

Reactions (HER) take place [ZAK85] [MOO09]:

in acid solution: (2-2)

in neutral or alkaline solution: (2-3)

During the HER, some hydrogen atoms are adsorbed to the surfaces of steel. Here, the hydrogen ion is

oxidized into atomic hydrogen which is adsorbed according to Volmer reactions [KAE03] [GEH94].

This kind of transfer reactions of hydrogen deposition take place on the interface corrosion

medium/steel substrate, as explained in Figure 2.8. After adsorption, a portion of adsorbed atomic

hydrogen is absorbed by the substrate, where the Hydrogen Absorption Reaction (HAR) takes place:

(2-4)

A portion is transferred into adsorbed molecular hydrogen according to the Heyrovsky reaction or the

Tafel reaction [REH13]. The recombination of atomic hydrogen into molecular hydrogen is impeded in


the process [KAE03] [GEH94]. Figure 2.8 shows the schematic illustration of possible reactions in the

process of hydrogen absorption.

Briefly, the process of hydrogen absorption in steel substrate combines several steps, which is

dependent on the factors like cathodic current density, type of the solution, solution composition,

surface condition of substrate, et al. The hydrogen absorbed in steel substrate in two formats: atomic

hydrogen and molecular hydrogen. The absorption is controlled by the rate of HER and HAR [CHA03].

(a) In acidic solution

(b) In neutral or alkaline solution

Figure 2.8: Illustration of the reactions in hydrogen absorption from: (a) acid corrosion mediums and (b) neutral or alkaline solution [REH13].

2.3.2 Hydrogen transport

Hydrogen transport in steels is enabled by interstitial diffusion or the movement of defects. The

generation of Hydrogen-Induced Cracking (HIC) requires the achievement of critical hydrogen

concentration. The dissolved hydrogen diffuses into the region where crack potentially occurs and leads

to the fracture [HWA94]. Hence the study of hydrogen solubility and diffusivity is essential for

understanding the mechanism of HIC.

Hydrogen solubility

Once the interstitial atom of hydrogen enters into the steel substrate, it is dissociated into proton and

electron. The repulsive force between the proton and iron nuclei causes the local lattice distortion.

According to the crystal structure, two typical lattice structures are mentioned here: body-centered cubic

(bcc) structure and face-centered cubic (fcc) structure. In Table 2.2 it shows the spacing and number of

tetrahedral/octahedral interstitial sites in the Fe-alloys with bcc and fcc lattice structure respectively

[BLE07] [GOT04].


Table 2.2: Description of interstitial sites in iron with bcc and fcc lattice structure [BLE07].

Lattice structure bcc fcc

Type interstitial sites Tetrahedral Octahedral Tetrahedral Octahedral

Number of interstitial sites 12 6 8 4

Size of interstitial sites 0.36 Å 0.19 Å 0.28 Å 0.53 Å

On the basis of the interstitial sites of iron lattice, the hydrogen atoms with radius of 0.32 Å can occupy

the tetrahedral interstitial sites in bcc structure (diameter 0.36 Å) and octahedral interstitial sites in fcc

structure (diameter 0.53 Å) without causing any lattice distortion. Moreover, the fcc and bcc lattice

structures display different solubility behaviors for hydrogen atoms. The solubility of hydrogen in fcc

iron is higher than that in bcc iron, because of the higher ability of octahedral interstitial sites to expand

[BLE07]. Once the hydrogen atom enters into iron substrate as an interstitial atom, the solubility in

atomic fraction cH can be described by Sievert’s law [MIS05]:

(2-5)

where stands for the pre-coefficient, presents for the partial molar heat, T presents for the

temperature in unite of K and is the hydrogen partial pressure.

Concerning the equilibrium based on Sievert’s law, the solubility of interstitial hydrogen is deeply

dependent on the temperature and partial pressure. Generally, hydrogen originally dissolved in pure

iron at high temperatures, in the state of liquid iron. During solidification the dissolved hydrogen has a

tendency to effuse from the substrate. As the temperature decreases, hydrogen solubility is decreased

at the same time, in particular when austenite-ferrite transformation occurs. At room temperature and

external hydrogen pressure of 105 Pa, the atomic fraction of hydrogen in pure iron is extremely low

with a value of 2×10-8 [HIR80].

Nevertheless, in real condition of commercial steels, the hydrogen solubility stays at a higher level due

to the presence of lattice defects [HIR80]. With the existence of crystallographic defects, a certain

amount of hydrogen is able to be retained, even below the specific temperatures. In some cases the

hydrogen contents in the steels after cold working process are many times higher than that after the hot

working process [ORI85]. The explanation is that during the cold working process, more lattice defects

were introduced than that in hot working process, resulting in a measurable increasing of hydrogen

content. The details of trapping sites of hydrogen are illustrated in the following section. Besides defects,

alloying elements also play a role in the effect of hydrogen solubility. Like lattice defects, the

introduction of mostly alloying elements such as Mo, W, Ni and Cr decrease the maximum solubility

of hydrogen in the steels through reducing the grain boundary areas [CHO81].


Hydrogen diffusivity

Once the hydrogen atoms occupy the interstitial sites, it can rapidly diffuse in the steel [STR06]. The

hydrogen concentration gradient is decisive to hydrogen diffusion, providing the driving force to

hydrogen atoms [GUO12]. Fick’s first law describes the diffusion flux density , which is defined as

the number of hydrogen atoms that passing through a unit area in the unit time, is proportional to the

concentration gradient (equation (2-6)).

(2-6)

where C stands for the hydrogen concentration and D is referred to hydrogen diffusion coefficient in

crystallographic lattice of metallic solid, which can be expressed by the Arrhenius equation (2-7):

(2-7)

where D0 is a pre-exponential factor, Q is the activation energy, which differs between different crystal

structures. R is the universal gas constant with a value 3.814×10-3 J·mol-1·K-1 and T is the temperature.

The diffusion coefficient is a characteristic factor for the materials and varies with different

microstructures. The difference between ferritic steel ( -Fe) and austenitic steel ( -Fe) is also

mentioned here. In austenitic steels the diffusion coefficient differs by several orders of magnitude

compared to that of ferritic steels. Austenitic steels display high hydrogen solubility but low hydrogen

diffusivity [FIG10]. The theory of trapping can be the explanation for hydrogen behavior in the various

microstructures. Table 2.3 summarizes several diffusion coefficients of different steels in previous

research. The diffusion coefficient is determined by multiple factors, regarding the amount and type of

hydrogen trapping sites in the materials. Lattice defects, such as dislocations, grain boundaries and

inclusions act as trapping sites for hydrogen. The trapping sites impede the migrating of hydrogen due

to the consequent trapping energy, resulting to the lower hydrogen mobility [FIS99] [GRA00].

Table 2.3: Diffusion coefficients of hydrogen in different steels in literature.

Material D [cm2·s-1] Literature

Pure iron 7.5 10-5 [ZAK99]

Carbon steel 2.5 10-6 [GEO97]

Austenitic stainless steel 2.15 10-12 [SMI00]

Dual-phase steel 2-2.7 10-7 [LOI14]

Complex-phase steel 5 10-7 [LOI14]

Martensitic steel 2.6 10-7 [LOI14]

Press-hardened steel 4.9 10-7 [LOI14]

The diffusion of hydrogen is determined not only by the crystallographic structures, but also by the

alloying elements. The interstitial alloying elements, such as C, have a significant influence on the


diffusion behavior of hydrogen and increase hydrogen diffusivity. Additionally, the introduction of

substitutional alloying elements such as Ni, Mn and Cr leads to a reduction of hydrogen diffusivity

[GRA00] [RIE85].

Considering the external conditions, temperature has the highest influence on the hydrogen diffusivity

in the materials. As expressed in equation (2-7), the diffusion coefficient is proportional to the

temperature [GOT04]. Another factor affecting hydrogen diffusivity is the stress state. Hydrostatic

stress is introduced in the area ahead of a crack tip, and consequently reaches a maximum within a short

distance from the crack tip. When a crack exists, hydrogen atoms in the metal bulk diffuse and

accumulated at the deformed site due to the expansion of the lattice [OLD09].

Hydrogen trapping

Usually, the distribution of hydrogen in a steel substrate is not homogenous. Many heterogeneous

structures (such as vacancies, dislocations, grain boundaries, voids and inclusions) are present in the

steels. These defect structures change the ideal hydrogen behavior and cause localized hydrogen

enrichment, acting as trapping site for hydrogen atoms. The presence of trapping sites increases the

hydrogen solubility and decreases the apparent hydrogen diffusivity. If the hydrogen is passing through

a trapping site, a longer residence time is required than that in the situation of normal interstitial lattice

site. As a consequence, the transport rate of hydrogen atoms through the material is reduced and the

hydrogen permeability in the steels is lowered [CHA03].

The trapping sites can be defined by the deepening of the associated potential well. As showed in the

sinusoidal curve (Figure 2.9), the activation energy for hydrogen diffusion in perfect crystal lattice is

expressed as Ea. The trapping site has a higher barrier energy E´ and ∆Ex than the normal sites. In the

diffusion process, hydrogen requires the energy higher than Ea+E´ to get into the trapping sites and

Ea+E´+∆Ex to escape from the trapping sits [VIY05].

Figure 2.9: Schematic illustration of interaction energy of atomic hydrogen and trapping sites [VIY05].

The trapping sites can be classified into reversible and irreversible trappings. Figure 2.10 schematically

demonstrates various trapping sites concerning the activation energy without concerning the barrier

energy [BIR95]. The reversible trapping sites (R-trap) are related to the low binding energy; the trapped


atomic hydrogen in reversible trappings is easily to be reactivated and becomes mobile. The trapped

hydrogen in reversible traps has tendency to participate in and enhance the HIC. On the contrary, the

irreversible traps (IR-trap) have much higher binding energy than reversible trapping sites. The resident

of hydrogen can be considered permanent. At ambient temperature the irreversible trapping sites do not

contribute to the hydrogen diffusivity and enrichment; when heated to very high temperatures, the

trapped hydrogen atoms can be released [ALB03].

The required activation energy for hydrogen interaction with trapping sites (∆Etrap), regarding trapping

sites and material systems is summarized in Table 2.4. Depending on the binding energy for hydrogen

dissolution in steel, four kinds of trapping sites are determined: I) point defects with ∆Etrap lower than

15 kJ·mol-1 are characterized as week traps, II) point defects with ∆Etrap over 20 kJ·mol-1 is characterized

as moderate trapping sites, III) dislocations with moderate to strong interaction is considered as strong

traps (∆Etrap in the range of 20-60 kJ·mol-1) and IV) grain boundaries or interface surfaces with strong

interaction (∆Etrap is 50-100 kJ·mol-1) [VIY05].

Figure 2.10: Schematic illustration of various trapping sites in terms of trapping energies [VIY05].

Table 2.4: Activation energy for hydrogen interaction with trappings, concerning trapping sites and material systems [VIY05] [GUO12].

Trapping site ∆E [kJ·mol-1] Material system

Lattice site 7-8 Fe-Ti alloy

Vacancy 22 bcc-steel

Dislocation 24-31 Fe-Ti alloy / bcc-steel

Grain boundary 53-59 Fe-Ti alloy / bcc-steel

Twin boundary 62 Fe-18Mn-1.5Al-0.6C steel

Micro void 35-48 carbon steel

MnS 40.5 bcc-steel

Al2O3 79 bcc-steel

TiC 95 bcc-steel


2.4 Hydrogen-induced cracking in steels

The term of Hydrogen Embrittlement (HE) is defined as “An important danger subject to tensile load

is unexpected brittle, instead of usually predictable ductile, fracture and a notorious variety is

embrittlement by uptake of atomic hydrogen as a solute in the metal lattice” [KAE03]. For AHSS, the

interaction of dislocations and existing hydrogen atoms aggravate the danger of cracking under

mechanical stress [KAE03]. The Hydrogen Induced Stress Corrosion Cracking (HISCC) is regarded as

the cooperation of hydrogen and mechanical loading. According to the source of affected hydrogen,

HISCC can be classified into Internal Hydrogen Assisted Cracking (IHAC) and Hydrogen Environment

Assisted Cracking (HEAC).

In IHAC, the internal dissolved atomic hydrogen is decisive for the cracking. An illustration of the

contribution of hydrogen to IHAC is shown in Figure 2.11: a macroscopic notched sample with internal

atomic hydrogen is under loading (Figure 2.11 (a)). In the region of notch the local stress has a high

value of σmacro and a stress gradient is obtained. The stress gradient acts as the driving force for the

diffusion of hydrogen, leading to the hydrogen re-contribution. The atomic hydrogen diffuses into the

high stress concentrated region like notched area and crack tips, where the high local concentration of

hydrogen is achieved (Figure 2.11 (b)). Under loading, once the critical hydrogen concentration to

crack is reached, HIC takes place, even when the applied environment is moderate and no

incursion of external hydrogen [INT05]. At the inhomogeneity, e.g., inclusions, a local stress

is superposed, which leads to also HIC [GER88] [SYM01].

(a) (b)

Figure 2.11: Distribution and diffusion tendency of hydrogen by IHAC: (a) presence of hydrogen in the steel before loading and (b) re-contribution of hydrogen after crack initiation [REH13].

In HEAC, the mechanical degradation and/or fracture is influenced by the environmental hydrogen.

The process of HEAC is always combined with the conjoint action of mechanical loading and chemical

reaction [ASTM08]. In corrosive surroundings the oxide film on the surface is firstly broken and

corrosion pits are created, then atomic hydrogen is absorbed around the corrosion pits. As shown in

Figure 2.12, on the bent sample in the aqueous solutions, the atomic hydrogen is primarily produced

2.5 Mechanism of hydrogen-induced cracking 17

by absorption from environment into the steel. Hydrogen atoms enrich at the crack tips due to the

enhanced hydrogen solubility and trappability by lattice distortion [GER88] [SYM01], and the bare

surface and subsurface areas (Figure 2.12 (a)). The absorbed hydrogen atoms diffuse in the direction

of the stress gradient, leading to local hydrogen concentration, which assists the final Stress Corrosion

Cracking (SCC) (Figure 2.12 (b)).

(a) (b)

Figure 2.12: Distribution and diffusion tendency of hydrogen by HEAC: (a) presence of hydrogen in the steel before loading and (b) re-contribution of hydrogen after crack initiation [REH13].

2.5 Mechanism of hydrogen-induced cracking

The fundamental mechanisms are intensively investigated and discussed to understand and explain the

phenomena of HE. The process of HE is due to the interaction between dislocations and atomic

hydrogen atoms on a very small scale, thus it is not easy to verify the proposed theories. So far,

following mechanisms are accepted in the research field: Hydrogen Pressure Theory (HPT), Hydrides

Formation Theory (HFT), Hydrogen Adsorption Theory (HAT), Hydrogen Enhanced Localized

Plasticity (HELP) and Hydrogen Enhanced De-cohesion (HEDE). It should be noted that, on the basis

of research until now, it is impossible to explain HE with one certain mechanism satisfactorily, since

the combination of several individual modes probably contributes [LOI14].

Hydrides Formation Theory (HFT)

In metallic materials like steel, the hydrides may exist in different forms, and different hydride potentials

cause different forms of metal-hydrogen bonds [NEL83]. The transition hydrides, which are formed by

the interaction of transition metals (e.g., V, Nb, Ti, etc.) and hydrogen, contribute to the HE [BIR95].

During loading and ductile failure, the precipitated brittle hydrides decrease the ductility and strength

of the materials. The hydrides are formed within the stress field at the crack tips, and cracks propagate

along the cleavage plane inside of hydrides [BIR95]. Figure 2.13 schematically depicts the hydride


theory [VIY05]: (a) the atomic hydrogen diffuses into the hydrostatic stress field (crack tips) by the

driving force of stress gradient (JH: diffusion flux of hydrogen); (b) local reduction of hydrogen

concentration leads to the formation of metallic hydride; (c) occurrence of cracking along the cleavage

plane inside of brittle hydride and (d) a repeat process causes the crack propagation. Thus a fracture

surface with brittle mode can be found on the fracture surfaces. The phenomenon is typical observed in

Ni-alloy and Al-alloy, especially when the hydride forming elements exist. Since in the standard

conditions hydrides cannot be formed in iron, the hydride formation theory is neglected in further

discussion of hydrogen-induced fracture in the steels.

(a) (b) (c) (d)

Figure 2.13: Schematic illustration of hydride theory showing crack propagation in the metallic hydride region ahead of crack tip by stress induced hydride formation [VIY05].

Hydrogen Adsorption Theory (HAT)

In Griffith criterion it is proposed that for the ideal brittle material the fracture stress is proportional to

square root of the surface energy [ENG69]. Hydrogen adsorption theory explains the HE by the loss of

surface energy due to adsorption of atomic hydrogen [PET52]. It is proposed that the adsorbed hydrogen

associates the reduction of surface tension. The adsorption theory is still contentious for its adequacy to

explain the crack propagation. Some researchers have pointed out, that the adsorption theory

underestimates the required energy to create fracture surfaces [HIR95].

Hydrogen Pressure Theory (HPT)

Hydrogen pressure theory explains HE on the basis of recombination of hydrogen atoms with

subsequent volume increasing. It is proposed that at the defect position the process of atomic hydrogen

recombination to molecular hydrogen takes place. The new produced hydrogen molecular precipitates

at local defects, resulting in the generation of high pressure and weakening of metal lattice [ZAP41].

Once the precipitation of hydrogen molecular is pronounced or the reinforcement of the lattice is weak,

nucleation and growth of micro-cracks may occur in the absence of externally applied stress [VIY05].

In the case of cathodic electrolysis, the pre-charged hydrogen can be great supersaturated. Once a certain

degree of supersaturation is achieved and the local hydrogen pressure is exceeding the elastic strength,

the lattice bonds is broken [VIY05]. The hydrogen pressure theory is significant to some cases such as

2.5 Mechanism of hydrogen-induced cracking 19

welding process, in which it is generally considered that the molecules are formed at high temperature

and high pressure is generated during cooling procedure [VIY05].

Hydrogen Enhanced Localized Plasticity (HELP)

The fundamental of Hydrogen Enhanced Localized Plasticity (HELP) is the interaction between

hydrogen and dislocations. The dissolved hydrogen atoms exist in the dislocation spacing and strain

fields around dislocations [TAB84]. The presence of hydrogen influences the dislocation mobility,

representing hardening effect or softening effect [NEL83].

The hardening effect can be explained by the formation of Cottrell-atmosphere in the high tensile stress

fields. The Cottrell-atmosphere hinders the movement of dislocations, making the deformation at crack

tip more difficult [VIY05]. As studied, the stabilizing of micro-cracks as well as the effect on work

hardening rate by hydrogen, a degradation of mechanical properties can be found. [LOU72]. By contrast,

in the effect of softening it is proposed that hydrogen reduces the fracture strength by facilitating the

movement of hydrogen on the crack tip surface; because the introduction of high concentrated hydrogen

increases the neighbor atoms on the surface, resulting in easier dislocation nucleation [LYN79].

The dissolved hydrogen reduced the interaction between the dislocations by shielding the elastic

interaction between dislocations with each other as well as with lattice defects. With the existence of

hydrogen, the pile-up dislocations move closer to each other [FER98]. In this manner, the elastic

energies of interaction between dislocations and other internal elastic stress fields are reduced by the

accumulation of hydrogen. Dislocation movement is eased by the elastic shielding effect and thus the

dislocation mobility is increased [BIR90]. The more easily overcome slip barrier makes the local

dislocation movement occur even at low shearing stress and thus deformation is facilitated [WAN01].

Due to supporting shielding calculations and microscopic observations, the HELP mechanism is well

established respect to the influence of hydrogen on the dislocation behavior [STR06] [TAH01].

Hydrogen Enhanced De-cohesion (HEDE)

In the mechanism of Hydrogen Enhanced De-cohesion (HEDE), the HE is related to the weakening of

the cohesion force between metallic atoms in the matrix due to the partial crucial accumulations of

hydrogen. In the elastic stress field the crystal lattice is dilated, providing spacing for hydrogen

accumulation. The trapped hydrogen diffuses into this region and is concentrated a short distance from

the crack tips. This reduces the atomic cohesive binding strength of metallic lattice [ORI72] [TRO60].

In non-hydride metals, the propagation of the crack is determined by the maximum tensile stress and

the maximum cohesive resistive force of the matrix metal. Once the maximum tensile stress at crack

tips is equal to the maximum cohesive force between the metallic atoms, crack propagation occurs


[ORI72] Thus, it is considered that hydrogen damage sites are located in the vicinity ahead of the crack

tip surface where tensile stress reaches maximum [GAN03].

HEDE explains delayed cracking in HE and is appropriate for the explanation of intergranular fracture.

It is considered that when the high concentration of hydrogen accumulates at energetically favorable

positions like grain boundaries, the cleavage fracture in intergranular mode occurs once the critical

concentration is reached [STR06]. However, the HEDE mechanism cannot explain the plastic parts on

the hydrogen embrittlement fracture surfaces. The theoretically predicted weakening on the metallic

matrix by interstitial dissolved hydrogen has not yet been experimentally confirmed [HIR87].

2.6 Important factors influencing hydrogen-induced cracking

The Hydrogen-Induced Cracking (HIC) is microscopically regarded as the consequence of interaction

between material, hydrogen and the stress, as showed in Figure 2.14.

Figure 2.14: Three main local factors for hydrogen-induced cracking.

Material

The influence of material-imminent factors on hydrogen-induced fracture has been investigated by

different researchers [KIM11] [TAK12] [GUO14]. Acting as hydrogen trapping sites, as mentioned in

section 2.3, lattice defects such as dislocations, grain boundaries, precipitates and inclusions play an

essential role on HIC. The binding energies with of different lattice defects decide the trapping types.

Dislocation mobility is increased by the presence of hydrogen [TAB84]. Takai et al. pointed out that in

the presence of hydrogen. The new vacancies generated by dislocation interaction contribute to the

hydrogen degradation [TAK08]. As studied by Fuchigami et al. [FUC06] and Takasawa et al. [TAK12],

decreasing dislocation density effectively facilitated the resistance to hydrogen embrittlement, owing

to the longer dislocation interval suppressing the formation of vacancies by dislocation dynamics.

Nevertheless, an increase in the dislocation density can increase the amount of tangled dislocations and

resulting in local hydrogen enrichment [LI14], which increases the cracking risk.

Grain refinement increases the area of grain boundaries and leads to enhance of the resistance to HE. It

leads to not only decrease the slip length of dislocations [TAK12], but also reduce the concentration of

2.6 Important factors influencing hydrogen-induced cracking 21

segregated hydrogen at grain boundaries [TAK12] [CHE14]. In TWIP steels, grain refinement plays an

essential role on HE, because it reduces the volume fraction of mechanical twins [KAN10] [PAR15]

[BOU11]. As investigated by Park et al. [PAR15], owing to the low volume fraction of twins, the

migration of hydrogen atoms in grain-refined TWIP steels is less prevalent than in coarse-grained

materials.

Du et al. [DU15] have reported that the HIC primarily nucleates at the non-metallic inclusions such as

MnS and Al2O3 in A537 Steel. Jin et al [JIN10] investigated in AP15L X100 steels and made the

conclusion that “The cracks are primarily associated with the Al- and Si-enriched inclusions, rather

than the elongated MnS inclusion”. In TWIP steel, the Nano-sized (Fe, Mn)3C cementite precipitates

along grain boundaries and reduce the DF resistance [HON13]. On the contrary, the homogeneously

distributed Ti(C,N) is reported as a beneficial trapping site for hydrogen, which can prevent the

formation of MnS and reduce diffusive hydrogen in the welded X70 pipeline steel [BEI09]. Generally,

the inclusions/precipitates play a detrimental role on HIC, since they act as affective hydrogen trapping

sites.

Hydrogen

In IHAC, the internal hydrogen dominates the cracking. As mentioned in Section 2.2, hydrogen entry

as well as residual stress cannot be eliminated in the steel manufacturing process. Inhomogeneity such

as hard second phases, inclusions, precipitates, micro-cracks on the surface or in the bulk enhance the

enrichment of redistributed hydrogen. In this case, HIC takes place easily [GER88] [SYM01].

Hydrogen-induced mechanical degradation and DF should be carefully considered before application.

In HEAC, however, DF is related to the overlap of internal dissolved hydrogen and the absorbed

hydrogen from the applied environment. Figure 2.15 illustrates the chemical reactions and hydrogen

migration processes from different hydrogen sources: gaseous hydrogen, water vapor and electrolyte

[GAN03].


Gaseous hydrogen Water vapor Electrolyte

(1) Molecular transport Molecular transport Oxygen depletion

(2) Physical adsorption H2 Physical adsorption H2O Cation dissolution, hydrolysis

(3) H2 migration, dissociation to atomic H Chemical reaction to FexOy·H2O, H H+ reduction to H

(4) Chemical adsorption Water reduction to H

(5) Solution to lattice, trap sides

(6) Bulk diffusion to FPZ

(7) Enhanced fatigue damage

Figure 2.15: The sequence of elemental processes that supply damaging H to the crack tip Fracture Process Zone (FPZ) [GAN03].

Stress

Both external and local residual stresses play an essential role in HE. During deformation, the plastic

incompatibility provides the crack initiation sites and the high stress concentration and deformation

strain distribute inhomogeneous. The stress gradient dominates the diffusion of atomic hydrogen.

Hydrogen assisted cracking takes place when the internal hydrogen atoms agglomerate at the crack

initiation sites, even in high purity material in which there is no segregation at grain boundaries [LIU08].

As assumed by Sidorenko et al. [SID75], the equilibrium concentration of hydrogen under stress c is

raised with local stress in (2-8):

) (2-8)

where c0 is the constant of hydrogen concentration, stands for mole volume of hydrogen and

stands for the hydrostatic stress.

The local stress dominates the crack initiation and propagation. Stress gradient functions as the driving

force of hydrogen diffusion and re-contribution. Rehrl has described the relationship between local

stress and concentration profile of hydrogen at the crack tip [REH13]. As illustrated in Figure 2.16, at

the crack tip both the hydrogen concentration and local stress are close to the peak value. The

concentrated hydrogen reacts with the matrix, by enhancing the plasticity (HELP mechanism) or

reducing the cohesion force of matrix (HEDE mechanism). Thereby the critical hydrogen content for

HIC is reduced because of the concentrated stress. As consequence, HIC takes place when the hydrogen

content achieves the threshold value.

2.6 Important factors influencing hydrogen-induced cracking 23

Figure 2.16: Distribution of hydrogen concentration and local stress at the crack tip [REH13].

In some cases, when pre-strain level is kept low, the uniform distribution of introduced dislocations

disperses hydrogen atoms and reduces the local hydrogen enrichment [LI14]. However, high strain

leads to hydrogen concentration increasing remarkably. Consequently, cohesion force between atoms

in matrix is deeply weakened, leading to the HIC.

Moreover, the stress and strain condition especially affects the DF behavior in TRIP-/TWIP steels

because it is decisive to the phase transformation or mechanical twinning mechanism, which influences

the hydrogen behavior in steels.

In summary, the HIC and its deformation mechanisms have been investigated in many scientific works.

Hydrogen atoms can entry and dissolve in the trapping sites in the steels, diffuse and re-contribute under

stress. Hydrogen embrittlement can be a challenge to apply the advanced high strength steels over

1000 MPa in automobile industry widely. A comparison of testing methods to evaluate the resistance

to hydrogen embrittlement is needed. The available knowledge focused on the structural steels or 1st

and 2nd generations of advanced high strength steels, concerning the degradation of properties and

delayed fracture behaviors. The new developed 3rd generation of advanced high strength steels

concerning their susceptibility to hydrogen embrittlement is still required to be clarified. Besides, the

influences of automotive production processes immanent parameters on HIC behaviors needs to be

further elucidated. The present work aims to investigate the hydrogen embrittlement in these highlights,

to conclude the guideline to prevent the car body from HIC, from material selection to the car

manufacture and functioning processes.

24

3 Materials and Experimental Methods

3.1 Investigated materials

In this study, series of advanced high strength steels have been investigated. As listed in Table 3.1, the

materials were classified according to their uniform tensile strength, into group 1180 MPa steels and

980 MPa steels. All the materials are industry-produced steel sheets for the applications in car bodies.

The materials were named by the suppliers according to the expected microstructures: Complex-Phase

(CP), Dual-Phase (DP), warm rolled Martensitic (MS-W) and Trip-aided Bainitic Ferrite (TBF).

Besides of MS-W1180 steel, other steel sheets supplied by various steel makers were cold rolled to

obtain thicknesses of 1.4 to 1.6 mm. After rolling, all the steel sheets were galvanized by hot-dip

galvanizing (GI) or commercial electro-galvanizing (EG).

Table 3.1: List of investigated steel grades.

Group Name Final rolling process

Galvanizing process

Thickness, in mm

1180 MPa Steels

CP1180+EG cold Electro-galvanized 1.6

DP1180A+EG cold Electro-galvanized 1.5

DP1180B+GI cold Hot-dip galvanized 1.4

MS-W1180+EG warm Electro-galvanized 1.5

TBF1180A+EG cold Electro-galvanized 1.5

TBF1180B+EG cold Electro-galvanized 1.5

TBF1180C+EG cold Electro-galvanized 1.6

980 MPa Steels

CP980+GI cold Hot-dip galvanized 1.6

DP980+EG cold Electro-galvanized 1.5

TBF980+EG cold Electro-galvanized 1.4

3.2 Material characterization

Chemical composition

The chemical compositions of the materials were measured by spark emission spectrometry with the

equipment manufactured by OBLF. The analysis of hydrogen concentration in as-delivered materials

and pre-charged samples was conducted by Thermal Desorption Spectroscopy (TDS), which is

introduced in details in section 3.3. Glow Discharge Optical Emission Spectrometry (GDOES) was also

applied to analyze the distribution of hydrogen, zinc and iron in the depth direction from the galvanized

sample.

3.2 Material characterization 25

Microstructure analysis

To characterize the microstructure of multi-phase steels, Light Optical Microscope (LOM) combined

with color-etching technology was widely applied [ANG06]. In the current work, the steel samples were

mechanically ground with conventional grinding paper and the final polishing step was conducted with

suspension of 1μm diameter silica particles. In DP and MS-W steels, etching was performed in 3%

Nital solution. In other steel grades, etching was carried out with Klemm solution. Due to different

solubility in the relevant etching agents, different phases are differed by varying colors under LOM

[ANG06].

The samples for Scanning Electron Microscope (SEM) were etched with 3% Nital solution. Under SEM

the C-enriched phases like austenite and martensite were stripped less than other phases.

Electron Back Scatter Diffraction (EBSD) based orientation microscopy was applied to quantitatively

analyze the phase fractions of the materials. The samples were electro-polished with commercial

electrolyte A2. The orientation microscopy was carried out on the RD-TD plane (RD: rolling direction,

TD: transverse direction) with a JEOL JSM 7000F SEM. EDAX-TSL camera “Hikari” was used as

EBSD system. The analysis was carried out with an acceleration voltage of 20 keV and a beam current

of 30 nA. The observed area was 100 μm×100 μm. The maps were measured with a step size of 100 nm,

in hexagonal grid. The measured data were collected and analyzed with the software OIM Data

Collection Analysis V6.2.

Because of differing crystallographic structures, the retained austenite was distinguished from the body

cantered cubic (bcc) phases. Martensite was differentiated from other bcc phases on the basis of Kikuchi

pattern relevant Image Quality (IQ) map. Crystalline defects such as dislocations have a distorted

Kikuchi pattern and cause a lower IQ value [WU05] [ZAE08]. The normalized IQ value (IQnormalized)

was used here to quantify the IQ of microstructure itself, preventing the possible influences from

operation processes. As described in equation (3-1) [WU05]:

(3-1)

where IQinitial stands for the absolute IQ value from the measuring, IQmin and IQmax represent the

minimum and maximum value of IQ respectively. Figure 3.1 shows a multi-peak model to analyze the

cold deformed microstructure [WU05]. The fraction of low IQ peak stands for the martensite phase and

the high IQ peak stands for the ferrite phase. In the normalized curve the grain boundary influence is

removed.

Both bainite and ferrite phases reveal low IQ values. During the bainite phase transformation process,

geometrically necessary dislocations forms, leading to a high misorientation gradient. Kernel Average

Misorientation (KAM) maps were analyzed to identify bainitic and ferritic constituents [ZAE08]. KAM

was calculated by determining the average of misorientation angles between all neighboring pixels of a

pixel [ZAE08]. In the current work, the third nearest neighbor kernel was adopted, so that the orientation


variation achieved 300 nm. A maximum value of KAM (KAMmax) of 0.6 was adopted as the boundary

value between ferrite and bainite. The microstructure with KAM value lower that 0.6º was defined as

ferrite while that with KAM value higher than 0.6º was identified as bainite and tempered martensite

structures.

Figure 3.1: The IQ distribution analysis using multi-peak model [WU05].

Quasi-static tensile test

The mechanical properties of the investigated materials were characterized by the quasi-static tensile

test according to SEP1240. The samples were manufactured according to ISO EN 10002, with a gauge

length of 80 mm. A universal extension machine Zwick Z100 was applied for quasi-static tensile test.

The strain rate for testing was set to a constant of 10-3 s-1. A laser extension sensor laserXtens was used

to measure the strain in the axial loading direction. The engineering stress-strain curves of each material

were recorded. The mechanical properties of each material were calculated as the average of three repeat

tests.

Hardness test

The local hardness on the edge surface and the cross section of welding point was measured. The

Vickers hardness test according to DIN EN ISO 6507-1 was carried out with a hardness machine

DuraScan from company Struers. A Vickers pyramid with surface angel 136° was applied. HV0.3 was

selected to characterize the edge surfaces and HV0.5 was selected for the analysis of welding point. The

hardness was analyzed by the software “ecos Workflow” automatically.

Roughness test

The roughness of the edge surface was measured using μsurf surface measurement. A confocal

microscope from the company NanoFocus was used for 3D measurement. The measured area was

320 μm × 320 μm on the sample. The data were analyzed by the software μsoft.

3.3 Hydrogen charging and analysis 27

3.3 Hydrogen charging and analysis

3.3.1 Galvano-static charging and electro-galvanizing

Galvano-static charging was used to modify the hydrogen concentration in the materials. The as-

delivered steel sheets were galvanized. The samples were firstly de-coated with an acid solution of

1.1g·l-1 HCl with 3.5 g·l-1 Hexamethylenetetramine for 2 min. The preparations of the sample differed

depending on the objective of the investigation. The samples for slow strain rate tests and forming tests

were ground to reach a surface roughness of Ra of 0.4 μm. The samples for U-bending were not ground

in order to maintain the original edge morphology.

Before charging, the samples were cleaned with an acid solution of 1.1 g·l-1 HCl with 3.5 g·l-1

Hexamethylenetetramine for 10 s, and then with distilled water and acetone to remove the impurities.

The procedures of hydrogen pre-charging and analysis are illustrated in Figure 3.2. The samples for

hydrogen measuring and mechanical testing underwent the same cleaning and hydrogen charging

processes. Additionally, the samples for mechanical tests were electro-galvanized immediately after

cathodic charging, and exposed in air for 24 h before the mechanical tests started.

Figure 3.2: Illustration of the procedure of hydrogen charging and following analysis.

A two-electrode charging cell was used for hydrogen charging. A solution of 3% NaCl with additional

0.3% NH4SCN was used. Ammonium-thiocyanate (NH4SCN) was added as poison agent to inhibit

hydrogen recombination and promote hydrogen entry rate effectively. The pH value of the solution is

close to neutral (pH 5.5-6), preventing the samples from corroding. The sample acted as the cathode

and a platinum-titanium net with a net factor 2.22 was used as the anode. The current density was set

to a constant value of 1 mA·cm-2. It means, for different sample geometries the required current density

was calculated based on its surface area.

The samples were pre-charged for varying lengths of time to obtain different grades of hydrogen

concentration. Table 3.2 summarizes the operation parameters of the samples used in different tests.

After pre-charging, the samples were cleaned with water to remove the residual solution on the surfaces.

Afterwards, they were exposed in air at room temperature for 5 min (minimum time for sample

preparation for the coming electro-galvanizing or measuring). The pre-charged samples were stored in

liquid nitrogen with a temperature of approximately -196°C to prevent from the dissolved hydrogen

escaping.


Table 3.2: Operated parameters of galvano-static charging for different kinds of samples.

Tests Solution Current density Duration Temperature

SSRT

3% NaCl +0.3% NH4SCN 1 mA/cm2

2.5/15/30 min (all the tests were carried out at room

temperature)

U-bend 10/30 min

Nakajima 30 min

VDA-bend 10/30 min

After pre-charging, the samples were electro-galvanized in order to effectively prevent hydrogen

effusion during following testing. The samples were activated with the HCl solution (1.1 g·l-1) with

3.5 g·l-1 Hexamethylenetetramine for 10 s and cleaned by distilled water and acetone. The commercial

acid zinc-chloride electrolyte Zylite ST HT (BC) (ZnCl2 concentration of 70 g·l-1) from Atotech was

applied for electro-galvanizing process. The samples functioned as the cathode and the anode was a

zinc bar. The parameters for galvanizing are listed in Table 3.3. The Zn-coated samples were exposed

in air for 24 h before the following mechanical testing, in order to homogenize the hydrogen in the

transverse direction.

Table 3.3: Operation parameters of electro-galvanizing for the different test samples.

Tests Solution Current density Duration Temperature

SSRT

Zylite ST HT (BC)

40 mA/cm2 10 min (all the tests

were carried out at room

temperature)

U-bend 40 mA/cm2 10 min

Nakajima 10 mA/cm2 30 min

VDA-bend 40 mA/cm2 10 min

3.3.2 Thermal desorption spectroscopy measurement

The hydrogen concentration was determined by Thermal Desorption Spectroscopy (TDS). The

apparatus of TDS is shown schematically in Figure 3.3. In the furnace, the samples were placed in a

glass pipe and the thermocouple was placed outside in another glass chamber in an equivalent location.

Before measuring the system was vacuumed so that the pressure inside the measuring system was in

the range of 10-8 to 10-2 Pa. The surrounding conditions of the sample and the thermocouple were quite

similar under high vacuum [ZEP03]. In the heating process the temperature of the radiation furnace was

measured by the thermocouple. The absorbed atomic hydrogen (Hab) desorbs from the substrate with

increased temperature and re-combination takes place at the sample surfaces. Thereby the amount of

desorbed molecule hydrogen was measured by quadruple spectrometers and recorded by computer. In

this study, the TDS apparatus was calibrated by standard TiH2. The heating rate was 0.1°C·s-1.

3.4 Mechanical tests to investigate HE behavior 29

Figure 3.3: Schematic apparatus of thermal desorption spectroscopy [ZEP03].

3.4 Mechanical tests to investigate HE behavior

3.4.1 Slow Strain Rate Test (SSRT)

Slow Strain Rate Test (SSRT) was carried out with the pre-charged and Zn-coated samples in the

atmosphere. In the SSRT a constant low strain rate was applied on the tensile samples until fracture

occurred. SSRT were carried out with a universal extension machine Zwick Z100 (as shown in Figure

3.4 (a)). The gauge length of the SSRT sample was 36 mm and the width of extension part was 5 mm,

as illustrated in Figure 3.4 (b). To eliminate the crack initiation related influences, the edges of the

sample were milled and after de-coating the surfaces of the sample were ground. The samples were pre-

charged with different amounts of hydrogen and afterwards electro-galvanized in lab (as described in

section 3.3.1). A constant strain rate of 10-5 s-1 was applied. A laser extension sensor laserXtens (Zwick

Company) was used to measure the strain in axial loading direction during the test. The engineering

stress-strain curves were automatically calculated and recorded. Each test repeated three times.

(a) (b)

Figure 3.4: (a) Universal extension machine Zwick Z100 with laser extension sensor; (b) geometry of slow strain rate test samples.

3.4.2 Incremental Loading Test (ILT)

The Incremental Loading Test (ILT) was carried out by the “Steinbeis-Transferzentrum Werkstoffe

Korrosion und Korrosionsschutz GmbH”. The samples were stretched in NaCl solution, under


incremental loading. The modified holed samples according to SEP1970 and samples with welding

points were tested. The sample geometries are shown in Figure 3.5: the holed samples were

manufactured with punched edges, while on the welded samples the welding points were firstly applied

by spot welding and afterwards the edges were machined to the test geometry.

The ILT was performed in a NaCl solution with a concentration of 3.5 % and start pH-value of 4

(adjusted by HCl). The maximum force until fracture (Fmax) of each material was examined in a quasi-

static tensile test. Based on the Fmax ILT was initiated at 75 % of Fmax and kept constant. After 48 h the

loading was increased by 5% of Fmax each 2 h until fracture or achieving 100 % of Fmax. The loading

curve is illustrated in Figure 3.6. The test duration and loading until occurrence of fracture were

recorded. Each test repeated five times.

(a)

(b)

Figure 3.5: Geometries of samples in the incremental loading tests: (a) holed samples with punched edges and (b) welded samples with machined edges.

Figure 3.6: Indication of loading in incremental loading test.

3.4.3 Constant Loading Test (CLT)

The Constant Loading Test (CLT) was carried out in surrounding of 3.5% NaCl solution with pH4

(adjusted by HCl). The samples were tested with the apparatus from the company Vetron. The sample

geometry for CLT is according to SEP1970, as illustrated in Figure 3.7. The samples were cut by


punching, and glued with the plastic container. For the testing the samples were fixed by bolts in the

loading rings (as shown in Figure 3.7), where the load was induced by applying a displacement by

adjusting the bolts and kept constant.

According to previous research, for the material with high yield ratio (Rp0.2/Rm) the notch effect on

sample is the lowest when the yield point (Rp0.2) was achieved. Therefore, for the material with yield

ratio over 0.9 a stress for testing of 80% of tensile stress (Rm) was selected [REH13]. The Rm and Rp0.2

of the material were determined by tensile tests with a strain rate of 10-3 s-1. In this work the loading

was kept a constant while the stress on the notch achieved Rp0.2 or 80 % Rm. The samples were subjected

to loading for 96 h. The time to fracture (tf) was recorded. In each material three samples were tested.

To determine the threshold stress of the SEP1970 samples in the NaCl surroundings, the applied stress

in the CLT was adjusted, from 90% of Rm and then lowered by 5% of Rm, until the fracture did not take

place within 96 h testing in the three repeat samples. The maximum allowed stress to fracture (σmax)

was recorded.

Figure 3.7: Apparatus and geometry of samples for constant loading test.

3.4.4 U-bending test

In order to produce the U-bend samples, sheet samples were bent by three-point-bending with a

universal testing machine. The sample geometry is shown in Figure 3.8 (a). The samples with three

kinds of edge conditions were tested: punched, milled and laser cut. In bending process the sample was

horizontally placed in the apparatus. The bending punch moved vertical to deform the samples, with a

moving rate of 2 mm·s-1. The sample was bent until the displacement of punch achieved 45 mm and the

two free legs were parallel to each other. Afterwards, the sample was fastened by plastic bolts, as

illustrated in Figure 3.8 (b).


(a) (b)

Figure 3.8: Illustration of U-bend sample: (a) sample geometry and (b) samples after bending.

The de-coated samples with punched edges were used to study the Internal Hydrogen Assisted Cracking

(IHAC). The samples were prepared with hydrogen pre-charging and electro-galvanizing as shown in

Figure 3.2. After a homogenizing in air for 24 h, the samples were bent with a punch radius (R) of

15 mm. The time to cracking initiation (tc) was recorded, and the cracking initiation and propagation

were observed. For each hydrogen concentration, at least three samples were tested.

In the U-bending test, the sheet samples with as-delivered Zn-coating were tested in the corrosive

medium concerning Hydrogen Environment Assisted Cracking (HEAC). Though the samples were Zn-

coated, the bare edges and the location where the protective layer is damaged display a potential

sensitivity to stress corrosion cracking [COR14]. The hydrogen pickup is promoted by the attack of

hydrogen from different corrosion mediums. The different parameters in each corrosion test are listed

in Table 3.4.

Table 3.4: Summary of the corrosion tests on U-bend samples.

Test Medium Concentration of NaCl pH Humidity Temperature Duration

Immersion test 100% solution 3.5% 4 n.c. RT 96 hours

Salt spray test 100% spraying 5% 6.5-7.2 n.c. 37-38°C 10 days

VDA corrosion test

Cyclically changed surroundings 1% n.c. in cycle in cycle 6 weeks

*n.c.: not controlled.

In the immersion test the U-bend samples were immersed in a 3.5% NaCl solution (pH4). The volume

of the solution met the requirement of 10 ml·cm-2 of the sample, according to DIN 50905-1. The solution

was adjusted with HCl solution twice per day, to control the pH-value. The samples were immersed in

the solution for 96 h or until failure. During the immersion, crack initiation and propagation on the

samples were observed and recorded.

The salt spray test was carried out according to standard DIN EN ISO 9227. The pre-bent samples were

placed in a salt spray chamber, where the NaCl solution with 5% concentration was homogenously,


continuously sprayed. The temperature in the chamber was controlled at 37-38°C. The salt spray test

took 10 days and during the test the samples were observed every 24 h.

In addition, the new VDA-corrosion test according to VDA233-102 was performed. In the testing

chamber salt spraying with 1% NaCl took place three times per week. The temperature and humidity

were varied during a cycle of 168 h (1 week), as illustrated in Figure 3.9. The duration of VDA-

corrosion was 6 weeks and chamber was opened every week to observe the samples.

Figure 3.9: Indication of charged parameters in VDA corrosion test for 1 week.

3.4.5 Four-point-bending test

In four-point-bending, the sheet samples in the as-delivered state were tested in salt solution. The

samples had a dimension of 115 mm × 40 mm. The bare Zn-Fe-Zn edges were cut by punching. The

apparatus for four-point-bending, as illustrated in Figure 3.10, were used. The samples were

horizontally placed in the apparatus, between the upper and lower points. The position of lower points

was adjusted by a bolt. The position, once the sample connected to all four points but not yet deformed

was set as the initiated position. Adjusted by bolt, the lower points went up and deformed the sample.

The displacement between initiated position and deformed position, as illustrated as y in Figure 3.10,

was measured by a dial indicator. The value of bending grade was equal the displacement y. The

samples were bent by four-point-bending to obtain different bending grades. Afterwards the apparatuses

with samples were immersed in 3.5% NaCl solution with pH4. The samples were observed up to 96 h

and the crack initiation and propagation were observed and recorded during immersion.


Figure 3.10: Photo and illustration of apparatus for four-point-bending.

3.5 Forming tests

3.5.1 Determination of the Forming Limit Curve (FLC)

Forming Limit Curve (FLC) is determined by Nakajima tests. Figure 3.11 illustrates the relationship

of major strain (φ1) and minor strain (φ2) in each strain state. The sample geometries in this work are

shown in Figure 3.12. The samples for the Nakajima tests were manufactured by machining. The

sample surfaces were ground to a roughness Ra of 0.4 μm. The samples were prepared by hydrogen pre-

charging and electro-galvanizing as described in section 3.3.1. Uncharged samples were firstly cleaned

and then directly electro-galvanized such as the pre-charged samples, in order to obtain same surface

condition to ensure the same friction.

Figure 3.11: Illustration of forming limit curve and the strain states in Nakajima test [BLE09].

3.5 Forming tests 35

Figure 3.12: Sample geometries of Nakajima tests.

Before the Nakajima tests, a characteristic stochastic black-and-white pattern was applied to the sample

surfaces by spraying. The punch velocity was 2 mm·s-1. During forming the patterns on the surface were

recorded by a GOM ARAMIS system in a frequency of 10 pictures per second. For each test group five

tests were repeated.

The recorded pictures were analyzed to detect the beginning of necking and fracture. The detection of

localized necking was performed by time based method “linear best fit” as mentioned by W. Volk and

P. Hora [VOL11], since it shows better performance and stability than the method according to

ISO12004-2 [FRI14]. Five elements in the necking areas were selected. The representative thinning rate

of last 20 pictures before the first occurrence of cracking was calculated as the average value from the

selected elements. As illustrated in Figure 3.13, the representative thinning rate initially increases

linearly with a low slope, being considered as stable deformation. After a transition period (curved area)

it increases linearly with a higher slope, representing instable deformation. The linear fitting curves

were calculated for both stable and instable areas. The intersection of the two linear fitting curves

represents the beginning of instable necking during the forming. According to the last picture number

in stable area, the major- and minor strains before the occurrence of necking were calculated as the

average values from selected elements. The forming limit in each strain state was determined by the

average value of critical major- and minor strains from five repeat tests. As well, the critical strain states

to fracture were calculated by using the last pictures before the occurrence of fracture.


Figure 3.13: Detection of localized necking in representative thinning rate diagram, uni-axial sample.

3.5.2 Determination of the Bending Limit Curve (BLC)

The Bending Limit Curve (BLC) is a strain-based criterion to predict failure under bending conditions.

To determine the BLC of a material, a drawing operation for pre-straining and a superposed bending

process are required [DEN12] [SHL09]. Figure 3.14 schematically describes how the BLC is derived.

A Marciniak test with a punch diameter of 340 mm was applied to introduce the pre-strain in the steel

sheets. The blank holder force was set as 400 kN and the punch velocity was 20 mm·s-1. The Marciniak

samples were drawn in uniaxial, plane strain and biaxial strain states, until the critical drawing depth

was reached. Afterwards, the bending samples with a dimension of 60 mm × 60 mm were cut from the

Marciniak samples, and prepared by pre-charging and electro-galvanizing.

Superposed bending loads were applied by the three-point-bending test according to VDA 238-100. A

universal testing machine with a bending device with low-friction mounted rollers was used, as

schematically shown in Figure 3.15. The distance between the two mounted rollers (L) depended on

the thickness of the samples (t), where . The bending punch had a bending-sword

radius of 0.4 mm. The edges of the sample were ground to remove any edge effect on crack initiation.

The sample was placed on the rollers and was pre-loaded by a force of 30 N with velocity of

10 mm·min -1. Afterwards the sample was bent with a punch velocity of 20 mm·min-1. A force drop of

30 N was defined as the failure criterion in the bending procedure. Five tests were repeated.

3.6 Investigation of the car component-similar samples 37

Figure 3.14: Schematic development of

BLC [DEN12]. Figure 3.15: Schematic illustration of bending device for

VDA238-100 bending test.

In addition, the strain distributions on the surfaces of Marciniak sample and bending sample were

measured by optical metrology: the stochastic black-and-white patterns before and after deformation

were recorded and analyzed by GOM ARAMIS system, thus the values of major and minor strains were

calculated.

3.6 Investigation of the car component-similar samples

90° V-bend samples

In the body-in-white structures, AHSS is usually applied in the tunnel components, where the sheet

samples are bent to an angle of 90°. To investigate if the susceptibility to HE is dependent on the

according radius, sheet samples with a dimension of 60 mm × 70 mm were bent with different punches

(bending radii 1 mm, 2 mm, 3 mm, 4 mm and 5mm), with a punch velocity of 2 mm·s-1. For each

material and punch radius, a pre-test and correction were carried out to ensure an angle of 90° between

the two legs of the sample. Because of the spring-back after bending, on the inner side of the bending

radius a localized tensile stress is present. The samples were tested in a cyclic immersion test, in which

they were alternatively immersed in solution (3.5%NaCl with pH4) for 10 min and exposed in air for

50 min. After 500 h cyclic immersion test, the samples were cleaned with mixed solution of HCl

(1.1 g·l -1) and Hexamethylenetetramine (3.5 g·l-1) to remove the corrosion product and Zn-coating.

Samples were taken from the bending edge, prepared and observed with an optical microscope.

Cross-die samples

The forming of the cross-die sample represents a classic deep drawing process. On the basis of

DIN 8584-3, the cross-die test involves a tensile-compress forming without changing the sheet


thickness [DOE10]. The stress distribution simulated in the cross-die model is similar to the stresses

generated by the drawing process of complex car body component.

As seen in Figure 3.16, the cross-die samples were produced with a blanking die and convex-concave

punch. The dimension of the sheet sample was 310 mm × 250 mm, with corners cut at an angle of 45°.

Deep drawing of the cross-die samples was carried out by a hydraulic press machine from the company

Müller-Weingarten. The lubricant OEST PLATINOL B804 was used to reduce the friction between the

sample and machine. The blank holder force was 400 kN, and the punch velocity was 21 mm·s-1.

Selected materials were investigated with the cross-die test. Depending on the drawability of each

material, different drawing depths were selected, as listed in Table 3.5.

Figure 3.16: Illustration of blank holder and sample geometry for cross-die test.

Table 3.5: Drawing depth of selected materials in the cross-die test.

Material Max. depth [mm] Applied depth [mm]

DP1180A+EG 17 17

TBF1180A+EG 21 21, 16

TBF1180B+EG 21 21, 16

TBF1180C+EG 23 23, 18

DP980+EG 20 20, 15

After deep drawing, the cross-die samples were prepared to obtain a similar state as in the real car

component. Figure 3.17 shows a cross-die sample with original punched edges. As illustrated in Figure

3.18, the edges of the sample were cut by laser cutting. A lid with laser cut edges was produced by

drawing and spot welded on the cross-die samples. The cross-die samples, directly drawn as well as

prepared by cutting and welding, were immersed in a salt solution of 3.5% NaCl with additional HCl

to obtain pH4 for 192 h. The crack initiation and the state after the immersion test were observed.

3.6 Investigation of the car component-similar samples 39

Figure 3.17: Illustration of the cross-die sample after deep drawing.

Figure 3.18: Illustration of the cross-die sample after cutting and welding.

40

4 Characterization of As-delivered Materials

4.1 Chemical composition

The chemical compositions of the investigated materials are listed in Table 4.1.

Table 4.1: Chemical compositions of the investigated materials, in wt.%.

Group Material C Si Mn Al P S Cr

1180 MPa steels

CP1180 0.150 0.033 2.013 0.042 0.011 0.001 0.237 DP1180A 0.155 0.213 1.900 0.024 0.024 0.001 0.216 DP1180B 0.123 0.255 2.570 0.024 0.010 0.001 0.600

TBF1180A 0.198 0.852 2.616 0.037 0.009 0.001 0.327 TBF1180B 0.200 1.511 2.590 0.047 0.012 0.001 0.320 TBF1180C 0.178 1.549 2.113 0.031 0.016 0.004 0.350 MS-W1180 0.105 0.080 1.404 0.031 0.010 0.001 0.233

Ni Mo V Ti Nb N Cu

1180 MPa steels

CP1180 0.013 0.004 0.006 0.026 0.004 0.002 0.006 DP1180A 0.012 0.002 0.005 0.023 0.004 n.d. 0.011 DP1180B 0.008 0.009 0.007 0.025 0.025 0.004 0.001

TBF1180A 0.001 0.004 0.006 0.004 0.026 n.d. 0.013 TBF1180B 0.011 0.080 0.005 0.005 0.004 0.006 0.010 TBF1180C 0.016 0.002 0.005 0.005 0.004 0.007 0.005 MS-W1180 0.022 0.003 0.004 0.041 0.003 n.d. 0.014

Group Material C Si Mn Al P S Cr

980 MPa steels

CP980 0.086 0.253 2.251 0.037 0.011 0.001 0.231 DP980 0.170 0.183 2.190 0.047 0.013 0.002 0.450

TBF980 0.190 0.881 2.356 0.038 0.011 0.001 0.275

Ni Mo V Ti Nb N Cu

980 MPa steels

CP980 0.013 0.005 0.006 0.030 0.005 0.003 0.008 DP980 0.040 0.009 0.005 0.002 0.022 0.015 0.013

TBF980 0.011 0.003 0.005 0.004 0.026 n.d. 0.009 *n.d.: not determined.

4.2 Microstructure characterization

The resulting micrographs from Light Optical Microscope (LOM) investigation are shown in Figure

4.1, in RD-TD plane (RD: rolling direction, TD: transverse direction). In the materials DP1180A,

DP1180B and MS-W1180 etched with 3% Nital solution, the martensitic structure is dark brown and

ferrite is white or light brown. In the other materials etched with Klemm solution, retained austenite

and martensite are white, while ferrite and bainitic structure are brown or blue [ANG06]. It is found

that in all the investigated materials, the grains are very fine and the grain boundaries cannot be

4.2 Microstructure characterization 41

determined. The microstructure in the materials is homogeneously distributed, displaying no

segregation. The determination of specific phase fraction of the materials via LOM is impossible.

(a) CP1180 (Klemm) (b) DP1180A (Nital)

(c) DP1180B (Nital) (d) TBF1180A (Klemm)

(e) TBF1180B (Klemm) (f) TBF1180C (Klemm)


(g) MS-W1180 (Nital) (h) CP980 (Klemm)

(i) DP980 (Klemm) (j) TBF980 (Klemm)

Figure 4.1: Light optical micrographs of the investigated materials, etched with Nital or Klemm.

Further microstructure analysis was conducted by Scanning Electron Microscope (SEM) and Electron

Back Scatter Diffraction (EBSD), as described in Section 3.2. In the CP1180, as shown in Figure 4.2,

the microstructure is fine and homogeneous. In the SEM image the characteristic sub-boundaries and

carbide distribution is observed. In the IQ-map the overall microstructure displays as the light region,

implying that quenched martensite is not present. Based on the analysis of IQ-map and KAM, it reveals

that CP1180 consists of mixed phase of bainite and tempered martensite.

(a) SEM image, etched with Nital (b) Details from (a)


(c) IQ-map (d) Phase-map

Figure 4.2: Micrographs of CP1180 in RD-TD plane, characterized by SEM and EBSD. (a) and (b): SEM image, etched with Nital, (c): IQ-map and (d): Phase-map.

In DP1180A as shown in Figure 4.3, ferrite is observed via SEM, featuring a uniform contrast inside

the grains. After tempering, the martensitic structure is differentiated with the typical, well-delineated

substructure and particular carbide. The IQ-map does not reveal any dark blocks with high dislocation

density, indicating that in DP1180A quenched martensite does not exist.



Figure 4.3: Micrographs of DP1180A in RD-TD plane, characterized by SEM and EBSD. (a) and (b): SEM image, etched with Nital, (c): IQ-map and (d): Phase-map.


Figure 4.4 shows the micrographs of TBF1180A. Under SEM both retained austenite and martensite

islands are smooth and featureless. Tempered martensite and bainite mix together. In the IQ-map after

EBSD analysis, the martensite appears dark because of its high IQ-value. According to the phase map,

retained austenite (red) and small amount of martensite (blue) are present in the material. Besides, small

fraction of ferrite (green) is determined, because both IQ and KAM have low values.



Figure 4.4: Micrographs of TBF1180A in RD-TD plane, characterized by SEM and EBSD. (a) and (b): SEM image, etched with Nital, (c): IQ-map and (d): Phase-map.


Figure 4.5 displays the characteristic microstructure of MS-W1180. The Nital-etched MS-W1180

reveals typical lath martensite structure under SEM. The IQ-map of MS-W1180 shows that the material

consists only of martensite.



Figure 4.5: Micrographs of MS-W1180 in RD-TD plane, characterized by SEM and EBSD. (a) and (b): SEM image, etched with Nital, (c): IQ-map and (d): Phase-map.

Furthermore, Figure 4.6 shows the micrographs of other steels with Rm of 1180 MPa. In the phase

analysis in EBSD the same threshold values were used to distinguish the phases in IQ-map and KAM-

map.

(a) (b)


(c) (d)

(e) (f)

Figure 4.6: Micrographs of other investigated 1180 MPa steels in RD-TD plane by SEM (etched with Nital) and EBSD. (a) (b): DP1180B, (c) (d) TBF1180B, and (e) (f) TBF1180C.

The volume fractions of each phase and grain size of the materials are listed in Table 4.2. It is worth

noting that the phases of bainite and tempered martensite are the main phases in all the investigated

materials except MS-W1180 (consists of 100% martensite). The retained austenite in CP-, DP- and MS-

W steels is in a very small fraction (below 1%). In contrast, the volume fractions of retained austenite

in all TBF steels stay in the range of 8.6% to 9.2%. The grain size was measured using the linear

intercepts according to ASTM E 112-12, with the EBSD phase-maps of the materials. In EBSD the

grain boundary was analyzed according to the misorientation. Once the misorientation is over 15°, the

structure was considered as grain boundary. Therefore, partial lath-boundaries are likely to be calculated

as the grain boundaries, resulting in an artificially fine calculated grain size. It should be noticed that

measuring of grain size is limited by the measured area by EBSD, which is 100 μm × 100 μm.

4.3 Mechanical properties 47

Table 4.2: Phase fractions and grain sizes of the investigated materials, in vol.%.

Group Material Residual austenite Ferrite Martensite

Bainite + Tempered martensite

Grain size [μm]

1180 MPa steels

CP1180 0.2 - - 99.8 1.2

DP1180A 0.4 16.4 - 83.2 1.4

DP1180B 0.3 12.1 13.8 73.8 0.8

TBF1180A 9.0 5.9 1.2 83.9 0.8

TBF1180B 8.6 15.0 5.3 71.1 1.1

TBF1180C 9.2 10.2 0.7 79.9 0.9

MS-W1180 - - 100 - 0.9

980 MPa steels

CP980 0.7 - 0.6 98.7 1.4

DP980 0.7 34.7 2.1 62.5 1.4

TBF980 8.9 18.8 5.2 67.1 1.2

4.3 Mechanical properties

The mechanical properties of the investigated materials were characterized by quasi-static tensile tests.

Table 4.3 summarizes the yield strength (Rp0.2), tensile strength (Rm), uniform elongation (Au) and total

elongation (Af). The yield ratios of the materials were calculated from Rp0.2/Rm.

It is revealed that MS-W1180+EG displays the highest strength but the lowest elongation. The TBF

steels reveal better elongation values (~8% Au and over 10% Af in 1180 MPa) than the other steel grades.

CP steel displays higher yield strengths but worse elongation properties than DP steels. In Figure 4.7

the distribution of strength and elongation illustrates the strength-ductility balance of the materials. The

TBF steels show better strength-ductility balance than the other steel grades.

Table 4.3: Mechanical properties in rolling direction of investigated materials after quasi-static tensile tests at strain rate of 10-3·s-1.

Group Material Rp0.2 [MPa]

Rm [MPa]

Au [%]

Af [%] Rp0.2/Rm

1180 MPa steels

CP1180+EG 1045 1194 3.3 6.2 0.88 DP1180A+EG 1024 1224 4.2 7.6 0.84 DP1180B+GI 925 1196 4.7 8.3 0.77

TBF1180A+EG 1035 1222 7.9 10.9 0.85 TBF1180B+EG 899 1194 9.2 12.9 0.75 TBF1180C+EG 1048 1229 8.1 12.1 0.85 MS-W1180+EG 1163 1263 2.4 4.8 0.92

980 MPa steel

CP980+GI 1054 1150 4.5 7.9 0.92 DP980+EG 753 991 7.1 11.4 0.76

TBF980+EG 742 1017 10.5 15.1 0.73


Figure 4.7: Distribution of tensile strength and total elongation of the investigated steels.

49

5 Characterization of Pre-charged Materials

5.1 TDS analysis of pre-charged samples

In order to investigate the hydrogen uptake and concentration during hydrogen charging, the SSRT

samples were pre-charged and measured by Thermal Desorption Spectroscopy (TDS). The hydrogen

desorption curves from TDS measuring provide information of diffusible and deeply trapped hydrogen.

Figure 5.1 shows the hydrogen desorption curves from four kinds of steels, before and after 30 min

hydrogen charging. The surface areas under the peaks represent the concentrations of desorbed

hydrogen. As illustrated, in the as-delivered state, the desorbed hydrogen at temperatures below 350°C

has a very small amount. Only 0.03 ppm hydrogen is detected in the steel of DP1180A. The desorbed

hydrogen at this low temperature is considered as diffusive hydrogen, standing for the hydrogen in the

reversible trapping sites. The second desorption peak on desorption curve is observed in temperatures

over 350°C. Hydrogen desorbs at this temperature accounts for carbides, oxides and the interfaces

between these phases and matrix. It is defined as metallurgical hydrogen or deeply trapped hydrogen.

After 30 min pre-charging, the first peaks in the curves rise visibly, occurring at temperatures in the

range of 154-198°C. The concentrations of diffusive hydrogen are obviously higher than that in the

initial condition. However, the second peaks rise slightly, and the maximum rise of 0.08 ppm occurs in

TBF and MS-W steels. It is deduced that when the investigated steels are charged with hydrogen in the

solution of NaCl with additional NH4SCN, the hydrogen absorbs mainly in the reversible trapping sites

which have a low binding energy. Compared to diffusive hydrogen, only a small amount of hydrogen

atom is absorbed in irreversible trapping sites. The trappability of the irreversible trapping sites in the

investigated steels is weak. Consequently, in the following study it is therefore assumed that the

diffusive hydrogen model is the most likely explanation for hydrogen induced cracking behaviors.

The dependence of hydrogen content on charging time is revealed by the calibration curves, as shown

in Figure 5.2. The samples were pre-charged for 2.5 min, 15 min and 30 min and measured up to 350°C

by TDS measurement. The content of diffusive hydrogen increases with increasing charging time. All

the materials in 1180 MPa and 980 MPa grades display same tendency in the calibration curves.


(a-1) CP1180, as-delivered (a-2) CP1180, pre-charged for 30 min

(b-1) DP1180A, as-delivered (b-2) DP1180A, pre-charged for 30 min

(c-1) TBF1180A, as-delivered (c-2) TBF1180A, pre-charged for 30 min

(d-1) MS-W1180, as-delivered (d-2) MS-W1180, pre-charged for 30 min

Figure 5.1: Hydrogen desorption curves in TDS measurement in different steels, in as-delivered and 30 min pre-charged samples, at a heating rate of 0.1°C·s-1.

5.2 GDOES analysis of electro-galvanized samples 51

(a) CP1180 and MS-W1180 (b) DP1180A and DP1180B

(c) TBF1180A, TBF1180B and TBF1180C (d) CP980, DP980 and TBF980

Figure 5.2: Calibration curves of content of diffusive hydrogen after different charging time, measured by TDS measurement, at a heating rate of 0.1°C·s-1.

5.2 GDOES analysis of electro-galvanized samples

In this work, the pre-charged samples were electro-galvanized to prevent hydrogen effusion during the

following mechanical test. However, hydrogen uptake can occur during the electro-galvanizing process,

as mentioned in section 2.2. A Glow Discharge Optical Emission Spectrometry (GDOES) was used to

analyze the hydrogen distribution through the depth of the steel sample. Figure 5.3 illustrates the

distribution of hydrogen (red), zinc (blue) and substrate steel (black) from the sample surface to a depth

of 30 μm after electro-galvanizing. Between the sample preparation and GDOES measurement the

sample was stored in liquid nitrogen. It can be seen in Figure 5.3 that after galvanizing, the thickness

of the Zn-layer displays as a function of the depth. An interface between Zn-layer and the substrate is

observed due to the sputtering effect during the deposition of zinc ions. In this interface the hydrogen

concentration is high and stays at a low concentration in both substrate and Zn-layer. It is assumed that

once the electrochemical reaction starts, the Zn deposition is associated with hydrogen uptake. Zinc can

be a barrier to protect the substrate from further hydrogen degradation [PAA11]. Once a certain Zn-

coating forms, the hydrogen uptake subsided.


The reference sample was electro-galvanized and stored in air for 24 h. When measured, it shows the

same distribution curves as illustrated in Figure 5.3. It provides evidence that the hydrogen trapped in

the interface is not diffusive at room temperature, since the hydrogen distribution in Zn, interface of Zn

and substrate and pure substrate displays no significant difference in both the liquid-nitrogen-stored

sample and the air-exposed sample. During the component production process, however, hydrogen

uptake may still occur through the bare edges which are not protected by the Zn-layer.

Figure 5.3: Distribution of hydrogen in the depth in DP1180A sample, pre-charged and electro-galvanized,

stored in liquid nitrogen before GDOES.

53

6 Effects of Hydrogen on Mechanical Properties

6.1 Results from Slow Strain Rate Test (SSRT)

Degradation of property

In the SSRT, the materials with different contents of hydrogen were strained in room atmosphere, at a

rate of 10-5 s-1 until failure. Figure 6.1 displays the selected engineering stress-strain curves on the

uncharged and pre-charged samples. The enriched hydrogen in the steels leads to earlier failure in the

tensile process. The reduction of strains depends on the hydrogen concentration. In CP1180 and MS-

W1180 steels, the degradations of mechanical properties are moderate compared to those in DP and

TBF steels with 1180 MPa tensile strength. The hydrogen-induced degradation of TBF1180 steels is

more obvious compared to other steel grades. Concerning the 980 MPa steels, in CP980 and DP980 the

influence of hydrogen on the stress-strain curves is not significant. In contrast, TBF980 degrades more

obviously than CP- and DP980 steels once the samples are pre-charged with hydrogen.

(a) CP1180

(b) DP1180A

(c) DP1180B

(d) TBF1180A


(e) TBF1180B

(f) TBF1180C

(g) MS-W1180

(h) CP980

(i) DP980

(j) TBF980

Figure 6.1: Engineering stress-strain curves of uncharged and pre-charged materials, at strain rate of 10-5 s-1.

To evaluate the susceptibility to HE, the percentage loss of mechanical property after hydrogen charging

was calculated from the equation (6-1):

(6-1)

where M represents the studied mechanical parameter (Rp0.2, Rm, Au and Af) of uncharged material and

MH indicates the same parameter of the charged material. A higher value of loss of property indicates

that the material is more susceptive to HE. A negative value of loss of property signifies that the property

is intensified due to increasing H contents. According to the hydrogen calibration curves as shown in

Figure 5.2, the degradations of strength of each material with different hydrogen concentrations are t,

6.1 Results from Slow Strain Rate Test (SSRT) 55

as shown in Figure 6.2. It can be seen that except for TBF steels, the changes in Rp0.2 and Rm of steel

are less than ±3%. In the TBF1180 and TBF980 steels the loss of Rp0.2 is not significant, while the Rm

degrades. The maximum degradation of Rm is found in TBF1180B, with a value closed to 10%.

Furthermore, the losses of elongation in each material are plot against the hydrogen concentration, as

shown in Figure 6.3. The degradation of elongation varies for different materials. In general, Af reveals

a stronger degradation and is more susceptive to HE than Au. In the 30 min pre-charged CP1180 samples,

the loss of Af is below 40%. In MS-W1180, the elongation is reduced by less than 30% because of lower

hydrogen concentration after 30 min pre-charging (according to Figure 5.2). The slopes of loss of

elongation in CP1180 and MS-W1180 are similar. The DP1180A displays a loss of Af of 58% with

1.45 ppm hydrogen comparing to 44% loss of Af in DP1180B when the hydrogen content reaches

1.21 ppm. While the TBF steels display a reduction of Af over 66%. TBF1180A, however, shows

highest value in loss of elongation, reaching 86% when the hydrogen content reaches 1.38 ppm.

The materials CP980 and DP980 show similar gradient of changes in elongation. However, even though

the Rm is also 980 MPa, after 2.5 min pre-charging the 0.53 ppm hydrogen causes a loss of Au of 30%

and Af over 40% in TBF980. Once the hydrogen content reaches 1.51 ppm, the loss of Au in the TBF980

as well as Af rise to 60%.


(a-1) (a-2)

(b-1) (b-2)

(c-1) (c-2)

(d-1) (d-2)

Figure 6.2: The loss of strength with increasing hydrogen content after SSRT, at strain rate of 10-5 s-1.


(a-1) (a-2)

(b-1) (b-2)

(c-1) (c-2)

(d-1) (d-2)

Figure 6.3: The loss of elongation with increasing hydrogen content after SSRT, at a strain rate of 10-5 s-1.


Fractography of SSRT samples

Fractography is vital in understanding the fracture mode of uncharged and pre-charged samples during

the SSRT. Compared to the uncharged sample, fracture surfaces of pre-charged samples displays two

different zones. The area located at the edges is named Surface Zone (SZ) and the area inside is called

Inside Zone (IZ), as shown in Figure 6.4. The microscopic fracture surfaces of the fractured samples

are shown in Figure 6.5 to Figure 6.8.

(a) uncharged (b) pre-charged

Figure 6.4: Macroscopic fracture surfaces of TBF1180A after SSRT: (a) uncharged and (b) 30 min pre-charged.

The fracture surfaces from CP1180 uncharged sample show overall ductile fracture mode with fine

dimples with the size of 1-2 μm or below, as displayed in Figure 6.5. After 30 min pre-charging, the

SZ on the fracture surface displays quasi-cleavage fracture mode with a small amount of dimples. The

IZ fracture surface still exhibits ductile fracture behavior. The size of the dimples in the IZ of pre-

charged samples is not larger than that in the uncharged sample.

Figure 6.6 shows the fracture surfaces from DP1180A after SSRT. The dimples on the fracture surfaces

from uncharged samples have an average size of 4-5 μm. The fracture surfaces in the SZ of pre-charged

DP1180A exhibit brittle fracture in transgranular mode with small amount of brittle facets. The IZ

fracture surfaces displays ductile fracture similar to the uncharged ones.

In the uncharged sample and the IZ of the pre-charged sample from TBF1180A, fine dimples with 1 to

2 μm diameters are observed (Figure 6.7 (a)). In the SZ of samples after pre-charging, the fracture

surfaces displays intergranular fracture with substantial amounts of brittle facets, as shown in Figure

6.7 (b). The IZ of fracture surface displays ductile fracture featured with dimples (Figure 6.7 (c)).

Figure 6.8 the fracture surfaces of the uncharged MS-W1180 reveal a great amount of substantial brittle

facets, resulting from an overload of plastic strain and brittle martensite phase. However, after pre-

charging, the SZ of fracture surfaces reveals quasi-cleavage fracture mode. The tearing under load

indicates that the hydrogen embrittled the material.

In general, except for MS-W1180, other uncharged materials show ductile fracture on the fracture

surfaces after SSRT. After H-charging, fracture surfaces in the SZ display quasi-cleavage fracture mode,


however, dimples can be found in the IZ, indicating that the hydrogen-induced cracking initiates from

the edge location. In the materials CP1180 and MS-W1180, transgranular fracture is the major fracture

mode when the crack initiates. In DP1180A, the brittle facets are found in the SZ, mixed with the

transgranular fracture, which is assumed as the accumulation of hydrogen at the grain boundaries. In

the case of TBF1180A, intergranular fracture is dominant, where the ESSRT displays the lowest value in

these four materials.

Generally, in the SSRT the strength of the materials degrades moderately compared to the degradation

of elongation, which is accordant to the results from other researchers [GUO12]. The Embrittlement

Index (EISSRT) is defined by using the loss of Af, as shown in equation (6-2):

EISSRT= (6-2)

The material with a higher strength displays a lower EISSRT, indicating a higher susceptibility to HE. It

should be noted that the gradient of change in EISSRT is not constant, but varies when the hydrogen

content increases. Table 6.1 summarizes the EISSRT of different materials when the hydrogen content is

in the range of 1.2–1.5 ppm. On the basis of EISSRT the HE-Susceptibilities of the investigated materials

are ranked in the following order:

For 1180 MPa steels:

TBF1180A > TBF1180B > TBF1180C > DP1180A > DP1180B > CP1180 > MS-W1180

And for 980 MPa steels:

TBF980 > DP980 > CP980

Table 6.1: EISSRT of investigated materials in the SSRT, when the hydrogen content is in range of 1.2-1.5 ppm.

Group Material H content [ppm] EISSRT [%]

1180 MPa steels

CP1180 1.30 68.9

DP1180A 1.45 41.9

DP1180B 1.21 56.4

TBF1180A 1.38 14.0

TBF1180B 1.45 26.1

TBF1180C 1.0* 33.7

MS-W1180 1.25 75.1

980 MPa steels

CP980 1.29 78.7

DP980 1.38 70.6

TBF980 1.51 35.8 * In TBF1180C, the hydrogen content can reach maximum 1 ppm, after 30 min pre-charging.


(a) uncharged (b) pre-charged, SZ

(c) pre-charged, IZ

Figure 6.5: Micrographs of fracture surface from CP1180 after SSRT: (a) uncharged, (b) SZ of 30 min pre-charged and (c) IZ of 30 min pre-charged.


(c) pre-charged, IZ

Figure 6.6: Micrographs of fracture surface from DP1180A after SSRT: (a) uncharged, (b) SZ of 30 min pre-charged and (c) IZ of 30 min pre-charged.



(c) pre-charged, IZ

Figure 6.7: Micrographs of fracture surface from TBF1180A after SSRT: (a) uncharged, (b) SZ of 30 min pre-charged and (c) IZ of 30 min pre-charged.


(c) pre-charged, IZ

Figure 6.8: Micrographs of fracture surface from MS-W1180 after SSRT: (a) uncharged, (b) SZ of 30 min pre-charged and (c) IZ of 30 min pre-charged.


6.2 Results of Incremental Loading Test (ILT)

6.2.1 Hydrogen-induced cracking on holed samples

Figure 6.9 shows the results of the Incremental Loading Test (ILT) on the holed samples. Each marker

in the diagram indicates when the sample breaks. In the ILT DP1180A+EG and TBF1180A+EG break

within a very short test time when the applied stress is 75% of Rm. The samples from CP1180+EG are

broken also in the constant period in the beginning 48 h. In contrast, the samples from MS-W1180+EG

stay generally a longer testing time until fracture.

Figure 6.9: Results of incremental loading test from CP1180+EG, DP1180A+EG, TBF1180A+EG and

MS-W1180+EG (punched edges, in NaCl solution).

To evaluate the susceptibility to HE, the applied stress (σ) and time to delayed fracture (tf) in different

materials are summarized in Figure 6.10, where σ was re-calculated as percentage of Rm and Rp0.2, and

tf was calculated from five repeated tests.

Delayed Fracture (DF) occurs in the investigated materials, except for MS-W1180+EG and CP980+GI,

when σ staying in 75% of Rm within 48 h. CP1180+EG displays a longer tf than DP- and TBF steels,

with the average value of 25.2 h. DP- and TBF1180 steels reveal short tf, in the range of 0.5 to 2.6 h.

Considering the yield behavior of the materials, the σ of 75% of Rm stays in around 90% of Rp0.2 in

DP1180A+EG, but 96% of Rp0.2 in DP1180B+GI. Thus DP1180A+EG reveals higher susceptibility to

HE than that of DP1180B+GI because of its shorter tf. TBF1180A+EG displays a longer tf than that of

TBF1180B+EG, because the lower hydrogen concentration in the as-delivered material. At the same

time, the start loading in TBF1180A+EG reaches 89% of Rp0.2 while in TBF1180B+EG reaches already

100% of Rp0.2. Compared to the other two TBF1180 steels, TBF1180C+EG has a longer tf. This is

assumed to be due to its lower hydrogen adsorption ability, which is revealed from the hydrogen

calibration curve in Figure 5.2. Moreover, it can be found that DP980+EG is more susceptive than

CP980+GI, because of the hydrogen content of as-delivered DP980+EG is 0.18 ppm, being much higher

that than in TBF980+EG (0.01 ppm). Thus the threshold value for crack initiation is easier to be

achieved.

6.2 Results of Incremental Loading Test (ILT) 63

Figure 6.10: Applied stresses and time to delayed fracture of different materials in incremental loading test

(punched edges, in NaCl solution).

The fracture surface macrographs of DP1180A+EG sample are illustrated in Figure 6.11. After the ILT

in the salt solution, the fracture surface in the triangular area close to the edge surface of the hole

(Surface Zone (SZ)) is perpendicular to the sample surface. In the Inside Zone (IZ) the fracture surface

is at an angle of 45° from the sample surface. It can be assumed that crack initiates from the punched

edges of the middle hole, where the stress concentrated.

Figure 6.11: Macroscopic fracture surface of sample from DP1180A+EG, failed incremental loading test (punched edges, in NaCl solution).

Figure 6.12 shows the microscopic fracture surface of sample from DP1180A+EG. In the SZ, the

fracture surface exhibits brittle fracture with an intergranular mode. The fracture surface in the IZ

displays ductile fracture. As a reference, the fracture surface in the sample after ILT in the air was also

examined. It displays also ductile fracture. It can be assumed, when the sample is tested under

incremental loading, the bare edge is attacked by the hydrogen atoms from the solution. The hydrogen

atoms concentrate at the area under stress. When the critical hydrogen concentration to crack initiation

is reached, DF takes place. Thus the crack initiates from the bare edge and propagates into the sample.


(a) macroscopic fracture surface, in NaCl solution

(b) SZ, tested in NaCl solution

(c) IZ, tested in NaCl solution

(d) fracture surface, tested in air

Figure 6.12: Micrographs of fracture surfaces of the sample from DP1180A+EG, after incremental loading test (punched edges): (a) to (c) in NaCl solution and (d) in air.

In ILT, the Embrittlement Index (EIILT) is defined as:

(6-3)

where σ is the applied stress when DF occurs, tf stands for the time to DF. Table 6.2 lists the value of

EIILT of the different steels. It is worth noting that a value of EIILT below 60 indicates that the DF is

observed during the constant loading process within 48 h at the beginning of the ILT. While a value of

74.8 represents that the material fails when the loading reaches tensile strength.

Regarding the EIILT, the susceptibilities to HE of the investigated materials are ranked in the following

order:


TBF1180B+EG > TBF1180A+EG > DP1180A+EG > TBF1180C+EG > DP1180B+GI > CP1180 +EG >

MS-W1180+EG


DP980+EG > TBF980+EG > CP980+GI


Table 6.2: EIILT of investigated materials in incremental loading test in NaCl solution.

Group Material EIILT [(%·h)1/2]

1180 MPa steels

CP1180+EG 44.3

DP1180A+EG 11.9

DP1180B+GI 14.0

TBF1180A+EG 9.6

TBF1180B+EG 6.1

TBF1180C+EG 13.7

MS-W1180+EG 62.1

980 MPa steels

CP980+GI 67.6

DP980+EG 16.2

TBF980+EG 17.4

Compared to SSRT, the steel grade of TBF1180A shows slightly better resistance to DF than

TBF1180B in ILT, and DP980 is more susceptive to DF than TBF980. The probable explanation is that

in the ILT samples the hydrogen content in TBF1180B is higher than that in TBF1180A, as described

in Figure 5.2; whereas, in the SSRT the hydrogen concentration was modified by hydrogen charging,

which is more comparable. Besides of the internal hydrogen, it is worth noting that in the SSRT the

loading is dynamically increased, until the fracture stress is reached. Hydrogen re-contribution is

dominant by the pre-charging and subsequent strain-induced phase transformation in TBF steels.

However, in the ILT, the loading is stepwise increased and the hydrogen uptake occurs during the

loading associated with corrosion from the surroundings. In this case the hydrogen concentration of as-

delivered state plays an essential role, by influencing the time of crack initiation. Under the same start

loading of 75% of Rm, the material has lower yield ratio, such as TBF1180B, have higher plastic strain

than the ones with lower yield ratios, e.g., TBF1180A, leading to probably stronger hydrogen entrance

from the edges. Additionally, the material of TBF1180C displays moderate susceptibility to HE in ILT

compared to that in SSRT. The possible reason can be the good resistance to corrosion, which will be

discussed in the Section 9.3.4.

6.2.2 Hydrogen-induced cracking on welded samples

To investigate the influence of welding, the tensile samples with spot welding point and milled edges

were tested by an ILT test. Figure 6.13 illustrates the results of σ and tf from the selected materials. The

steels with 1180 MPa tensile fail at the applied stress of 75% of Rm, during the initial constant load

period. CP1180+EG shows a tf of 24.6 h, while DP1180A+EG and TBF1180A+EG display a shorter tf

of less than 1 h. After welding the samples from CP980+GI fail when the stress reaches 95% of Rm.

The fracture in the welded DP980+EG occurs when the stress is increased to 100% of Rm, indicating

the highest resistance to HE of the investigated welded samples.


Figure 6.13: Applied stresses and time to delayed fracture of welded samples in incremental loading test (in

NaCl solution).

The welded samples after the ILT display different fracture surfaces macroscopically. In CP1180+EG,

DP1180A+EG and TBF1180A+EG, the samples crack along the boundary of the welding point when

they are subjected to load in air (Figure 6.14 (a-1)). When tested in the NaCl solution, the crack in the

sample goes through the welding point, as shown in Figure 6.14 (a-2), and the fracture surface is approx.

90° to the loading direction. In the samples from CP980+GI and DP980+EG, the samples loaded in air

show characteristic shear fracture. The fracture surface is 45-60° to the loading direction and near the

fracture surface typical necking areas can be observed (Figure 6.14 (b-1)). When the samples were

tested in the solution of NaCl, the fracture occurred in in the region near the boundary of welding point

Figure 6.14 (b-2).

DP1180A+EG reveals a high susceptibility to HE considering its short fracture time under stress in the

value of 75% of Rm. Under LOM, several sub-cracks in the welding region are found. As illustrated in

Figure 6.15, the severe sub-cracks are at 90° to the direction of loading. The arrested region of the sub-

cracks displays a branch-like shape. It can be concluded that stress corrosion cracking dominates the

earlier failure when the sample is loaded in the NaCl solution.

The fracture surfaces of the welded samples were examined with SEM. Figure 6.16 shows the

fractography of DP1180A+EG. It can be found that in the welding region the fracture surface displays

intergranular fracture mode. The second micro-cracks between the grains can be seen. While in the area

outside of the welding point, the fracture surface displays ductile fracture, analogical to the as-delivered

materials after tensile test. The microscopic fracture surface provides the evidence that on the welded

sample, the DF initiates from the welding region and propagates into the area around the welding point,

due to hydrogen embrittlement.


(a-1) DP1180A+EG, in air (a-2) DP1180A+EG, in NaCl

(b-1) DP980+EG, in air (b-2) DP980+EG, in NaCl

Figure 6.14: Illustration of fractures of the welded samples after incremental loading test, tested in air and NaCl solution: (a) DP1180A+EG and (b) DP980+EG.

Figure 6.15: Sub-cracks in the welding region near the fracture surface via LOM, from DP1180A+EG after

incremental loading test in NaCl solution.

To characterize the welded region, the local hardness was tested using the Vickers hardness test.

Figure 6.17 illustrates the distribution of local hardness across the welded region in TBF1180A+EG.

In the base material of TBF1180A+EG, the hardness is in the range of 400-430 HV0.5. In the welding

zone, the hardness increases greatly, into the range of 520-570 HV0.5. In the welded zone the

microstructure is directly influenced by the current density during welding, where martensite is formed,

displaying a very high hardness [WAN13]. The increase of hardness in the welded zone is attributed to

the carbon equivalent of the material [LOI14]. Between the base material and the welded zone, in the

so-called Heat Affected Zone (HAZ), the local hardness drops to approx. 350 HV0.5. The rapid drop of

hardness in HAZ is probably due to the heat effect in welding process. Thus the microstructure in the

vicinity of welded zone is affected and coarse-grained microstructure can be found [WAN13].


(a) (b)

(c) (d)

Figure 6.16: Microscopic fracture surface of sample from DP1180A+EG after incremental loading test in NaCl solution: (a) (b) welding region and (c) (d) outside of welding.

Figure 6.17: Hardness profile across the welding point in TBF1180A+EG.

6.3 Results of Constant Loading Test (CLT) 69

6.3 Results of Constant Loading Test (CLT)

In the CLT the Zn-coated samples with bare cutting edges were immersed in salt solution and subjected

to constant stress. Figure 6.18 illustrates the time to fracture (tf) of the investigated materials, in

minimum value of time to fracture (gray) and the average value of time to fracture from 5 repeat tests

(blue). Because the bare edges had been cut by punching, the tf in the CLT reveals the resistance to HE

combined with material edge sensitivity.

When the applied stress was Rp0.2 (when Rp0.2/Rm<0.9) or 80% Rm (when Rp0.2/Rm>0.9), all 1180 MPa

steels fail within 96 h of testing time. CP1180+EG reveals a longer tf compared with other materials,

indicating a better resistance to HE. DP1180A and DP1180B have tf in an average value of 5.2 h and

9.4 h respectively. TBF1180A+EG and TBF1180B+EG show a short tf, while TBF1180C+EG performs

better with a tf of 8.6 h. The samples from MS-W1180+EG fail after 13.5 h testing. In the group of

980 MPa steels, DF does not occur in CP980+GI after 96 h testing, demonstrating an excellent

resistance to HE. The materials of DP980+EG and TBF980+EG fail after 6.4 h and 8.9 h respectively.

Figure 6.18: Comparison of time to fracture in constant loading test in the NaCl solution. Applied stress σ=80%

Rm (MS-W1180+EG and CP980+GI, with Rp0.2/Rm>0.9) or σ=Rp0.2 (other materials, with Rp0.2/Rm<0.9).

Generally, the result of CLT under certain stress is consistent with that of ILT. The shorter incubation

time to DF of MS-W1180 in CLT compared with that in ILT can be explained by the different loading

grades. Similar to ILT, on the fracture surface the area close to the punched edge reveals a brittle fracture

mode, as shown in Figure 6.19. In the SZ the fracture surface from CP1180+EG (Figure 6.19 (a))

reveals quasi-cleavage fracture with transgranular mode. In the sample from DP1180A+EG (Figure

6.19 (b)) and TBF1180B+EG (Figure 6.19 (c)) the brittle facets and second cracks can be observed.

The fracture surface of MS-W1180+EG in Figure 6.19 (d) demonstrates also transgranular fracture and

micro-cracks, characterized as the hydrogen-induced fracture. DF initiates at the bare Zn-Fe-Zn edges

before the maximum tensile stress had been reached, because of hydrogen uptake in the aqueous

surroundings (as mentioned in section 2.3.1).


(a) CP1180+EG

(b) DP1180A+EG

(c) TBF1180B+EG (d) MS-W1180+EG

Figure 6.19: Fracture surfaces in surface zone after constant loading test (in 3.5%NaCl). Applied stress σ=80%Rm (MS-W1180+EG with Rp0.2/R m>0.9) or σ=Rp0.2 (other materials, with Rp0.2/Rm<0.9).

For industrial production it is important to learn the critical loading for the material so as to ensure

suitable applications. To identify the theoretical allowed applied loading, each material was additionally

tested in salt solution under different loads. The maximum allowed stress (σmax) was defined when all

three repeat tests run 96 h without showing DF. Figure 6.20 illustrates the σmax for the investigated

materials, expressed as in percent of Rm.

Figure 6.20: Maximum allowed applied stress of the investigated materials in constant loading tests in NaCl

solution.

6.4 Results of U-bending tests 71

The σmax indicates the HE-susceptibility of galvanized materials with punched edges, limiting further

application. The Embrittlement Index in CLT (EICLI) can be expressed by σmax:

EICLT= (6-4)

Table 6.3 lists the EICLI of the various materials. The lower value of EICLI stands for shorter time to DF,

indicating a higher susceptibility to HE. In general, the ranking of HE-susceptibility for 1180 MPa steels

is:

TBF1180A+EG / TBF1180B+EG > TBF1180C+EG / DP1180A+EG /DP1180B+GI > CP1180+EG

>MS-W1180+EG

And for 980 MPa steels: DP980+EG / TBF980+EG > CP980+GI

Table 6.3: EICLT of investigated materials in constant loading test in NaCl solution.

Group Material EICLT [% of Rm]

1180 MPa steels

CP1180+EG 65

DP1180A+EG 55

DP1180B+GI 55

TBF1180A+EG 50

TBF1180B+EG 50

TBF1180C+EG 55

MS-W1180+EG 70

980 MPa steels

CP980+GI 85

DP980+EG 65

TBF980+EG 65

6.4 Results of U-bending tests

6.4.1 U-bending on pre-charged samples

When the pre-charged and Zn-coated samples were bent via three-point-bending test and fixed by bolts,

DF can be observed after a certain incubation time, thus the Internal Hydrogen Assisted Cracking

(IHAC) occurs. Figure 6.21 shows exemplarily the observed normal crack initiation in the U-bend

samples. On the pre-charged samples, the crack initiates at the edges and then propagates into the

sample, perpendicular to the sample edges. The cracks from both edges can stop after some hours, but

most frequently they meet and lead to fracture.


(a) directly after bending (b) 1 h after bending

Figure 6.21: Crack initiation in a delayed fractured U-bend sample from TBF1180A (punched edges, 30 min pre-charged, R=15 mm): (a) directly after bending and (b) 1 h after bending.

Different materials show varying characteristics of crack initiation and propagation behaviors. Figure

6.22 describes the crack propagation in the samples from DP1180A and TBF1180A. In both materials,

the incubation period is short. The cracks in DP1180A grow rapidly, and after as long as 10 h they

propagate and meet together, thus the cracks go through the U-bend sample. In contrast, in TBF1180A,

after initiation the propagation is more moderate than that in DP1180A.

(a) DP1180A (b) TBF1180A

Figure 6.22: Crack propagation and divided propagation phases in delayed fractured U-bend samples (punched edges, 30 min pre-charged, R=15 mm): (a) DP1180A and (b) TBF1180A.

The average value of time to crack incubation (tc) for various materials is illustrated in Figure 6.23.

After 30 min pre-charging, tc of U-bend samples is obviously shorter than that of the 10 min pre-charged

samples due to the higher hydrogen concentration (Figure 5.2). In some materials such as CP steels and

MS-W1180, DF does not take place. After 10 min pre-charging, DF does not occur in DP1180A samples.

While when the pre-charging time is increased to 30 min, the DF initiates in the sample after 2 to 3 h.

DP1180B shows also a short tc in the 30 min pre-charged samples, but only 2 from 5 samples show DF

when the pre-charging time is 10 min. In the TBF steels, tc from TBF1180A and TBF1180B display a

low value. In TBF1180C tc is longer than that in TBF1180A and TBF1180B, but the range error in

TBF1180C is higher than in other TBF steels. All the TBF steels with 1180 MPa are susceptive to HE.


Compared to 1180 MPa steels, the 980 MPa steels show lower HE-susceptibility in pre-charged U-bend

samples. There exists only one exception from TBF980 thus cracks approximately in 4 h after bending.

Figure 6.23: Comparison of incubation time to delayed fracture (tc) in U-bend samples (Punched edges,

R=15 mm). i/j: the amount of cracked sample is i and the total tested sample in this variant is j.

Aside from tc, the length of crack propagation (ltot) of each sample was measured until the crack went

through the sample or up to 96 h. Figure 6.24 summarizes the average values of ltot in the pre-charged

U-bend samples from different materials. The amount of cracks (n) in each sample depends on both

material and charging time, as compared in Figure 6.25. When the DF does not take place within 96 h

observation, ltot and n have a value of 0. Generally, the material that has a shorter tc displays a higher

value of ltot and n. The longer charging time leads to a higher value of ltot. In TBF1180A and TBF1180B,

ltot reach over 50 mm.

Figure 6.24: Comparison of total lengths of cracks (ltot) in the U-bend samples (Punched edges, R=15 mm).

i/j: the amount of cracked sample is i and the total tested sample in this variant is j.


Figure 6.25: Comparison of amounts of cracks (n) in the U-bend samples. (Punched edges, R=15 mm). i/j: the

amount of cracked sample is i and the total tested sample in this variant is j.

The fracture surfaces of the U-bend samples were investigated via SEM. The fracture surface is divided

into initiation region, middle region and arrested region, according to the crack initiation and

propagation processes, as displayed in Figure 6.26.

Figure 6.26: Macrographs of Fracture surface of delayed fractured U-bend sample.

Figure 6.27 shows the fracture surfaces in the U-bend sample from DP1180A. In the crack initiation

region, brittle facets and second micro-cracks can be seen, indicating high susceptibility to HE. In the

middle region, the fracture surface transforms into transgranular and ductile mixed fracture mode. Here

stress corrosion cracking dominates the fracture of the sample. In the arrested region, ductile fracture

with dimples is observed, which is considered as a feature of ductile overload under stress.

Furthermore, the fracture surfaces in the initiation region of the samples from other fractured materials

are shown in Figure 6.28. In pre-charged material of DP1180B, the fracture surface is in transgranular

fracture mode, associated with small amount of brittle facets and second cracks, as shown in Figure

6.28 (a). In the TBF steels, as revealed in Figure 6.28 (b) to (d), intergranular fracture is predominant.


(a) initiation region (b) middle region

(c) arrested region

Figure 6.27: Micrographs of fracture surface of delayed fractured U-bend sample from DP1180A (punched edges, 30 min pre-charged, R=15 mm).

(a) DP1180B

(b) TBF1180A

(c) TBF1180B

(d) TBF1180C

Figure 6.28: Micrographs of fracture surface of delayed fractured U-bend sample (punched edges, 30 min pre-charged, R=15 mm).


In general, ltot and n are inversely proportional to tc. The Embrittlement Index in the U-bending test on

the pre-charged samples (EIU-H) is characterized as tc, as:

EIU-H=tc (6-5)

The EIU-H of investigated materials is given in Table 6.4, the ranking of HE-susceptibility in U-bend

test with the pre-charged samples is ordered from high to low:


TBF1180A > TBF1180B > DP1180A > DP1180B > TBF1180C > CP1180 / MS-W1180


TBF980 > DP980 / CP980

Table 6.4: EIU-H of the investigated materials in U-bending test on pre-charged samples.

Group Material EIU-H [h]

1180 MPa steels

CP1180 >96*

DP1180A 2.3

DP1180B 5.4

TBF1180A 1.7

TBF1180B 2.0

TBF1180C 6.9

MS-W1180 >96*

980 MPa steels

CP980 >96*

DP980 >96*

TBF980 4** * The delayed fracture was not found in 96 h observation. ** 1 from 3 samples cracked.

The trend is consistent with the previous section as in SSRT, except the material TBF1180C+EG. It is

likely that after 30 min charging the hydrogen concentration in TBF1180C is much lower than that of

other steels, which is confirmed in the TDA-measuring in section 5.1. Contrary to SSRT, the strain in

U-bending with bending radius of 15 mm is at a low level. The strain-induced transformation in TBF

steels may not be completed, therefore the supersaturated hydrogen concentration in new strain-induced

martensite has not yet been achieved. Thus the DF in TBF1180C takes place even later than DP steels.

However, when the samples were pre-charged for 10 min, the TBF1180C displays an even earlier DF

than DP1180 steels, indicating the dominance of hydrogen re-contribution due to strain-induced phase

transformation.


6.4.2 U-bend samples in salt solution

Punched samples

The U-bend samples were immersed in 3.5%NaCl (pH4, adjusted with HCl) to investigate the Hydrogen

Environment Assisted Cracking (HEAC) in the materials. In the salt solution, a great amount of

hydrogen and hydroxyl ions are present. All the samples were used in the as-delivered state with original

Zn-layer, with bare Zn-Fe-Zn edges. The samples were bent with punch radius of 15 mm and fixed with

bolts. All the edges of the substrate steels were bare to the solution, which allows directly attack by

hydrogen ions and chlorine ions.

Figure 6.29 illustrates the comparison of tc of different materials, which was calculated as the average

value from five repeat tests. The edges were cut by punching. It shows that in CP1180+EG, MS-

W1180+EG, CP980+GI and DP980+EG, DF is not observed within a 96 h immersion tests. The

TBF1180 and DP1180 samples start to crack within the first 8 h after immersion.

Figure 6.29: Comparison of incubation time to delayed fracture on the U-bend samples (punched edges,

R=15 mm, in NaCl solution). i/j: the amount of cracked sample is i and the amount of total tested sample in this variant is j.

During the immersion the cracks propagate rapidly after initiation. The sample was taken out of the

solution once completely fractured. The total length of cracks (ltot) and number of cracks (n) were

measured. Figure 6.30 illustrates the ltot and n of the investigated materials. Briefly, the comparisons

of ltot and n are in accordance with the ranking by tc.


Figure 6.30: Comparison of total length of cracks and amount of cracks on the U-bend samples (punched

edges, R=15 mm, in NaCl solution). i/j: the amount of cracked sample is i and the amount of total tested sample in this variant is j.

The microscopic fracture surfaces of the fractured samples were characterized via SEM, as displayed

in Figure 6.31 and Figure 6.32. In DP1180A+EG, the fracture surface close to punched edge reveals a

quasi-cleavage fracture mode. Secondary cracks and brittle facets are found and are considered to be

the result of hydrogen attack from the NaCl solution. In the middle region of the fracture surface, both

intergranular fracture and shear ridges can be observed. However, in the middle of fracture surface

ductile fracture is observed, which is considered as a feature of ductile overload due to high stress. In

TBF1180A+EG, the fracture surface in the edge region reveal intergranular fracture mode, and the

brittle facets is finer than that of DP1180A+EG. This is probably due to finer grain size of

TBF1180A+EG, as analyzed by the phase maps of EBSD. In the middle region, the friction of

characteristic intergranular fracture in TBF1180A+EG is higher than that of DP1180A+EG. Indicating

that under the attack of hydrogen in HEAC, degradation of de-cohesion force of TBF1180A+EG matrix

is more severe. Furthermore, in the crack arrested region of TBF1180A+EG samples, the quasi-cleavage

fracture mode is found, which is mostly likely that stress corrosion cracking dominates the fracture.


(a) initiated region (b) middle region

(c) arrested region

Figure 6.31: Fracture surfaces of U-bend samples from DP1180A+EG (punched edges, R=15 mm, in NaCl solution).

(a) initiated region (b) middle region

(c) arrested region

Figure 6.32: Fracture surfaces of U-bend samples TBF1180A+EG (punched edges, R=15 mm, in NaCl solution).


The Embrittlement Index of U-bending samples in NaCl solution (EIU-NaCl) is identified as tc:

EIU-NaCl=tc (6-6)

Table 6.5 gives the values of EIU-NaCl of the investigated steel grades regarding to the punched edges.

The HE-susceptibilities of different steels are ranked as follows:


TBF1180A+EG > TBF1180B+EG > TBF1180C+EG > DP1180B+GI > DP1180A+EG > CP1180+EG

/ MS-W1180+EG


TBF980+EG > DP980+EG / CP980+GI

Table 6.5: EIU-NaCl of investigated materials in U-bending samples in the NaCl solution.

Group Material EIU-NaCl [h]

1180 MPa steels

CP1180+EG >96*

DP1180A+EG 7.6

DP1180B+GI 5.3

TBF1180A+EG 3.8

TBF1180B+EG 4.8

TBF1180C+EG 4.9

MS-W1180+EG >96*

980 MPa steels

CP980+GI >96*

DP980+EG >96*

TBF980+EG 26.1** * The delayed fracture was not found in 96 h observation. ** 3 from 5 samples cracked.

The HE-susceptibilities of the investigated material by using the U-bend samples in the NaCl solution

surroundings is similar as that by ILT and CLT. In the study of U-bend samples in the salt solution, the

crack initiation in DP1180B+GI takes place earlier than that in DP1180A+EG. It is assumed probably

due to two reasons. Firstly, the lower yield ratio of DP1180B+GI results in the higher plastic strain than

that of DP1180A+EG when the bending radius is the same. Secondly, the DP1180B+GI is hot-dip

galvanized material, thus the Zn-coating is different from the one by electro-galvanizing of

DP1180A+EG. The Zn-Fe intermetallic phase after hot-dip galvanizing is likely to reduce the corrosion

resistance of the material, therefore promotes the possibility of hydrogen pickup from the salt solution.


Milled and laser cut samples

Aside from samples with punched edges, the laser cut and milled U-bend samples were tested in the

same solution. Table 6.6 summarizes the characteristic tc, ltot and n of the U-bend samples in different

edges conditions. Compared with the results of punched samples, the samples with milled edges from

all the steel grades do not reveal any DF. When the edges are cut by laser cutting, DF takes place only

in the materials of TBF1180 steels. It is observed that the samples from TBF1180C+EG display longer

tc but comparable ltot and n than that from materials TBF1180A+EG and TBF1180B+EG.

Table 6.6: Comparison of incubation time to delayed fracture (tf), total length of cracks (ltot) and number of cracks (n) of U-bend samples (R=15 mm, in NaCl solution).

Material Punched edge Lasered edge Milled edge

tc [h]

ltot [mm]

n [-]

tc [h]

ltot [mm]

n [-]

tc [h]

ltot [mm]

n [-]

1180 MPa steels

CP1180+EG > 96 0 0 > 96 0 0 > 96 0 0

DP1180A+EG 7.6 21.4 2.7 > 96 0 0 > 96 0 0

DP1180B+GI 5.3 22.5 2.7 > 96 0 0 > 96 0 0

TBF1180A+EG 3.8 26.1 4.3 4.2 23.1 2.7 > 96 0 0

TBF1180B+EG 4.8 26.0 4.6 5.5 23.0 3.2 > 96 0 0

TBF1180C+EG 4.9 24.0 3.7 14.5 22.3 2.7 > 96 0 0

MS-W1180+EG > 96 0 0 > 96 0 0 > 96 0 0

980 MPa steels

CP980+GI > 96 0 0 > 96 0 0 > 96 0 0

DP980+EG > 96 0 0 > 96 0 0 > 96 0 0

TBF980+EG 26.1 19.6 2.0 > 96 0 0 > 96 0 0

The fracture surfaces of the fractured TBF1180A+EG samples with lasered cut edges (Figure 6.33)

differ from that with punched edges (Figure 6.32). On the fracture surface of punched sample,

intergranular fracture mode can be found in the edge region, where the crack initiates. However, when

the edges were laser cut, in the location of 40-50 μm below the edge surface a transit area in fracture

surface is observed, where the fracture mode changes (intergranular–transgranular-intergranular). The

assumption for this observation is the microstructure transformation due to the heat effect during laser

cutting.


(a) overview of fracture surface (b) Micrographs P1 of (a)

(c) Micrographs P2 of (a) (d) Micrographs P3 of (a)

Figure 6.33: Fracture surfaces of U-bend samples from TBF1180A+EG (laser cut edges, R=15 mm, in NaCl solution).

6.4.3 U-bend samples in other corrosion mediums

The U-bend samples with punched edges were additionally tested by salt spray test and the VDA102-

233 corrosion test. The samples in salt spray test were observed once per day, thus the evolution of

corrosion is shown in Figure 6.34. In salt spraying surroundings, the Zn-coating protects the steel

substrate from corrosion in the beginning. The corrosion product occurs after 3 days on both sample

surface and bare edges. After 10 days the rust is found, indicating the severe corrosion of steel substrate.

Figure 6.34: Photos of U-bend samples in salt spray test (punched edges, R=15 mm).

The incubation time to crack initiation (tc) of U-bend samples in salt spray test is illustrated in Figure

6.35. Because of the observation frequency, once the DF is observed the samples are already completely

broken. Therefore, the initiation and propagation behaviors cannot be characterized. Hence, here tc is

only an approximate parameter, which cannot indicate the susceptibility to HE very accurately. Briefly,


the U-bend samples from CP1180+EG and MS-W1180+EG display only corrosion without DF during

10 days salt spray test, which is consistent with the results of U-bend samples in the salt solution. The

samples from DP and TBF steels with Rm of 1180 MPa fail in the first three days.

Figure 6.35: Comparison of fracture time on the U-bend samples in salt spray test (punched edges, R=15 mm, salt spray tests).

Figure 6.36 shows the photos of U-bend samples during VDA corrosion test. The Zn-coating protects

the surfaces of the sample from corrosion in the beginning; the corrosion product is formed on the bare

Zn-Fe-Zn edges. After 3 weeks, the surface area of the U-bend sample starts to corrode, where several

corrosive pitting can be seen. The pitting enlarges with increasing time, meanwhile new pitting

generates on the sample surfaces. By the end of corrosion test after 6 weeks, the Zn-layers are

completely destroyed and the samples are heavily corroded.

Figure 6.36: Photos of U-bend samples in VDA233-102 corrosion test (punched edges, R=15 mm).

Figure 6.37 shows the results of U-bend samples in VDA233-102 corrosion test. The edges of sample

were cut by punching. It can be seen that only partial samples from DP1180A+EG, DP1180B+GI and

TBF1180A+EG fail during this test. The error amount in the diagram displays a high value, indicating

the low reproducibility of the repeat tests. It is considered that in the VDA corrosion test, hydrogen

uptake is limited. Corrosion is more dominant than HE.


Figure 6.37: Comparison of fracture time on the U-bend samples in corrosion test (edges punched, R=15 mm,

VDA corrosion tests). i/j: the amount of cracked sample is i and the total tested sample in this variant is j.

The Embrittlement Indexes of U-bend samples in salt spray test (EIU-SS) and VDA corrosion test (EIU-

VDA) are expressed by tc, as shown in equation (6-7):

EIU-SS, EIU-VDA=tc (6-7)

Table 6.7 lists the values of EIU-SS and EIU-VDA of the investigated materials. In the salt spraying

surroundings, the HE-susceptibilities of different steels are ranked from high to low:


TBF1180A+EG / DP1180A+EG > TBF1180B+EG / DP1180B+GI > TBF1180C+EG > CP1180+EG /

MS-W1180+EG

The 980 MPa steels cannot be ranked, because DF is not observed.

In the case of the VDA corrosion test, however, only small amount of samples cracked. The

susceptibility to HE of the investigated material is not enough to be compared here.


Table 6.7: EIU-SS and EIU-VDA of the investigated materials of U-bend samples in corrosion tests.

Group Material EIU-SS [d] EIU-VDA [w]

1180 MPa steels

CP1180+EG >10* >6*

DP1180A+EG 2 3.5**

DP1180B+GI 2.3 4**

TBF1180A+EG 2 2**

TBF1180B+EG 2.3 >6*

TBF1180C+EG 3 >6*

MS-W1180+EG >10* >6*

980 MPa steels

CP980+GI >10* >6*

DP980+EG >10* >6*

TBF980+EG >10* >6* * Delayed fracture is not observed. ** Delayed fracture occurs on partial samples.

6.4.4 Characterization of different edges

Depending on different cutting methods, the edge surfaces after cutting display different morphologies,

as illustrated in Figure 6.38. On the punched edge surface, a characteristic flush-cut surface can be

observed (Figure 6.38 (a) and (b)). Close to the flush-cut surface, there is the fractured region after

punching. In contrast, the laser cut surface is quite homogeneous, as displayed in Figure 6.38 (c) and

(d). The molten metal on the edge surface can be observed, which is confirmed to consist of Zn by

Energy-Dispersive X-ray (EDX)-detection. The reason is assumed to be that during laser process the

Zn-coating on the sample surfaces is melted and flows along the edge surface. Figure 6.38 (e) and (f)

display the edge surface after milling. It can be observed that the fine scratches from the milling process

presents on the milled surface. The values of roughness in Ra of edge surfaces by different cutting

methods are shown in Table 6.8.

Table 6.8: Roughness of the edge surfaces by different cutting methods, TBF1180A+EG.

Cutting method Punching

Laser cutting Milling flush-cut region Fracture region

Ra [μm] 0.34 0.47 0.31 0.32


(a) (b)

(c) (d)

(e) (f)

Figure 6.38: Morphology of edge surfaces of TBF1180A+EG after: (a) (b) punching, (c) (d) laser cutting and (e) (f) milling.

The geometry of the punched edge has been described in a publication by E. Deoge [DOE10]. As shown

in Figure 6.39 (a), after punching the edges from the free side display strong deformation, with a burr

in the horizontal direction beneath the top of the steel sheet. In all test samples the shear cutting sides

with ridge were used, to obtain the same critical state in the bending test. At the edge, a grid with step

of 0.1 mm was set, and the local hardness on the intersection of grid was measured accordingly. Figure

6.39 (b) illustrates the hardness map of punched edge in TBF1180A sample. The change of color from

blue-green to yellow reveals an increase of hardness especially in the region close to the burr. The edges

after laser cutting and milling show no shear cut character, but rather smooth surfaces without a burr.


The hardness of lasered edges, as shown in Figure 6.39 (c), reveal a significant increase in hardness.

The maximum hardness reaches 580 HV0.3 on the overall edge surface. In contrast, the distribution of

hardness in the milled samples is homogenous and no significant increase in hardness at the edge can

be found (Figure 6.39 (d)).

(a) Measured area, punched edge (b) Punched edge

(c) Laser cut edge (d) Milled edge

Figure 6.39: Illustration of measured area and the hardness map of different edges.

The values of hardness of the tested materials with different edges are plot in Figure 6.40 additionally.

In the punched samples, the measurements were made from the top point from burr, along the horizontal

direction (Figure 6.40 (a)). As illustrated in Figure 6.40 (b), the hardness is highest in the first

measurement near the edge surface, and then decreases as the measurement location departs from the

edge surface. TBF1180A steel shows the highest maximum localized hardness on the burr region of

punched edge, with a value of 492 HV0.3. DP1180A and CP1180 show similar values of hardness in the

overall measured field. The hardness of MS-W1180 is the lowest of the four steel grades, with the

maximum hardness of 423 HV0.3.


(a) Measuring direction, punched edge (b) Punched edges

(c) Laser cut edges (d) Milled edges

Figure 6.40: Hardness distribution in the transversal surface, from cut edge into the material.

The values of hardness measured from the edge surface along the horizontal direction in the samples

after laser cutting are shown in the diagrams in Figure 6.40 (c). In all the materials, the first

measurements close to the lasered edge display a significantly higher value of hardness. Along the depth

from edge surface in the sample, the hardness drops to a bottom value when the distance from edge

surface reaches 0.15 mm and increases again, to have a value as basic material. In TBF1180A the

maximum hardness close to edge surface has the highest value of the four materials. However, the

hardness of the other three materials, CP1180, DP1180A and MS-W1180 show still a local hardness

increase to a value below 500 HV0.3 through the influence of the laser.

In contrast, in the edge areas after milling the change of hardness is not found in all investigated

materials, as shown in Figure 6.40 (d).

The relative increase of hardness (∆HV) in the materials is calculated by equation (6-8):

(6-8)

where HV0.3, max is the maximum hardness on punched edge, HV0.3, basic is the hardness of basic material,

without the effect of cutting. Based on the measured hardness, the values of ∆HV0.3 due to different

cutting method in different materials are summarized in Table 6.9.

6.5 Results of four-point-bending test 89

Table 6.9: Values of local hardness in the edge regions from four kinds of materials, measured from edge surface in the depth direction.

Material Hardness of

basic material [HV0.3]

Punched edge Lasered edge Milled edge Max.

hardness [HV0.3]

∆HV0.3 [%]

Max. hardness [HV0.3]

∆HV0.3 [%]

Max. hardness [HV0.3]

∆HV0.3 [%]

CP1180 406 438 7.9 498 22.7 409 0.7

DP1180A 410 444 8.3 479 16.8 418 2.0

TBF1180A 407 492 20.9 557 36.9 412 1.2

MS-W1180 410 438 6.8 448 9.3 418 2.0

6.5 Results of four-point-bending test

The stress state in the U-bend sample is determined by the bending radius. To identify the crucial stress

to DF, four-point-bending test combined with immersion test in the NaCl solution was carried out.

Figure 6.41 displays the time to crack initiation tc of four steel grades. In CP1180+EG and MS-

W1180+EG, DF is not observed even the bending grade reaches maximum bending grade (16) of the

apparatus. In DP1180A+EG and TBF1180A+EG, DF does not occur when the bending grade is 10;

when the bending grade is increased to 12, DF is observed in partial four-point-bend samples. However,

when the bending grade is 16, cracks initiate after a short tc.

Figure 6.41: Comparison of incubation time to delayed fracture under different bending grades in four-point-

bend test (punched edges, in NaCl solution). i/j: the amount of cracked sample is i and the amount of total tested sample in this variant is j.

The stress and strain states of the four-point-bending samples were calculated based on Finite Element

(FE) method. In the calculation, the elastoplastic two-dimension symmetrical modelling by four-point

elements was used in software LS-Dyna. The mechanical properties were obtained from quasi-static

tensile tests with strain rate of 10-3 s-1. The elastic modulus was 206 GPa. The true stress as a function

of true plastic strain in the range of yield strength to tensile strength was used in the simulation. The

maximum strain and stress occurred in the middle of outside surface. The values of effective strain and


effective stress in each bending grade of CP1180 (t=1.6 mm) and TBF1180A (t=1.5 mm) are plot in

Figure 6.42. The effective strain increases linearly with increasing bending grade. The higher strain

leads to higher value of stress, associated with flow-curve.

(a) CP1180 (b) TBF1180A

Figure 6.42: Effective stress and strain on the outside surface of four-point-bend sample, calculated by 2-dimension FE-modelling in LS-Dyna: (a) CP1180 and (b) TBF1180A (bending grades from 10 to 16).

The maximum allowed bending grade (ymax) was determined when no crack is observed in three repeat

tests after 96 h immersion. The maximum allowed stress (σmax) was calculated according to the ymax, by

FE method as introduced above, and expressed as percentage of Rp0.2. In some materials such as

CP1180+EG and MS-W1180+EG, ymax cannot be detected, because DF is not observed. Figure 6.43

illustrates the values of ymax and σmax of various materials. It can be found that DP1180 steels

TBF1180A+EG reveal a ymax of 11, while both TBF1180B+EG and TBF1180C+EG display a lower

ymax in the value of 10. The σmax in the both DP1180 steels are above 102% of Rp0.2, while in TBF1180

steels σmax distributes in the range of 100% to 102% of Rp0.2.

Figure 6.43: Comparison of maximum allowed bending grades (ymax) and the simulated value of maximum

allowed stress (σmax) in the four-point-bend samples (punched edges, NaCl solution for 96 h).

6.5 Results of four-point-bending test 91

The Embrittlement Index (EI4PB) in the four-point-bending test, is defined by ymax:

EI4PB= ymax (6-9)

Table 6.10 shows the EI4PB of different steels. Thereby the HE-susceptibilities are ranked as following,

from high to low:


TBF1180B+EG / TBF1180C+EG > TBF1180A+EG / DP1180A+EG / DP1180B+GI > CP1180 / MS-

W1180

And for 980 MPa steels, the samples did not crack even when the maximum bending grade was reached.

The susceptibility to HE is not evaluated here.

Table 6.10: EI4PB of investigated materials in four-point-bending test.

Group Material EI4PB[-]

1180 MPa steels

CP1180+EG ≥16*

DP1180A+EG 11

DP1180B+GI 11

TBF1180A+EG 11

TBF1180B+EG 10

TBF1180C+EG 10

MS-W1180+EG ≥16*

980 MPa steels

CP980+GI ≥16*

DP980+EG ≥16*

TBF980+EG ≥16* * The delayed fracture is not found in testing time.

92

7 Effects of Hydrogen on Forming Behavior

7.1 Forming Limit Curve (FLC)

Degradation of forming limit

According to the hydrogen calibration curve in Figure 5.2, the concentration of diffusive hydrogen in

DP1180A after 30 min pre-charging reaches 1.45 ppm. The forming limits in uni-axial, plane strain and

biaxial strain states on the uncharged and 30 min pre-charged samples were determined by Nakajima

tests, as shown in Figure 7.1. It is revealed that the forming limits in different strain states drop by

different amounts. In the uni-axial strain state, the obvious distinction between uncharged and pre-

charged samples can be found. The major strain drops from 0.11 to 0.10 while the minor strains stay

the same. In contrast, in strain path of plane strain, the forming limit changes to a lower level when the

samples are pre-charged for 30 min. Due to hydrogen charging the major strain in plane strain state

decreases from 0.09 to 0.06. In biaxial strain state, the forming limit in uncharged material is 0.18 of

major strain and 0.16 of minor strain. It is worth noting that after pre-charging the localized necking

period does not occur in biaxial strain state. As illustrated in Figure 7.2 (a), on the uncharged material

the representative thinning rate diagram reveals both stable and unstable deformation periods. Therefore,

the necking is determined by the intersection of fitting curves. Conversely, after pre-charging the

representative thinning rate increases linearly, indicating only stable deformation (Figure 7.2 (b)),

implying that the samples underwent stable deformation until failure [GAO16a].

Figure 7.1: Forming limits of uncharged and charged DP1180A in Nakajima tests [GAO16a].

7.1 Forming Limit Curve (FLC) 93

(a) uncharged (b) pre-charged

Figure 7.2: Representative thinning rate diagrams at strain state of biaxial from (a) uncharged and (b) pre-charged DP1180A samples in Nakajima tests [GAO16a].

To study the strain state effect on the degradation of forming limit, the critical strains to fracture in the

sample without and with 1.45 ppm diffusive hydrogen are compared in Figure 7.3. It can be seen that

in all three investigated strain states, the threshold strains leading to fracture are decreased due to the

hydrogen enrichment. The degradation of critical strain to fracture is more significant than that of

forming limit. In Table 7.1 the values of critical effective strains to necking and fracture on the

uncharged and pre-charged samples are compared. After hydrogen charging the critical effective strains

to fracture is reduced from 0.26 to 0.18 in the uni-axial state, from 0.17 to 0.10 in the plane strain state

and from 0.46 to 0.30 in the biaxial state respectively [GAO16a].

Figure 7.3: Critical strains to fracture of uncharged and charged DP1180A in Nakajima tests [GAO16a].


Table 7.1: Comparison of critical effective strains to necking and fracture on uncharged and pre-charged DP1180A in Nakajima tests.

Before necking Before fracture

Strain state Uni-axial Plane strain Biaxial Uni-axial Plane strain Biaxial

uncharged 0.11 0.12 0.36 0.26 0.17 0.46

H pre-charged 0.10 0.09 - 0.18 0.10 0.30

Fractography

Figure 7.4 shows the fracture surfaces of the fractured Nakajima samples. In the uncharged samples,

as shown in Figure 7.4 (a-1), (b-1) and (c-1), the fracture surfaces exhibit overall ductile fracture with

characteristic dimples, indicating void nucleation and coalescence. After hydrogen pre-charging, the

fracture surfaces display quasi-cleavage fracture in all investigated strain states. The transgranular

fractures in vicinity of the aluminum-oxide inclusion (Figure 7.4 (a-2)) and inside of large dimples

(Figure 7.4 (b-2) and (c-2)) can be observed. The other areas outside displays still ductile fracture mode.

The occurrence of quasi-cleavage fracture indicates the cracks due to the enrichment of hydrogen. As

explained by Hydrogen Enhanced De-cohesion (HEDE) mechanism, the concentrated hydrogen

degrades the cohesion force of matrix, causing transgranular fracture. Besides, the earlier failure of

Nakajima sample is probably due to the effect of hydrogen on the motion of dislocations, according to

Hydrogen Enhanced Plasticity (HELP) mechanism.

7.1 Forming Limit Curve (FLC) 95

(a-1) uni-axial, uncharged (a-2) uni-axial, pre-charged

(b-1) plane strain, uncharged (b-2) plane strain, pre-charged

(c-1) biaxial, uncharged (c-2) biaxial, pre-charged

Figure 7.4: Fracture surfaces of uncharged and pre-charged samples after Nakajima test in the strain state of (a) uni-axial, (b) plane strain and (c) biaxial.


7.2 Bending Limit Curve (BLC)

Bending Limit Curve (BLC) was obtained by superimposing of pre-strain by Marciniak test and bending

load via VDA 238-100 bending test. Hydrogen concentration in the samples was modified by electro-

charging before the bending test. The contents of diffusive hydrogen in the pre-strained samples are

listed in Table 7.2. Though the hydrogen pre-charging processes of all the samples were same, the

samples with pre-strain adsorbed more hydrogen than the samples without pre-strain. The similar results

have been studied in the previous investigations from Li et al. and Ronevich et al. [LI14] [RON12]. The

rising of hydrogen content can be explained by the introduction of dislocations and the extension of

lattice structures during pre-strain, which provide more trapping sites for hydrogen atoms [LI14].

Table 7.2: Diffusive hydrogen content of VDA bend samples after pre strain via Marciniak test. Pre-charged for 30 min.

Pre-strain Without pre-strain Uni-axial Plane strain Biaxial

Content of diffusive hydrogen [ppm] 1.45 1.52 1.60 1.63

The strain states after pre-strain in uni-axial, plane strain and biaxial states and VDA bending test,

namely enhanced Bending Limit Curve (eBLC) [DEN11], are shown in Figure 7.5. The pre-strain

processes in three investigated strain states provide the distribution of major and minor strains. After

pre-strain the effective strain in the sample in three investigated strain states reached 0.02. As exhibited

in Figure 7.5, a major strain is superimposed in the pre-strained samples. By comparing the eBLC in

the uncharged and pre-charged samples, it can be found that the enrichment of hydrogen lowers the

eBLC of the material.

Figure 7.5: eBLC of DP1180A by pre-straining in Marciniak test and superimposed VDA-bending test

[GAO16b].

7.2 Bending Limit Curve (BLC) 97

The effective strains to bending fracture in the samples via different pre-strain paths are calculated and

listed in Table 7.3. The degradation of effective strain is defined as an Embrittlement Index (EIBL) to

evaluate susceptibility to hydrogen, which was calculated according to equation (7-1):

(7-1)

where is the maximum allowed effective strain without bending fracture. The higher represents

higher susceptibility to HE during bending process. For the material without pre-strain, the is 27.6 %

after 10 min pre-charging and 31.5% after 30 min pre-charging. The degradation of critical effective

strain of pre-strained material has a higher value compared to the uncharged material. The samples pre-

strained in biaxial strain state displays the highest degradation of critical effective strain after hydrogen

pre-charging. After 30 min pre-charging the effective strain to bending fracture reduces from 0.177 to

0.116. After pre-straining in plane strain the effective strains degrade in 32.5% and 33.9% after 10 min

and 30 min pre-charging respectively. The bend samples pre-strained in uni-axial strain state show

lower degradation of bending limit compared to other two strain states.

Table 7.3: Comparison of critical effective strains to bending fracture on as-delivered and pre-strained DP1180A samples VDA238-100 bending test.

Charging time [min]

As-delivered

Pre-strained

Uni-axial Plane strain Biaxial

Effective strain [-]

0 min 0.159 0.184 0.188 0.177

10 min 0.115 0.131 0.127 0.119

30 min 0.109 0.127 0.124 0.116

EIBL [%] 10 min 27.6 28.8 32.5 32.9

30 min 31.5 31.3 33.9 34.3

98

8 Behaviors of Component-similar Samples

8.1 Results of 90° V-bend samples

After the 500 h immersion tests, corrosion products mixed with salt were formed on the surface of the

V-bend samples. After cleaning and de-coating, macroscopic crack is not observed. However, in the

ground samples, the microscopic cracks located in the inner radius are found, as shown in Figure 8.1

(a). Figure 8.1 (b) and (c) exhibit the micrographs of cracks in the samples from CP1180+EG with

bending radius (R) of 4 mm and 5 mm respectively. When R is 4 mm, as shown in Figure 8.1 (b),

cracks from the inner surface are present. The maximum depth of the crack reaches 20 μm. The arrested

side of the cracks exhibits the branch-like structures. However, when R is 5 mm, (Figure 8.1 (c)), the

sample surface is still smooth, without crack.

(a)

(b) (c)

Figure 8.1: Micrographs of 90° V-bend samples from CP1180+EG after 500 h immersion test: (a) location of crack, (b) R=4 mm, and (c) R=5 mm.

The occurrence of cracks in V-bend samples is attributed to the co-operation of hydrogen attacking

through corrosion process and the spring-back of the material. In this study the bending process was

corrected to ensure the final angle of 90° in the sample. It means that the samples were bent to get an

angle smaller than 90°, and a spring-back occurred as the force was removed. After spring-back, the

residual stress on the inner side of bending radius exists. In the salt solution, the Zn-layer is firstly

dissolved. Once the substrate steel contacts to the corrosion medium, uptake of hydrogen is meanwhile

started. The residual tensile stress on the inner radius provides local expansion and facilitates the

absorption of hydrogen, leading to the initiation and propagation of cracks. In the automobile

8.2 Results of cross-die samples 99

components, the structure of tunnel is in the same situation. The initiation of hydrogen-induced cracks

is not allowed. According to the assumption above, the minimum allowed ratios of R/t (R: bending

radius, t: thickness) are calculated, which can be used as the Embrittlement Index (EIV). As summarized

in Table 8.1, different steel grades display similar values of minimum R/t. It can be concluded that the

1180 MPa steels are allowed to be bent until the angle of 90° in car body constructions, when R/t is

over 3.

Table 8.1: Comparison of minimum allowed punch radius (R) and thickness (t) of the investigated materials.

Material t [mm] Min. R [mm] Min. ratio of R/t

CP1180+EG 1.6 5 3.1

DP1180A+EG 1.5 5 3.4

DP1180B+GI 1.4 5 3.6

TBF1180A+EG 1.5 5 3.4

TBF1180B+EG 1.5 5 3.4

TBF1180C+EG 1.6 5 3.1

8.2 Results of cross-die samples

Table 8.2 gives the results of cross-die samples after immersion tests. It reveals that after 192 h

immersion in the 3.5%NaCl solution (pH4), DF is not detected in the lasered and welded samples

(Figure 3.18). However, in the samples with original punched edges from three Zn-coated TBP1180

steels, the cracks from free edges are observed after 17 h immersion, and then propagated in the location

of angle, as shown in Figure 8.2. However, the samples from DP1180A+EG and DP980+EG exhibit

no DF.

Table 8.2: Time to DF of the cross-die samples, immersion in 3.5%NaCl solution for 192 h.

Material drawn sample (free edges) drawn sample + cutting + welding

drawing depth [mm] time to DF [h] drawing depth

[mm] time to DF [h]

DP1180A+EG 17 n.d. 17 n.d.

TBF1180A+EG 21 17 21 n.d.

16 n.d. 16 n.d.

TBF1180B+EG 21 17 21 n.d.

16 n.d. 16 n.d.

TBF1180C+EG 23 17 23 n.d.

18 n.d. 18 n.d.

DP980+EG 20 n.d. 20 n.d.

15 n.d. 15 n.d. *n.d.: DF was not detected after 192 immersions.

100 Behaviors of Component-similar Samples

(a) Cross-die sample after immersion test (b) Detail from (a)

Figure 8.2: Delayed fracture in the cross-die sample from TBF1180B+EG (original punched edges, in NaCl solution).

The fracture surfaces from the failed cross-die samples from TBF1180 steels were investigated, as

shown in Figure 8.3. On the fracture surface the sub-cracks, mixed quasi-cleavage and intergranular

fracture mode are found, indicating the typical hydrogen-induced fracture. The intergranular facets

imply the significant hydrogen segregation on the grain boundaries, while the quasi-cleavage fracture

is probably initiated inside of new generated martensite and martensite/austenite interfaces.

(a) Cross-die sample after immersion test (b) Detail from (a)

Figure 8.3: Micrographs of fracture surface of cross-die sample from TBF1180A+EG (original punched edges, in NaCl solution).

The surfaces of cross-die samples are protected with Zn-coating, except the bare Zn-Fe-Zn edges. It is

understandable that DF occurs only on the bare edge. To explain the location of the onset of cracking,

a FE-model in LS-Dyna was analyzed. Thus the stress condition on the sample surface during forming

was simulated. Figure 8.4 displays the results from simulation on TRIP-assisted AHSS steel with a

tensile strength of 800 MPa. The blue area is the location of highest effective stress, where is most

crucial during forming process. On the vicinity of bare edges, especially on the shorter edges, a low

stress band colored with red/orange can be found. In this narrow band area the effective stress displays

below 400 MPa. An obvious stress gradient presents in the vicinity of the band, where the effective

8.2 Results of cross-die samples 101

stress drops from approximately 900 MPa to 200 MPa. Additionally, the vector of stress indicates that

the material in the low-stress bands is drawn perpendicular to the length of the bands.

In the NaCl solution, absorbed hydrogen atoms enter the sample through the bare edges and start

diffusion. It can be assumed that the local hydrogen concentration on the edges is much higher than that

in the inside parts. The hydrogen atoms can diffuse and redistribute in the sample. The stress gradient

plays an essential role in this redistribution process, as the driving force. The hydrogen accumulates at

the expanded region and degrades the critical stress to the DF. When the crack initiates, it extends along

the hydrogen concentrated band and arrests when the local stress is released by the crack.

Figure 8.4: Stress distribution on the cross-die sample of TRIP-assisted AHSS.

In the laser cut and additionally welded samples, DF is not observed. The main reason could be that in

the cutting process these stress gradient areas were cut off and the stress distribution on the vicinity of

edges is homogeneous. The after welding there are gaps between the cross-die and the lip, but the local

stress is not high enough to lead DF.

102

9 Discussion

9.1 Evaluation of testing methods to investigate HE of AHSS

In the present study different mechanical tests were carried out to characterize the Hydrogen

Embrittlement (HE) of Advanced High Strength Steels (AHSS). In addition, forming tests to study the

impact of strain state were conducted. Finally the mechanical tests were validated with car component-

similar samples. In this section, these tests are evaluated by concerning the main advantages (+) and

disadvantages (-) regarding procedure of the test and the significance of the results.

Slow strain rate test

In each steel in both the 1180 MPa group and the 980 MPa group, a hydrogen-induced

degradation of the strain property can be detected, which is considered as an Embrittlement

Index (EISSRT) to evaluate the susceptibility of the material to HE.

The hydrogen content can be modified by hydrogen charging. Combining with hydrogen

calibration, the inclining to HE at varying hydrogen contents can be evaluated.

PC controlled operation ensures the accuracy of the measurement.

The requirements of sample preparation are very rigid. The geometry of the sample strongly

influences the results. Even a 0.01 mm error in the width outside of the gauge length may cause

an unexpected localized break.

The Slow Strain Rate Test (SSRT) evaluates the Internal Hydrogen Assisted Cracking (IHAC) in the

materials. The hydrogen concentration was modified by the hydrogen charging process, and the samples

were loaded with a strain rate of 10-5 s-1, to ensure the diffusion and re-contribution of hydrogen in the

materials. In many investigations SSRT was used to evaluate the impact of hydrogen on the mechanical

properties [SUN15] [LI14] [LAU15] [REH14]. Concerning the strict requirements of sample

preparation, it is recommended that notched samples be used instead of standard tensile samples in the

future. The presence of the notch causes local stress concentration to control the position of fracture

during SSRT, however, increases the complexity of sample preparation.

Incremental loading test

Varying sample geometries are available. Hence, the as-delivered material and welded material

can be examined.

The critical stress to HE can be determined in 56 h, in which time three consecutive tests can

be conducted.

The sample preparation is simple.

9.1 Evaluation of testing methods to investigate HE of AHSS 103

In the current work, many materials failed in the constant loading procedure, where the stress

was held at 75% of Rm.

The Incremental Loading Test (ILT) evaluates the Hydrogen Environment Assisted cracking (HEAC)

in the materials. The ILT was carried out in the solution of 3.5%NaCl with pH4. Due to a strong reaction

with hydrogen in the test surroundings, Delayed Fracture (DF) occurred within a short amount of time.

Therefore, the threshold stress to delayed fracture in the ILT had been exceeded. In this case, the

embrittlement Index EIILT was calculated based on the applied stress and time to delayed fracture.

Thereby the HE behavior in the investigated AHSS was clearly evaluated. In the current work, the initial

load was 75% of Rm, and most 1180 MPa steels broke in a short time when subjected to this load (Figure

6.10). In order to investigate the threshold stress to DF in the ILT and obtain a better differentiation of

the results, a lower starting load is recommended. A start loading for future tests could be reduced to

50% or 60% of Rm of the materials.

Constant loading test

Threshold stress to hydrogen-induced DF of each material can be determined.

Sample preparation is simple.

The loading is achieved by adjusting the fixed bolt to control the distance between free sides of

the samples. This loading process is operated by hand. Hence errors may occur during the

testing procedure.

To determine the threshold stress the sample should be tested under different loads. The

required time is several weeks.

The Constant Loading Test (CLT) is considered as a simplified test, which separates each loading step

into an individual test. CLT evaluates the HEAC as well as ILT. In this work the time to DF and the

threshold stress in the corrosive medium (3.5%NaCl solution with pH4) were studied. Actually, in some

literature CLT has been usually used to determine the critical fracture stress of the materials with

varying hydrogen contents, for instance, after hydrogen charging [WAN06] [TAK14]. The results of

CLT are consistent with the results from ILT, but the testing time is much longer than that of ILT.

U-bending test

Sample preparation is simple.

The sample geometry is simple, thus different edge conditions are available.

The bending process is easy.

The investigation of the U-bend samples is very flexible. The samples can be tested in different

corrosive mediums to study the HEAC, and be available to examine the IHAC associated with

hydrogen pre-charging process. Additionally, many samples can be tested at the same time.

104 Discussion

The stress distribution of the U-bend sample is limited by the bending radius. The critical stress

to DF cannot be directly detected.

The results from pre-charged samples (referring to IHAC) and immersed sample (referring to HEAC)

in the U-bending test displayed a similar ranking of the different materials. The time to crack initiation

(tc) is used as the EI to evaluate the susceptibility to HE of different AHSS. In the research from Takagi

et al., U-bending tests were carried out to evaluate HE for two kinds of 1180 MPa steels [TAK10]. But

in their work the strain state on the bent surface was not mentioned. Rehrl reported the stress distribution

through the thickness, but did not compare the DF behavior with localized stress [REH13]. Actually, to

clearly evaluate the susceptibility of the materials to HE, a good understanding of the stress distribution

in the U-bend sample is very important.

Four-point-bending test

The sample has a simple geometry.

The four-point-bending test procedure is easy.

The bending grade of the sample can be adjusted continuously. Thereby the stress distribution

on the investigated sample is no longer limited by the regular bending radii.

The use of a four-point-bending apparatus limits the variation of the test conditions (for

example corrosive surroundings are not possible).

To determine the critical bending grade to delayed fracture, several series of test are needed.

The testing time is up to several weeks.

To understand the stress distribution in the sample, further analysis like simulation is required.

Four-point-bending test is considered as an additional test to the U-bending test. The main difference

between four-point-bending and U-bending is the way to tighten the sample. In four-point-bending test,

the critical bending grades on the investigated materials displayed similar values. Concerning the local

stress on the bent surface, the maximum allowed stresses reflect the susceptibility to HE. The ranking

of the investigated material is consistent with the EI in U-bending tests.

90° V-bend test

Component-similar test, in which the test condition is close to reality.

The sample preparation is complex, including the production of sample, and the 90° bending.

The testing time is up to 500 h.

To analyze the results, the samples need to be additionally prepared for microscopy.

In some materials, like MS-W1180+EG, the maximum bending radius of 5 mm is not possible.


The 90° V-bend samples were tested in a cyclic immersion test for 500 h. In the bent sample, the spring-

back force provides the stress, and the hydrogen uptake is associated with the corrosion process. After

the immersion test the samples need to be prepared before observation in a microscope. The minimum

allowed ratio of R/t (R: bending radius, t: thickness of the sheet sample) provides guidelines with radii

less sensitive to delayed fracture.

All the investigated methods to evaluate the HE of AHSS for automotive applications enabled a

differentiation of the various materials concerning their susceptibility to HE. In order to identify the

most suitable method, the differentiation of the results of the various tests has to be taken into account.

For each test method, the Embrittlement Index (EI) is defined, as summarized in Table 9.1. All the EIs

are inversely proportional to the susceptibility to HE.

Table 9.1: Summary of test methods and defined EI to investigate HE.

Test method Test condition Embrittlement Index (EI) Difficulty of preparation

/test

Required time of test

Slow Strain Rate Test (SSRT)

pre-charged (IHAC)

(loss(Af): loss of fracture strain

due to hydrogen) high 2 days

Incremental Loading Test (ILT)

NaCl solution (HEAC)

( : applied stress, tf: incubation

time to DF) low 3 days

Constant Loading Test (CLT)


(σmax: maximum allowed stress) low 2-3 weeks

U-bending

pre-charged (IHAC)

(tc: time to crack initiation) med 4 days


(tc: time to crack initiation) low 4 days

salt spraying surroundings

(HEAC)

(tc: time to crack initiation) low 10 days

corrosion surroundings

(HEAC)

(tc: time to crack initiation) low 6 weeks

Four-point-bending (4PB)


EI4PB = ymax

(ymax: maximum allowed bending grade)

low 2-3 weeks

90° V-bending NaCl solution (HEAC)

EIV = (R/t)min ((R/t)min: threshold ratio of R/t) high 3 weeks

Figure 9.1 illustrates the distribution of EISSRT (referring to IHAC) and EIILT (referring to HEAC). It is

found that the distribution of both EISSRT and EIILT display a similar ranking. The differences in materials

5 and 9 are probably due to the hydrogen concentration in the as-delivered states. In SSRT, the EISSRT

implies the degradation of total elongation as a function of hydrogen concentration. Although the ILT

106 Discussion

investigated the impact of environmental hydrogen on DF, the internal hydrogen in the as-delivered

state influences the EI by reducing the critical hydrogen content to DF.

Moreover, the required testing time is an important factor in the industry. The required time listed in

Table 9.1 is the running time of the test procedure, including the observation of DF under or after

loading, and the required time for repeat tests. The time for sample preparation was not considered.

Slow strain rate test Incremental loading test

Figure 9.1: Illustration of distribution of (a) EISSRT and (b) EIILT. Each number represents a steel grade, listed in Table 9.2.

To study the statistics of EI from each method, equation (9-1) is introduced to define the Index (A):

(9-1)

EIA is the value of EI of each investigated material by using the selected method. EImin and EImax

represent the minimum and maximum value from all the investigated materials respectively. The range

of Index (A) from 0 to 100 is divided into 10 intervals. The number of intervals, which are captured by

Index (A), represents the differentiation of the method. Table 9.2 shows the example to calculate the

differentiation of the SSRT. The differentiation of the test of U-bend samples in VDA corrosion

surroundings has a value 2 because only small amount of samples revealed DF; most of them underwent

the tests without showing DF.


Table 9.2: Calculation of differentiation of slow strain rate test.

Number 1 2 3 4 5

Material CP1180 DP1180A DP1180B TBF1180A TBF1180B

EISSRT 68.9 41.9 56.4 14 26.1

Index(A) 84.9 43.1 65.5 0.0 18.7

Captured Intervals 80-90 40-50 60-70 0-10 10-20

Number 6 7 8 9 10

Material TBF1180C MS-W1180 CP980 DP980 TBF980

EISSRT 33.7 75.1 78.7 70.6 35.8

Index(A) 30.4 94.4 100.0 87.5 33.7

Captured Intervals 30-40 90-100 90-100 80-90 30-40

Differentiation 7

The simplicity of the test method is evaluated in the terms of (I) simplicity of sample preparation and

(II) simplicity of the test procedure. For both factors (I) and (II), the simplicity is evaluated in three

grades: 1 represents the most complex, 2 stands for moderate and 3 means easy. In simplicity of the test

procedure, the efficiency of the test is concluded. The simplicity of the test method is calculated by the

equation (9-2):

Simplicity (test) = Simplicity (I) × Simplicity (II) (9-2)

In this case, the range of simplicity of test is 1-9. For example, the SSRT has a high requirement of

sample manufacture, edge condition and surface condition. An additional hydrogen charging process is

needed. The simplicity (I) is 1, indicating complex. The test procedure of SSRT is easy and computer-

controlled. SSRT can be carried out within several hours, but only one test can be carried out at a time.

Thus the simplicity (II) is evaluated as 2 (moderate). The simplicity of SSRT according to equation (9-

2) is 2.

Figure 9.2 illustrates the differentiation and simplicity of the investigated methods. The differentiation

is evaluated as 0 (poor) to 10 (good). The differentiation as calculated in this work strongly depends on

the number of investigated steels. In the current work, the differentiation is calculated based on ten

investigated AHSS alloys. But in the 90° V-bending test, only 6 kinds of 1180 MPa steels were

investigated. Thus the differentiation of immersion tests with V-bend samples has the lowest value of

all test methods; however this is not directly comparable to the differentiation of the other tests. The

simplicity, which is shown in x-axial, is in the range from 1, meaning the most complex, to 9, indicating

the easiest. The test methods which stay in the blue area are the suggested methods, displaying good

balance between differentiation and simplicity in this work.

108 Discussion

Figure 9.2: Illustration of differentiation and simplicity of test methods to evaluate HE of AHSS.

* In the V-bend + immersion test, only 6 materials with 1180 MPa Rm were tested.

Another important factor to evaluate the test methods is the reproducibility of the results. The deviations

of the results in different tests are compared. When the samples are produced accurately and the loading

is carried out by computer-controlled processes, the reproducibility of the tests is higher. In the SSRT,

the deviation in the loss of strain due to hydrogen charging stays at a small level, indicating good

reproducibility. In the ILT, there is a wide range in time to fracture of the samples from CP1180+EG

and MS-W1180+EG (CP1180+EG: 10-48 h, MS-W1180+EG: 33-53 h). The reason is assumed to come

from the sample geometry after punching. When the loading is carried out manually and the samples

are produced by punching, for example in CLT and the U-bending test with VDA condition, HEAC

takes place. It may occur that only partial samples crack due to HE even when the samples are in the

same test condition. However, in the NaCl solution and pre-charged condition, the U-bend samples

crack reproducibly because of the strongly attack by hydrogen from the solution and the high hydrogen

content after charging.

It can be identified that the U-bend samples combined with immersion tests in NaCl solution, in which

the susceptibility concerning HEAC is investigated, displays a good balance of a moderate

differentiation of 5 and a high simplicity of 9. The U-bending test with pre-charged samples concerning

IHAC also displays a good differentiation of 6 and a moderate simplicity of 6. The ranking of

susceptibility to HE of the investigated materials in U-bending test with pre-charged sample and in the

NaCl solution is similar in both tests. Due to this reason, as well as the easy sample preparation, the

immersion test with U-bend sample is recommended for evaluating the HE-susceptibility of AHSS in

industry. Additionally, in this test, different edge conditions can be investigated. To optimize the test

method and obtain a relationship between DF and the stress distribution, it is recommended that in

9.2 Evaluation of susceptibility to HE in different materials 109

future investigations different bending radii are chosen to enable different stress levels in the U-bend

samples. The simulation by FE-modelling can be done to provide the stress-strain condition in the

deformed samples.

Another test method to evaluate the HEAC worth paying attention to is the Incremental Loading Test

(ILT). ILT displays also a good simplicity of 6 but a more moderate differentiation of 4 than U-bend

samples in the immersion tests. Considering the short time of testing, ILT could be also a good selection

to evaluate the susceptibility to HE. Based on the current results, the ILT can be further improved by

lowering the loading at the beginning of the test to enable the study of critical stress of the material. In

order to improve the reproducibility of the results of ILT, the sample can be produced in an optimized

way to increase the accuracy in geometry. For instance, the outside contour of the sample can be firstly

cut by milling or water jet cutting, and then the hole in the middle can be produced by punching.

9.2 Evaluation of susceptibility to HE in different materials

Different AHSS grades with values of Rm of 980 MPa and 1180 MPa were investigated. According to

the characterization of microstructure via LOM and EBSD, the materials are divided into four

classifications: so-called CP steels with homogenous microstructure of bainite and tempered martensite;

DP steels with 10-35% ferrite and martensite or mixed phase of tempered martensite and bainite; TBF

steels with approx. 9% residual austenite, 6-19% ferrite, small amounts of martensite (maximum 5.3%

in TBF1180B+EG) and bainite; MS-W1180+EG is the hot rolled martensitic steel with 100% lath-like

martensite.

By comparing steels in the same classification, it is found that the 980 MPa steels are less susceptive to

HE than the 1180 MPa steels, implied by less strain loss in SSRT, a longer time to crack initiation in

bent samples and a higher threshold stress to DF in the four-point-bending tests. Figure 9.3 shows the

relationship between Rm· and EISSRT when the concentration of diffusive hydrogen in the steels is in the

range of 1.2 to 1.5 ppm. The EISSRT here illustrates the susceptibility to HE of different material, where

higher EISSRT indicates lower susceptibility/better resistance to HE. As an exception, the maximum H

content in TBF1180C+EG is only 1 ppm, even when pre-charged for 30 min. A possible explanation

could be that the low diffusion coefficient of TBF1180C+EG and oxide-layers on the sample surface

hinders the uptake of hydrogen.

From Figure 9.3 it is concluded that within a steel group, the susceptibility to HE increases with

increasing Rm. In the previous work from Loidl and Rehrl the same observation was made [LOI11]

[REH13]. One of the reasons for the higher strength of 1180 MPa compared to 980 MPa is the higher

strengthening accompanied by a higher dislocation density. Dislocations influence the susceptibility to

HE in two ways. Firstly during the slip of dislocations new vacancies will be created in the presence of

hydrogen, resulting in the annihilation of dislocations, which likely contributes to the property

110 Discussion

degradation [TAK08]. Secondly the dislocations travel in the steels, and pile up when they meet obstacle

or boundaries. The hydrogen will accumulate in the vicinities of the obstacle and boundaries, where the

high dislocation density localizes, and reduce the cohesion of the matrix [TAK12] [BAI16] [PAR15].

Hence, when the dislocation density in the steel is in a higher level after strengthening, the susceptibility

to HE is clearly increased.

Figure 9.3: Embrittlement Index in SSRT against tensile strength of the materials. Diffusive hydrogen

concentration is in the range of 1.2-1.5 ppm. (*the H-content TBF1180C is 1 ppm). Higher EI means better resistance to HE.

By comparing the results from investigations concerning Internal Hydrogen Assisted Cracking (IHAC)

(SSRT and pre-charged samples in U-bending test), the ranking of HE-susceptibilities of four

classifications of materials from high to low is:

TBF steel > DP steel > CP steel > MS-W steel

The results of the investigations concerning the behavior of Hydrogen Environment Assisted Cracking

(HEAC) are consistent with that of the IHAC-related tests. There are only some exceptions to this trend.

In ILT and CLT DP980+EG reveals higher susceptibility to HE than TBF980+EG, because of the higher

hydrogen concentration of DP980 in as-delivered state. Furthermore, when the sheet samples were

deformed to a certain deformation level and tested in moderate corrosive mediums, such as U-bend

sample in salt spray test or VDA corrosion test, some TBF1180 steel are slightly more resistance to DF

than the DP1180 steels. This is assumed to be due to either the differing stress level after bending or

the higher resistance to corrosion.

To discuss the susceptibility of different steels to HE, possible influences such as the volume fraction

of retained austenite and the grain size are discussed in this work. As investigated, the content of C


influences the HE-behavior by changing the diffusion coefficient and solubility of hydrogen [GUO12].

Equivalent Carbon content (CE) is discussed due to its effect on hydrogen-induced cracking. The

equation by Bleck is used to calculate the CE of the materials based on the weight percentages of the

elements listed [BLE12]:

(9-3)

In the HEAC mode, the resistance to corrosion determines the hydrogen uptake from the surroundings.

Therefore Pitting Resistance Equivalent Number (PREN) is used to evaluate the resistance to pitting

corrosion, which is calculated from the equation (9-4) [SED96]:

(9-4)

Table 9.3 summarizes the possible influences on HE on the investigated steels, including retained

austenite, grain size, CE, PREN and the content of Ti and Nb.

Table 9.3: Factors influencing the HE susceptibility of AHSS.

Material

Influence Factors Classification Retained

austenite [vol. %]

Grain size [μm]

CE [wt.%]

PREN [wt.%]

Ti+Nb [wt.%]

MS steel MS-W1180 - 0.9 0.39 0.24 0.044

CP steel CP1180 0.2 1.2 0.54 0.29 0.034

CP980 0.7 1.4 0.51 0.30 0.035

DP steel

DP1180A 0.4 1.4 0.52 0.22 0.027

DP1180B 0.3 0.8 0.68 0.69 0.050

DP980 0.7 1.4 0.63 0.71 0.024

TBF steel

TBF1180A 9.0 0.8 0.70 0.34 0.030

TBF1180B 8.6 1.1 0.66 0.38 0.009

TBF1180C 9.2 0.9 0.57 0.45 0.009

TBF980 8.9 1.2 0.64 0.29 0.030

MS-W steel

The results from the mechanical test to investigate the susceptibility to HE reveal that MS-W1180 steel

is less susceptive to HE, even than the CP1180 steel. As shown in Figure 4.5, MS-W1180 steel consists

of 100% lath-like martensite, thus the microstructure and the lath-boundary distribution are

homogeneous. As listed in Table 9.3, considering the sub-boundaries the grain size of MS-W1180

reaches 0.9 μm. The sub-boundaries hinder the dislocation movement during plastic deformation. As a

consequence the hydrogen atoms pile up at the boundaries [NAG14]. In many scientific works it has

been reported that the grain boundaries can suppress the local accumulation of hydrogen in hydrogen

diffusion process, resulting in less susceptibility to HE [TAK12] [BAI16] [PAR15]. The sub-boundary

112 Discussion

has a lower misorietation angle than grain boundary, but is assumed to influence the interaction with

slip system and reduce the susceptibility to HE as well. Nagao found that when the slip band interacts

with the sub-boundary-like lath structure the fracture surface reveals quasi-cleavage mode [NAG14],

which is in accordance with our observation.

According to the hydrogen calibration curves under electro-charging, as illustrated in Figure 5.2, after

the same charging time the hydrogen concentration in MS-W1180 is lower than that in CP1180. It can

be deduced that in MS-W1180 the hydrogen diffuses more slowly than that in CP1180. The research

of Loidl confirmed that the diffusion coefficient of hydrogen in Martensitic steel is much lower than

that in complex-phase steel [LOI14]. Additionally, the CE value of MS-W1180 is the lowest in the

investigated material (0.39%, as shown in Table 9.3), which can be also prove that the diffusion of

hydrogen is at a slow velocity.

CP steels

CP1180 displays a higher EI in the mechanical tests of this work, indicating a lower susceptibility to

HE compared to DP and TBF steels. Loidl et al. have reported similar observation when the steel

samples are loaded in H2 atmosphere [LOI11].

As illustrated in Figure 4.2, in the SEM images in the Nital-etched samples from CP1180, mixed bainite

and tempered martensite phases with boundaries (grain boundary and sub-boundary) and carbides have

been observed. The grain size of CP1180 is 1.2 μm and a sub-boundary structure has been found. The

grain boundaries and sub-boundaries in CP steels play an important role in HIC processes. Usually the

hydrogen segregated to the grain boundary causes intergranular fracture [PAR15]. But in the current

work, on the fracture surfaces of CP1180 quasi-cleavage fracture mode can be found Figure 6.5 and

Figure 6.19 (a). Because of the function of boundaries, CP1180 displayed higher EI in the mechanical

test indicating HE. The dispersed carbides are also seen as a reason for the low susceptibility to HE in

CP steels. According to the chemical composition, as shown in Table 4.1, it can be deduced that the

carbides in CP1180 are TiC. TiC acts as a beneficial trapping site for hydrogen atoms during the

hydrogen diffusion and trapping processes, due to its high required activation energy for hydrogen (

[VIY05][BEI09]. The homogeneously distributed TiC in CP1180 effectively prevents the local

hydrogen accumulation process. In the mechanical tests investigating both IHAC and HEAC, CP1180

displayed low sensitivity to HE.

DP steels

DP1180A and DP1180B display a higher susceptibility to HE than the CP1180 and MS-W1180 steels,

shown by a value of EISSRT in the range of 40-60% and a lower threshold stress to DF in the CLT and

the four-point-bending test.


The microstructures of DP1180A and DP1180B are shown in Figure 4.3 and Figure 4.6. In these two

materials, the volume fraction of ferrite reaches 16.4% and 12.1% respectively. The presence of ferrite

determines the heterogeneity of the microstructure. Although the dislocation density in ferrite is much

lower than that in the martensitic phases, a significant amount of lattice defects exist in the vicinity of

the martensite-ferrite boundaries. Thereby these localized lattice defects can cause a heterogeneous

distribution of diffusive hydrogen in ferrite. This fracture mode has been discussed by M. Calcagnotto

et al. and Koyama et al. [CAL10] [KOY14]. In the traditional dual-phase steel, the cracks initiate usually

at the martensite-martensite or martensite-ferrite-martensite boundary [RAM16]. Figure 9.4 illustrates

the crack initiation and propagation processes in DP steel due to enrichment of hydrogen: the crack

nucleates at the location of martensite-martensite boundary and then propagates and opens in ferrite

phase. Koyama et al. have researched the damage in hydrogen pre-charged DP steel and pointed out

that “hydrogen not only decreases the critical strain for the de-cohesion in martensite regions through

the Hydrogen Enhanced De-cohesion (HEDE) mechanism, but also promotes ferrite/martensite

cracking, ferrite/martensite boundary sliding and cracking of the ferrite grains.” [KOY14].

Figure 9.4: Schematic of crack nucleation, propagation and opening in dual phase steel. F: ferrite, M:

martensite, LS: localized slip [KOY14].

DP1180B consists of 13.8% quenched martensite, but displays lower susceptibility to HE than

DP1180A. By comparing the chemical compositions of both materials, it can be found that aside from

a similar content of Ti (0.025% in DP1180B and 0.023% in DP1180A), the content of Nb in DP1180B

is much higher than that in DP1180A (0.025% in DP1180B and 0.004% in DP1180A). The addition of

Nb promotes the formation of Nb(C, N), which is beneficial in decreasing the mobility of hydrogen

atoms in the steel, by trapping them in the Nb(C, N) irreversible trapping sites [ZHA15].

As reported, the addition of Cr can increase the resistance to pitting corrosion by the formation of a

passive oxide layer on the substrate [SPE06]. The content of Cr in DP1180B is 0.600%, compared to

that of 0.216% in DP1180A (as shown in Table 4.1), leading to a higher value of PREN. In the corrosive

medium, DP1180B is more resistant to pitting corrosion than DP1180A. In the corrosive medium,

pitting corrosion firstly occurs on the substrate of the steel or on the interface between Zn-coating and

steel, where the oxidation reaction on the iron occurs (Fe→Fe2+ +2e-). In the region close to the

114 Discussion

steel/solution interface, a great amount of Cr ions are generated. The hydrogen ions are enriched in the

vicinity of the locations of pitting corrosion. Thereby the hydrogen uptake is enhanced locally. A better

resistance to pitting corrosion weakens the accompanied hydrogen penetration into the material. As a

consequence, the DF due to HEAC in DP1180B occurs more moderately than that of DP1180A, which

is confirmed by ILT and U-bending test in NaCl solution.

In addition, the grain boundaries in DP1180B are assumed to be to the reason for its lower susceptibility

to HE than DP1180A. The grain size of DP1180B (0.82 μm) is finer than that of DP1180A (1.38 μm).

The explanation could be that the presence of more grain boundaries suppresses the local concentration

of hydrogen under deformation, thereby increasing the resistance to HE.

TBF steels

TBF steels have a total elongation over 10% when the tensile strength is 1180 MPa and over 15% when

the tensile strength is 980 MPa. The ductility of the material is due to the retained austenite and ferrite

phases. Meanwhile the susceptibility to HE of TBF steels is higher than that in other steels. The crystal

structure of austenite, a face-centered cubic (fcc) lattice, dictates the high solubility and low diffusivity

of hydrogen [BLE07] [LAU15]. The hydrogen effect in Transformation-Induced-Plasticity (TRIP)

steels has been investigated in many scientific works [GUO12] [LAU15]. It has been found that the

hydrogen solubility in austenite is much higher than that of phases with body-centered cubic (bcc)

structures like martensite and ferrite [ZHU14] [LAU15]. The phase transformation of retained austenite

decomposing into martensite can occur under plastic deformation [GUO12], leading to an increase of

hydrogen diffusivity [KIM12]. Thus the newly generated martensite is supersaturated with hydrogen

because of its lower hydrogen solubility than that of austenite, thus explaining the high susceptibility

of TRIP steels to HE [LAU15] [ZHU16] [ZHU13]. In this damage mechanism, ferrite does not play an

essential role, contrary to DP steel. Evidence suggesting the segregation of hydrogen to the grain

boundaries are the intergranular facets on the fracture surfaces [PAR15]. According to the research of

Zhu et al., in TRIP-assisted AHSS the supersaturated diffusive hydrogen segregates quickly to the misfit

strained locations between transformed martensite and austenite or inside of martensite blocks [ZHU14],

as illustrated in Figure 9.5.


Figure 9.5: Schematic description of initiation of sub-cracks in TRIP-assisted steels. Under loading

the hydrogen atoms segregate on the misfit locations between generated martensite and retained austenite, resulting in localized cracks.

The volume fraction of retained austenite in all the investigated TBF materials is in the range of 8.6%

to 9.2%. The TBF1180C shows a better resistance to HE than the other two TBF1180 steels, concerning

both IHAC and HEAC. The CE of TBF1180C is 0.57%, which is much lower than that of TBF1180A

(0.70%) and TBF1180B (0.66%). The PREN of TBF1180C is 0.45%, indicating a better resistance to

pitting corrosion comparing than TBF1180A (PREN=0.34%) and TBF1180B (PREN=0.38%). As a

consequence, TBF1180C reveals the longest time to delayed fracture in HEAC of the TBF steels.

The difference in HE susceptibility between TBF1180A and TBF1180B is insignificant due to a balance

between grain size, CE and PREN. In most cases, TBF1180B shows a better or similar resistance to the

degradation of properties and to DF. When the materials were examined to investigate their IHAC

resistance, the lower CE of TBF1180A appeared to be the main factor causing the lower diffusivity of

hydrogen. Therefore the re-concentration process occurs more slowly than that in TBF1180B. By

comparing the PREN, TBF1180B also indicates a better resistance to pitting corrosion when the

materials are exposed to corrosive mediums. In the CLT and ILT, TBF1180B reveals a lower resistance

to HE than TBF1180A. It is probably due to the high concentration of hydrogen in the as-delivered

material of TBF1180B. In the hydrogen enriched environment the hydrogen pick-up takes place locally

on the edges. The higher internal hydrogen content makes it easier for the critical fracture hydrogen

concentration to be reached. It is worth noting that the yield ratio (Rp0.2/Rm) of TBF1180B is 0.75, which

is lower than that of TBF1180A (0.85). It means that under the initial load of 75%Rm in the CLT, the

plastic deformation in TBF1180B is at a higher level, causing more free space in the lattice for hydrogen.

The strong hydrogen uptake from the NaCl solution is dominant; thus weakening influence from the

carbon equivalent and resistance to pitting corrosion.

116 Discussion

In summary, the susceptibility to HE is usually proportional to the tensile strength of AHSS, because

strengthening is usually associated with an increase in dislocation density. The susceptibility to HE is

influenced by various material-based parameters. To compare the different steel groups, various factors

influence the HE behavior of AHSS, like volume friction of retained austenite, chemical compositions

and grain size of the microstructure. Martensitic steel shows the lowest sensitivity in the four

investigated steel groups, which is attributed to its homogeneous fine lath-like structure and low CE.

The CP steel also displays low susceptibility to HE due to its homogenous microstructure of mixed

tempered martensite and bainite. Moreover, the dispersed TiC in CP1180 is assumed to reduce the

susceptibility to HE by the trapping effect on the mobility of hydrogen, thereby the resistance to DF of

CP1180 is increased. The DP steels show a higher susceptibility to HE than CP and MS-W steels. The

presence of ferrite in the martensite/tempered martensite structure increases the heterogeneity of the

microstructure. Aside from the phase fraction, the most dominant factor contributing to the higher HE

susceptibility of DP1180A is the coarser grain size than that of other materials. Meanwhile, the high

CE of DP1180B reduces its resistance to HE. The TBF steels reveal the highest susceptibility to HE of

the four steel groups. The explanation is the presence of retained austenite and a high CE value of the

material.

9.3 Influence of process-immanent parameters

9.3.1 Influence of welding

In the ILT in a NaCl solution, DF occurs in the welded samples when the loading was below the Rm of

the welded material. In DP1180A+EG and TBF1180A+EG steels, the time to fracture is very short,

even when the load stays at 75% of Rm, indicating high susceptibility to HE in the tested medium. This

result is consistent with previous work in the literature, in which it has been reported that after welding

the risk to HE of the material is increased [STO94] [PER15] [LOI14].

In the fracture surface analyzed by SEM it can be seen that the hydrogen-induced crack initiates from

the welded zone (Figure 6.16). The high susceptibility of the welded material can be explained by the

phase transformation during the welding process, which leads to an increase in hardness. As shown in

Figure 6.17, the hardness in welding zone of TBF1180A+EG is increased from 400-430 HV0.5 to 520-

570 HV0.5, and in the Heat Affected Zone (HAZ) the hardness is only 350 HV0.5. The increase in

hardness depends on the C content of the material, as illustrated in Figure 9.6. A higher C content leads

to a higher hardness in the welded zone. When the tensile test is carried out in air, the softened

microstructure in the HAZ results in fracture initiation locally. In contrast, when the welded samples

are exposed to a hydrogen-containing environment, reversible hydrogen embrittlement causes

unpredictable failure of the material under low loads [SUN15] [SHV00]. The diffusivity and solubility

of hydrogen are discretized in terms of microstructures in the divided zones after welding.

9.3 Influence of process-immanent parameters 117

Figure 9.6: Average values of hardness in the base material and the welded zone, relating to carbon content and

carbon equivalent.

In spite of a high local hardness/strength on spot welded zone, Liquid Metal Embrittlement (LME) also

contributes to the degradation of the properties of the welded material in hydrogen-assisted

environments. In resistance spot welding, the electrical current passes through the material leading to

internal heating by the Joule effect, causing the melting of the substrate steel and the zinc [RAO12].

The welded zone can be considered as a consequence of the solidification of the molten nugget. During

this heating-solidification process, the liquid zinc penetrates into the austenite grain boundaries. As

reported by L. Cho et al., the Zn-coating reacts with the steel and forms a mixed phase of the bcc solid

solution of Zn in α-Fe (α-Fe(Zn)) and Г-Fe3Zn10 phase during austenitization [CHO14]. Figure 9.7

schematically illustrates the formation of low strength structure α-Fe(Zn). In the current work, direct

cracking due to LME did not occur, but the presence of α-Fe(Zn) at the prior austenite boundaries is a

factor contributing to the high susceptibility of 1180 MPa steels to HE. As known, zinc has more

negative chemical potential than iron, which is prone to electron loss. Combined with a strength increase

in the welded zone as described above, once the welded material is loaded in the presence of hydrogen,

the hydrogen accumulation at the welded zone, especially at the prior austenite boundaries is more

pronounced than at the basic material, resulting a higher susceptibility to HE.

Figure 9.7: Schematic illustration of Zn grain boundary diffusion-mitigated phase transformation during resistance spot welding [CHO14]. γ-austenite, α´-martensite.

118 Discussion

In addition, after austenitizing, the solubility of hydrogen in the material increases. The amount of

hydrogen absorption during welding process depends on the hydrogen potential level. The hydrogen

dissociation from atmospheric moisture takes place, where the moisture converts to potential hydrogen

[VIY05]. After cooling, the hydrogen concentration of the welded zone has already changed. A higher

local hydrogen content can enhance the susceptibility to HE because the localized critical hydrogen

content required to induce the crack is more likely to be reached.

By comparing the rankings of susceptibility to HE of the welded material (Table 6.2) and as-delivered

material (Figure 6.13) in the same test surroundings (3.5%NaCl solution, pH4), it can be found that in

both tests the 980 MPa steels displayed lower sensitivity to HE than 1180 MPa steels. In the 1180 MPa

steels, the ranking of the sensitivity to HE of welded material is consistent with that of the basic

materials. Thereby the susceptibility to HE from high to low is: TBF1180A+EG > DP1180A+EG >

CP1180+EG. As an exception, the welded DP980+EG is less susceptive to HE than CP980, which is

opposite to the results of ILT with basic material. This is probably due to the high hydrogen

concentration in the as-delivered DP980+EG. In the ILT, the critical hydrogen content to induce the

cracking in the basic material is more easily reached. However, during the welding process, the

hydrogen uptake probably occurs in all the materials. The hydrogen content in welded DP980+EG is

therefore no longer outstanding. Other parameters such as PREN thus become more significant. The

high value of PREN of DP980+EG (0.71%) leads to a better resistance to HE after welding than

CP980+EG (PREN=0.30%).

9.3.2 Influence of the edge condition

Regardless of whether environmental hydrogen or pre-charged hydrogen is present, the initiation of DF

is always located on the edged because of the lower required energy at the free edges. When the sample

is completely coated with Zn-layer, as analogous to pre-charged and galvanized samples in SSRT or U-

bend test, Internal Hydrogen Assisted Cracking (IHAC) dominates. The crack is attributed to the

hydrogen redistribution along the stress gradient. But this is not enough to explain the DF in corrosive

surroundings. As reported, when the sample is completely coated with Zn-layer, the substrate steel is

protected from the hydrogen rich environment because the Zn-layer is a barrier to hydrogen absorption;

consequently, a property degradation has not been detected [MUE14]. However, in the situation of Zn-

coated sample with bare edges, the sandwich-like edges should be taken into account. When there is a

defect in the Zn-layer on the steel, or at the Zn-Fe-Zn sandwich-like edge, the hydrogen absorption is

enhanced in the beginning and reduced to a lower level after 4 days, when the metallic Zn disappears

[OTS13]. It is concluded that when galvanic coupling of Zn and steel is present, the steel is more

susceptive to HE than the blank steel due to the fact that zinc has a lower potential which causes much

hydrogen desorption on steel substrate [OTS13].


In the immersion test of the U-bend samples, the EIU-NaCl of DP and TBF 1180 MPa steels with punched

edges has a value below 8 h, indicating the high susceptibility to HE in the NaCl solution (Figure 6.29).

The edge condition impacts the HE behavior of the sample, primarily in terms of morphology. After

punching, the edge surface displays flush-cut surface and a fractured region (Figure 6.38). The

roughness of the flush-cut region reaches 0.34 μm while the roughness of the fractured region is

0.47 μm. The roughness of the edge surface is decisive to the local stress distribution at the crack tips

even when the samples are under the same loading. A higher roughness promotes micro-crack

nucleation. The presence of micro-cracks on the edge surfaces can cause a higher localized stress; thus

there is a stronger hydrogen re-contribution due to the higher driving force of the stress gradient. The

critical stress to DF is easier to be reached.

The punched edges are produced in a short time, but the resulting increase in temperature is not enough

to cause phase transformation. The hardening on the punched edge is due to the localized mechanical

loading. On the punched edges, as illustrated in Figure 6.39, the edge surfaces are hardened by work

hardening. According to Table 6.9, in TBF1180A+EG the increase in hardness due to punching has the

highest value of all the investigated steels (∆HV0.3 is 20.9%). The susceptibility to DF is also the highest.

The material MS-W1180+EG, in which DF has not been observed in the immersion test with U-bend

samples with punched edges, has the lowest value of ∆HV0.3. It can be concluded that the hardening on

the edge is related to the carbon equivalent (CE) of the material. The susceptibility to HE of the punched

sample is correlated to the hardening due to punching. This is likely to be due to the increase of

dislocation density during the hardening process, thus the amounts of hydrogen trapping sites are

increased. Both micro-cracks on the edge surface and the high local stress contribute to causing the

earlier failure or greater degradation of the property due to HE. When the edges are produced by

punching, the EI of the U-bend samples (EIU-NaCl) from different materials is reciprocal to the hardening

because of punching, where:

HV0.3, MS-W1180+EG < HV0.3, CP1180+EG < HV0.3, DP1180A+EG < HV0.3, TBF1180A+EG

EIU-NaCl, MS-W1180+EG /EIU-NaCl, CP1180+EG > EIU-NaCl, DP1180A+EG > EIU-NaCl, TBF1180A+EG

The laser-cut edges of AHSS have a smooth edge surface because the high temperatures causes the edge

of the metals to melt. However, it is worth noting that the U-bend samples with laser-cut edges from

TBF1180 steels also show a short time to DF in the NaCl solution (Table 6.6). The high susceptibility

to HE in the laser-cut U-bend samples can be attributed to the hardening on the edge surface, which is

probably due to phase transformation to martensite. The maximum hardness attained depends on the

CE, whereby the higher CE causes a higher hardness at the edge (Figure 9.8). Since the morphology of

laser-cut edges is already in the ideal condition, the susceptibility to HE of the laser-cut sample is

determined by the local hardening. The resistance to HE decreases with increasing strength or hardness

of the material. Once the stress state and hydrogen concentration reach the threshold value to fracture,

DF occurs, similar to U-bend samples from TBF1180A+EG in the immersion test. After laser-cutting,

120 Discussion

the maximum hardness on the edge of TBF1180A+EG reaches 557 HV0.3, with an increase in hardness

of 36.9%. In contrast, DP1180A+EG also showed a high sensitivity to HE in the punched samples. But

after laser-cutting the hardness of the edges in DP1180A+EG increased only to 479 HV0.3. Combined

with the low roughness of the edge surface, DP1180A+EG with laser-cut edges showed good resistance

to HE.

Conversely, the milled edges have a moderate roughness. During milling, a cooling agent is used to

prevent high temperatures on the sample and the miller. Thereby the heating effect during milling is

minimized; no phase transformation occurs. The milled edges do not display local hardening on the

edge surfaces. This is consistent with the observation that samples with milled edges demonstrate the

lowest susceptibility to HE.

Figure 9.8: Maximum hardness due to cutting and its relationship with the carbon equivalent of the materials.

In summary, the edge condition affects the HE behavior of the material in terms of the

morphology/roughness of the edge surface and the hardening during cutting. The roughness of the edges

is determined by the cutting process. However, the hardening on the cut edges is dominated by the

chemical composition of the material, or more specifically, the carbon equivalent. After cutting, if the

edge surface shows high roughness or significant hardening, the sample is more sensitive to hydrogen-

induced delayed fracture.

9.3.3 Influence of strain state

In the studies of the Forming Limit Curve (FLC), the strain state plays a critical role in the degradation

of formability. The FLC is one of the most commonly used processing criteria, representing the maximal

in-plane formability [DEN12]. In this criterion, the rupture is caused by the voids generated inside the

material [MUE06] [BLE09]. The voids are firstly nucleated at vulnerable locations such as phase

boundaries, grain boundaries and inclusions as a result of dislocation motion and piling up. Under stress,

the void nucleation can continue to grow and merge through different mechanisms. No matter which


mechanism, the local stress in the vicinity of the voids is not constant but rather a gradient. The absorbed

hydrogen is activated from the reversible trapping sites, diffuses along the stress gradient from high to

low stress, accumulates at the tips and influences the formability through Hydrogen Enhanced Plasticity

(HELP) or Hydrogen Enhanced De-cohesion (HEDE) mechanisms.

The strain state is determined by the major deformation degree (φ1) and minor deformation degree (φ2)

of the sheet sample. In the uni-axial strain state the sheet material is in the uniaxial tension stress, where

φ1= -2φ2. In the state of plane strain, the sheet sample does not deform in the direction of minor

deformation (φ2=0). In the biaxial strain state, the sheet material is equally stretched in both major and

minor stress directions; thereby the φ1 is equal to φ2, in which the sheet is deformed with uniform sheet

thickness reduction [BLE09]. It is noteworthy that the sample is deformed in Nakajima test with a high

deformation velocity (the punch velocity 2 mm·s-1). The hydrogen re-contribution can occur only

locally, in limited locations. The hydrogen diffusion through the thickness of the material when the

material is drawn in a biaxial strain state is assumed to be stronger. Thus the degradation of formability

is most obvious in the biaxial strain state [GAO16a] [GAO16b].

The influence of the strain state on the bendability of pre-strained samples was studied by the

determination of the Bending Limit Curve (BLC). In the previous studies, it is concluded that the

hydrogen concentration increases with increasing pre-straining [LI14]. The explanation is probably due

to the generation of defects, which act as trapping sites for hydrogen atoms; and/or the interaction

between hydrogen and dislocations. The strain in the tensile direction is more significant than that in

the compressive direction, since the expanded lattices provide the space for hydrogen atoms. The

hydrogen concentration of the pre-strained samples reveals that the biaxial pre-strained samples have a

higher hydrogen concentration. It is assumed that pre-strain in the biaxial state provides more defects

for hydrogen because of the expanded lattice. The bending fracture is dominated by the shear stress on

the outside of bending edge, where the local stress and strain display the maximum values in the

thickness direction. Because of the punch velocity, (20 mm·s-1), the diffusion of hydrogen occurs within

a limited distance. The higher hydrogen concentration leads to a stronger degradation of the bendability

[GAO16b].

9.3.4 Influence of corrosive surroundings

U-bend samples were tested in three different corrosive surroundings: 3.5% NaCl solution (pH 4), 5%

NaCl spray surroundings, and VDA233-102 corrosive surroundings in cyclically changed 1% NaCl

spraying, humidity and temperature (Table 3.4). The EI is defined as the time to crack imitation (tc) to

evaluate the susceptibility to HE, concerning the HEAC. In the U-bend samples in the NaCl solution,

the Hydrogen-Induced Cracking (HIC) initiated in DP and TBF 1180 MPa steels after 3-8 h immersion

(Figure 6.29). However, in the salt spray surroundings, HIC was found the samples from the same

122 Discussion

materials after 1-3 days (Figure 6.35). In the VDA corrosive surroundings, the DF was observed after

1-2 weeks in small amount of samples (Figure 6.37).

In the NaCl solution, a high concentration of Cl- is present, which can destroy the oxide-layer on the

substrate, resulting in pitting corrosion. The oxidation reaction of Fe takes place. As a consequence,

hydrogen absorption occurs (Figure 2.8), leading to HIC. In the solution without blowing gas, a very

small fraction of O2 dissolves in the solution, hence the surroundings are considered be anaerobic. As

investigated, the substrate of the steel becomes more susceptible to HE when the surroundings are

anaerobic [MAY15]. When the U-bend samples are placed in the salt spray chamber, the surroundings

are regarded as aerobic; only when the O2 is exhausted does it becomes anaerobic. Though the Cl- is

enough in the salt spray surroundings, the reduction reaction 2H2O +O2 +4e- 4OH- is preferred,

thereby the Volmer reaction in hydrogen adsorption (H2O +e- Had +OH-) is compressed, resulting in

the longer time to HIC than in the NaCl solution. In the VDA corrosion test, the amount of Cl- is at a

small level due to the cyclic spraying of 1% NaCl solution (1 h, 3 times per week, Figure 3.9). In this

case, the destruction of the oxide-layer on the substrate surface takes a longer time than that in the NaCl

solution and spraying conditions. Additionally, the large chamber of VDA test allows that the test

surrounding is aerobic for a long time. All these reasons lead to a longer tc in the VDA corrosion test.

Changes in humidity and temperature are likely to affect the HE behaviors slightly.

In general, the HE of the materials is related to the corrosive surroundings. The corrosive ions and

aerobic/anaerobic conditions are the predominant factors influencing the HEAC related HE-

susceptibility of the materials.

9.4 Validation of mechanical tests with component-similar tests

As described in Section 3.6, HE-related tests on component-similar samples were conducted to validate

the results gained on regular sample geometries with geometries which reflect the shape of car body

components. The 90° V-samples were chosen to represent the conditions that are present on radii found

on automotive components, like tunnel parts, which are produced mainly by bending operations. The

choice of cross die samples is related to deep drawn parts like B-pillars. The testing of these samples

took place in corrosive mediums which were supposed to reflect realistic or even worse conditions than

what is typical during the production and life of the car.

The V-bend sample enables the identification of the minimum ratio of R/t (R: bending radius, t:

thickness of sheet material) up to which DF takes place. This information can be used to design the radii

of component parts providing higher safety against DF. In the cross-die samples, the maximum allowed

drawing depth in each steels was reached, then the samples were immersed in a NaCl solution. The

TBF1180 steels with original punched edges displayed delayed fracture, while in the other steels

(DP1180A+EG and 980 MPa steels) delayed fracture was not observed.

9.5 Consequences for using AHSS for car body applications 123

Although the differentiation of the HE behavior of the different steels in component-similar tests is not

as good as in the regular mechanical tests to investigate the HE-susceptibility, the results of the cross-

die samples are consistent with mechanical tests like the ILT and the U-bending test, where the

TBF1180 steels showed the highest susceptibility in the corrosive environment.

The corrosive conditions of component-similar samples are comparable to that in the HE-related

mechanical tests, which enables a link between the results from HE-related mechanical tests to the

impact of hydrogen on real components. The results from regular HE-related mechanical tests showed

that the effect of hydrogen on the investigated steels is strong, e.g., in U-bending test in the condition

of NaCl solution the HIC occurred in 2 h in DP1180A+EG. The tests on the cross-die samples, however,

only displayed failure on the edges on the samples from TBF1180 steels. It can be concluded that the

regular HE-related mechanical tests in the current work are more stringent than the real components.

The results from regular HE-related mechanical tests display similar tendency but significantly higher

differentiation than component-similar tests.

In summary, the HE-related mechanical tests investigated in the current work are appropriate to evaluate

the HE-behavior of real parts. In the material release process, a material which is already safely applied

in a component can be selected as a reference to validate the applicability of the new material by

comparing the HE-susceptibility. The threshold value in the HE-related mechanical tests can be

validated by the reference material. Additionally, the cross-die samples can be also investigated to

ensure the safe usage of advanced high strength steels in the car bodies.

9.5 Consequences for using AHSS for car body applications

The factors contributing to the occurrence of Hydrogen-Induced Cracking (HIC) can be classified into

three groups: material, hydrogen and stress. Based on the results and discussion in the current work,

Figure 9.9 illustrates a summary of the impact factors in each group.

Figure 9.9: Summary of the factors contributing to hydrogen embrittlement.

As a consequence, the proposed methods to improve the resistance to Hydrogen Embrittlement (HE)

and thereby increase the safety of usage AHSS for car body applications are summarized below:

124 Discussion

Optimize material selection

The application of AHSS in the body-in-white structure usually depends on the strength, ductility and

weldability of the material. However, is also important to take into account the susceptibility to HE of

high strength steels. When the strength and ductility levels are comparable, the material which is less

susceptive to HE takes precedence.

The material-related factors influence the susceptibility to HE in various ways. The presence of retained

austenite ensures a good performance in ductility, but meanwhile increases the susceptibility to HE, as

observed in the TBF steels in this study. To optimize the material selection for body-in-white

construction, the fraction of retained austenite should be minimized. The susceptibility to HE can be

reduced by increasing the number of effective grain boundaries using a fine-grained microstructure.

Hence in the material selection procedure, an AHSS with a homogenous, fine grain size takes

precedence. Additionally, the HE-susceptibility is related to the chemical composition of the material.

Once the steels are in the same classification, e.g., DP steel, the material with a lower CE has a higher

resistance to HE. The addition of alloying elements, such as Ti and Nb, can reduce the susceptibility to

HE, because the formation of carbides of TiC or NbC hinders hydrogen diffusion and re-contribution

by the trapping effect. In most cases the hydrogen uptake occurs accompanied with corrosion processes.

Thus the addition of Cr, Mo, and N can decrease the HE-susceptibility by decelerating the adsorption

of hydrogen with increased resistance to pitting corrosion.

Prevent hydrogen uptake

When the hydrogen content in the as-delivered material is increased, the potential to HE is also

increased because the critical hydrogen concentration to HIC becomes more easily to be reached. The

results of the investigations on IHAC (mechanical tests on pre-charged materials) also provide evidence

that higher hydrogen content leads to higher susceptibility to HE. Therefore, the first solution to improve

the resistance to HE is minimization of the hydrogen content in as-delivered material from the steel-

making process.

Aside from the internal hydrogen in the steels, hydrogen from the environment can also cause a strong

enrichment of hydrogen in the material and lead to HIC. As discussed in section 9.3.4, the hydrogen

concentration is always accompanied with corrosion reactions. Once the corrosive ions enrich the

surroundings, especially when the corrosive medium is anaerobic, HE becomes critical to the

components (section 9.3.4). The protective layer in the body-in-white component plays an important

role in increasing of resistance to HE. In most vehicles Zn-coated steel sheets are used. The Zn-coating

can provide the first protective layer to suppress the uptake of hydrogen and protect the steel substrate

from corrosion. The following E-coating and painting can also enhance the protection to the substrate.

However, during the car making process the edges and welded parts are exposed to air and the pre-

treatment solution before E-coating. Thus external methods, such as sealing the exposed edges with


wax or protecting with oil can protect the susceptive parts from direct contact with the possible corrosive

surroundings, thus improving the resistance to HE of components.

Another vulnerable factor in car components is the gap between the components. Moisture can easily

be retained in these small gaps. The corrosion in this kind of location is strong. In the case of defects in

the Zn-coating, hydrogen uptake may occur. In the construction process the amount of gaps in the car

body design should be reduced to avoid local corrosion. Sealing fine gaps can protect the components

from the corrosive surroundings, thus the hydrogen uptake can be reduced and the susceptibility to DF

of the body-in-white is decreased.

Avoid stress concentration

It has been confirmed that a stress gradient acts as the driving force of the diffusion of absorbed

hydrogen and the consequent re-contribution leads to HIC. To improve the resistance to HE, the stress

distribution should be paid attention to. Based on Figure 2.6, the maximum increase of stress takes

place during the forming of the component parts. To avoid critical stress leading to DF, the local stress

concentration needs to be kept below the yield strength. For the bending and hemming parts, high local

stresses can be reduced by avoiding a small bending radius. As studied in the 90° V-bend samples, the

steel sheet has an allowed minimum bending radius, with which the V-bend sample does not reveal HIC

on the inner radius, when the sample is exposed to corrosive mediums.

In some special regions like cut edges and welded parts, which are susceptive to HE, the applied stress

needs to be minimized. The parts with edges or welding should be designed far away from the highly

deformed regions, particularly under tensile stress, to avoid high local hydrogen concentration. In some

cases, it is not possible to change the location of the edge. Here it is particularly important to optimize

the cutting method. As investigated, punched edges reveal the highest susceptibility to HE compared to

laser-cut and milled edges. Avoiding punched edges and preferring laser-cut or milled edges can

improve the condition of the edge surface, thus reducing the local stress concentration at the micro-

crack tips on the edge surfaces. It is worth noting that different materials reveal different DF behavior

on the laser-cut edges. TBF1180 steels, with laser--cut edges show also high susceptibility to delayed

fracture because of their microstructure and chemical compositions. To improve the resistance to HE,

theoretically, the edge of TBF1180 should be milled to minimize work-hardening and ensure a

homogeneous surface condition, although it is acknowledged that milling is not feasible in an industrial

setting.

Tensile stress is more harmful than compressive stress in most materials subject to HIC, as investigated

by Chu et al [CHU84]. In this work it is found that when the strain state is in biaxial, where the major

and minor stresses are both tensile stresses, the degradation of the properties due to hydrogen

enrichment is more significant than other in other strain states. It is assumed that avoiding high local

126 Discussion

tensile stresses, or the modification the tensile stress to become compressive stress, can prevent the

component from cracking due to HE.

In summary, a component part will be at risk of HIC once the mentioned three factors (material,

hydrogen and stress) are in the critical level at the same time as illustrated in Section 9.5. To reduce the

susceptibility to HE to a safe level, at least one of the three factors should be minimized. This is the

most essential guideline for enabling the safe, widespread use of AHSS in car-making, enabling

reduction in weight, and contributing ultimately to the goal of reducing CO2 emissions.

127

10 Conclusion

The current work investigated the Hydrogen Embrittlement (HE) behaviors of different Advanced High

Strength Steels (AHSS) by various hydrogen-related mechanical tests and component-similar samples.

On the basis of the obtained results the following conclusions can be drawn:

Suitable test method to evaluate the susceptibility to HE

Various mechanical tests to evaluate the reduction in properties due to hydrogen enrichment were

compared in the terms of differentiation of the results and the simplicity of the test method and sample

preparation. The reproducibility of the test was also taken into account. The tests related to Hydrogen

Environment Assisted Cracking (HEAC) (e.g., immersion test with U-bend samples) and Internal

Hydrogen Assisted Cracking (IHAC) (e.g., slow strain rate test) displayed a similar ranking in the

susceptibility to HE. This demonstrates that it is not necessary to test both HEAC and IHAC, but that

one of these types of test is sufficient. In an industrial setting it is recommended to conduct a HEAC-

related test rather than an IHAC-related test, due to the simplicity of sample preparation.

The immersion test of U-bend samples in NaCl solution is identified as a suitable method to investigate

HE of AHSS. This method has the highest value of simplicity in all investigated methods, combined

with a good differentiation of 5 and high reproducibility. In addition, the investigation of different edge

conditions is possible.

Another appropriate method is the Incremental Loading Test (ILT). ILT requires a short time to evaluate

the susceptibility to HE (max. 56 h), and meanwhile displays a high value of simplicity. Although the

differentiation of ILT was moderate, further optimization in future work could be done. It is

recommended to reduce the initial load from 75% to 50% or 60% of Rm. The reproducibility of ILT can

be further improved by optimizing sample preparation, such as cutting the outside contour with other

methods such as milling and water jet cutting instead of punching.

The results from hydrogen-related mechanical tests are consistent with that of the investigations on

component-similar samples, although the occurrence of HE was less pronounced in the component-

similar samples. The test conditions of the hydrogen-related mechanical tests were more sensitive to

HE than the component-similar samples, which is significant for the material release process. The

selection of reference material or identification of the threshold value in the HE-related mechanical tests

can validate the HE-susceptibility for further applications, ensuring the safe usage of AHSS in car body

construction.

128 Conclusion

Material-associated influences on HE

The 1180 MPa steels are more susceptible to HE than 980 MPa steels, owing to the higher dislocation

density after strengthening. The ranking of susceptibility to HE of the 1180 MPa steels from high to

low is:

TBF steel > DP steel > CP steel > MS-W steel

In the TBF steels the presence of approx. 9% retained austenite is the main factor causing the highest

HE-susceptibility. In the DP steels, the ferrite increases the heterogeneity of the microstructure and

meanwhile hinders crack propagation, decreasing the susceptibility to HE. The higher resistance to HE

in CP steels is attributed to the homogenous fine microstructure and dispersions of carbides. The highest

HE-resistance in MS-W1180 is due to the lath-like microstructure and the low C content of the material.

The material-related influences on the susceptibility to HE can be classified into two groups, depending

on their positive or negative effect on the HE-susceptibility. Briefly, HE-susceptibility is proportional

to the volume fraction of retained austenite (γ), grain size and the content of C, as well as the Carbon

Equivalent (CE) of the material. Once retained austenite is present, the strain-induced phase

transformation may take place, generating martensite (α'). Owing to the difference of hydrogen

solubility in γ and α' the hydrogen supersaturates in α', and concentrates in the misfit region between

these two phases, resulting in HIC. The grain size influences the HE-susceptibility by suppressing

hydrogen diffusion and accumulation through the irreversible trapping in grain boundaries and sub-

boundaries. Increasing C content as well as CE reduces the hydrogen solubility and increases the

hydrogen diffusibility, thereby increasing the HE-susceptibility.

In contrast, the addition of Ti and Nb in the steel can improve the resistance to HE, since TiC and NbC

function as irreversible trapping sites and are favorable for hindering hydrogen re-contribution. The

addition of alloying elements Cr, Mo and N are also advantageous in decreasing the susceptibility to

HE. They increase the resistance to pitting corrosion and delay the hydrogen uptake when the material

is exposed to corrosive surroundings.

Production-associated influences on HE

After spot welding, the welded material is particularly susceptible to HE when the sample is exposed

to corrosive mediums like NaCl solutions. The unpredictable failure may initiate in the welded zone,

where the microstructure is strengthened by welding. The formation of Zn-Fe mixture phase at the prior

austenite grain boundaries due to liquid metal embrittlement is also a reason for the high susceptibility

to HE. Additionally, the welding process itself can cause further hydrogen absorption, thereby

increasing the risk of the hydrogen content reaching the critical concentration for HE.

HE-susceptibility is related to the edge condition of the sample. The bare Zn-Fe-Zn edge is sensitive to

the initiation of delayed fracture once anaerobic corrosive surroundings are present. The punched edges


are susceptible to HE, since the work hardening on the edges increases the dislocation density and the

coarse edge surfaces enables micro-crack initiation. TBF 1180 samples with laser-cut edges are also

critical to HIC because the laser-cutting process hardens the edges by phase transformation. Therefore,

the milled edges offer the highest resistance to HE, owing to a smooth edge surface and homogenous

hardness distribution in the transverse surface.

A forming path in the biaxial strain state is more susceptible than that in uni-axial and plane strain states,

because the expanded lattice facilitates the re-contribution of hydrogen. Similarly, pre-straining in the

biaxial strain state facilitates hydrogen adsorption, resulting in more degradation of bendability.

Consequences for using AHSS for car body applications

To enable the safe application of AHSS in car bodies, the possible following three factors must be taken

into account:

i) Material: microstructure, chemical composition, and carbides (NbC, TiC)

ii) Hydrogen: internal hydrogen and hydrogen from surroundings

iii) Stress: parts subject to local tensile stress

According to the requirements on the material strength and ductility, at least one of the three factors

should be optimized to reduce the susceptibility of component parts to hydrogen embrittlement.

In summary, this research work evaluated different testing methods to investigate hydrogen

embrittlement of advanced high strength steel sheets and gave recommendations to select and improve

the suitable methods. The results identified the material-associated factors and clarified the critical

factors in car manufacture and during the life of the car which could lead to hydrogen embrittlement,

whether from internal or environmental hydrogen. The current thesis made recommendations for

materials selection and car manufacture. The results of the work provide a scientific basis for the

industrial use of advanced high strength steels in car body constructions, including essential guidelines

to avoid the occurrence of hydrogen embrittlement in safety-relevant components.

130

Reference

[ALB03] S.K. Albert, V. Ramasubbu, Influence of alloying on hydrogen assisted cracking and diffusible hydrogen content in Cr-Mo steel welds. Academy Proceedings in Engineering Sciences 28 (2003) 383-393.

[ANG06] J. Angeli, E. Füreder, M. Panholzer, A.C. Kneissl, Ätztechniken für die Phasencharakterisierung von niedriglegierten Dual-Phasen- und TRIP-Stählen. Practical Metallography 43 (2006) 489-504.

[ASTM08] ASTM G5-94 Standard reference test method for making potentiostatic and potentiodynamic anodic polarization measurements. ASTM, 2008.

[BAI16] Y. Bai, Y. Monotan, M.C. Chen, A. Shibata, N. Tsuji, Effect of grain refinement on hydrogen embrittlement behaviors of high-Mn TWIP steel. Materials Science and Engineering A 651 (2016) 935-944.

[BHA05] D. Bhattacharya, Developments in advanced high strength steels, Proceeding of the fifth International Conference on HSLA Steels, 2005, Sanya, China, 1-7.

[BEI09] B. Beidokhti, A. Dolati, A.H. Koukabi, Effects of allying elements and microstructure on the susceptibility of the welded HSLA steel to hydrogen-induced cracking and sulfide stress cracking. Materials Science and Engineering A 507 (2009) 167-173.

[BIL13] E. Billur, T. Altan, Three generations of advanced high-strength steels for automotive applications, Part 1 - The first generation. Stamping Journal, November-December (2013) 16-17.

[BIL14a] E. Billur, T. Altan, Three generations of advanced high-strength steels for automotive applications, Part 2 - The second generation. Stamping Journal, January-February (2014) 12-13.

[BIL14b] E. Billur, T. Altan, Three generations of advanced high-strength steels for automotive applications, Part 3 - The third generation. Stamping Journal, March-April (2014), 12-13.

[BLE07] W. Bleck, Material Science of Steel. Textbook for RWTH students, 2007.

[BLE09] W. Bleck, Material Characterisation. Texbook for RWTH students, 2009.

[BLE12] W. Bleck, Spezielle Werkstoffkunde der Stähle für Studium und Praxis, 2012.

[BLE14] W. Bleck, Advanced high strength steels for the automotive industry-from microstructures to nanostructures. Stahl und Eisen 134 (2014) 25-33.

[BIR90] H.K. Birnbaum, Mechanism of hydrogen related fracture of metals. In Hydrogen Effects on Material Behavior, Eds. N. R. Moody and A. W. Thompson, Warrendale PA and The Metallurgical Society Inc, 1990, 629-660.

[BIR95] H.K. Birnbaum, Hydrogen related second phase embrittlement of solids. In Hydrogen Embrittlement and stress Corrosion Cracking, Eds. R. Gibala and R. F. Hehemann, ASM International, 1995, 153-177.

[BOU11] O. Bouaziz, S. Allain, C.P. Scott, P. Cugy, D. Barbier, High manganese austenitic twinning induced plasticity steels: a review of the microstructure properties relationships. Current Opinion in Solid State & Materials Science 15 (2011) 141-168.

[CAL10] M. Calcagnotto, D. Ponge, E. Demir, D. Raabe, Orientation gradients and geometrically necessary dislocations in ultrafine grained dual-phase steels studied by 2D and 3D EBSD. Materials Science and Engineering A 527 (2010) 2738-2746.

131

[CHA03] G. Chalaftris, Evaluation of aluminum-based coatings for cadmium replacement. PhD thesis, Granfield University, 2003.

[CHE14] S. Chen, M. Zhao, L. Rong, Effect of grain size on the hydrogen embrittlement sensitivity of a precipitation strengthened Fe–Ni based alloy. Materials Science and Engineering A 594 (2014) 98-102.

[CHO81] W.Y. Choo, J.Y. Lee, C.G. Cho, Hydrogen solubility in pure iron and effects of alloying elements on the solubility in the temperature range 20 to 500°C. Journal of Materials Science 16 (1981) 1285-1292.

[CHO14] L. Cho, H. Kang, C. Lee, B.C. De Cooman, Microstructure of liquid metal embrittlement cracks on Zn-coated 22MnB5 press-hardened steel. Scripta Materialia 90-91 (2014) 25-28.

[CLA11] O. Clauzeau, N. Le Bozec, D. Thierry, Hydrogen embrittlement and delayed fracture of advanced high-strength steels for the automotive industry. Research report, Institut de la Corrosion 2011: 4.

[COO11] B.C. De Cooman, K. Chin, J. Kim, High Mn TWIP steels for automotive applications. In New Trends and Developments in Automotive System Engineering, Eds. M. Chiaberge, InTech, 2011, 101-128.

[COR14] D. Cornette, C. Allely. P. Dietsch, B. Weber, T. Sturel et al. No detrimental impact of car manufacturing process and simulation of vehicle in-service conditions on DP1180 hydrogen embrittlement. Proceedings of 2nd Steely Hydrogen Conference, 2014, Gent, 45-58.

[CUR10] S. Curtze, V.T. Kuokkala, Dependence of tensile deformation behavior of TWIP steels on stacking fault energy, temperature and strain rate. Acta Materialia 58 (2010) 5129-5141.

[DEN11] R. Denninger and M. Liewald, C. Held, M. Sindel, Investigation on bendability of lightweight materials for various load paths. Proceedings of ESAFORM2011, 1601-1604.

[DEN12] R. Denninger, M. Liewald, M. Sindel, Failure prediction in sheet metal forming depending in pre-straining and bending superposition. Proceedings of ESAFORM 2012, 101-106.

[DEP11] T. Depover, D. Perez Escobar, E. Wallaert, M. Verhaege, In-situ mechanical evaluation of hydrogen embrittlement for TRIP, FB, DP, and HSLA steels. Proceedings of 1st Steely Hydrogen Conference, 2011, Ghent, 31-42.

[DU15] X.S. Du, W.B. Cao, S.J. Li, J.Y. Zhao, Y.F. Sun, Effect of microstructures and inclusions on hydrogen-induced cracking and blistering of A537 steel. Materials Science and Engineering A 642 (2015) 181-186.

[DUM08] A. Dumay, J. Chateau, S. Allain, S. Migot, O. Bouaziz, Influence of addition elements on the stacking-fault energy and mechanical properties of an austenitic Fe-Mn-C steel. Materials Science and Engineering A 483-484 (2008) 184-187.

[DOE10] E. Deoge B.A. Behrens, Handbuch Umformtechnik: Grundlagen, Technologien, Maschinen. Springer-Verlag Berlin Heidelberg 2010.

[ENG69] H.J. Engell, M.O. Speidel, Ursachen und Mechanismen der Spannungsrisskorrosion. Material and Corrosion 20 (1969) 281-300.

[FER98] P.J. Ferreira, I.M. Robertson, H.K. Birnbaum, Hydrogen effects on the interaction between dislocations. Acta Materialia 46 (1998) 1749-1757.

[FIG10] D. Figueroa, M.J. Robinson, Hydrogen transport and embrittlement in 300M and AerMet100 ultra high strength steels. Corrosion Science 52 (2010) 1593-1602.

132 Reference

[FIS99] D.J. Fisher, Hydrogen diffusion in metals. Defect and Diffusion Forum (1999) 167-168.

[FRI14] H. Friebe, Optical 3D metrology for sheet metal analysis and material parameters determination. FLC workshop, Zurich, 2014.

[FUC06] H. Fuchigami, H. Minami, M. Nagumo, Effect of grain size on the susceptibility of martensitic steel to hydrogen-related failure. Philosophical Magazine Letters 86 (2006) 21-29.

[FUJ85] F.E. Fujita, The iron-hydrogen phase diagram. In Hydrogen degradation of ferrous alloys, Eds. R.A. Oriani, J.P. Hirth, M. Smialowski, Park Ridge, N.J.: Noyes Publications, 1985, 1-15.

[GAN03] R.P. Gangloff, Hydrogen assisted cracking of high strength alloys. Comprehensive Structural Integrity 6 (2003) 131-101.

[GAO16a] Q. Gao, F. Han, D. Wortberg, W. Bleck, M. Liewald, The impact of hydrogen on the formability of AHSS in Nakajima tests. Proceedings of 19th ESAFORM, 2016, Nantes.

[GAO16b] Q. Gao, F. Han, D. Wortberg, W. Bleck, M. Liewald, Influence of hydrogen on formability and bendability of DP1180 steel for car body application. Proceedings of IDDRG2016, 2016, Linz.

[GEH94] F. Gehrmann, Einfluß der Nitride von V, Nb und Ti auf die Diffusion und Löslichkeit von Wasserstoff in Eisen. PhD thesis, Universität Dortmund, 1994.

[GEO97] J.S. Georgiev, L.A. Anestiev, Influence of the surface processes on the hydrogen permeation through ferric steel and amorphous Fe-40Ni-40Mo-4B alloy specimens. Journal of Nuclear Material 249 (1997) 133-141.

[GER88] W.W. Gerberich, S. Chen, Environment-induced cracking of metals fundamental processes: micromechanics. Proceedings of 1st International Conference on Environment-induced Cracking of Metals, 1988, NACE, Houston, 167-187.

[GOT04] G. Gottstein, Physical foundations of materials science. Springer-Verlag Berlin Heidelberg, 2004.

[GRA00] H.J. Grabke, E. Riecke, Absorption and diffusion of hydrogen in steels. Materials and Technology 331 (2000) 331-342.

[GUO12] X. Guo, Influences of microstructure, alloying elements and forming parameters on delayed fracture in TRIP/TWIP-aided austenitic steels. PhD. thesis, RWTH-Aachen University, 2012.

[GUO14] X. Guo, Q. Gao, W. Bleck, S. Kim, Mechanical properties and delayed fracture behavior of a hot dip galvanized Fe-15Mn-0.6C-2Al TWIP steel. Proceedings of 4th International Conference on Steels in Cars and Trucks, 2014, Braunschweig, 605-612.

[HEB08] T. Hebesberger, A. Pichler, H. Pauli, S. Ritsche, Dual-phase and complex-phase steels: AHSS Material for a wide range of applications. Proceedings of 2nd International Conference on Steels in Cars and Trucks, 2008, Wiesbaden, 456-463.

[HEB14] T. Hebesberger, A. Pichler, D. Kriyan, F. Winkelhofer, C. Walch et al., High Ductility AHSS Grades: Dual-Phase and Complex-Phase Grades with Improved Formability. Proceedings of 4th International Conference on Steels in Cars and Trucks, 2014, Braunschweig, 79- 86.

[HIL04] E.M.K. Hiller, M.J. Robinson, Hydrogen embrittlement of high strength steel electro-plated with zinc-cobalt alloys. Corrosion Science 46 (2004) 715-727.

[HIR80] J.P. Hirth, Effects of hydrogen on the properties of iron and steel. Metallurgical TransActions A 11 (1980) 861-890.

133

[HIR87] J.P. Hirth, Hydrogen Induced Fracture. In Chemistry and Physics of Fracture. Eds. R. M. Latanison, R. H. Jones, Martinus Nijhoff Publishers, 1987, 538-551.

[HIR95] J.P. Hirth, Theories of hydrogen induced cracking of steels. In hydrogen embrittlement and stress corrosion cracking, Eds. R. Gibala, R.F. Hehenmann, ASM, 1995, 29-41.

[HON13] S. Hong, J. Lee, B.-J. Lee, H.S. Kim, S.-K. Kim et al., Effects of intergranular carbide precipitation on delayed fracture behavior in three Twinning Induced Plasticity (TWIP) steels. Materials Science and Engineering A 587 (2013) 85-99.

[HWA94] L.C. Hwang, T.P. Perng, Hydrogen transport in ferritic stainless steel under elastic stress. Materials Chemistry and Physics 36 (1994) 231-235.

[IML04] K.-P. Imlau, T. Heller, U. Etzold, O. Hoffmann, Neue Stähle für Fahrzeugkarosserien. Praktische Metallographie Sonderband 36, Frankfurt, Werkstoffinformations- gesellschaft, 2004, 9-20.

[INT05] C.G. Interrante, L. Raymond, Hydrogen Damage. In Corrosion tests and standards: Application and Interpretation, Eds. R. Baboian, West Conshohocken, ASTM International, 2005, 322-340.

[JIN10] T.Y. Jin, Z.Y. Liu, Y.F. Cheng, Effect of ono-metallic inclusions on hydrogen-induced cracking of AP15LX100 steel. International Journal of Hydrogen Energy 35 (2010) 8014-8021.

[KAE03] H. Kaesche, Corrosion of Metals: Physical chemical Principles and current problems. Springer-Verlag Berlin Heidelberg, 2003.

[KAN10] S. Kang, Y.S. Jung, J.H. Jun, Y.-K. Lee, Effects of recrystallization annealing temperature on carbide precipitation, microstructure, and mechanical properties in Fe-18Mn-0.6C-1.5Al TWIP steel. Materials Science and Engineering A 527 (2010) 745–751.

[KIM11] J. Kim, S.J. Lee, B.C. De Cooman, Effect of Al on the stacking fault energy of Fe-18Mn-0.6C twinning-induced plasticity. Scripta Materialia 65 (2011) 363-366.

[KIM12] S.J. Kim, D.W. Yun, D.W. Suh, and K.Y. Kim, Electrochemical hydrogen permeation measurement through TRIP steel under loading condition of phase transition. Electrochemistry Communications 24 (2012) 112-115.

[KOY14] M. Koyama, C.C. Tasan, E. Akiyama, K. Tsuzaki, D. Raabe, Hydrogen-assisted decohesion and localized plasticity in dua-phase steel. Acta Materialia 70 (2014) 174-187.

[KUZ08] R. Kuyiak, R. Kawalla, S. Waengler, Advanced high strength steels for automotive industry. Archives of Civil and Mechanical Engineering 8 (2008) 103- 117.

[LAH14] C. Lahaijie, A. Blowey, J. Loisaux B. Ennis, Mass Saving Potential of dual phase steels with improved formability, Proceedings of 4th International Conference on Steels in Cars and Trucks, 2014, Braunschweig, 103-108.

[LAU15] A. Laureys, T. Depover, R. Petrov, K. Verbeken, Characterization of hydrogen induced cracking in TRIP-assisted steels. International Journal of Hydrogen Energy 40 (2015) 16901-16912.

[LI14] X. Li, Y. Wang, P. Zhang, B. Li, X. Song et al., Effect of pre-strain on Hydrogen embrittlement of high strength steels. Materials Science and Engineering A 616 (2014) 116-112.

[LIU08] X.Y. Liu, J. Kameda, J.W. Anderegg, S. Takaki, K. Abiko et al., Hydrogen-induced cracking in a very-high-purity high-strength steel. Materials Science and Engineering A 492 (2008) 218-220.

134 Reference

[LEE12] Y.K. Lee, Microstructural evolution during plastic deformation of twinning-induced plasticity steels. Scripta Materialia 66 (2012) 1002-1006.

[LOI11] M. Loidl, O. Kolk, S. Veith, T. Göbel, Charaterization of hydrogen embrittlement in automotive advanced high strength steels. Materialwissenschaft und Werkstofftechnik 42 (2011) 1105-1110.

[LOI14] M. Loidl, Entwicklung einer Prüfmethodik zur Charakterisierung höchstfester Karosseriestähle hinsichtlich des Risikos zur Wasserstoff induzierten Rissbildung. PhD thesis, Universität Stuttgart, 2014.

[LOU72] M.R. Louthan, G.R. Caskey, J.A. Donovan, D.E. Rawl, Hydrogen embrittlement of metals, Materials Science and Engineering 10 (1972) 357-346.

[LYN79] S.P. Lynch, Mechanisms of hydrogen assisted cracking. Metals Forum 2 (1979) 189-200.

[MA11] M. Ma, H. Yi, Lightweight car body and application of high strength steels. In Advanced Steels, Eds. Y. Wenig, H. Dong, Y. Gan, Springer-Verlag Berlin Heidelberg, 2011, 187-198.

[MAA08] P. Maaß, P. Peißker, Handbuch Feuerverzinken, Wiley-VCH Verlag, 2008.

[MAT09] D.K. Matlock, J.G. Speer, Third Generation of AHSS: Microstructure design concepts. In Microstructure and Texture in Steels, Eds. A. Haldar, S. Suwas, D. Bhattacharjee, Springer-Verlag London, 2008, 185-205.

[MAY15] R. Mayrhofer, Wasserstoffsensitität durch Korrosion bei UHSS, Workshop, Sindelfingen, 2015.

[MIS05] B. Mishra, D.L. Olson, F. Sanchez, An assessment of magnetization effects on hydrogen cracking for thick weld pipelines. Academic report to US Department of Transportation, Report-No.: CSM/MT/CWJCR/2005/009.

[MOL14] L. Moli-Sanchez, Z. Zermout, L. Duprez, L. Malet, Hydrogen embrittlement of 4 martensitic steels with strength levels above 1000 MPa. Proceedings of 2nd Steely Hydrogen Conference, 2014, Ghent, 70-84.

[MOO09] A. Moon, R. Balasubramaniam, Determination of Hydrogen Diffusivity in Rail Steels by Sub-Scale Microhardness Profiling Technique. Defect and Diffusion Forum 293 (2009) 41-45.

[MOR02] O. Moriau, L.T. Martinez, P. Verleyzen, J. Degrieck, Dynamic mechanical properties of cold rolled TRIP steels for crash-relevant applications. Proceedings of International Conference on TRIP Aided High Strength Ferrous Alloys, 2002, Ghent, 247-251.

[MUE06] S. Münstermann, Numerische Beschreibung des duktilen Versagensverhaltens von hochfesten Baustählen unter Berücksichtigung der Mikrostruktur, Berichte aus dem Institut für Eisenhüttenkunde, Band 4/06, Verlag Shaker, 2006.

[MUE14] T. Müller, Hydrogen embrittlement resistant electro galvanized martensitic steels for the automotive industry. Proceedings of 4th International Conference on Steels in Cars and Trucks, 2014, Braunschweig, 410-417.

[NAB10] N. Mabho, Determination of diffusible and total hydrogen concentration in coated and uncoated steel, Ph.D thesis, Universität Duisburg-Essen, 2010.

[NAG14] A. Nagao, C.D. Smith, M. Dadfarni, P. Sofronis, I.M. Robertson, Interpretation of hydrogen-induced fracture surface morphologies for lath martensitic steel. Procedia Materials Science 3 (2014) 1700-1705.

[NEL83] H.G. Nelson, Hydrogen Embrittlement. Treatise on materials science and technology 25 (1983) 275-359.

135

[OLD09] V. Olden, C. Thaulow, R. Johnsen, E. Østby, T. Berstad, Influence of hydrogen from cathodic protection on the fracture susceptibility of 25% Cr duplex stainless steel-Constant load SENT testing and FE-modelling using hydrogen influenced cohesive zone elements. Engineering Fracture Mechanics 76 (2009) 827-844.

[OLI07] S. Oliver, T.B. Jones, G. Fourlaris, Dual phase versus TRIP strip steels: microstructural changes as a consequence of quasi-static and dynamic tensile testing. Materials Characterization 58 (2007), 4, 390-400.

[ORI72] R. Oriani, A mechanistic theory of hydrogen embrittlement of steels. Berichte der Bunsengesellschaft für physikalische Chemie 76 (1972) 848-857.

[ORI74] R. Orini, P. Josephic, Equilibrium aspects of hydrogen-induced cracking of steels. Acta Metallurgica 22 (1974) 1065-1074.

[ORI85] R.A. Oriani, J.P. Hirth, M. Smialowski, Hydrogen degradation of ferrous alloys. Park Ridge, N.J., U.S.A.: Noyes Publications, 1985, 799-813.

[OTS13] S. Otsuka, Y. Sugimoto, S. Fujita, E. Tada, A. Nishikata et al., Behavior of hydrogen absorption on Zn-coated steel sheet under dry-wet cyclic corrosion environment. Proceedings of 9th International Conference on Zinc and Zinc Alloy Coated Steel Sheet & 2nd Asia-Pacific Galvanizing Conference, 2013, Beijing, China, 679-683.

[PAA11] W. Paatsch, Wasserstoffbestimmung und Wasserstoffversprödung – Sinn und Nutzen. In Galvanotechnik, Eds. G. Eugen, Leuze Verlag, 2011. 48-55.

[PAR15] I. Park, S. Lee, H. Jeon, Y. Lee, The advantage of grain refinement in the hydrogen embrittlement of Fe-18Mn-0,6C twinning-induced plasticity steel. Corrosion Science 93 (2015) 63-69.

[PER15] P.A.S. Pereria, C.S.G. Franco, J.L.M. Guerra Filho, D.S. dos Santos, Hydrogen effects on the microstructure of a 2.25Cr-1Mo-0.25V steel welded joint. International Journal of Hydrogen Energy 40 (2015) 17136-17143.

[PET52] N.J. Petch, Delayed fracture of metals under static load. Nature 169 (1952) 482-483.

[PUN05] A. Pundt, R. Kirchheim, Hydrogen in metals: microstructural aspects. Annual Review of Materials Research 36 (2006) 555-608.

[RAO12] R. Raoelison, A. Fuentes, Ph. Rogeon, P. Carre, T. Loulou et al., Contact conditions on nugget development during resistance spot welding of Zn coated steel sheets using rounded tip electrodes. Journal of Materials Processing Technology 212 (2012) 1663-1669.

[RAM16] A. Ramazani, M. Abbasi, S. Kazemiabnavi, S. Schmauder, R. Larson et al., Development and application of a microstructure-based approach to Characterize and model failure initiation in DP steels using XFEM. Materials Science and Engineering A 660 (2016) 181-194.

[RAT07] E. Ratte, Wasserstoffinduzierte verzoegerte Rissbildung austenitischer Stähle auf CrNi(Mn)- und Mn-Basis. PhD thesis, RWTH-Aachen University,2007.

[REH13] J. Rehrl, Wasserstoffversprödung in hochfesten, mikrolegierten Stählen. PhD Thesis, TU-München, 2013.

[REH14] J. Rehrl, K. Mraczek, A. Pichler, E. Werner, The impact of hydrogen on the mechanical properties of AHSS/UHSS grades at low- and high strain rates. Proceedings of 2nd Steely Hydrogen Conference, 2014, Gent, 21-32.

[RON12] J A Ronevich, B.C. De Cooman, J.G. Speer, E. De Moor, D.K. Matlock, Hydrogen Effects in Prestrained Transformation Induced Plasticity steel. Metallurgical and Materials Transaction A 43A (2012) 2012-2293.

[SED96] A.J. Sedriks, Corrosion of stainless steels. Wiley-Interscience, 1996.

136 Reference

[SCH01] T. Schlueter, Wasserstoffaufnahme während des Beizens und des Glühens von warmgewalztem Stahlband vor der Feuerverzinkung. PhD thesis, Universität Dortmund, 2001.

[SCO06] C. Scott, S. Allain, M. Faral, N. Guelton, The development of a new Fe-Mn-C austenitic steel for automotive applications. Revue de Metallurgy Paris 103 (2006) 293-302.

[SHL09] R. Shleich, C. Held, M. Sindel, M. Liewald, Investigation on the effect of curvature and sheet thickness on forming limit prediction for aluminium sheet metal alloys. International Journal of Material Forming 2 (2009) 411-414.

[SHV00] V.I. Shvachko, Cold cracking of structural steel weldments as reversible hydrogen embrittlement effect. International Journal of Hydrogen Energy 25 (2000) 473-480.

[SID75] V.M. Sidorenko, B.F. Kachmar, N.S. Borisova, Effect of deformation and stress on the diffusion characteristics of hydrogen in metals. Materials Science 9 (1975) 500-504.

[SMI00] L. Smith, M. Celant, A. Pourbaix, A guideline to the successful use of duplex. Proceedings of Duplex America 2000 Conference, Houston, Texas, 2000.

[SPA00] H. Spähn, Druckwasserstoffangriff auf unlegierte und niedriglegierte Stähle im Temperaturbereich oberhalb 200°C. In Wasserstoff und Korrosion, Eds. Kuron, Verlag Irene Kuron Bonn, 2000, 127-183.

[SPE06] M.O. Speidel, Nitrogen containing austenitic stainless steels. Material Wissenschaft und Werkstofftechnik 37 (10): 875-880.

[STO94] M.R. Stoudt, R.E. Ricker, Characterizatijon of the hydrogen induced cold cracking susceptibility at simulated weld zones in HSLA-100 steel. Academic report for Office of Naval Research, U.S. Department of Commerce, 1994.

[STR06] M.E. Stroe, Hydrogen embrittlement of ferrous materials. PhD Thesis, Universite Libre de Bruxelles, 2006.

[SUN15] Y. Sun, J. Chen, J. Liu, Effect of hydrogen on ductility of high strength quenched and tempered (QT) Cr-Ni-Mo steels. Materials Science and Engineering A 625 (2015) 89-97.

[SUG08] K. Sugimoto, M. Murata, Y. Mukai, Stretch-Flangeability of C-Mn-Si-Al-Nb Ultra High-Strength Cold Rolled Sheet Steels. Proceedings of Advanced High Strength Sheet Steels for Automotive Applications, 2008, Orlando, Fla., USA, 57-67.

[SYM01] D.M. Symons, A comparison of internal hydrogen embrittlement and hydrogen environment embrittlement of X-750. Engineering Fracture Mechanics 68 (2001) 751-771.

[TAB84] T. Tabata, H.K. Birnbaum, Direct observations of hydrogen enhanced crack propagation in iron. Scripta Metallurgica 18 (1984) 231-236.

[TAH01] A. Taha, P. Sofronis, A micromechanics approach to the study of hydrogen transport and embrittlement. Engineering Fracture Mechanics 68 (2001) 803-837.

[THI11a] R.G. Thiessen, T. Heller, K. Mraczek . Nitschke, A. Pichler, Influence of microstructure on the susceptibility to hydrogen embrittlement. Proceedings of 1st. Steely Hydrogen Conference, 2011 Ghent, 43-51.

[THI11b] R.G. Thiessen, K. Bergers, T. Heller, Influence of electro-galvanizing parameters on absorbed hydrogen content in AHSS. Proceedings of 1st. Steely Hydrogen Conference, 2011 Ghent, 249-259.

[TAK08] K. Takai, H. Shoda, H. Suzuki, M. Nagumo, Lattice defects dominating hydrogen-related failure of metals. Acta Materialia 56 (2008) 5158-5167.

137

[TAK10] S. Takagi, Y. Toji, K. Hasegawa, Y. Tanaka, Evaluation methods of hydrogen embrittlement for ultra-high strength steel sheets for automobiles. Materials Science and Technology October (2010) 17-21.

[TAK12] K. Takasawa, R. Ikeda, N. Ishikawa, R. Ishigaki, Effects of grain size and dislocation density on the susceptibility to high-pressure hydrogen environment embrittlement of high-strength low-alloy steels. International Journal of Hydrogen Energy 37 (2012) 2669–2675.

[TAK14] S. Takaji, Y.Toji, M. Yoshino, K. Hasegawa, N. Wada et al., Hydrogen embrittlement resistance of 1180 MPa grade automotive steel sheet evaluated by various methods. Proceedings of 2nd Steely Hydrogen Conference, 2014, Ghent, 13-20.

[TAV15] S. Tavares, I. Bastos, J. M. Pardal, T.R. Montenegro, M.R. da Silva, Slow strain rate tensile test results of new multiphase 17% Cr Stainless steel under hydrogen cathodic charging. International Journal of Hydrogen Energy 40 (2015) 16992-16999.

[THI11] R.G. Thiessen, T. Heller, K. Mraczek, A. Nitschke, A. Pichler, Influence of Microstructure on the susceptibility to hydrogen embrittlement. Proceedings of 1st Steely Hydrogen Conference, 2011, Ghent, 43-51.

[TRO60] A. Troiano, The role of hydrogen and other interstitials in the mechanical behavior of metals. Transactions of the American Society of Metals 52 (1960) 54-80.

[VIT06] L. Vitos, J.O. Nilsson, B. Johansson, Alloying effects on the stacking fault energy in austenitic stainless steels from first-principles theory. Acta Materialia 54 (2006) 3821-3826.

[VIY05] E. Viyanit, Numerical simulation of hydrogen assisted cracking in supermartensitic stainless steel welds. PhD thesis, Federal Institute for Materials Research and Testing Berlin, 2005.

[VOL11] W. Volk, P. Hora, New algorithm for a robust user-independent evaluation of beginning instability for the experimental FLC determination. International Journal of Material Forming 4 (2011) 339-346.

[WAN01] J.S. Wang, The thermodynamics aspects of hydrogen induced embrittlement. Engineering Fracture Mechanics 68 (2001) 647-669.

[WAN06] M.Q. Wang, E. Akiyama, K. Tsuzaki, Determination of the critical hydrogen concentration for delayed fracture of high strength steel by constant loading test and numerical calculation. Corrosion Science 48 (2006) 2189-2202.

[WAN13] L. Wang, J.G. Speer, Quenching and Partitioning Steel Heat Treatment. Metallography, Microstructure, and Analysis 2 (2013) 268-281.

[WOR14] www.WorldAutoSteel.com, Advanced High-Strength Steels Application Guidelines: Section 1 – Introduction to advanced high-strength steels (AHSS), Version 5.0, 2014.

[WU05] J. Wu, P.J. Wray, C.I. Garcia, M. Hua, A.J. Deardo, Image quality analysis: a new method of characterizing microstructures. ISIJ International 45 (2005) 254-262.

[YAK02] O. Yakubovsky, N. Fonstein, D. Bhattacharya, Stress-strain behavior and bake hardenability of TRIP and TRIP-aided multiphase steels. Proceedings of International Conference on TRIP Aided High Strength Ferrous Alloys, 2002, Ghent, 263-270.

[ZAE08] S. Zaefferer, P. Romano, F. Friedel, EBSD as a tool to identify and quantify bainite and ferrite in low-alloyed Al-TRIP steels. Journal of Microscopy 230 (2008) 499-508.

[ZAK85]

T. Zakroczymski, Entry of hydrogen into iron alloys from the liquid phase. In Hydrogen degradation of ferrous alloys, Eds. R. Oriani, J. Hirth, M. Smialowski, Park Ridge, N.J., U.S.A.: Noyes Publications, 1985, 215-250.

[ZAK99] T. Zakroczymski, Elektrochemical Determination of Hydrogen in Metals. Journal of Electroanalytical Chemistry 82-88 (1999).

138 Reference

[ZAP41] C.A. Zappe, C.E. Sims, Hydrogen embrittlement, internal stress and defects in steel. Transactions of the Metallurgical Society of AIME 145 (1941) 225-259.

[ZEP03] F. von Zeppelin, M. Haluska, M. Hirscher; Thermal desorption spectroscopy as a quantitative tool to determine the hydrogen content in solids. Thermochimica Acta 404 (2003) 251-258.

[ZHA15] S. Zhang, Y. Huang, B. Sun, Q. Liao, H. Lu et al., Effect of Nb on hydrogen-induced delayed fracture in high strength hot stamping steels. Materials Science and Engineering A 626 (2015) 136-143.

[ZHO14] Q.J. Zhou, L. Wang, J.X. Li, Experimental study on delayed fracture of TS 980 MPa grade steels for automotive applications. Proceedings of 2nd Steely Hydrogen Conference, 2014, Ghent, 35-44.

[ZHU13] X. Zhu, W. Li, H. Zhao, X. Jin, Effects of cryogenic and tempered treatment on the hydrogen embrittlement susceptibility of TRIP-780 steels. International Journal of Hydrogen Energy 38 (2013) 10694-10703.

[ZHU14] X. Zhu, W. Li, H. Zhao L. Wang, X. Jin, Hydrogen trapping sites and hydrogen-induced cracking in high strength quenching & partitioning (Q&P) treated steel. International Journal of Hydrogen Energy 39 (2014) 13031-13040.

[ZHU16] X. Zhu, K. Zhang, W. Li, X. Jin, Effect of retained austenite stability and morphology on the hydrogen embrittlement susceptibility in quenching and partitioning treated steels. Materials Science and Engineering A 658 (2016) 400-408.

139

Curriculum Vitae

PERSONAL DATA

Family name, given name: Gao, Qi

Date and place of birth: 19.05.1986, in Nei Mongol, China

Gender: Female

Nationality: Chinese

EDUCATION

10.2009-09.2012 Master of Science in Metallurgical Engineering RWTH-Aachen University, Germany

09.2005-06.2009 Bachelor of Engineering in Material Science and Engineering

University of Science and Technology Beijing, China

RELEVANT EXPERIENCE

08.2013-12.2016 Ph.D student in Group Material Engineering of Metals

Process Development and Materials Daimler AG, Sindelfingen, Germany

04.2012-09.2012 Master thesis in Group of Material Characterization

Department of Ferrous Metallurgy RWTH-Aachen University

05.2010-08.2012 Student Assistant in Group of Advanced Material Simulation

Department of Ferrous Metallurgy RWTH-Aachen University

Documents

Investigations on Hydrogen Embrittlement of Advanced High