mechanical and material characterization of high strength steel … · i would also like to thank...

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Mechanical and material characterization of high strength

steel welded joints

João Pedro Ribeiro Marques

Thesis to obtain the Master of Science Degree in

Mechanical Engineering

Supervisors: Prof. Virgínia Isabel Monteiro Nabais Infante

Prof. Ricardo Miguel Gomes Simões Baptista

Examination Committee

Chairperson: Prof. Luís Filipe Galrão dos Reis

Supervisor: Prof. Virgínia Isabel Monteiro Nabais Infante

Member of the Committee: Prof. Ana Mafalda Saldanha Guedes

December 2017

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Abstract

The improvement and development of technology lead to a necessity of new materials with specific

properties. As consequence, a comprehensive study is necessary in order to understand their strength,

limitations, and behaviours under several conditions.

The purpose of this work is to evaluate the fatigue resistance of a high strength steel with a specific

welding geometry and parameters, undergoing a 0.1 stress ratio. The Paris Law is determined for each

case, using the Digital Image Correlation Method.

Additionally, it is assessed the variation on hardness and Yield strength along the weld zone and heat

affected zone and it is studied the influence of 4 different sets of welding parameters on those properties.

It is found that the increase in the heat input has as consequence an overall reduction of hardness and

yield strength.

Moreover, it is also studied the steel microstructure in different zones, which appears as a result of the

welding process, and it is compared with the microstructure of the base material. The increase of heat

input promoted the occurrence of ferrite in the distinct zones examined.

Finally, a numerical analysis is performed for 3 distinct cases: a specimen without a weld, another with

a transverse weld and one with a non-centred vertical weld, in order to compare the different Stress

Intensity Factor evolutions.

It is found that the average SIF value is almost constant for all cases, however, the SIF evaluation

through the specimen thickness differ, depending on the case considered.

Keywords: High strength steel (class 700 MPa), Fatigue, Digital Image Correlation, Hardness, Yield

strength, Microstructure

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Resumo

Com a contante evolução da tecnologia, a necessidade de fabricar novos materiais torna-se prioritária.

Deste modo é necessário quantificar as propriedades mecânicas de novos materiais, de forma a

compreender as limitações e vantagens do mesmo.

Neste trabalho é avaliada a resistência à fadiga, através da determinação dos parâmetros da lei de

Paris, de um aço de alta resistência (classe 700 MPa) sujeito a uma razão de tensões de 0.1 para uma

geometria de soldadura e determinados parâmetros de soldadura, utilizando o método Digital Image

Correlation. Adicionalmente, é estudada a variação de dureza, das tensões de cedência, bem como a

variação da microestrutura, para diversas amostras, de forma a observar a influência de 4 diferentes

conjuntos parâmetros de soldadura nos diferentes casos. Observou-se que um aumento da entrega

térmica promoveu uma redução geral no valor de dureza e tensão de cedência. Da análise da

microestrutura constatou-se um aumento de ferrite com o aumento da entrega térmica.

Por fim, uma análise numérica é realizada para 3 casos distintos, um provete sem soldadura, um

provete com um cordão de soldadura perpendicular à direção da propagação da fenda e um provete

com uma soldadura vertical descentrada, de forma a comparar as diferentes evoluções do factor de

intensidade de tensões. Apesar das evoluções em função do tamanho de fenda terem sido bastante

próximas em todos os casos, as evoluções ao longo da espessura demonstraram-se diferentes

dependendo do provete e tamanho de fenda considerados.

Palavras-chaves: Aço de alta resistência (classe 700 MPa), fadiga, Digital Image Correlation, dureza,

tensão de cedência, microestrutura

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Acknowledgments

I would like to express my gratitude for Dr. Virginia Infante and Dr. Ricardo Batispta for proposing this

multifaceted work and for their guidance, availability and support.

I am grateful to Dr. Mafalda Guedes for the assistance given on the hardness tests and microstructure

analysis.

To my friends and family who accompanied me during this project, especially to José Lopes, Letícia

Carvalho, Diogo Nascimento, Tiago Lopes, Rodrigo Martins, Pedro Esteves, Gabriel Maciel, Rui

Parente, João Mendonça and Mariana Raposo

A special word to Mr. Daniel Jesus for the assistance given in polishing of the samples.

I would also like to thank the collaboration of Junaid Akhtar of Aalto University for the assistance given

on the Digital Image Correlation method.

I would like to express my gratitude for my parents and brothers who demonstrated a massive support,

not only during this project, but also during my entire life.

This work was supported by FCT, through IDMEC, under LAETA, project UID/EMS/50022/2013.

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Table of Contents

Abstract .................................................................................................................................................. ii

Resumo.................................................................................................................................................. iv

Acknowledgments ................................................................................................................................ vi

Table of Contents ............................................................................................................................... viii

List of Figures ...................................................................................................................................... xii

List of Tables ..................................................................................................................................... xviii

Nomenclature ....................................................................................................................................... xx

Glossary ............................................................................................................................................. xxii

1. Introduction .................................................................................................................................... 1

1.1. Topic Overview ...................................................................................................................... 1

1.2. Motivation .............................................................................................................................. 3

1.3. Objectives .............................................................................................................................. 3

1.4. Thesis Outline ........................................................................................................................ 4

2. Literature Review ........................................................................................................................... 5

2.1. Fracture Mechanics ............................................................................................................... 5

2.1.1. Loading Modes ......................................................................................................... 5

2.1.2. Stress Intensity Factor .............................................................................................. 6

2.1.3. Plane Strain and Plane Stress .................................................................................. 9

2.1.4. J Integral ................................................................................................................. 11

2.1.5. J and K evolutions along the thickness .................................................................. 12

2.2. Fatigue ................................................................................................................................. 13

2.2.1. Fatigue stages ........................................................................................................ 14

2.2.2. Microstructure influence on fatigue crack propagation ........................................... 20

2.2.3. Stress ratio influence on the fatigue crack propagation ......................................... 21

2.2.4. Influence of welding and welding residual stresses on fatigue crack propagation . 22

2.2.5. Crack propagation laws .......................................................................................... 23

2.3. MIG/MAG (Metal Inert Gas/ Metal Active Gas) Welding ..................................................... 25

2.3.1. Arc physics ............................................................................................................. 25

2.3.2. Consumables .......................................................................................................... 26

2.3.3. Welding process parameters .................................................................................. 27

2.4. Steel Microstructure ............................................................................................................. 28

2.4.1. High strength steels ................................................................................................ 30

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2.4.2. Heat Affected Zone microstructure ......................................................................... 31

2.4.3. Heat input influence ................................................................................................ 33

2.4.4. S700MC .................................................................................................................. 34

2.5. Young’s Modulus ................................................................................................................. 35

2.6. Digital Image Correlation ..................................................................................................... 37

2.6.1. Introduction ............................................................................................................. 37

2.6.2. Sobel method and Fast Fourier Transform............................................................. 38

2.6.3. Method developed to measure crack lengths ......................................................... 40

2.6.4. Speckle pattern ....................................................................................................... 41

3. Material and Welding Procedure ................................................................................................ 43

3.1. Material ................................................................................................................................ 43

3.2. Welding procedure .............................................................................................................. 43

4. Laboratorial part .......................................................................................................................... 47

4.1. Fatigue tests ........................................................................................................................ 47

4.2. Micro hardness tests ........................................................................................................... 49

4.2.1. Samples preparation ............................................................................................... 49

4.2.2. Micro hardness test conditions ............................................................................... 49

4.3. Metallographic Analysis ....................................................................................................... 52

5. Numerical Analysis ...................................................................................................................... 53

6. Results and Discussion ............................................................................................................. 55

6.1. Fatigue results ..................................................................................................................... 55

6.2. Micro hardness results ........................................................................................................ 56

6.3. Yield Stress results .............................................................................................................. 61

6.4. Microstructure analysis ........................................................................................................ 64

6.5. Numerical simulation results................................................................................................ 70

7. Conclusions and Recommendations ........................................................................................ 77

7.1. Conclusions ......................................................................................................................... 77

7.2. Future works ........................................................................................................................ 78

References ........................................................................................................................................... 79

Annexs .................................................................................................................................................. 85

A1 Fatigue data ............................................................................................................................ 86

A2 Temperature profiles and Modelled molten and austenittized zones ............................... 87

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A3 Microstructure ........................................................................................................................ 91

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List of Figures

Figure 1.1: Axial loading S-N at various mean stresses for unnotched specimens of an aluminium alloy

[2] ............................................................................................................................................................. 1

Figure 1.2: Failure origin of Comet G-ALYU around square windows [3] ............................................... 2

Figure 2.1: Loading modes types [10] ..................................................................................................... 6

Figure 2.2: Stresses near the crack tip [11] ............................................................................................. 7

Figure 2.3: Stress normal to the crack plane in Mode I [9] ...................................................................... 8

Figure 2.4: Specimen geometry and respective f (a/W) factor[9] ............................................................ 9

Figure 2.5: Plane stress fracture. Plastic zone diameter r0 comparable to or greater than sample

thickness. Modified from [13] ................................................................................................................... 9

Figure 2.6: Plane strain fracture. Plastic zone diameter r0 much less than sample thickness [13] ....... 10

Figure 2.7: Effect of specimen thickness on fracture toughness [14] .................................................... 10

Figure 2.8: Path for the calculation of J-integral ..................................................................................... 11

Figure 2.9: Local J normalized by far-field/remote J, plotted vs position along the half-crack front, at

various loading levels [17] ..................................................................................................................... 12

Figure 2.10: Kj based evolution through the thickness for different crack front shapes [18] ................. 13

Figure 2.11: Fatigue cycle [19] .............................................................................................................. 13

Figure 2.12: Primary fracture mechanisms in steels associated with sigmoidal variation of fatigue crack

propagation rate (da/dN) with alternating stress intensity (ΔK)[20] ....................................................... 14

Figure 2.13: Typical propagation of a fatigue crack [21] ....................................................................... 15

Figure 2.14: Development of extrusions and intrusions during fatigue [21] .......................................... 16

Figure 2.15: Schematic of stages I (shear mode) and II (tensile mode) transcrystalline microscopic

fatigue crack growth Adapted from [14] ................................................................................................. 16

Figure 2.16: A model for the creation of saw-tooth type striations [26] ................................................. 17

Figure 2.17: A higher magnification SEM photomicrograph of the transition from ductile overload fracture

(dimples) of the weld HAZ to brittle overload (cleavage) fracture of the base metal [27] ..................... 18

Figure 2.18: a) Ductile striation growth through tempered martensite during fatigue (unembrittled steel),

b) Brittle intergranular cracking during striation growth in fatigue of embrittled steel [28] .................... 18

Figure 2.19: Variation of regression lens and values of the slope m for fatigue crack growth rates of

unembrittled steel for several stress ratios R [28] ................................................................................. 19

Figure 2.20: Fatigue crack propagation data. a) Martensitic steels; b) Ferritic-pearlite steels ............. 20

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Figure 2.21: A family of curves showing effect of stress ratio on fatigue crack propagation of JES SM58Q

steel [33] ................................................................................................................................................ 21

Figure 2.22: Fatigue crack growth rate as a function of stress intensity range and effective stress

intensity range a) Base Material; b) Weld Material and Cross-Bond [35] ............................................. 22

Figure 2.23: Fatigue crack growth rate versus crack length for Transversal specimen, adapted from [35]

............................................................................................................................................................... 23

Figure 2.24: MIG/MAG process [33] ...................................................................................................... 25

Figure 2.25: Comparison of melting rates in welding with a covered electrode, a solid wire and a cored

wire, respectively [42] ............................................................................................................................ 26

Figure 2.26: Influence of shielding gas in MIG/MAG welding processes [42] ....................................... 27

Figure 2.27: Iron diagram phase [45] .................................................................................................... 28

Figure 2.28: Different structures obtained in steel after cooling from austenite by different cooling rates

[46] ......................................................................................................................................................... 28

Figure 2.29: Upper Bainite (UB), Lower Bainite (LB) and Polygonal Ferrite (PF) microstructures [47] 29

Figure 2.30: a) Isothermal transformation diagram for a eutectoid steel [46]; b) Ferritic-martensitic

structure [48] .......................................................................................................................................... 29

Figure 2.31: Microstructure evolution during thermal processing of high advanced high strength steels

[49] ......................................................................................................................................................... 30

Figure 2.32: Different rolling stages of TMCP [6] .................................................................................. 30

Figure 2.33: Microstructural control by different cooling routes after hot rolling [7] .............................. 31

Figure 2.34: Typical domains of heat affected zone in welds for steel [50] ........................................... 31

Figure 2.35: Vickers hardness measurements in the HAZ of a structural steel and grain size variation

[53] ......................................................................................................................................................... 32

Figure 2.36: Acicular ferrite structure [48] ............................................................................................. 32

Figure 2.37: HAZ overlapping in multi-pass welding [55] ...................................................................... 33

Figure 2.38: a) Structural changes in the S700 MC steel in terms of welding CCT-W ; b) Distribution of

hardness as a function of the cooling time [59] ..................................................................................... 35

Figure 2.39: Maximum temperatures (oC) in the middle part area of cross-section [63] ....................... 36

Figure 2.40: Heat-affected zone and thermal cycle for P point [63] ...................................................... 36

Figure 2.41 - Young's modulus versus grain size for ultra-fine grain sized mild steel [64] .................... 37

Figure 2.42: Characteristic microstructural zones of the HAZ of a double-submerged arc-welded joint

[65] ......................................................................................................................................................... 37

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Figure 2.43: Crack length vs no. of cycles obtained by travelling microscope and DIC using sobel edge

detection method [68] ............................................................................................................................ 39

Figure 2.44: Example of a displacement field ....................................................................................... 40

Figure 2.45: Crack tip location (initial assumption) ............................................................................... 40

Figure 2.46: Crack surface plot ............................................................................................................. 41

Figure 2.47: Influence of the speckle pattern on strain measurements [72] ......................................... 42

Figure 3.1: Clamping system and dimensions of the plates .................................................................. 44

Figure 3.2: Schematic presentation of joint design and welding sequence for all welds ...................... 44

Figure 4.1: Fatigue experimental apparatus.......................................................................................... 47

Figure 4.2: Test specimen ..................................................................................................................... 47

Figure 4.3: Specimen dimensions. Adapted from [74] .......................................................................... 48

Figure 4.4: Speckle pattern applied on specimen test ......................................................................... 48

Figure 4.5: Fatigue test images ............................................................................................................. 49

Figure 4.6: Vickers hardness test schematics [75] ................................................................................ 50

Figure 4.7: Hardness measurements .................................................................................................... 51

Figure 4.8: Scanning electron microscope at IST ................................................................................. 52

Figure 4.9: Location of the regions where the microstructures were analysed ..................................... 52

Figure 5.1: BM case (left), T case (centre), HAZ case (right)................................................................ 53

Figure 5.2: Mesh at the crack tip (crack tip is located at the centre of the circumferences) ................. 54

Figure 5.3: Application of static forces ................................................................................................... 54

Figure 6.1: a) Crack propagation rate for HAZ 6A case 1; b) case 2 .................................................... 56

Figure 6.2: Hardness profile 5A sample ................................................................................................ 57




Figure 6.6: Average hardness measured for each sample and respective standard deviation ............ 59

Figure 6.7: Yield stress profile sample 5A ............................................................................................. 62



Figure 6.10: Yield stress profile sample 8A ........................................................................................... 63

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Figure 6.11: Different zones due to the welding process ...................................................................... 64

Figure 6.12: Weld geometries for the distinct samples. Sample 5A (case A), sample 6A (case B), sample

7A (case C), sample 8A (case D) .......................................................................................................... 65

Figure 6.13: Microstructures position A 2000X magnification ............................................................... 66

Figure 6.14: Microstructures position B 2000X magnification ............................................................... 66

Figure 6.15: Microstructures position C 2000X magnification ............................................................... 67

Figure 6.16: Microstructures position D 2000X magnification ............................................................... 67

Figure 6.17: Microstructures position E 2000X magnification ............................................................... 68

Figure 6.18: Microstructures position F 2000X magnification ............................................................... 68

Figure 6.19: Microstructures position G 2000X magnification ............................................................... 68

Figure 6.20: Microstructures position H 2000X magnification ............................................................... 69

Figure 6.21: Microstructures position I 2000X magnification, apart from sample 7A which was performed

with a 5000X magnification.................................................................................................................... 69

Figure 6.22: Evolution of the middle SIF value for a=40 mm ................................................................ 71

Figure 6.23: Comparison between numerical simulation SIF values and standard SIF values ............ 72

Figure 6.24: SIF evolution for the distinct cases ................................................................................... 73

Figure 6.25: SIF relative variation between BM case and the other cases ........................................... 73

Figure 6.26: SIF evolution along the thickness T case a=40mm .......................................................... 74

Figure 6.27: Location of the crack tip in both cases .............................................................................. 74

Figure 6.28: SIF evolution along the thickness for T case a=13mm ..................................................... 74

Figure 6.29: SIF evolution along the thickness for T case a=18mm ..................................................... 75

Figure 6.30: Crack tip location for a=23mm .......................................................................................... 75

Figure 6.31: SIF evolution along the thickness T case for a=23mm ..................................................... 76

Figure 6.32: SIF evolution along the thickness HAZ case for a=12mm ................................................ 76

Figure A.1: Temperature profiles sample 5A - Face side (top image) Root side (bottom image) ......... 88

Figure A.2: Temperatures profile sample 6A - Face side (top image) Root side (bottom image) ......... 88



Figure A.5: Modelled molten and austenitized zones overlapped with actual macrographs for root pass

and for all second passes ...................................................................................................................... 90

Figure A.6: Position A, Top images 2000X magnification; Bottom images 5000X magnification .......... 91

xvi

Figure A.7: Position B, Top images 2000X magnification; Bottom images 5000X magnification .......... 92

Figure A.8: Position C, Top images 2000X magnification; Bottom images 5000X magnification ......... 93

Figure A.9: Position D, Top images 2000X magnification; Bottom images 5000X magnification ......... 94

Figure A.10: Position E, Top images 2000X magnification; Bottom images 5000X magnification ........ 95

Figure A.11: Position F, Top images 2000X magnification; Bottom images 5000X magnification......... 96

Figure A.12: Position G, Top images 2000X magnification; Bottom images 5000X magnification ....... 97

Figure A.13: Position H, Top images 2000X magnification; Bottom images 5000X magnification ....... 98

Figure A.14: Position I, Top images 2000X magnification; Bottom images 5000X magnification ......... 99

xviii

List of Tables

Table 2.1: C and m in Base Material, Weld Material and Cross-Bond specimens [35] ......................... 22

Table 2.2: Mechanical properties of heat treated and untreated NST 37-2 steel [61] ........................... 35

Table 3.1: Steel chemical composition .................................................................................................. 43

Table 3.2: Mechanical properties ........................................................................................................... 43

Table 3.3: Welding parameters for each t8/5 ......................................................................................... 45

Table 6.1: Comparison between measured Yield Strength and material Yield strength ....................... 59

Table 6.2: Softest and Hardest point’s location ..................................................................................... 60

Table 6.3: Comparison between the initial yield strength and the lowest and the highest values ........ 64

Table 6.4: SIF values for several number of points along the thickness a=11 mm ............................... 70

Table 6.5: SIF values for several number of points along the thickness a=40mm ................................ 70

Table 6.6: SIF values provided by standard and numerical analysis for a=11 and a=40 mm ............... 72

Table A.1: Case 1 .................................................................................................................................. 86

Table A.2: Case 2 .................................................................................................................................. 87

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Nomenclature

𝑨 Total elongation

𝒂 Crack length

𝑨𝟏, 𝑨𝒄𝟏 Starting temperature for austenitic transformation at heating

𝑨𝟑 𝑨𝒄𝟑Temperature at which the full austenitization reached at heating

𝒂𝒄 Critical crack length

𝑨𝒓𝟏 Finishing temperature for formation of ferrite at cooling

𝑨𝒓𝟑 Temperature at which the ferrite formation starts at cooling

𝑩 Specimen’s thickness

𝑪 Paris law Constant

𝑫 Standard deviation

𝒅 Grain size

𝑬 Young’s Modulus

𝒆 Relative variation

𝑭𝟐 Joint type factor for 2-Dimensional Heat Flow

𝒇𝒊𝒋, 𝒀 Factors dependent of loading conditions and specimen geometry

𝒇𝑹, 𝒇𝒕𝒉, 𝒇𝑪 Correction factors

𝑯𝒗 Hardness Vickers

𝑱 Contour integral

𝑲 Stress intensity factor

𝑲𝒄 Fracture Toughness

𝒌 Thermal Efficiency Factor

𝒌𝒚, 𝝈𝟎 Materials constants

𝒎 Paris law exponent

𝑵𝒇 Fatigue life

𝑷 Applied force

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𝑸 Energy arc

𝑹 Stress ratio

𝑹𝟐 Correlation coefficient

𝒓 Distance to the crack tip

𝑺 Strain-energy density

𝑻 Tension vector

𝑻𝟎 Room Temperature

𝒕 Workpiece thickness

𝒕𝟖/𝟓 Cooling time from 800oC to 500oC

𝒖𝒙, 𝒖𝒚 Displacement field

𝑾 Specimen’s width

𝒛 Coordinate along the specimen’s thickness

Greek Variables

𝜶 Quotient between the crack length and the specimen’s width

𝜞 Path surrounding the notch tip

𝜺 Strain tensor

∆𝑲 Stress intensity Factor variation

∆𝝈 Stress variation

𝝈𝒙, 𝝈𝒚, 𝝈𝒛 Normal stress

𝝈𝟏, 𝝈𝟐, 𝝈𝟑 Principal stresses

𝝈𝒀𝑺 Yield stress

𝝈𝑼𝑻𝑺 Ultimate tensile stress

𝜽 Angle

𝝁 Shear modulus

𝝂 Poisson Coefficient

𝝉𝒙𝒚, 𝝉𝒚𝒛, 𝝉𝒙𝒛 Shear stresses

xxii

Glossary

AWS – American Welding Society PSB – Persistent Slip Bands

BM – Base Material SCHAZ – Subcritical Heat-Affected Zone

CB – Cross-Bond SIF – Stress Intensity Factor

CGHAZ – Coarse Grained Heat-Affected Zone SRCG – Subcritically Reheated CGHAZ

DCEN – Direct-Current Electrode Negative TMCP – Thermo-Mechanically controlled process

DCEP – Direct-Current Electrode Positive TRIP – Transformation Induced Plasticity

DFT – Discrete Fourier Transform UB – Upper Bainite

DIC – Digital Image Correlation WZ – Weld Zone

EPFM – Elastic-Plastic Fracture Mechanics ZOI – Zone Of Interest

FCHAZ – Fine Grained Heat-Affected Zone

FEA – Finite Element Analysis

FFT – Fast Fourier Transform

FL – Fusion Line

HAZ – Heat-Affected Zone

HSLA – High Strength Low Alloy steel

ICHAZ – Inter Critical Heat-Affected Zone

IRCG – Intercriticallly Reheated CGHAZ

LB – Lower Bainite

LEFM – Linear Elastic Fracture Mechanics

MAG – Metal-Active Gas welding

MIG – Metal-Inert Gas welding

MVC – Microvoid Coalescence

1

1. Introduction

1.1. Topic Overview

The advances in technology, create new engineering challenges, particularly in the search of new

materials, where the necessity of materials with better properties becomes critical. One important aspect

relies on the continuous improvement of fatigue strength of newer materials since fatigue phenomenon

is the major responsible for the failure of components in engineering (around 80%). In order to

emphasize the magnitude of this undesired process, it is worth to mention that the costs related to

fatigue failure may reach 3% of the GNP (gross national product) [1].

The fatigue process occurs, due to crack propagation, when a dynamic loading is applied to the

component. In general, it occurred for a low stress level, which can be considerable inferior to the yield

strength of the material. The fatigue failure usually occurs suddenly and unexpectedly contrary to what

occurs, for example, in components subjected to wear, where a reshape of the part is observed.

Therefore, due to the latter information, the study of this phenomenon reveals extremely important in

order to predict the component’s life.

This topic has been analysed through the years and one of the first studies was performed by August

Wӧhler in the 1850s. This study centred on rail axes which presented an inferior durability to the

expected by static analyse criterions. In this case, although the ductile nature of the material used, the

fracture presented no plastic deformation [1]. After performing several fatigue tests, it was found that not

only the cyclic stresses had an effect on fatigue, but also that the mean stress affected the fatigue

process (Figure 1.1).

Figure 1.1: Axial loading S-N at various mean stresses for unnotched specimens of an aluminium alloy [2]

The work from Wӧhler was a starting point served for several fatigue studies, such as the works

performed by Gerber and Goodman, which had the purpose of predicting the mean stress effects.

Since the fatigue process is promoted near regions where the stress intensity factor is higher, it was

important to introduce fracture mechanics concepts to evaluate the fatigue phenomenon. This was the

2

approach of Wells in order to justify the occurrence of failures at the “Comet” airplane (Figure 1.2). The

crack initiation was found at the plane windows, where the sharpened geometry of the window promoted

a local increase on stress. The material failure occurred as the crack reached a critical size.

Figure 1.2: Failure origin of Comet G-ALYU around square windows [3]

A few years later, Paris proposed a law, known as Paris Law [4], which relates the crack propagation

rate with the stress intensity range. This law allowed to characterize materials in terms of fatigue

strength, and it is used to characterize the material evaluated in this work.

The material analysed is a S700 MC-1, which is a high strength steel recently developed by SSAB [5].

It is characterized by the low presence of carbon and alloy elements. The characteristic high strength of

this steel is due to the grain refining process, phase transformation and precipitation hardening

(precipitates formation of Ni, Ti or V), which the steel undergoes [6]. This steel is manufactured by TMCP

(Thermo-Mechanically Controlled Process) which is composed by two main stages: the rolling stage at

high temperature, and the cooling phase where depending on the cooling rate, certain phases are

achieved [7]. Apart from presenting high strength, S700 MC-1 shows good weldability and high

toughness as well. The selection of this steel allows two possibilities, which are weight reduction, since

for a specific loading, the amount of steel needed is lower than for a steel which presents an inferior

strength. The second possibility is an increase in the durability of the structure and an increase in loading

capacity as well, maintaining the same overall structure weight.

Due to these properties, the S700 MC-1 reveals as an interesting material, which can be found in the

automotive industry, building cranes and ships.

3

1.2. Motivation

In engineering, most of the projects required permanent assemblies, which are obtained using mainly

fusion welding processes. The fusion welding process is a high energy process, which severely affects

the steel structure, due to high temperatures cycles in which steel undergoes. At the same time, the

mechanical properties are also affected, thus, the evaluation of the effect of the welding process on

mechanical properties such as hardness, fatigue strength and yield strength reveals extremely

important.

The focus of this work is the determination of the Paris law constants experimentally and evaluate the

hardness variation for different welding conditions for this material. Since it is a recent steel, the

information is scarce about it.

The evaluation of the best set of welding parameters is extremely important in this case, since, as

mentioned before, this steel reveals high strength which is crucial in several applications and the welding

process severely affects the strength of the materials. Moreover, the welding is performed with two weld

runs since the parts to weld present a considerable thickness which introduces a higher effect on the

steel properties. Therefore, the best set is achieved for the case where the strength is less affected,

allowing the use of this material for the defined applications without a considerable reduction on

mechanical performance.

Furthermore, the importance of performing experimental tests relies on the fact that, for example, the

Paris law constants determined experimentally for a specific condition can be used numerically and

treated by Finite Element Analysis (FEA) to simulate the behaviour of the same material for different

geometries. Therefore, there is no longer necessary to perform the same experimental test for each

condition, which results in a reduction in the number of specimens manufactured and consequently

reducing the overall costs.

1.3. Objectives

This thesis has the main objective of performing the material characterization of the high strength steel

(S700 MC-1) developed by SSAB, and this main objective can be divided into several objectives which

are the following:

Determination of the Paris law constants, when it is welded with a specific set of welding

parameters and for a specific welding geometry;

Evaluation of the influence of different welding parameters on hardness and to comprehend its

evolution along the heat affected zone and weld zone of the steel analysed. From the

hardness evolution, another objective is revealed, which is to perform an identical analysis of

the influence of different welding parameters, however, in this case, on Yield strength.

Analyse the same influence, which was referred in the previous point, on the microstructure, in

certain regions of the steel in question.

Examine the influence of the different weld geometries on the stress intensity factor for several

cracks lengths and also along the specimen’s thickness. These results are compared with the

unwelded scenario.

4

1.4. Thesis Outline

The present dissertation is divided into several chapters, which are enumerated below:

In Chapter 1 a brief introduction to the work performed is made and the several objectives of this work

are presented.

In Chapter 2 are reviewed the most important topics considered in this work. A brief introduction to

Fracture Mechanics is made. Then, a solid theoretical background about the distinct fatigue phases,

especially phase I and II, is performed. Similar fatigue propagation tests, which studied the influence of

parameters, such as stress ratio, microstructure and grain size on crack propagation rate, are meticulous

analysed as well. An introduction to the MIG/MAG process is performed and subsequently a detailed

analyse to the welding influence on the steel microstructure is done. Similar tests (hardness tests)

carried out with a similar steel are scrutinized as well. It is also discussed some aspects that can

influence the Young’s modulus. In the end, the Digital Image Correlation method used is explained in

detail.

Following the Literature review, both chemical and mechanical properties of the material analysed are

introduced, as the welding procedure and the different welding parameters used to perform the welds

are presented in Chapter 3.

In Chapter 4, the fatigue tests conditions are presented, as the preparation stages, in which the four

samples were subjected in order to perform the hardness tests and the microstructure analysis. The

conditions and decisions taken for each case are also explained in this chapter. It is introduced the

correlation used to determine the Yield strength along the distinct samples as well.

Chapter 5 consists on a detailed explanation of the Numerical simulation carried out to evaluate the

evolution of the Stress Intensity Factor (SIF) value for different welding geometries, where is mentioned

the loading conditions and mesh properties used in the simulation.

The results are presented and discussed in chapter 6. It is divided into five groups; the fatigue results,

hardness results, Yield strength results, microstructure analysis and at last the results of the numerical

analysis.

In the end, in Chapter 7, an overall conclusion and recommendation for future works are discussed.

5

2. Literature Review

In this chapter, a discussion is made about the main topics of this study. Initially, Fracture Mechanics

concepts are introduced. Afterwards, an approach on fatigue process is made, and it is analysed the

influence of several parameters on fatigue crack propagation rate. Posteriorly, some aspects of

MIG/MAG welding and the typical characteristics of the distinct zones which are formed due to the

welding process are presented. Furthermore, the influence of some aspects, such as the grain size and

phase transformation on Young’s modulus is analysed. In the end, a brief introduction relative to Digital

Image Correlation method is made.

2.1. Fracture Mechanics

Fracture Mechanics can be divided into two main approaches, depending on the physical

considerations: Linear Elastic Fracture Mechanics (LEFM) and Elastic-Plastic Fracture Mechanics

(EPFM).

The LEFM was developed by Griffith and Irwin [8] and it is applied to materials, which present a linear-

elastic behaviour and the inelastic behaviour occurs in a small region. Therefore, with an inelastic zone

relatively small, the influence of that region is also small and LEFM principals can be applied to

characterize the material behaviour. In this case, the fracture surface presents a shining and granular

aspect.

In EPFM, the plastic deformation is considerable, due to the application of high stresses. In order to

increase the plastic deformation, the specimen thickness can also be reduced. Therefore, in this case,

the fracture occurs due to a ductile process. The fracture surface is rough and shows a dull aspect.

From Fracture Mechanics, an important concept was developed, in order to evaluate the crack

propagation, when a loading is applied on a component, which is the stress intensity factor, 𝐾. This

parameter measures the raising effect on stress, due to the presence of stress concentration regions.

𝐾 depends on several conditions, such as loading mode, crack size and specimen dimensions. When

the stress intensity factor reaches a critical value, 𝐾 acquires the value of 𝐾𝑐, toughness, and the

specimen failure takes place [9].

2.1.1. Loading Modes

Taking into account a solid specimen with an arbitrary crack subjected to a determined loading, different

crack opening modes can occur. Mode I (in-plane loading), Mode II (in-plane sliding) and Mode III (out

of plane loading) are present in Figure 2.1.

6

Figure 2.1: Loading modes types [10]

In Mode I, the crack growth direction is perpendicular to the loading plane, due to the occurrence of

tension stresses at the crack tip. Mode II, also defined as sliding mode, is characterized by shear

stresses at the crack tip since, the crack propagation and the loading have the same direction. Finally

mode III, known as tearing mode, consists in a loading, parallel to the crack front, which also promotes

the occurrence of shear stresses at the crack tip.

2.1.2. Stress Intensity Factor

The stress intensity factor (SIF), 𝐾, as the name indicates, quantifies the stress field intensity, near the

crack tip, when a loading is applied to a cracked body. This Fracture Mechanic parameter is affected by

several aspects, as mentioned previously. Depending on the loading mode, 𝐾 presents different values:

𝐾𝐼 for mode I, 𝐾𝐼𝐼 for mode II and finally 𝐾𝐼𝐼𝐼 for mode III. In this work, the analysed mode was the opening

mode (Mode I) and for that reason, only the relations for this mode are discussed in this document.

Authors as Westergaard and Irwin [8], established the relation between the stress intensity factor and

the resulting stress field, when a cracked part is subjected to an external load. Defining a polar

coordinate axis with the origin at the crack tip (Figure 2.2), the stress field, in a cracked body, assuming

isotropic linear material behaviour, can be described by the following equation:

𝜎𝑖𝑗(𝑟, 𝜃) =

𝐾

√2𝜋𝑟𝑓𝑖𝑗(𝜃) + ∑ 𝐴𝑚𝑟

𝑚2

∞

𝑚=2

𝑔𝑖𝑗(𝑚)(𝜃) (2.1)

where 𝑟 expresses the distance to the crack tip, 𝜃 is the angle between the crack propagation direction

and the horizontal direction, as represented in Figure 2.2, while 𝑓𝑖𝑗 represents a dimensionless function,

which depends of the applied loading and geometry. 𝐴𝑚 is the amplitude and 𝑔𝑖𝑗 is a dimensionless

function of 𝜃 [9].

7

Figure 2.2: Stresses near the crack tip [11]

The field stress near the crack tip is given when 𝑟 → 0. From Equation 2.1, and considering the previous

condition, the first term, presenting an 1/√𝑟 evolution, tends to infinite, while the other terms tend to

zero. As a result, regardless the crack body geometry, the stress evolves with a 1/√𝑟 factor. The crack

tip is a singularity point since, stress presents an asymptote at 𝑟 = 0.The stress field, for Mode I, is

described by:

𝜎𝑥𝑥 =

𝐾𝐼

√2𝜋𝑟cos (

𝜃

2) [1 − sin (

𝜃

2) sin (

3𝜃

2)] (2.2)

𝜎𝑦𝑦 =

𝐾𝐼

√2𝜋𝑟cos (

𝜃

2) [1 + sin (

𝜃

2) sin (

3𝜃

2)] (2.3)

𝜏𝑥𝑦 =

𝐾𝐼

√2𝜋𝑟cos (

𝜃

2) sin (

𝜃

2) cos (

3𝜃

2) (2.4)

𝜏𝑥𝑧 = 𝜏𝑦𝑧 = 0 (2.5)

At direction z, stress can be given by two different expressions, depending on the plane state.

𝜎𝑧𝑧 = {

0 𝑓𝑜𝑟 𝑝𝑙𝑎𝑛𝑒 𝑠𝑡𝑟𝑒𝑠𝑠

𝜈(𝜎𝑥𝑥 + 𝜎𝑦𝑦) 𝑓𝑜𝑟 𝑝𝑙𝑎𝑛𝑒 𝑠𝑡𝑟𝑎𝑖𝑛 (2.6)

where 𝜈 is the material´s Poisson coefficient.

The crack-tip displacement field was also expressed and the respective functions, for x and y direction,

are presented in Equations (2.7) and (2.8).

𝑢𝑥 =

𝐾𝐼2𝜇√𝑟

2𝜋cos (

𝜃

2) [𝜅 − 1 + 2 sin2 (

𝜃

2)] (2.7)

𝑢𝑦 =

𝐾𝐼2𝜇√𝑟

2𝜋sin (

𝜃

2) [𝜅 + 1 − 2 cos2 (

𝜃

2)] (2.8)

where 𝜇 represents the shear modulus while, 𝜅 = 3 − 4𝜈 and 𝜅 = (3 − 𝜈)/(1 + 𝜈) for plane strain and

plane stress condition, respectively.

8

Assuming 𝜃 = 0, the stress state is only given by normal stresses since, all shear stresses are null,

resulting in a crack plane coincident with the principal stress plane. Moreover, 𝜎𝑥𝑥 and 𝜎𝑦𝑦 present the

same evolution equal to 𝐾𝐼

√2𝜋𝑟 .

In Figure 2.3, it is represented the evolution of 𝜎𝑦𝑦, as a function of the crack tip distance. It is clear to

see that, near the crack tip, the stress is well adjusted by the previous evolution. For a far point, the

same does not occur and the stress is equal to the applied stress, 𝜎∞.

Figure 2.3: Stress normal to the crack plane in Mode I [9]

For mode I, 𝐾 is related with the associated stress value by:

𝐾𝐼 = lim𝑟→0

√2𝜋𝑟 ∙ 𝜎𝑖𝑗(𝐼)(𝑟, 𝜃)

𝑓𝑖𝑗(𝐼)(𝜃)

= lim𝑟→0

√2𝜋𝑟 ∙ 𝜎𝑦𝑦(𝑟, 𝜃) (2.9)

The stress intensity factor, in general, is described, also as:

𝐾 = 𝑌𝜎√𝜋𝑎 (2.10)

𝑌 represents a dimensionless factor which, depends on the specimen geometry but also, on the crack

size, 𝜎 the stress applied, while 𝑎 is the crack length. Since mode I is the most analysed mode, for

specific specimens subject to this mode, it was defined Equation (2.12) to evaluate the SIF value on

those specific specimens [12]:

𝐾𝐼 =𝑃

𝐵√𝑊𝑓 (

𝑎

𝑊) (2.11)

Where P is the applied force, B the specimen thickness, W is the useful width while, 𝑓 (𝑎

𝑊) is a

dimensionless factor, which depends of the specimen geometry and loading mode. It is presented in

Figure 2.4, the schematic of the specimen used in this work and the respective 𝑓 (𝑎

𝑊) factor.

9

Figure 2.4: Specimen geometry and respective f (a/W) factor[9]

2.1.3. Plane Strain and Plane Stress

As seen before, the field stress differs and either a plane strain or a plane stress can take place. This

fact is, typically, due to the specimen thickness. In a smaller thickness specimen, a triaxial stress state

is less noticed, being only noticed near the thickness middle point. In a rough approximation and

considering a specimen with an infinitesimal thickness, it is possible to justify the previous sentence.

Considering the square, Q, (Figure 2.5), the surface is a discontinuous stress line. Therefore, normal

and tangential forces, at the surface, have to match. For that reason, the normal stress to the surface, 𝜎𝑧,

is null and the tangential shear stress to the surface, is also null. For equilibrium reasons, the opposite

shear stresses are null as well, leading to a plane stress state. The previous approximation considers a

limit case, however, the triaxial stress state is always present.

In a thicker specimen, the presence of a triaxial stress state is more noticed, as represented in Figure

2.6. When the stress is applied in direction y, at the crack tip the material tends to contract in the other

directions. On the other hand, the surrounding material oppose against that tendency, constraining the

crack tip material, causing then a triaxial stress state. The approximation to a plane strain relates to the

fact that, in the central position, the deformation in the thickness direction is negligible, compared to the

plate thickness.

Figure 2.5: Plane stress fracture. Plastic zone diameter r0 comparable to or greater than sample thickness. Modified from [13]

10

Figure 2.6: Plane strain fracture. Plastic zone diameter r0 much less than sample thickness [13]

Each stress state lead to a different fracture surface as seen in Figure 2.5 and 2.6. In plane stress

condition, the fracture occurs by shear lips (45o shear fracture). The formation of this type of fracture

(shear lips) requires a great amount of energy, leading to a higher value of 𝐾𝐼𝐶, as shown in Figure 2.7.

Although the plain strain condition also has the presence of shear lips, these are located only near the

surface, and consequently the fracture surface is flat

In Figure 2.7 it is illustrated the evolution of toughness, regarding the specimen thickness and it is seen

that, as thickness increases, the fracture toughness decreases, until reaches a minimum value, 𝐾𝑐 ,

which occurs for the plane strain condition. After 𝐾𝑐 is reached, the fracture toughness remains constant

and independent of the thickness value. It is important to notice that, the represented beach marks are

due to fatigue precracking at a low cyclic stress, to guarantee a sharp crack tip. Therefore, it is promoted

crack growth at that location [14].

Figure 2.7: Effect of specimen thickness on fracture toughness [14]

11

2.1.4. J Integral

Integral J represents a contour in which it is possible to evaluate a stress field and the respective

deformation in the vicinity of a crack tip, for linear and non-linear elastic materials. It was developed by

Rice [15] and it was based on the “energy momentum tensor” concept formulated by Eshelby [16], which

has the purpose of characterizing generalized forces on dislocations in elastic fields.

Defining an arbitrary counterclockwise path surrounding the notch tip (𝛤, Figure 2.8), and starting from

the lower and moving to the upper flat notch surface, integral 𝐽 is expressed by:

𝐽 = ∫ (𝑆𝑑𝑦 − 𝑇 ∙𝜕𝑢

𝑑𝑥𝑑𝑠)

𝛤

(2.12)

where 𝑇 represents the tension vector which is characterized by 𝑇𝑖 = 𝜎𝑖𝑗𝑛𝑗 with the same direction of 𝑛

(outward normal along 𝛤), represented in Figure 2.8. 𝑢 is the displacement vector, while 𝑑𝑠 is an

infinitesimal arc length element along 𝛤. 𝑆 is the strain-energy density defined by (Eq. (2.14)):

𝑆(휀) = ∫ 𝜎𝑖𝑗𝑑𝜖𝑖𝑗

𝜖

0

(2.13)

where 𝜖 expresses the strain tensor, while 𝜎𝑖𝑗 represents all stresses present.

Figure 2.8: Path for the calculation of J-integral

Rice also observed that 𝐽 can be defined as the energy release rate, which occurs in a nonlinear elastic

body presenting a crack (Eq. 2.15). For a linear elastic body, 𝐽 is equivalent to 𝐺.

𝐽 = −(𝑑𝑈

𝑑𝑎) (2.14)

Where 𝑈 represents the strain energy stored in the body and 𝑎 is the crack area with a unitary thickness.

Moreover, the crack propagation takes place when the energy release rate reaches a critical value. This

value occurs when the stress intensity factor is also achieved at the crack tip.

A relationship between 𝐾 and 𝐽 was founded, as well, when it is considered a homogeneous, isotropic

and linear elastic material. For mode I and for a plane stress state the relation is:

𝐽𝐼 = 𝐾𝐼2

𝐸 (2.15)

12

For plane strain state, the relation is given by the following equation:

𝐽𝐼 = 𝐾𝐼2 (1 − 𝜈2

𝐸) (2.16)

2.1.5. J and K evolutions along the thickness

In section 2.1.3 it was observed the thickness influence on the fracture toughness. However, the stress

intensity factor is not constant along the thickness. The Finite Element Analysis (FEA) introduced an

important role to evaluate the SIF evolution along the thickness. Several works were developed in order

to study this evolution. Nakamura [17], defined two 𝐽 integrals, the local, 𝐽𝑙𝑜𝑐𝑎𝑙, obtained by the domain

integral method, and 𝐽𝑓𝑎𝑟, which is equal to the 𝐽 definition (𝐾𝐼2/𝐸) for the plane stress condition. In order

to evaluate the evolution of local J, 𝐽𝑙𝑜𝑐𝑎𝑙 was normalized by 𝐽𝑓𝑎𝑟. This evolution was analysed for

different load levels and it is presented in Figure 2.9. It is important to refer that z is the distance to the

centre of the specimen and t is the specimen thickness.

Figure 2.9: Local J normalized by far-field/remote J, plotted vs position along the half-crack front, at various loading levels [17]

As it is possible to observe, the maximum value occurred at the centre of the specimen, while the

minimum value took place at the surface. The 𝐽𝑙𝑜𝑐𝑎𝑙 tended to be higher than the 𝐽𝑓𝑎𝑟 in the majority of

the thickness, especially for lower load levels. Near the surface, the 𝐽𝑙𝑜𝑐𝑎𝑙 value presented an abrupt

reduction. The increase on the load level promoted an increase on the difference between the maximum

and minimum value of local 𝐽. For smaller load levels, it was also seen that the opening displacement

presented the plane strain solution along the crack front, except near the surface which, indicated the

existence of a plane strain state along the crack front. As the load level increased, the plane stress

solutions took place which is in agreement with the remote boundary condition, since the defined

specimen thickness was small. From [18], it was observed a similar evolution to the one presented in

the previous case, for normalized SIF value (straight front condition, Figure 2.10), where 𝐾𝑗 is the local

13

SIF value. As happened, in the case above, the evolution was only evaluated for half thickness since,

the specimen was symmetric and hence, the evolution was also symmetric.

Figure 2.10: Kj based evolution through the thickness for different crack front shapes [18]

2.2. Fatigue

Fatigue can be defined as a permanent structural changing process, progressive and localized that

occurs in a material subjected to stresses and cycled deformations in a point or in several points, which

can lead to cracks or the complete fracture, after a determined number of cycles.

The process is progressive due to the fact that occurs along a certain period of time and localized since,

tends to appear in small areas, where geometric constrains are found (stress concentration factor),

residual stress or material defects.

Linear Elastic Fracture Mechanics, LEFM, considers that all materials contain flaws. Those flaws

reduces the strength of a component, since these are intensity stress raisers. On an initial phase, the

size of those flaws is small and insufficient to cause the component failure. However, the crack growth,

as occurs in most engineering cases, is promoted due to fatigue. The fatigue process can also promote

the formation of cracks. An arbitrary fatigue cycle is shown in Figure 2.11.

Figure 2.11: Fatigue cycle [19]

14

The variable 𝜎𝑎, stress amplitude and 𝜎𝑚, stress midrange, presented in Figure 2.11, are defined by

the following equations:

𝜎𝑎 =𝜎𝑚𝑎𝑥 − 𝜎𝑚𝑖𝑛

2 (2.17)

𝜎𝑚 =𝜎𝑚𝑎𝑥 + 𝜎𝑚𝑖𝑛

2 (2.18)

Where 𝜎𝑚𝑎𝑥 and 𝜎𝑚𝑖𝑛 are the maximum and minimum stress applied. In fatigue, it is also defined 𝑅,

as the stress ratio, which is given by:

𝑅 =𝜎𝑚𝑖𝑛𝜎𝑚𝑎𝑥

(2.19)

In most engineering cases, the stress range (∆𝜎 = 𝜎𝑚𝑎𝑥 − 𝜎𝑚𝑖𝑛) is significantly inferior to the material’s

Yield stress, however, the fatigue failure of a structure due to crack growth still occurs. The stress

intensity factor also varies between a maximum and minimum value and ∆𝐾 is given by:

∆𝐾 = 𝐾𝑚𝑎𝑥 − 𝐾𝑚𝑖𝑛 = (𝜎𝑚𝑎𝑥 − 𝜎𝑚𝑖𝑛)𝑌√𝜋𝑎 (2.20)

In Figure 2.21, it is represented the typical evolution of the fatigue crack propagation on steels, which is

divided into three main regimes: in Regime A, the crack nucleation process occurs and the crack

propagation rate is smaller at this stage. Regime B represents the stable crack growth phase, while in

Regime C the crack growth is unstable, leading to a catastrophic failure.

Figure 2.12: Primary fracture mechanisms in steels associated with sigmoidal variation of fatigue crack propagation rate (da/dN) with alternating stress intensity (ΔK)[20]

2.2.1. Fatigue stages

Applying a cyclic loading on a component can initiate the nucleation of a microscopic fatigue crack. This

crack continues to propagate due to the continuous application of the cyclic loading. As a result, the

cross-section area of the component is reducing, increasing the stress, until reaches a value that is

15

higher than the one that the component can support, leading to the component failure. In Figure 2.13

can be seen the crack evolution along a loading period.

Figure 2.13: Typical propagation of a fatigue crack [21]

Until material failure the fatigue process can be divided into two different periods, crack initiation and

crack growth period. Therefore, the component total life is the sum of the number of cycles occurred in

crack initiation the number of cycles in crack growth period. Depending on the stress level, the fatigue

process tends to occur differently. When a low stress level is applied, the majority of the fatigue life

occurs at nucleation and microcrack growth stages. For a high stress level, the fatigue life is mainly

spent at microcrack and macrocrack stages.

Crack initiation phase

From Figure 2.12, the fatigue process can be divided into 3 stages. Starting with the crack initiation

phase, this phase is characterized by microcrack development. The size of fatigue cracks, at the

beginning of the initiation period, can be smaller than 0.5 µm according to Fine [22]. The surface

condition has a large effect in this period since, surface grains are less constrained which, associated

with the presence of a defect, can promote the crack initiation for a lower stress level. The presence of

discontinuities in the material which, can be an inclusion, a secondary phase or even a pore, can also

promote the crack initiation since, the stress intensity is increased in those regions. In order to explain

crack initiation and comprehend the mechanical response of metals, it is important to mention the

influence of dislocations on material deformation. The dislocation movement promotes fatigue damage

and its density is increased, as inelastic deformation takes place. Between dislocations and inelastic

deformation, a relationship can be established and it is defined by “slip”. Slip consists on a shear

deformation, which occurs due to the movement of the dislocations along the crystallographic planes

within the individual grains. Depending on the grain, the movement of dislocations is more or less eased

since, every grain has its own preferential direction and specific directional properties. In a grain where

the planes are oriented with the direction of the maximum applied shear stress, the slip phenomenon is

eased, promoting inelastic deformation at that same location. Fatigue crack initiation occurs due to cyclic

plastic deformation and formation of Persistent Slip Bands (PSB). The designation ‘persistent’ is due to

16

the fact that, in a fatigue test where the specimen was polished and subjected to fatigue once again, the

slip activity occurred at the same location in both situations. The PSB are zones of high cyclic slip activity.

This activity leads to the occurrence of surface extrusions and intrusions, as illustrated in Figure 2.14.

Since the slip process implies some strain hardening in the slip band and despite the reversed slip

occurs, preferably on the same slip band, the strain hardening in the slip band is not fully reversible,

forming extrusions and intrusions depending on the formation direction. A surface intrusion is the start

of a fatigue microcrack. [23][24].

Figure 2.14: Development of extrusions and intrusions during fatigue [21]

At this stage, crack growth results, essentially, due to a shear process and it is affected by the material

grain size and the slip characteristics of the material (Figure 2.12).

Crack growth phase

The second fatigue stage is known for a larger crack growth than the existent at the nucleation stage,

where the crack growth occurs only for some grains (Figure 2.15, Stage I). The growth occurs, mainly,

perpendicular to the loading direction, as shown in Figure 2.15 and, in this case, the crack growth results

of a tensile process.

Figure 2.15: Schematic of stages I (shear mode) and II (tensile mode) transcrystalline microscopic fatigue crack growth Adapted from [14]

The reason behind this occurrence relies on the fact that, as the microcrack grows into the material, the

slip displacement tends to be constraint on account of the grains in the vicinity. This promotes the

occurrence of slip displacement on more slip planes, leading to a deviation from the initial orientation.

The general tendency is the growth to occur perpendicular to the loading direction, as mentioned before.

17

For this stage can arise three different fatigue crack growth modes, such as striation formation, Microvoid

Coalescence, (MVC), and microcleavage [14].

Striation formation is characteristic of fatigue cyclic damage. Some authors consider that one striation

is associated to one load cycle. Laird and Smith [25] exhibited that on the uploading moment, crack tip

blunting took place while, on the downloading moment, a striation and a crack tip re-sharpening was

witnessed on a high strains loading cycle. In order to understand the crack propagation phenomenon

due to striation formation, in Figure 2.16 is illustrated a model of saw-tooth type striations. This saw-

tooth pattern was observed by several researchers [26].

Figure 2.16: A model for the creation of saw-tooth type striations [26]

MVC, as the initials stand for, occurs during plastic deformation due to the coalescence of microvoids,

which leads to dimples, Figure 2.17. This process is a high-energy process, which takes place at high

crack growth rates. It is responsible for the cracking between inclusions and the surrounding material.

Finally, Cleavage is the most brittle form of fracture in crystalline materials. The crack growth requires

low energy and tends to appear in notched components or where plastic flow is constrained. The crack

develop along determinate crystallographic planes, forming cleavage facets as presented in Figure 2.17,

that are, in general, flat.

18

Figure 2.17: A higher magnification SEM photomicrograph of the transition from ductile overload fracture (dimples) of the weld HAZ to brittle overload (cleavage) fracture of the base metal [27]

The fatigue propagation mechanisms for materials, which present a ductile behaviour, is the ductile

striation growth. The growth is controlled by the alternating plastic strain per cycle, leading to a little

effect of mean stress. In a material, which presents a more brittle behaviour, the crack growth occurs by

the combination of ductile fatigue striations and intergranular cracking. The two propagation

mechanisms can be observed in Figure 2.18 (unembrittled case), where ductile striation growth is

presented and where a brittle intergranular cracking is shown (embrittled case). The arrow, in both

figures, indicates the crack propagation direction [28].

Figure 2.18: a) Ductile striation growth through tempered martensite during fatigue (unembrittled steel), b) Brittle intergranular cracking during striation growth in fatigue of embrittled steel [28]

From the information provided by Ritchie’s work [28], it is possible to conclude that the crack propagation

rate was higher for the embrittled steel case, as demonstrated in Figure 2.19.

19

Figure 2.19: Variation of regression lens and values of the slope m for fatigue crack growth rates of unembrittled steel for several stress ratios R [28]

From Figure 2.19 it was also seen, that the embrittlement of the steel, besides increasing the growth

rate, increased also the influence of the mean stress on growth rate. The mean stress influence on

growth rate is a consequence of the occurrence of ‘static’ fracture modes, associated with the striation

growth mechanism [28].

Paris [4] applied fracture mechanics concepts to evaluate crack propagation, due to the fatigue

phenomenon, defining the Paris law, Equation (2.21), for regime B.

𝑑𝑎

𝑑𝑁= 𝐶∆𝐾𝑚 (2.21)

Constants 𝐶 and 𝑚 are obtained experimentally and can vary depending on the stress ratio and the

microstructure. The influence of these aspects is presented further in this chapter. 𝑚 is typically found

between a range of 2-4 [28].For high strength steels, 𝑚, can present higher values, especially in steels

with a low fracture toughness.

Apart from DIC method, which was applied in this work, other methods have recently been developed,

in order to establish the Paris law [29]–[31].

Fatigue life,𝑁𝑓, considering only regime B, can be determined by integrating the Paris law between the

limits of initial and final crack size as presented below, considering ∆𝐾 = 𝑌∆𝜎√𝜋𝑎,

∫ 𝑑𝑁𝑁𝑓

0

= 1

𝐶(∆𝜎)𝑚(𝜋)𝑚/2𝑌𝑚∫

𝑑𝑎

𝑎𝑚/2

𝑎𝑓

𝑎𝑖

(2.22)

Integrating Eq. 2-22, fatigue life is obtained:

𝑁𝑓 =

{

𝑎

𝑓

(−𝑚2 )+1 − 𝑎

𝑖

(−𝑚2 )+1

(−𝑚2+ 1)𝐶(∆𝜎)𝑚(𝜋)

𝑚2𝑌𝑚

1

(𝐶∆𝜎𝑌√𝜋)2 ln (

𝑎𝑓

𝑎𝑖) , 𝑓𝑜𝑟 𝑚 = 2

, 𝑓𝑜𝑟 𝑚 ≠ 2 (2.23)

20

Where 𝑎𝑓 is the final crack length (critical), while 𝑎𝑖 is the initial crack length.

Although the extrapolation of the Paris law to region III is nonconservative, due to the shape of the curve

in that period, the integration of the Paris law over the entire crack growth may be considered reasonable

since, the crack growth life in region C is smaller than the one occurred in region A and B.

Instable crack growth and component failure phase

The third stage is characterized by an exponential increase in the crack growth rate and when the crack

length reaches a critical value, 𝑎𝑐, the material can no longer support further damage, leading to the

failure of the material. The failure mechanism at this stage is, in general, cleavage, as mentioned in

Figure 2.12.

2.2.2. Microstructure influence on fatigue crack propagation

As showed before, Figure 2.12, the fatigue crack propagation can be affected by different aspects.

Several works about the effect of microstructure, stress ratio and welding residual stresses on fatigue

crack propagation, especially for the second phase of the fatigue process are presented.

Starting with the microstructure influence, Barsom [32], studied the fatigue crack growth of several

martensitic and ferritic-pearlitic steels, at region B, and the results are presented in Figure 2.20.

Figure 2.20: Fatigue crack propagation data. a) Martensitic steels; b) Ferritic-pearlite steels

21

Analysing both cases, the C value, for ferritic-perlitic steels is much smaller compared to the respective

value for the martensitic steels. This fact may be explained by the crack branching phenomenon.

Contrary of what occurred in martenstic steels, A36 steel, a ferritic-perlitic steel, showed severe

secondary cracking at the crack tip. This promoted a reduction in the stress intensity factor, leading to a

smaller fatigue crack growth rate compared to the occurred in a single crack front situation. For the

matersitic case, Barsom also witnessed that, despite the different tested steels presenting a large

properties variation, the fatigue crack rate data was founded within a relatively narrow scatter band [32].

2.2.3. Stress ratio influence on the fatigue crack propagation

From Ohta and Sasaki work [33], it is possible to observe that stress ratio (𝑅) influence the crack

propagation curve, as demonstrated in Figure 2.21 and previously in Figure 2.19. The tests were

performed for a 𝑅 range of -1 to 0.8.

Analysing Figure 2.21, the threshold value was higher when it was applied a negative stress ratio.

Compressive loads introduced residual compressive stresses, which inhibited the crack growth, leading

to a higher threshold value. The opposite occurred, as 𝑅 increased, the threshold value decreased. It is

important to mentioned that the information given by Figure 2.21 was coherent with the information

presented by Cooke and Beevers [34] which mentioned that, for a 𝑅 range of [0.05; 0.73], at higher

crack growth rates, the 𝑅 variation became, practically, insignificant and similar crack propagation rates

were observed. Contrarily, for lower crack growth, the 𝑅 increment had an important role, as the

threshold level was reached. It is important to mention that, only Region A and Region B are represented

in the figure.

Figure 2.21: A family of curves showing effect of stress ratio on fatigue crack propagation of JES SM58Q steel [33]

22

2.2.4. Influence of welding and welding residual stresses on

fatigue crack propagation

Since, in this work, welded specimens are evaluated, it is important to comprehend the influence of the

welding effect on fatigue crack propagation. In Shi’s work [35], the welds were produced by submerged

arc welding, a MIG/MAG derivation, and from this work, it was seen that the crack growth rate was lower

for WM (weld metal) and CB (cross-bond, in this case the crack front was presented between the weld

metal and parent metal regions) comparing to the non-welded specimen (Figure 2.22 and Table 2.1,

∆𝐾). The reason behind these results relies in the residual stress field which occurred. Both residual

stresses, longitudinal and transverse led to a compressive stress field, which promoted the crack

closure, hindering the crack growth. It is important to notice that, the difference between the two BM

specimens was the rolling direction, which was used to manufacture both specimens. However, and

considering an effective stress intensity range, ∆𝐾𝑒𝑓𝑓 , in order to remove the residual stress effect, the

crack growth rate increased significantly in both cases (WM and CB). The difference between results

can be explained by the different microstructure occurred at the weld zone, which indicated poorer

fatigue strength. The CB case presented an intermediate value comparing to the remaining cases

(Figure 2.22 and Table 2.1, ∆𝐾𝑒𝑓𝑓).

Figure 2.22: Fatigue crack growth rate as a function of stress intensity range and effective stress intensity range a) Base Material; b) Weld Material and Cross-Bond [35]

Table 2.1: C and m in Base Material, Weld Material and Cross-Bond specimens [35]

A transverse weld was also considered and for a constant value of ∆𝐾, the crack growth variation along

the different regions (BM, HAZ, CB and WM) is illustrated in Figure 2.23.

23

Figure 2.23: Fatigue crack growth rate versus crack length for Transversal specimen, adapted from [35]

From the analysis of Figure 2.23, it is clear to observe that, in a first phase, as the crack grows towards

the HAZ, the crack growth rate tended to decrease, as the crack length increases, reaching a minimum

value at the boundary between heat affected zone and CB zone. Then, an increase on the crack growth

rate occurred. This increase was more accentuated as the distance to the weld material increased. This

occurred since, in a first phase, a compressive state was present, leading to a reduction on the growth

rate. As the crack tip reached the weld centre, a tensile state took place instead. Initially, the increase

on crack growth rate was relatively small due to the occurrence of crack closure, as a result of a residual

stress redistribution. Posteriorly, with the crack length increment, the crack growth rate tended to

increase significantly since, the residual stresses presented lower values [35].

2.2.5. Crack propagation laws

Paris law was a starting point on Fatigue propagation laws. Several laws were developed after Paris

studies, such as Forman, Walker and Nasgro laws, which are presented below.

Forman law

Forman witnessed that, for high load ratios, the Paris law was inadequate. Furthermore, he recognized

that the Paris law did not consider two different effects which were the stress ratio, 𝑅 in the crack growth

rate and the crack growth instability, when the maximum stress-intensity factor (SIF) reaches the

material fracture toughness, 𝐾𝑐. In 1967, R. Forman [36] proposed a new crack propagation law:

𝑑𝑎

𝑑𝑁=

𝐶′(∆𝐾)𝑚′

(1 − 𝑅)𝐾𝑐 − ∆𝐾 (2.24)

Thus, the mean stress effect was introduced, as the fatigue crack growth behaviour of region C on the

crack propagation law. However, since it is difficult to determine 𝐾𝑐, the application of this law becomes

complex. Nevertheless, this law presents good results for Al-Alloys and high strength steels.

Walker law

Although the stress ratio effect was also consider on Walker’s law (Equation (2.27)), C and m vales,

were defined as equal to the Paris law values for 𝑅 = 0 case [37].

24

𝑑𝑎

𝑑𝑁= 𝐶 (

∆𝐾

(1 − 𝑅)1−𝑤)𝑚

(2.25)

Where 𝑤 is a material constant.

Nasgro law

While the previous laws were defined for 𝑅 ≥ 0, this law took into account compressive loading. Apart

from that, it considers regime A and regime C, as well [38].

𝑑𝑎

𝑑𝑁= 𝐶 (

(1 − 𝑓)

1 − 𝑅∆𝐾)

𝑛 (1 −∆𝐾𝑡ℎ∆𝐾

)𝑝

(1 −𝐾𝑚𝑎𝑥𝐾𝑐

)𝑞 (2.26)

Where the crack opening function, 𝑓, satisfies the following conditions [39]:

𝑓 = {𝐴0 + 𝐴1𝑅 + 𝐴2𝑅

2 + 𝐴3𝑅3 𝑅 ≥ 0

𝐴0 + 𝐴1𝑅 − 1 ≤ 𝑅 < 0 (2.27)

𝐴0 = (0.825 − 0.34𝛼 + 0.05𝛼2) (cos(𝜋 (

𝑆𝑚𝑎𝑥2𝜎0

)))

1𝛼

(2.28)

𝐴1 = (0.415 − 0.071𝛼)𝑆𝑚𝑎𝑥𝜎0

(2.29)

𝐴2 = 1 − 𝐴0 − 𝐴1 − 𝐴3 (2.30)

𝐴3 = 2𝐴0 + 𝐴1 − 1 (2.31)

Where C, n, p, q, ∆𝐾𝑡ℎ, 𝐾𝑐, 𝛼 and 𝑆𝑚𝑎𝑥/𝜎0 are material parameters.

More recently, Dhondt proposed a new crack propagation law [40], which has the aim of avoid the

redetermination of the Paris law parameters, when are varied several effects in the experimental test,

such as the stress ratio, 𝑅. In order to minimize the corrections influence on the Paris-type law, the

author used the multiplicative formulation to introduce the effect of those corrections. The respective law

is:

𝑑𝑎

𝑑𝑁= [(

𝑑𝑎

𝑑𝑁)𝑟𝑒𝑓

(∆𝐾

∆𝐾𝑟𝑒𝑓)

𝑚

]𝑓𝑅 ∙ 𝑓𝑡ℎ𝑓𝐶

(2.32)

Where 𝑓𝑅 is the R-correction factor, 𝑓𝑡ℎ is the threshold correction, 𝑓𝐶 is the correction factor at high

fatigue crack growth rates which are presented following:

𝑓𝑅 =1

(1 − 𝑅)(1−𝑤)𝑚 (2.33)

𝑓𝑡ℎ = (1 −∆𝐾𝑡ℎ∆𝐾

)𝑝

(2.34)

𝑓𝐶 = 1 − exp [𝛿 (𝐾𝑚𝑎𝑥𝐾𝐶

− 1)] , 𝐾𝑚𝑎𝑥 < 𝐾𝐶 (2.35)

The term in square brackets represents the Paris law where C is given by:

𝐶 = (𝑑𝑎

𝑑𝑁)𝑟𝑒𝑓

(1

∆𝐾𝑟𝑒𝑓)

𝑚

(2.36)

Where, the reference values used were determinate in previous experimental test.

25

The previous laws referred are purely empirical and none was defined considering the mechanisms of

fatigue crack propagation. Nevertheless, the different laws presented reasonably results for a specific

interval or region. The Paris law was the chosen to compute the several crack propagation laws,

presented in this work, due to its simplicity to apply and acceptable accuracy

2.3. MIG/MAG (Metal Inert Gas/ Metal Active Gas)

Welding

The welding process can be described as a joining process that produces coalescence of materials by

heating them to the welding temperature, with or without the application of pressure or by the application

of pressure alone, and with or without the use of filler metal.

There are many welding processes and they can be classified due to the physic state of the different

parts which compose the weld: fusion welding, solid welding or liquid-solid phase welding. In fusion

welding (MIG/MAG case) occurs the partial fusion of the base material and if it exists, the fusion of the

filler material in order to join the different parts. In MIG/MAG, Figure 2.24, the heat source is the electric

arc. The arc is established between two electrodes where, one is a consumable wire, which is fed into

the system with a continuous speed and the other is the workpiece. A gas or gas mixture, designated

as shield gas is also used. It has objective the protection of both electrodes, the arc and the weld pool

from the atmospheric contamination. The process is designed as MAG when the shield atmosphere

contains a percentage of an active gas (CO2 and O2 are more frequently used), even if the percentage

is much lower than the inert gas. On the other hand, the process is considered as MIG when the

protective atmosphere is only composed of inert gases (the more common are Helium and Argon) [41].

Figure 2.24: MIG/MAG process [33]

2.3.1. Arc physics

Before moving forward, it is necessary to understand basic principles of the electric arc. The arc can be

defined as a conductor gas, which transforms electric energy into heat. An electrical discharge between

the two electrodes creates the electrical arc. One way to create an electric arc is to promote the contact

between the two electrodes. Then, it is created a very large short circuit current that is enough to melt

and vaporize some metal from the electrodes, due to the high contact resistance. The presence of

metallic atoms reduces the ionization energy, leading to an inferior ionization voltage compared to the

26

available voltage in the welding machine and subsequently the arc is established. Afterwards, the

negative electrode (cathode) emits electrons to the positive electrode (anode) and the positive ions

dislocate to the cathode, as the welding process takes place [41].

2.3.2. Consumables

Electrode-wire

The electrode used in MIG/MAG welding is a filler wire. The electrode suitability is defined by the

metallurgical, physical and chemical compatibility with the base material. The filler wire can have a

different chemical composition comparing to the base material, in order to improve mechanical

properties, such as ductility or even to avoid the hot cracking phenomenon, increasing the content of

manganese. The electrode wire can be solid or tubular. The tubular electrode, can have metallic or non-

metallic core powder. From Figure 2.25, it is possible to visualize that cored-wire presents a higher

melting rate compared to the solid-wire and covered electrode cases. The reason relies on the fact that,

the energy density is higher on the cored-wire case since, the core of the cored wire is not conductive.

For the two scenarios (cored-wire and solid-wire), the melting rate is higher in smaller wires since, the

energy density is higher as well [42].

Figure 2.25: Comparison of melting rates in welding with a covered electrode, a solid wire and a cored wire, respectively [42]

Shield gas

Reaching the fusion temperature, most metals show a considerable propensity to oxidize, forming

oxides when the weld pool is exposed to the atmosphere. This lead to poorer properties at the weld.

The main function of the gas shield is to avoid the contact between the weld pool and the oxidant

atmosphere. In addiction the shield gas also influences other weld properties, as those presented in

Figure 2.26.

It is important to mention that in a shield gas, where the main composition is an inert gas, the addition

of active gas, as CO2, even in low concentration, allows a more stable arc since, the ionization energy

decreases with the presence of an active gas, improving the welding process. However, a higher

percentage can promote the weld oxidation, leading to poorer properties [41].

27

Figure 2.26: Influence of shielding gas in MIG/MAG welding processes [42]

2.3.3. Welding process parameters

Although shield gas has an influence on the welding process, there are parameters which can also

modify the quality and shape of the weld bead, such as the current, voltage, polarity and welding speed

[41][43]:

Current Intensity

Current intensity is a very important parameter since influence the metal transfer mode. An increase on

current promotes a wider and deeper weld bead. The opposite occurs when a decrease on the current

value takes place.

Arc voltage

Increasing this parameter can promote a larger weld seam, however, it may reduce penetration and an

excessing value of voltage can result in porosity and spattering. On the other hand, an arc formed with

a low voltage can promote porosity but also produces a sharply convex weld bead.

Welding speed (linear speed of the torch)

The increase of this parameter results on a decrease of the welding bead penetration and it leads to a

narrower weld bead, by allowing less material deposition. However, with a lower welding speed, it is

promoted a higher material deposition, leading to a wider bead.

Type of current and Polarity

In order to create an electric arc on MIG/MAG process, it is used direct current since, with an alternate

current the arc becomes unstable when current intensity reaches zero. In most cases, the workpiece is

connected to the negative pole and the torch is connected to the positive pole (DCEP- Direct current

electrode positive). The arc formed is more stable, allowing a more regular metal transfer. Direct current

electrode negative (DCEN) may increase the deposition rate since, in this case, higher temperatures

are reached at the electrode-wire. However, it produces a more irregular metal transfer and thus, a

poorer weld bead. Tests have been performed to enable the application of alternate current to capitalize

the benefits of using DCEP and DCEN [44].

28

2.4. Steel Microstructure

Steel is known for the distinct phases which can present at different temperatures and with different

carbon percentage, below 2.14%. For a higher carbon percentage, it is considered a cast iron. Figure

2.27 shows the iron diagram, which illustrates the phase evolution through steel undergoes, at different

temperatures for different carbon percentages.

Figure 2.27: Iron diagram phase [45]

Below the eutectoid temperature (A1 in Figure 2.27), steel can present 3 different structures, depending,

mainly, on the cooling rate which steel, in Austenite phase, is exposed (Figure 2.28).

Figure 2.28: Different structures obtained in steel after cooling from austenite by different cooling rates [46]

Slow cooling rate promotes a more efficient carbon redistribution. At temperatures close to the eutectoid

temperature, the carbon diffusion process is promoted, resulting in the formation of thick layers of ferrite

and pearlite. Fine pearlite is formed at lower temperatures where the carbon diffusion rate decreases,

leading to thinner layers [46]. As the cooling temperature decreases, the carbon diffusion rate decreases

as well. Two different microstructures can be found, upper bainite (UB) and lower bainite (LB), Figure

2.29. In upper bainite, carbon concentrates at the phase boundaries. Therefore, ferrite grows rapidly

29

and cementite (Fe3C) precipitates between the ferrite grains. For lower bainite, the carbon fusion rate is

even lower promoting the cementite formation inside the ferrite grains (lath shape) [46].

Figure 2.29: Upper Bainite (UB), Lower Bainite (LB) and Polygonal Ferrite (PF) microstructures [47]

In the fast cooling rate case, the temperature gradient is extremely accentuated. As consequence, the

transformation is a diffusionless process and martensite grains nucleate and grow rapidly, presenting a

lath shape (Figure 2.30 b)). In Figure 2.30 a), the martensite transformation is demonstrated by line b.

Figure 2.30: a) Isothermal transformation diagram for a eutectoid steel [46]; b) Ferritic-martensitic structure [48]

The previous isothermal transformation diagram is characteristic of a steel presenting a eutectoid

composition. For a different composition, the diagram can vary, depending on the elements present in

the steel. For a high strength steel, the diagram presents a typical evolution, as the illustrated in Figure

2.31. From the figure, it is possible to see that different microstructures may occur, either purely

martensitic, dual-phase, TRIP (transformation induced plasticity) or complex phase.

30

Figure 2.31: Microstructure evolution during thermal processing of high advanced high strength steels [49]

2.4.1. High strength steels

The studied steel in this work is a high strength steel. This type of steels are produced by a Thermo-

Mechanically Controlled Process, TMCP. This process is divided, essentially, into two phases: hot rolling

and cooling stage, allowing to achieve a fine grain size, which improves the strength and toughness of

the steel. The improved strength is explained by the Hall-Petch effect, which indicates that the strength

is related to the grain size by [46]:

𝜎𝑦 = 𝜎0 +𝑘𝑦

√𝑑 (2.37)

Where d is the average grain diameter while 𝜎0 and 𝑘𝑦 are material constants.

According to Figure 2.32, the rolling phase is composed of three steps.

Figure 2.32: Different rolling stages of TMCP [6]

At the first step, the rolling is performed at a temperature which the grain recrystallization occurs. The

rolling process tends to elongate the material grains. However, at this temperature, the grains

recrystallize, resulting in a considerable grain size reduction, maintaining a non-elongate shape.

31

The second step is executed below the recrystallization range, avoiding the grain recrystallization. In

this case, the austenite grains are deformed, presenting an elongated shape. At this stage, the

nucleation of ferrite occurs inside the austenite grains.

Finally, at the third stage, the grains are deformed and elongated once again, hardening even more the

steel. The formation of subgrains structures, ferrite, occurs inside the grains.

After the hot rolling stage is completed, the cooling stage takes place. Depending on the cooling rate,

the microstructure can vary, as mentioned before. The different microstructures, which can be achieved

by distinct cooling rates, are presented in Figure 2.33.

Figure 2.33: Microstructural control by different cooling routes after hot rolling [7]

2.4.2. Heat Affected Zone microstructure

The fusion welding process is high temperatures process. Depending on the reached peak temperature,

the microstructure can present significant variations between zones, presenting different properties.

In Figure 2.34 is illustrated the different domains and the approximate temperature that leads to each

domain formation.

Figure 2.34: Typical domains of heat affected zone in welds for steel [50]

Solidified weld is characterized by the presence of elongated grains, oriented towards the weld centre.

The grain starts to solidify from the melted surface of the base metal and it acquires the same orientation

32

of the respective base material grain. It is important to also mention the fact that the grains tend to

present typically considerable dimensions, demonstrating poor properties. It is possible to improve the

properties by decreasing the grain size and promote the formation of acicular ferrite since, it improves

the weld toughness [51]. Inclusions (mainly oxides) tend to provide favourable regions for intragranular

nucleation of acicular ferrite. The ideal oxygen concentration range for acicular ferrite formation is

between 150-350 ppm [52].

Solid-liquid transition zone, also known by fusion line, consists in a discontinuity separation between the

weld metal and the unmolten base material. From Figure 2.35, it is possible to observe that, at the fusion

line, there is a significant gradient in terms of grain size. As the peak temperature falls, the grain size

decreases and a hardness reduction is witnessed. The maximum hardness point is typically found

adjacent to the fusion boundary as showed in Figure 2.35.

Figure 2.35: Vickers hardness measurements in the HAZ of a structural steel and grain size variation [53]

Next to the fusion line is located the Coarse Grained Heat Affected Zone (CGHAZ). It occurs when the

peak temperature is above 1100 0C according to Figure 2.34. At this zone, the austenite grain size of

the base material is severely increased. When the cooling rate is high, it is promoted the formation of

hard and brittle phases as martensite. Due to this fact, the hardness at this zone can be higher than the

initial hardness of the base material. The CGHAZ can be the more problematic zone due to the presence

of brittle and hard phases which reduce the material toughness and increases the probability of cold

cracking [53]. The presence of acicular ferrite enhances the toughness of this region as well [51].

Figure 2.36: Acicular ferrite structure [48]

33

Fine Grained Heat Affected Zone (FGHAZ) is the following zone to be analysed. The peak temperature

is above 900oC. According to Figure 2.35, the grain size reaches lower values, which may lead to better

mechanical properties than the initial material properties, since the number of grain boundaries

increases. Thus, the dislocations have more obstacles to overcome [53].

At Inter-Critical Heat Affected Zone (ICHAZ) the reached peak temperature is found between A1 and A3

temperatures presented in Figure 2.34 (approximately 900oC and 750oC, respectively). Part of the

material is transformed into austenite and recrystallizes. The transformation begins at regions where a

carbon rich domain, as pearlite, is found [54]. The cooling occurs rapidly, forming martensite-austenitic

islands. Due to this fact, the impact toughness is deteriorated. Vanadium and carbon tend to increase

the austenite hardenability, promoting the formation of hard and brittle phases as martensite [55]. It is

important to mention that in low carbon steels, the quantity of pearlite is low resulting in smaller

martensitic-austenitic islands. The presence of islands on the microstructure can provide more barriers

to the dislocation propagation and subsequently, in these cases, the hardness is improved at this zone.

The last zone of the HAZ is the Sub-Critical Heat Affected Zone (SCHAZ). The peak temperature is

between A1 (eutectoid temperature) and 600oC. Below A1, transformation and recrystallization

processes do not occur. Instead, at these temperatures, a tempering process takes place, promoting

the grain refining. Although this process may lead to a strength and hardness reduction [56], the

precipitation of some elements, as Cu, can promote an increase in hardness and strength [56].

The welds which are analysed were performed with two weld passes and according to Figure 2.37 it is

possible to see that due to the multi-pass welding, two different zones are formed, Intercritically

Reheated CGHAZ (IRCG) and Subcritically Reheated CGHAZ (SRCG).

Figure 2.37: HAZ overlapping in multi-pass welding [55]

2.4.3. Heat input influence

In this work, one of the main objectives is to evaluate the heat input influence on the steel hardness and

consequently on the Yield strength. From Prasad’s study [57], for a HSLA steel, it was observed that the

peak harness was reached at the coarsened zone. The reduction of the average hardness value on both

34

weld and HAZ, as heat input increased, was also concluded. The justification presented by Prasad was,

with the heat input increase, the grain coarsening was promoted, increasing the fraction of soft phases,

due to a lower cooling rate.

It was also seen the occurrence of different microstructures at the weld seam. Near the fusion line, the

grain presented a very fine size. Closer to the root face, the grain tended to be finer than the remaining

grains at the weld centre, which showed a coarser size and a cellular structure. As both heat input and

welding speed increased, it was promoted the formation of an equiaxed grains band, which extended

along the weld centreline, instead of the formation of a coarse columnar grain structure which was

presented in the adjacent areas. This structure tended to be finer for low heat input values, as higher

values promoted the grain coarsening [57].

2.4.4. S700MC

Considering the particular case of the steel studied in this work, several authors attended to comprehend

the influence of different thermal cycles on the steel microstructure, by simulating a welding cycle. This

steel presents an elongated and fine ferritic-bainitic structure, characteristic of the thermo-mechanically

controlled process, which is used to fabricate the steel. One of these works was performed by Górka

[58], where the welding cycle was simulated using resistive heating, subjecting the sample to different

maximum temperatures. 3 distinct structures were seen, depending on the temperature range. Between

400-700 0C, occurred a partial recrystallization, as a slight grain growth took place, leading to a reduction

on toughness (from 80 J/cm2, parent material, to 37 J/cm2) and hardness (Parent material-280 HV1,

actual value- 260 HV1). For a temperature range of 800-900 0C, a fine grained ferritic-bainitic structure

took place instead. In this case, toughness was largely improved (200 J/cm2) while, the hardness was

slightly inferior to the parent material hardness (265 HV1). This fact may be explained by a matrix

strengthening, as a result of the coagulation of precipitates. Finally, for an interval of 1000-1300 0C, a

grain coarsening occurred, where the structure is, manly, constituted by bainite, reducing largely the

toughness (5 J/cm2), due to the dissolution of MX phases, such as NbN and TiN, which occurs at those

temperatures and the uncontrolled secretion, during the cooling period, led to a weaker matrix. The

hardness presented a value of 230 HV1. Górka [59] also studied the influence of the cooling time on the

microstructure and hardness value. The real maximum temperature cycle was, in every case, higher

than 1300oC. Figure 2.38 shows the influence of cooling time on the microstructure and hardness,

respectively.

It was witnessed that, for a cooling time inferior to 3 s, a martensitic structure was formed while, between

3 and 10s the microstructure showed a mixture of martensite and bainite. A mainly bainitic structure

occurred for a cooling time interval between 10 and 14 s. From 14 s, the increase of the cooling time

promoted the occurrence of ferrite on the microstructure. Considering now, the hardness evolution, it is

possible, once again, to see that higher cooling times led to a hardness reduction, as referred by Prasad

[57] since the content of ferrite increased and the grain coarsening also occurred.

35

Figure 2.38: a) Structural changes in the S700 MC steel in terms of welding CCT-W ; b) Distribution of hardness as a function of the cooling time [59]

2.5. Young’s Modulus

Apart from the several properties which were mentioned in previous chapters, it is worth mentioning the

influence of specific parameters, which are related to the welding process, on Young’s Modulus. The

Young’s Modulus is a linear material property, which measures the stiffness of a material and can be

affected by several conditions such as plastic prestrain [60], chemical composition, heat treatment [61]

and environment temperature. However, for this case, the influence of heat treatment is a more

interesting topic to mention since, the environment temperature was unchanged and the steel used was

always the same.

From the work of Fadare [61], it is possible to observe the variation of the Young’s Modulus when the

same steel was subjected to different heat treatments. As illustrated in Table 2.2, the influence of the

heat treatment on Young’s modulus is substantial with the exception of the hardened treatment.

Table 2.2: Mechanical properties of heat treated and untreated NST 37-2 steel [61]

Mechanical properties

Heat treatment

Tensile strength (N/mm2)

Hardness (BHN)

Toughness (J)

Elongation (%)

Reduction (%)

Yield strength (N/mm2)

Young’s modulus (N/mm2)

Untreated 343.80 100.10 58.88 21.16 63.23 217.31 465.78

Annealed 325.42 95.95 64.10 23.24 71.94 209.47 562.00

Normalised 422.30 188.00 57.26 20.38 71.81 232.75 534.85

Hardened 678.70 460.50 24.67 8.42 41.14 288.05 1235.31

Tempered 385.42 131.00 60.70 21.00 76.92 228.52 535.17

36

The hardened condition promoted the formation of Martensite, which led to a higher Young’s modulus.

This information empowers the conclusion made by Pramanik [62], which indicated that phase

transformation influence the elastic modulus of steels, despite the lack of data to evaluate that same

influence on the distinct transformation phases

From the work of Winczek [63], a typical temperature field for steels during a multi-pass GMAW is

presented. The maximum temperatures reached are presented in Figure 2.39.

Figure 2.39: Maximum temperatures (oC) in the middle part area of cross-section [63]

Making a more localized evaluation of the temperature evolution, it was considered a specific point P

and the respective thermal cycle is shown in Figure 2.40.

Figure 2.40: Heat-affected zone and thermal cycle for P point [63]

Analysing Figure 2.40 and although the temperatures achieved were higher than the fusion temperature,

in this case, it occurred for a short period of time, opposed to the one occurred on heat treatment

processes, where steel is maintained at a specific temperature for a determined period of time, larger

than the showed in this case. Furthermore, this type of evolution may be compared to a quenching

process, in stages 3 and 4, due to the rapid cooling rate which occurred. If it is added stage 5, then, the

overall process may be compared to a quenched and tempering process, although the tempering period,

in this case, was also significantly smaller than the occurred in the conventional process.

Moreover, a parameter, which is also important to examine is the grain size, since the welding process

promotes the occurrence of distinct zones, each presenting different grain sizes. The influence of the

grain size on the Young’s Modulus (for mild steels) can be observed in Figure 2.41.

37

Figure 2.41 - Young's modulus versus grain size for ultra-fine grain sized mild steel [64]

As mentioned before, the welding process promotes a significant variation in grain size. For a

thermomechanically rolled high-strength steel, welded by double-submerged arc process, the different

grain sizes for the distinct zones are presented in Figure 2.42.

Figure 2.42: Characteristic microstructural zones of the HAZ of a double-submerged arc-welded joint [65]

From the figure analysis, it is notorious the grain variation for the distinct zones, which empowers the

statement that the Young’s modulus may vary with the introduction of the welding effect.

2.6. Digital Image Correlation

2.6.1. Introduction

The Digital Image Correlation, DIC, was proposed in the 1980s by researchers of the University of South

Carolina [66] and it is a non-contact optical technique. It is applied to determine displacement fields by

matching grey levels between an image taken before loading period (as a reference image) and an after

the un-loading of a specimen (as a deformed image). Images are composed by pixies and for each pixel

is associated a grey level.

In order to determinate the displacement field, the image is divided into sub-regions, also known as the

Zone Of Interest (ZOI). The objective of the correlation method is to match the ZOIs of the two images.

38

The size of the subsets is defined in order to establish the condition that the deformation can be

considered as homogeneous. The deformed position of an arbitrary point in a subset, (𝑥𝑑𝑒𝑓)𝑖, can be

given by:

(𝑥𝑑𝑒𝑓)𝑖= 𝑥𝑖𝑛𝑖 + (𝑢)𝑖 +

𝜕{𝑢𝑖(𝑥)}

𝜕𝑥𝑗 𝑑𝑥𝑗 𝑖, 𝑗 = 1, 2 (2.38)

Where 𝑥𝑖𝑛𝑖 is the initial point position, (𝑢)𝑖 corresponds to the point vector displacement and 𝜕𝑢𝑖/𝜕𝑥𝑗

consists in the deformation gradients components.

In order to obtain the displacement and deformation gradient terms of a local subset, a correlation

coefficient 𝑋 was defined, where 𝐴 represents the reference image and 𝐵 the deformed image [66]:

𝑋 (𝑢𝑖 ,𝜕𝑢𝑖𝜕𝑥𝑗

) =∬[𝐴(𝑥𝑖𝑛𝑖) − 𝐵(𝑥𝑑𝑒𝑓)]2 𝑑𝑥 𝑖, 𝑗 = 1, 2 (2.39)

This parameter, also called global correlation residual, provides a simple quality evaluation of the

matching between both images. It is defined since, the brightness conservation cannot be totally

achieved due to image imperfections or due to the fact that the actual displacement field may not be

exactly part of the chosen ZOI. The ideally condition is achieved when the correlation coefficient is zero

[66].

2.6.2. Sobel method and Fast Fourier Transform

Through the years several algorithms were developed, such as Sobel Edge Detector and Fourier

Transform with the purpose of detecting and measuring crack lengths.

The Sobel Edge Detector is a first order derivative method and consists in the application of a derivative

approximation in order to find edges (in this case cracks). The presence of edges is determinate, at

regions, where the gradient is maximum for a considered image. A threshold value is defined and only

above this value, the existence of a crack is considered [67].

A practical application of the Sobel edge detection method to measure crack lengths is presented. In

this case, it was performed a similar fatigue test as the one performed in this report and the crack length

was determinate by two different processes: travelling microscope and DIC by using the Sobel method.

Based on the analysis of Figure 2.43, it is possible to observe that the crack length values obtained for

the two processes were fairly close even though, from 100000 cycles onwards, approximately, the

results started to spread apart slightly, until regime C was reached.

39

Figure 2.43: Crack length vs no. of cycles obtained by travelling microscope and DIC using sobel edge detection method [68]

The Fast Fourier transform (FFT) consists in a mathematical technique, which converts a time-domain

digital signal (in this case, an image) into the frequency domain:

𝐹(𝑢, 𝑣) =1

𝑀𝑁∑ ∑𝑓(𝑥, 𝑦)𝑒−2𝜋𝑗(

𝑥𝑢𝑀+𝑦𝑣𝑁)

𝑁−1

𝑦=0

𝑀−1

𝑥=0

(2.40)

Where f(x,y) is the image in spatial domain with a 𝑀 ×𝑁 size.

The main concept of the FFT is to divide the DFT (Discrete Fourier Transform) time-domain sequence

of length N into smaller DFTs with the purpose of reducing the number of arithmetic operations [69]. The

crack definition is similar to the Sobel method since, for a defined threshold value, the existence of a

crack is determined.

In spectre domain, the application of the FFT allows to reduce computation costs. The relative speed

between the correlation based on the FFT and the spatial domain correlation was evaluated according

to the number of multiplications and additions needed in each algorithm. Considering quadratic subsets

(n=m) and searching areas 𝑆 = 𝑆𝑥 = 𝑆𝑦 in order to simplify the demonstration, the number of operations

for the spatial domain, 𝑁𝑠𝑝𝑎𝑡𝑖𝑎𝑙 (Equation 2.41) and for the spectral domain, 𝑁𝑠𝑝𝑒𝑐𝑡𝑟𝑎𝑙 (Equation 2.42) are

presented below [70]:

𝑁𝑠𝑝𝑎𝑡𝑖𝑎𝑙 = 𝑛2𝑆2 (2.41)

𝑁𝑠𝑝𝑒𝑐𝑡𝑟𝑎𝑙 = 3𝑛2 log2 𝑛 (2.42)

Although FFT-based correlations are faster than the spatial domain correlations, the results achieved

with FFT based tended to present poorer results [71].

40

2.6.3. Method developed to measure crack lengths

Due to the difficulties to create a robust program from the edge detection methods, such as Sobel

method and the lack of precision of the FFT, none of these methods were applied to determine the crack

length in this work. Instead of applying a method based on the first derivate, as in the Sobel method, the

method used was based on the second order derivate. For this method, the test images were captured

and imported to the DaVis software, in order to acquire a displacement field, as represented in Figure

2.44. For the determination of the displacement field, a reference image was needed.

Figure 2.44: Example of a displacement field

Thereafter, the Fast Fourier Transform was used to reduce, as possible, the level of noise to ensure a

better evaluation of the crack length. The crack edge detection method was based on a 2nd order

derivative method, as mentioned before. The vertical displacement was differentiated two times and the

location where the minimum value occurred was recorded as a crack point. Ideally, the second derivate

would be zero, however, due to the fact that the fracture surface was not flat, zero condition was not

achieved. Although a 1st order derivative method would also be considered, however, in that case the

maximum value would be searched instead. For shorter cracks, the 2nd order derivative method showed

better results. Furthermore, it was observed that the differential method led to an issue since, the

searching process of a minimum value continued, even if it was clear that those points were off the crack

domain, as depicted in Figure 2.45.

Figure 2.45: Crack tip location (initial assumption)

41

For that reason, it was needed to define a threshold value. For example, considering a threshold value

of 9 pixels, for the cases where the vertical distance between two consecutive points was higher than 9

pixels, the location of the crack tip was considered at the previous point.

In the next step, in order to improve the crack tip location, it was recorded the displacement values for

the upper and lower crack surface, usually, 6 pixels above and below the crack edges. The displacement

values were fitted by:

𝑌𝑡 = {𝑎 𝑥 < 𝑐

𝑎 + 𝑏𝑡 √(𝑥 − 𝑐) + 𝑑(𝑥 − 𝑐) 𝑥 ≥ 𝑐 (2.43)

𝑌𝑏 = {𝑎 𝑥 < 𝑐

𝑎 + 𝑏𝑏√(𝑥 − 𝑐) + 𝑑(𝑥 − 𝑐) 𝑥 ≥ 𝑐 (2.44)

Where c is the crack tip location. The subscripts t and b referred to top and bottom, respectively. The

displacement for the lower and upper super surface as the respective fitting curves are illustrated in

Figure 2.46. Therefore, as the crack tip location was fitted, the crack length was obtained.

Figure 2.46: Crack surface plot

2.6.4. Speckle pattern

An important aspect which is also important to mention is the need to create a speckle pattern to ensure

a better accuracy on data capture, in order to obtain the displacement field results. A good speckle

pattern allows a noise reduction and enables a better identification of subsets on the deformed image.

A speckle pattern has to have specifics characterizes, such as high contrast, a small variation in the

speckles size, and a random pattern. This latter aspect reveals important since a false matching can

occur, leading to inaccurate results.

The influence of the type of speckle was evaluated by Barragner [72]. It was tested black paint, white

paint and spread powder speckle patterns. For each one of the speckle patterns displacement and strain

measurements were performed. The displacement results were similar for the all cases, however, the

same did not occur on the strain results. The results for the powder speckle patterns underestimate the

42

strain values obtained with a mark tracking technique (assumed as the correct values) as presented in

Figure 2.47 which is more accentuated for strains >10%. This can be explained by the fact that powered

particles remained rigid during the test, leading to a non-uniform deformation of the grey level

distribution. In the painted speckle patterns, the results presented slight differences between the mark

tracking results, though these were smaller than the results for the powered case, due to the occurrence

of cracks on the speckle patterns.

Figure 2.47: Influence of the speckle pattern on strain measurements [72]

43

3. Material and Welding Procedure

In this chapter it is presented the chemical composition and the mechanical properties of the material

evaluated in the project. In addition, it is also described the welding procedure to perform the several

welded joints and the four different welding parameters that were used to perform the welds.

3.1. Material

The material in question is a hot-rolled structural steel produced by TMCP (Thermo-Mechanically

Controlled Process), known also as S700MC-1 and developed by SSAB. In Table 3.1, the respective

chemical composition of the steel is shown.

Table 3.1: Steel chemical composition

C

(max %)

Si

(max%)

Mn

(max%)

Al

(min%)

S

(max%)

P

(max%)

Nb*

(max%)

Ti*

(max%)

V*

(max%)

0.12 0.25 2.10 0.015 0.010 0.020 0.09 0.15 0.20

*The sum of Nb, Ti and V equals or lower than 0.22 wt-%

In Table 3.2 the mechanical properties of the S700MC-1 steel are enumerated. The tensile test was

carried out with a 20 mm width specimen. The result of the Charpy-V impact test was obtained for a

specimen with a thickness of 7.5mm. Both properties were provided by the manufacturer

Table 3.2: Mechanical properties

Tensile tests Charpy-V impact test

Test direction 𝜎𝑌𝑆 (MPa) 𝜎𝑈𝑇𝑆 (MPa) A (%) Test direction Impact energy at -

400C (J)

Transverse 785 867 13.4 Transverse 66

3.2. Welding procedure

The welds were performed with a Kemppi Promig 530 automatized system at the welding laboratory of

SSAB’s Research centre, situated in Raahe, Finland.

In order to perform the welds, two plates were joined and four longitudinal 16mm thickness bars were

used to support each one of planes. A clamping system with eight clamping elements (four at each side)

was used, as presented in Figure 3. This fixed the two plates, avoiding any plate displacement. The

thickness of the plates was 8mm (Figure 3.2).The width and the length were 200 mm and 1000 mm,

respectively, as indicated in Figure 3.1.

44

Figure 3.1: Clamping system and dimensions of the plates

The heat input was determined by calculative method using Equation 3.1, where this cooling time, t8/5,

was defined as the time, which the weld seam and the adjacent HAZ cooled from 800oC to 500oC. The

defined cooling times were 5 s, 10 s, 15 s, 20 s [73].

𝑡8/5 = (4300 − 4.3𝑇0) ∙ 105 ∙𝑘2 ∙ 𝑄2

𝑡2∙ [(

1

500 − 𝑇0)2

− (1

800 − 𝑇0)2

] ∙ 𝐹2 (3.1)

Where 𝑇0 is the working temperature (oC), 𝑘 is the thermal efficiency of the welding procedure (for MAG

welding process is considered 0.8), 𝑄 is the arc energy (kJ/mm), 𝑡 is the workpiece material thickness

(mm) and 𝐹2 is the joint type factor in two-dimensional heat conduction (considered 0.9 in butt welds).

The room temperature for each weld was 23oC and the inter-pass time needed was, at least, 40 min.

The shielding gas used was Mison 8, which contains Ar (Argon) + 8% CO2 (carbon dioxide) + 0.03%

NO (nitrogen monoxide). In Figure 3.2 is presented the joint design and the welding sequence.

Figure 3.2: Schematic presentation of joint design and welding sequence for all welds

As mentioned before, the welds were produced with four different heat inputs which were determined

by using Equation 3.1. The welding parameters, for each weld, are presented in Table 3.3.

45

Table 3.3: Welding parameters for each t8/5

Welding parameters

Run Current

(A)

Voltage

(V)

Wire

feed

(m/min)

Travel

speed

(cm/min)

Arc energy

(kJ/mm) Tag

t8/5=5s

1 253 25.8 11.6 56.0 0.7

5A

2 268 27.8 14.0 63.8 0.7

t8/5=10s

1 253 25.8 11.7 55.7 0.7

6A

2 270 28.6 12.7 46.3 1.0

t8/5=15s

1 253 25.8 11.4 55.9 0.7

7A

2 269 29.0 12.8 39.0 1.2

t8/5=20s

1 253 25.8 11.5 56.0 0.7

8A

2 271 29.0 12.7 33.6 1.4

It is important to notice that the root pass (first run) of each weld was performed with the heat input

calculated for a cooling time, 𝑡8/5, of 5s, changing only the heat input of the second run in each sample.

47

4. Laboratorial part

In this Chapter the experimental testing methods and testing conditions are described.

4.1. Fatigue tests

All fatigue tests were carried out with a servohydraulic MTS 810 Material Testing System machine at

Aalto University. The tests were performed at room temperature and according to the E 647-00 standard

[74]. In Figure 4.1 is showed the experimental apparatus used to perform the fatigue tests.

Figure 4.1: Fatigue experimental apparatus

The fatigue tests were performed using a stress ratio, 𝑅, of 0.1 for the specimen presented in Figure 4.2

and the generic specimen dimensions are shown in Figure 4.3. The maximum force was 10kN, while

the minimum force was equal to 1kN. A loading frequency of 5 Hz was used to perform the fatigue tests.

Figure 4.2: Test specimen

48

Figure 4.3: Specimen dimensions. Adapted from [74]

As indicated previously, it was necessary to apply a speckle pattern on the specimen surface, which

promotes an easier and more accurate measurement of the crack length when the DIC method is used.

In this case, the speckle pattern was performed with spray paint, as illustrated in Figure 4.4.

Figure 4.4: Speckle pattern applied on specimen test

The spray paint is known for its easiness to implement and for the fact that the waiting time needed

before the test is small. Despite the previous fact, the speckle pattern process was a time-consuming

process, since it was necessary to reach the optimal condition in order to proceed with data acquisition.

The image capture was executed with a LaVision Imager pro X camera.

As mentioned before, the images of the tests were processed according to the method described

previously in Chapter 2.6 (DIC), in order to proceed with the crack length determination. This data was

treated and then it was used to obtain the Paris Laws for the different tests. The different Paris laws

were determine based on the seven point incremental polynomial technique, presented in E 647-00

standard [74]. The fatigue test can be observed in Figure 4.5.

49

Figure 4.5: Fatigue test images

4.2. Micro hardness tests

4.2.1. Samples preparation

For the hardness tests, it was necessary to proceed with the samples polishing. Several sanding papers

(class 240, 320, 600, 800, 1000, 2400 and 4000) were used, submerged in water, during 3 minutes for

each one of them, to ensure that the existent scratches presented the same size of the grain size of the

respective sanding paper used.

Posteriorly, two more polishing steps were used. In these cases, the polishing was performed with

diamond, presenting a grain size of 3µm and 1µm, respectively, and lubricant during 1 minute in both

cases. All samples were polished with the assistance of a Streuers DAP-7 wet polisher. It is important

to mention that after every polish stage, the several samples were clean with ethanol, in order to avoid

corrosion. Afterwards, the samples were etched with 2% Nital (2% nitric acid HNO3 and 98% ethanol

C2H5OH) for a time period between 10 and 15 seconds.

4.2.2. Micro hardness test conditions

Hardness test allows quantifying the material’s resistance to present small indentations when a localized

compressive load is applied. In order to perform the hardness test, a small indenter is pressed, with a

constant load, against the surface of the sample which pretends to analyse.

Several indenter shapes can be used to measure the hardness value, depending on the hardness scale

that is more convenient to work. For each case, the indenter leaves a specific pattern mark on the

surface. In this work, it was performed a Vickers hardness test, which the characteristic indenter is a

Diamond pyramid. In Figure 4.6 is illustrated the indenter geometry and the top view mark shape as

well.

50

Figure 4.6: Vickers hardness test schematics [75]

The Hardness value is determine according to:

𝐻𝑉 = 1.854

𝑃

(𝑑1 + 𝑑22

)2 (4.1)

Where P is the load applied in kgf (kilograms force) and 𝑑1 and 𝑑2 are the indentation dimensions, which

are in mm. Hardness measurements were performed at IST, Lisbon with a Struers Duramin hardness

measure. Initially, an HV 0.025 scale (25 grams) was defined, however, soon was seen that the

indentations formed had approximately the size of the grains. Consequently, the measured value was

equivalent to the hardness grain, leading to a significant overestimation of the hardness value. This

result may be explained by the grain boundary effect [76]. Nevertheless, the hardness tests were carried

out using a relatively small scale (HV 0.1, 100 grams) in order to ensure a more detailed analysis, but

also, on the same time, guarantee a hardness evaluation without the grain boundary effect. In order to

determine the hardness profiles, a horizontal interval of 0.5 mm between indentations was defined,

ensuring that the hardness measurements were not influenced by the previous indentations (Figure 4.7).

Lines C, D, E, and F were performed with a vertical interval of 2 mm between them while line G was

1mm below line F. The distance between the indention marks of line C and G and the sample’s boundary

was 0.5 mm in order to prevent any influence of the resin on the hardness results. Additionally, a

minimum horizontal distance of 0.5mm from the resin was defined so that the hardness measurements

do not suffer any kind of effect caused by it as well.

51

Figure 4.7: Hardness measurements

Furthermore, from the values of hardness, it is possible to estimate the yield stress for the steel. Based

on several hardness and strength values acquired through the years, Pavlina [77], proposed several

correlations in order to relate tensile and yield stress with the hardness value.

It is important to mention that, all hardness values used to perform this correlation were converted to

the same hardness scale, which in this case was Vickers hardness.

The correlation determined for yield stress, 𝜎𝑌𝑆, is given by:

𝜎𝑌𝑆 = −90.7 + 2.876𝐻𝑣 (4.2)

Where 𝐻𝑣 is the Vickers hardness value.

The study showed that the determination coefficient, 𝑅2, for equation 4.2, was 0.92 and the correlation

valid range was between 129 and 632 (Vickers Hardness). 165 points were used to obtain the

correlation. Furthermore, a correlation for tensile strength was also established in that work, however, a

comparison was made between predict and actual stress for yield and tensile correlations. Although the

yield strength-hardness regression presented no systematic deviations, the same cannot be said about

the tensile strength-hardness regression, which under predicts the strength value at higher hardness

values.

Moreover, from the data analysis, it was reckoned that for lower hardness values, strength might have

a nonlinear correlation with hardness and for that reason, it is not recommended the application of these

correlations for low hardness values.

52

4.3. Metallographic Analysis

The four samples were analysed using a scanning electron microscope (SEM) JEOL JSM-7001F at IST

Lisbon, Portugal (Figure 4.8) in order to evaluate the different microstructures formed due to the distinct

welding parameters and to examine the multi-pass effect on the microstructure, as well.

Figure 4.8: Scanning electron microscope at IST

Firstly, the proceedings used to perform the sample polishing and the sample etching were the same of

the ones selected to prepare the samples to the hardness tests. Thereafter, samples were washed by

ultrasounds to remove impurities in order to allow a better evaluation of the steel microstructure.

In order to perform the microstructure analysis and evaluate the influence of the different welding

parameters used to perform the welds specific regions were defined. The distinct locations where the

microstructure of each sample was analysed are presented in Figure 4.9.

Figure 4.9: Location of the regions where the microstructures were analysed

It is important to mention that, positions I, H and G were localized near the fusion line (the boundary

between the weld zone and the heat affected zone).

53

5. Numerical Analysis

One of the aims of this work is to evaluate the geometric influence of the weld on the SIF value. For that

purpose 3 different specimens, based on the ones sent by Aalto university were modelled, as shown in

Figure 5.1. The first model was an unwelded model (BM case), which served as benchmark to evaluate

the influence of the different weld geometries. The other cases evaluated were a specimen with a weld

transverse to the crack propagation direction (T case) and a specimen with a vertical weld, deviated

from the centre (HAZ case).

It was defined a 200 GPa Young’s modulus and a Poisson coefficient of 0.3 for each specimen. The

general dimensions were the same for all specimens and equal to the selected for the fatigue test

specimen.

Figure 5.1: BM case (left), T case (centre), HAZ case (right)

The linear elastic simulations were performed using Abaqus software and the stress intensity factor, for

all cases, was determined using the contour integral J method. Regarding the finite element mesh, it

presents a direct influence on the obtained results and the existence of notches, as occurred in this

case, emphasising the need of an adequate mesh in order to obtain a good estimation of the parameters

of interest.

The mesh was performed using quadratic elements along the specimen domain, however, in each case,

it was needed to define two circumferences, centred at the crack tip. Since the crack tip is sharp, the

result is the occurrence of a singularity at the crack tip.

The inclusion of the singularity improves the accuracy of the contour integral and therefore, in order to

introduce the singularity effect (since, as seen before, the stress near the crack tip increases by a 1/r

factor) and simulate linear elasticity, in the smaller circumference, the midside nodes were moved

towards the crack tip. The midside node parameter, t, was considered equal to 0.25 instead 0.5 (the

crack tip is localized at t=0). It was also necessary to collapse hexahedral elements to wedges elements,

with multiple nodes at each location, at the crack tip. The nodes presented at the same location (at the

crack tip) were constrained to move together as a single one node. The second circumference had the

purpose of smooth the transition between the interior mesh with the remaining mesh part (Figure 5.2).

54

Figure 5.2: Mesh at the crack tip (crack tip is located at the centre of the circumferences)

A tension state was promote using a 10 KN static loading as showed in Figure 5.3. Both force vectors

were applied at the thickness middle point, to ensure that the main opening crack mode observed was

mode I. The stress intensity factor was calculated along the thickness, for a crack length range of 11 to

40 mm. The crack length was defined as indicated in E647-00 standard [74] and showed in Figure 4.3.

Figure 5.3: Application of static forces

For the BM specimen, several meshes were selected with different numbers of points along the

thickness. This number ranged between 15 and 33 with this variation being explained by the purpose of

evaluating the solution convergence. Although the evolution was better addressed for the case with

more points, as expected, the medium and maximum values of SIF presented a small variation for the

several cases evaluated, therefore, the increase on the number of points did not compensate the

increase of the computation time. From the converge analysis (presented further and performed for

crack lengths equal to 11mm and 40mm), it was decided to use 15 points along the thickness in order

to evaluate the SIF evolution, as the crack length increases.

For the other specimens, the number of points used was the same in order to perform the comparison

between the distinct specimens. Finally, it is also important to mention that for the BM case, the number

of contours used was 10. The introduction of a higher number of contours improves the estimation of

the SIF value. However, it was seen that by using the contour integral method, the SIF presented a

constant value for the 10 contours used. For that reason, both in the T and HAZ cases, 5 contours were

used instead.

55

6. Results and Discussion

In this chapter the results for the fatigue and hardness tests are discussed. Furthermore, from the

hardness data, the Yield Strength profiles were determined and are presented as well. The

microstructure analysis is also presented as the numerical simulation results as well.

6.1. Fatigue results

The Paris laws presented below were determine according to the following steps:

In case 1 the initial number of crack length values was 108, while in case 2 this number was 36. The

crack length data obtained by the DIC method showed an inconstant evolution, since, for smaller crack

lengths, the crack size either increased or decreased as the number of cycles increased. This occurred

due to the difficulty in measuring a crack with a relatively small size. In order to avoid this issue, for the

data treatment of the fatigue test, it was considered that the crack length only increased with the increase

on fatigue cycles and for that reason, when the measured crack length was inferior to the previous value,

that value and the respective number of cycles were not considered. Furthermore, the last point of case

1 and the last two points in case 2 were removed from the analysis, since those points already belong

to the third fatigue phase. Therefore, from the data treatment, the final number of points in both cases

was 51 (case 1) and 31 (case 2). The treated data is presented below for the different cases studied, in

56

Figure 6.1, in which the specimens evaluated were the HAZ case with a weld performed with the 6A

welding parameters case and a stress ratio of 0.1 (Pmax=10kN and Pmin=1kN).

Figure 6.1: a) Crack propagation rate for HAZ 6A case 1; b) case 2

Comparing the two cases, coefficient 𝐶 and 𝑚 present similar values, which was expected since both

cases were performed with the same conditions despite the total number of cycles was different: 970000

and 778000, respectively. The two cases presented a correlation coefficient near to the unit, which

indicates that the Paris law for the two cases was well addressed. In terms of the C and m values, it is

only possible to mention that these results are found within the typical interval of values for steel.

6.2. Micro hardness results

First of all, and for the sake of the hardness analysis, it is important to introduce information from the

previous work [78]. In [78], the temperature profiles were determined at root side and face side for the

four samples. Furthermore, and to complement the information provided by the temperatures profiles, it

was presented, as well, the molten areas (yellow) and the areas where a full austenitization occurred

(pink) for the different welding parameters. These information can be found in Annex A2 and is extremely

important to comprehend and to analyse the effects of the different welding parameters on the hardness

profiles, and on the steel microscruture, which are discussed further in this chapter.

The hardness profiles were performed for the different samples and are shown in Figures 6.2, 6.3, 6.4

and 6.5.

57

Figure 6.2: Hardness profile 5A sample


58



In order to evaluate the measurement precision, it was considered four hardness points in each line,

two on each side of the sample. These were sufficiently far from the HAZ to ensure that the welding

59

effect was practically inexistent. Therefore a comparison between the average value measured, and the

“theoretical” value, given in Table 3.2 was made. The average value was determined from the average

value measured for each line, converted afterwards to a yield strength value, using the Eq. 4.2. It was

also computed the standard deviation for all the samples taking into consideration the same points,

however, in this case, in hardness unit.

In Table 6.1, are presented the results of the different samples.

Table 6.1: Comparison between measured Yield Strength and material Yield strength

Yield Strength

(table 3.2) (MPa)

Average Yield Strength

measured (MPa)

Difference (MPa)

Standard deviation (Hv)

Sample 5A 785 785.4 0.4 4.3

Sample 6A 785 770.2 14.8 7,3

Sample 7A 785 813.8 -28.8 5.8

Sample 8A 785 768.8 16.2 9.4

From the information provided in Table 6.1, it is clear to observe that the highest variation took place at

the 7A sample. This indicates that the values measured for sample 7A may present slight higher values

than the expected since the average yield strength measured was slightly higher than the material yield

strength. Nevertheless, the variation between the average measured values and actual value was small

for all cases, which validates the performed measurements. In Figure 6.6 are presented the average

hardness value measured and the respective standard deviation. As observed in Table 6.1 and

illustrated in the Figure 6.6, the highest standard deviation occurred in sample 8A, which indicates in

that sample that the values measured presented a higher dispersion. On the other hand, the lowest

standard deviation occurred in sample 5A, however, the standard deviation values were rather small.

Figure 6.6: Average hardness measured for each sample and respective standard deviation

0

50

100

150

200

250

300

350

5A 6A 7A 8A

HV

0.1

SAMPLES

60

Taking into account now the different hardness profiles at a region near the weld zone, a tendency can

be detected, as the heat input increased, an average reduction in the hardness value occurred. This

fact confirmed the information provided by Prasad’s work [57]. It was also observed, from the

microstructure analysis, that a higher cooling time promoted the presence of ferrite, which is a softer

structure compared to bainite or martensite.

Comparing the different hardness profiles, a common evolution was observed as, at the weld seam, the

hardness value presented relatively small variations when was considered the same line. Between

different lines, the variation was more noticed. Moving further away from the weld zone, the hardness

tended to decrease as the distance to the weld centre increased, until a minimum value was reached.

From that, the hardness increased and, in some cases, a maximum hardness point was achieved.

Finally, at a considerable distance of the weld seam, the thermal effect was less observed, since the

hardness tended to present a more constant value, closer to the base material hardness. In sample 8A

this tendency was less evident which may indicate that the welding effect reached a higher distance

than the expected, which can also explain a higher standard variation for this sample.

From the hardness evolution, it was observed that the hardness variation was more evident at line F

and G, which were located at the first weld run for the different hardness profiles. This may be explained

by the fact that, at those zones, the microstructure was affected by the two weld runs, instead of one

weld run only, which occurred at line C and D. Line E presented also a significant variation, compared

to the occurred at C and D since it is located near the boundary between the two welds. At the weld

zone, the lines which hardness presented higher values were line F and G, in most cases.

Although the two weld centres were slightly misaligned, it is possible to see that the hardness profiles

for the different lines were almost symmetric. The softest and hardest points locations are presented in

Table 6.2, as well as the respective hardness value and the line at which occurred.

Table 6.2: Softest and Hardest point’s location

Sample Softest point

(Hv 0.1)

Location to the weld centre

(mm) Line

Hardest point (Hv 0.1)

Location to the weld centre

(mm) Line

5A 242 -3.5 G 364 1 F

6A 250 4 G 368 5.5 E

7A 239 3.5 F 366 -6 E

8A 231 -1.5 C 359 8 E

From the information given in table 6.2, it is possible to observe that, in most cases, the hardest point

was located in line E. In sample 5A, the hardest point was found in the vicinity of the weld (fusion line)

which was referred before (Chapter 2.4 Figure 2.43) as a plausible scenario. The reason for that was

the lower cooling time imposed which promoted the formation of harder structures, as martensite, at that

location. At line F, microstructure was affected by the two weld runs, however, since the heat input was

low, the second run enhanced the hardness at that position. For the other samples, the cooling times

associated were sufficiently high, occurring the formation of softer structures at that location. Instead,

61

the hardest points were found relatively far from the weld centre. The peak temperatures at these points

were within a range between 700-750oC, which led to a partial recrystallization, maintaining the

characteristic structure induced by the rolling process. The second weld run promoted an enhancement

of the hardness at those points. Apart from that, aspects such as the presence of inclusions, TiN,

associated with a recovery process due to the temperature reached, might lead to higher hardness

values than the hardness observed for the base material. Furthermore, it was also possible to witness

that the hardness of the hardest point decreases with the increase of the heat input, as occurs when the

hardest values are compared for the same line (line E), as expected.

Regarding the softest points, an interesting outcome took place in sample 8A. Contrary to the other

samples, in which the softest point occurred at similar zones (regions where the peak temperature range

was between 900-950oC for the different samples), in sample 8A the softest point was detected at the

weld seam. This occurrence may be related to the high energy input which tended to promote the grain

coarsening, and this fact associated with the existence of softer structures but also with the elongated

shape which is characteristic of the grains at the weld seam, led to an inferior hardness value. For the

other samples, as mentioned before, the softest point took place at a region, where the peak

temperatures were within 900-950oC range. At these temperatures, the structure presented grains with

a small size. However, at this area, the effect of the induced hardening provided by the rolling process

was undone, since a full recrystallization occurred. Thus, the mechanical properties were poorer and

hence the respective hardness value was also lower at this region. This information confirms the

information provided by the previous work [78].

It is important to mention that, despite the general tendency on the reduction of hardness with the

increase of heat input, the hardness of the softest point of sample 6A was higher than the presented in

sample 5A. The reason behind this fact was the peak temperature reached, for sample 6A, at that

location, was higher and promoted the enhancement on the hardness.

6.3. Yield Stress results

Yield stress profiles were performed using the correlation mentioned previously (Eq. (4.2)). As occurred

for the hardness profiles, the data was evaluated separately for each sample. The correspondent results

for the distinct Yield stress profiles are presented below (Figures 6.7, 6.8, 6.9 and 6.10).

62

Figure 6.7: Yield stress profile sample 5A


600

650

700

750

800

850

900

950

1000

-9 -8 -7 -6 -5 -4 -3 -2 -1 0 1 2 3 4 5 6 7 8 9

Yie

ld S

tre

ss (

MP

a)

Distance to the weld centre (mm)

C D E F G

600

650

700

750

800

850

900

950

1000

-9 -8 -7 -6 -5 -4 -3 -2 -1 0 1 2 3 4 5 6 7 8

Yie

ld s

tre

ss (

MP

a)

Distance to the centre weld (mm)

C D E F G

63



The Yield stress profiles showed a similar evolution to the one presented for the hardness case, for the

respective sample. This is in agreement with the fact that the correlation used to determine the Yield

Strength from the hardness value was linear. Table 6.3 indicates the highest and lowest values of Yield

strength for the samples analysed. The variation percentage in relation to the initial Yield strength value

was also calculated and the obtained results are presented in Table 6.3 as well.

550

600

650

700

750

800

850

900

950

1000

-10 -8 -6 -4 -2 0 2 4 6 8 10

Yie

ld S

tre

ss (

MP

a)


C D E F G

550

600

650

700

750

800

850

900

950

-10 -8 -6 -4 -2 0 2 4 6 8 10

Yie

ld S

tre

ss (

MP

a)


C D E F G

64

Table 6.3: Comparison between the initial yield strength and the lowest and the highest values

Sample Initial Yield

Strength (MPa)

Lowest Yield Strength

(MPa)

Variation %

Highest Yield Strength

(MPa)

Variation %

5A 785 605.3 -22.9 956.2 21.8

6A 785 628.3 20.0 967.7 23.3

7A 785 596.7 -24.0 961.9 22.5

8A 785 573.7 -26.9 941.8 20.0

Analysing the information provided by Table 6.3, it is possible to detect that the sample which presented

the lowest variation when it was considered the lowest Yield strength was sample 6A, however, it

presented the highest variation when was considered the highest yield strength. The opposite scenario

occurred in sample 8A.

Sample 6A presented the highest lowest yield strength and the highest yield strength leading to a better

strength overall when it was compared to the remaining samples.

6.4. Microstructure analysis

Before the microstructure analysis, the macroscopic effect of the different welding parameters used in

each weld is evaluated. The welding process affects largely the steel microstructure as referred in

Chapter 2. Apart from the base material area (also known as the unaffected zone), two more main zones

may be considered, at a first sight and in every sample: a weld zone, WZ, and a heat affected zone,

HAZ, as depicted in Figure 6.11.

Figure 6.11: Different zones due to the welding process

Starting with the WZ zone, 3 different subzones can be seen. At the bottom and presenting a darker

coloration is found the first run weld seam (position D, Figure 4.9). High temperatures of the 2nd weld

65

pass affected the previous weld, leading to a darker appearance. The intermediate part shows a lighter

colour tone and the reason may be related to the effect of the first weld pass in the second welding pass

(position E, Figure 4.9). Finally, the top of the 2nd weld seam (positon F, Figure 4.9) presents the lightest

colour tone of the weld seam and at this subzone and the intermediate subzone, the grains tended to

grow from the first subzone, contrary of what occurred at the bottom, where the grains tended to grow

towards the centre. Nonetheless, the grains, at the WZ zone demonstrate a sharpened and elongated

shape, which is shown further in this chapter.

Moving to the HAZ, four sub zones can be observed, especially on the 8A sample where the different

zones are more notorious. The first zone (positions H and G in figure 4.9) exhibits a similar colour tone

to the one presented at the bottom part of the weld zone. This coloration results of the high temperatures

reached. Next, a thin region presented a colour tone almost identic with the base material. This may

lead to believe that the microstructure can present some similarities. As the hardness tests were carried

out with samples etched it was possible to visualize that, at this location, the grains were fine, as

occurred in the base material although with a different geometry (more rounded in this case). The two

final sub-zones, adjacent to the last subzone, are difficult to distinguish due to the closer colour tone.

Figure 6.12: Weld geometries for the distinct samples. Sample 5A (case A), sample 6A (case B), sample 7A (case C), sample 8A (case D)

Taking into consideration all the samples (Figure 6.12), it was understood that the heat input increase

(arc energy) resulted in a larger effect on the steel structure. In order to justify the previous sentence, it

is enough to consider the lowest heat input case (sample 5A) and highest heat input case (sample 8A).

Comparing the two cases, the affected area on case 8A was significantly larger than the respective area

on the 5A case. Furthermore and regarding the two cases where were applied the lowest values of heat

input, few aspects can be noticed, apart from the already mentioned difference between affected areas.

66

Relatively to the second run, the weld seam, in case 6A, was larger and deeper than the observed in

5A case, as expected. This occurred, since the current and voltage values were slightly higher for 6A

and additionally, due to an inferior travel speed. As mentioned previously, a current and an arc tension

increase promote the formation of a wider weld seam. Although a voltage increase causes a penetration

reduction, the current increase and the reduction of the travelling speed overlapped that effect, leading

to a deeper weld beam in the 6A case. A similar scenario occurred when it was compared the two cases

with higher values of arc energy, 7A and 8A. In this situation, the welds were performed with the same

voltage value. Nevertheless, the difference between the weld seams depth was less prominent than the

one occurred in the previous comparison, despite the increase in the current value and the reduction of

the travelling speed. This relates to the fact that the travelling speed reduction was inferior, resulting in

a less notorious difference.

Performing now a microscopic analysis, the several microstructures analysed for the different defined

positions (Figure 4.9) and for each sample are evaluated. These microstructures can be observed, on

a greater scale, in Annex A3. The microstructures were captured with a 2000X and 5000X magnification.

The base material microstructure can be seen at positions A, B and C since, for all samples, the

temperatures reached, due to the welding process, are relatively insignificant, which is proven by the

similarity of the different samples. The microstructure at these positions is characterized by a fine and

elongated ferrite-bainite grain structure, as a result of successive rolling stages at high temperatures.

(Figure 6.13, 6.14 and 6.15)

Figure 6.13: Microstructures position A 2000X magnification

Figure 6.14: Microstructures position B 2000X magnification

67

Figure 6.15: Microstructures position C 2000X magnification

Considering now position D (weld zone) and sample 5A, the low heat input value promoted the formation

of a mainly bainitic structure (lath shape). Ferrite can also be seen since the peak temperature resultant

of the second run was near 800oC, allowing a partial austenitization and thus the formation of ferrite.

For sample 6A, a bainitic structure took place instead, since the cooling time was higher in this case.

Once again ferrite occurred, however in this case on a higher scale since the temperature peak provided

by the second run was also high which allowed a higher recrystallization. For 7A and 8A samples, the

zone near the resin was analysed. Nevertheless, a ferritic-bainitic structure is present in sample 7A.

Sample 8A revealed a coarser grain structure, mainly composed by ferrite (Figure 6.16).

Figure 6.16: Microstructures position D 2000X magnification

Attending position E, the temperatures achieved at this position are less known, since the temperature

profiles (in annex A2) were associated with the face side and root side, respectively, and E is found in

the middle of the weld. Nevertheless, the microstructures for the different samples (Figure 6.17)

presented a similar structure, as presented in sample 5A, position D. In sample 6A is clear to observe a

coarsening of the ferrite phase, when compared to the 5A sample microstructure. Samples 7A and 8A

presented, once again, higher ferrite content as a result of a higher cooling rate. It is important to be

aware that, in sample 7A, corrosion was presented in the microstructure.

68

Figure 6.17: Microstructures position E 2000X magnification

At position F, due to the high temperatures reached, the bainitic structure tended to appear in great

amount. Once again, in sample 7A and 8A the microstructures were analysed near the resin and a

bainitic-ferrite structure was found (Figure 6.18).

Figure 6.18: Microstructures position F 2000X magnification

Regarding position G, in sample 5A, the main structure is given by the first weld since the second weld

led to relatively low temperatures (about 400oC) promoting only a refining of the grain. A ferritic bainitic

structure was present, however, the bainite laths are in less evidence. In sample 6A, the second weld

pass caused temperatures around 700oC, leading to a partial recrystallization and a slight grain growth.

However, a ferritic-bainitic structure took place as well, although the ferrite content had increased. For

sample 7A and 8A due to a higher cooling rate, a structure mainly composed by ferrite occurred (Figure

6.19).

Figure 6.19: Microstructures position G 2000X magnification

69

For position H, the temperature profiles are also unknown, as occurred for position E. However, the

microstructures of samples 6A and 7A showed similarities with the ones presented in sample 7A and 8A

at position G. Considering the microstructures of the 7A sample at position H and G, it is possible to

observe that a finer grain structure occurred at H, when compared with G. This fact may be explained

due to the occurrence, at position H, of a higher and closer peak temperature, caused by the second

weld run, to the temperature in which the formation of a finer structure is promoted. In the 6A case, it

appears to have a recrystallization compared to the structure presented at G. Relatively to sample 8A,

the differences between position G and H are patents. At H, the laths of bainite are evident in the ferritic

matrix. Despite the higher cooling rate, the peak temperature reached was also higher, since occurred

an increase in heat input. This led to the formation of a bainitic structure and a grain coarsening of the

ferritic matrix (Figure 6.20).

Figure 6.20: Microstructures position H 2000X magnification

At position I the reached temperature was around 1400oC, which promoted the formation of a ferritic-

bainitic structure, as the laths of bainite are easily observed. The sample 8A, due to a higher cooling

time, presented a more dominant ferrite structure. For sample 7A, it was only possible to present the

microstructure using a 5000X magnification, however, a similar structure to the one found in sample 8A

was observed (Figure 6.21)

Figure 6.21: Microstructures position I 2000X magnification, apart from sample 7A which was performed with a 5000X magnification

70

6.5. Numerical simulation results

As mentioned previously for the base material specimen, several meshes with different numbers of

points along the thickness were applied, to evaluate the convergence of the evolution on the SIF values.

The number of points ranged between 15 and 33. The convergence analysis reveals extremely

important since it allows to determine the suitable mesh in which, from the iterative process, the solution

converges to a constant value.

The evaluation was performed for an 11 mm crack length and a 40 mm crack length. The main results

in both analysis are presented in Table 6.4 and 6.5, where z is the coordinate along the specimens’

thickness. The origin of the axis is located at the surface of one of the specimen’s side. For T and HAZ

specimens, the origin is located on the side where the weld is not present.

Table 6.4: SIF values for several number of points along the thickness a=11 mm

Nº points

SIF_middle (z=4mm) 𝑴𝑷𝒂√𝒎

SIF_average 𝑴𝑷𝒂√𝒎

SIF (z=3mm) 𝑴𝑷𝒂√𝒎


15 29.4 28.6

17 29.5 28.6 29.5 29.3

19 29.5 28.6

21 29.5 28.6 29.1

23 29.5 28.6

25 29.5 28.6 29.4 29.2

27 29.5 28.6

29 29.5 28.6 29.1

31 29.5 28.6

33 29.5 28.6 29.4 29.2

Table 6.5: SIF values for several number of points along the thickness a=40mm

Nº points

SIF_middle (z=4mm) 𝑴𝑷𝒂√𝒎

SIF_average 𝑴𝑷𝒂√𝒎



15 245.9 236.7

17 246.6 236.9 246.0 243.8

19 246.0 237.0

21 246.5 237.2 242.5

23 246.1 237.3

25 246.4 237.4 245.3 243.2

27 246.1 237.4

29 246.3 237.5 242.7

31 246.1 237.6

33 246.3 237.6 245.6 243.0

In both cases, the value at the middle of the specimen converge to a determined value. Similarly, to z=3

and z=2. However due to the smaller data obtained, it was only possible to compute the value for specific

71

meshes. Despite in the previous cases the amount of data retrieved was smaller, since it was only

possible to obtain the value at those coordinates in certain meshes. For that reason, the convergence

of the SIF values at those positions was less noticed. For the middle coordinate, the SIF value evolution

as the number points increased is presented in Figure 6.22.

Figure 6.22: Evolution of the middle SIF value for a=40 mm

From Table 6.4 and 6.5 it is possible to conclude that the average values stabilize with the increase of

the number of points. The evaluation of the average SIF value revealed to be important, since the

average SIF value was used to compare to the SIF value obtained from the standard. An interesting

aspect can be noticed since, as the number of points used increased, the average value tended to

deviate from the values obtained from the E 647-00 standard [74], Eq. (6.1).

∆𝐾 =∆𝑃

𝐵√𝑊

(2 + 𝛼)

(1 − 𝛼)32

(0.88 + 4.64𝛼 − 13.32𝛼2 + 14.72𝛼3 − 5.6𝛼4) (6.1)

Where 𝛼 is the coefficient between the crack length (𝑎) and specimen width (𝑊), ∆𝑃 represent the load

variation (𝑃𝑚𝑎𝑥 − 𝑃𝑚𝑖𝑛), while 𝐵 is the specimen thickness.

In figure 6.23, it is presented the SIF evolutions as a function of crack length obtained from numerical

simulation and from the standard equation.

72

Figure 6.23: Comparison between numerical simulation SIF values and standard SIF values

The data provided by the numerical simulation presented a higher deviation from the standard data, as

the crack length increases. The highest relative error and the lowest relative error are shown in Table

6.6 and also the SIF values obtained from the simulation and from the standard definition for the lowest

and highest crack length analysed.

Table 6.6: SIF values provided by standard and numerical analysis for a=11 and a=40 mm

Crack length

(mm)

SIF Abaqus

𝑴𝑷𝒂√𝒎

SIF standard

𝑴𝑷𝒂√𝒎

Difference

𝑴𝑷𝒂√𝒎

Relative error

%

11 28.6 25.3 3.3 13.0

40 236.7 230.3 6.4 2.8

Analysing both cases, and as seen in Figure 6.23, the SIF value obtained from numerical simulation

was higher than the obtained from the standard definition.

Afterwards, a comparison between the SIF for the remaining cases and the BM case was made. The

different evolutions of the average SIF value, as the crack length increases, are presented below (Figure

6.24) for the different specimens analysed.

It is possible to see that, at a first phase, for lower crack length values, the medium stress intensity factor

is lower for the SIF_T specimen. This fact can be related to the local increase of thickness. As mentioned

before, the increase of thickness leads to an inferior toughness value.

In Figure 6.25, it is presented the percentage variation between the BM specimen and the other

specimens, in order to observe more clearly the difference between all cases. The relative variation, 𝑒,

was calculated according to Eq. 6.2.

𝑒 =𝑆𝐼𝐹𝐵𝑀 − 𝑆𝐼𝐹𝑖

𝑆𝐼𝐹𝐵𝑀 × 100 (6.2)

0

50

100

150

200

250

10 15 20 25 30 35 40

SIF

(𝐌𝐏𝐚√𝐦

)

Crack length a (mm)

SIF Abaqus SIF Standard

73

Where, 𝑆𝐼𝐹𝐵𝑀 represents the SIF value for the BM case and 𝑆𝐼𝐹𝑖 is the correspondent value for the other

two cases (T and HAZ case).

Figure 6.24: SIF evolution for the distinct cases

Figure 6.25: SIF relative variation between BM case and the other cases

As mentioned before, the SIF value for the T case for lower crack lengths is clearly inferior to the values

obtained for the other cases. From a crack length of 20 mm, SIF value for T case was higher than the

respective SIF values obtained for the other cases. Regarding the HAZ case, it can be said that the

relative variation was almost constant and the SIF value for the HAZ case was always lower than the

value for the BM case.

At the end, the two charts, SIF_T and SIF_BM, overlapped (Figure 6.24), since the weld geometry effect,

at a considerable distance, was seen as negligible. The SIF evolution in both cases tended to behave,

as represented in Figure 6.26.

20

70

120

170

220

10 15 20 25 30 35 40

SIF


)

Crack lenght a (mm)

SIF_T_case SIF_HAZ_case SIF_BM_case

-2

0

2

4

6

8

10 15 20 25 30 35 40 45

Re

lati

ve v

aria

tio

n (

%)

Crack length a (mm)

HAZ case T case

74

Figure 6.26: SIF evolution along the thickness T case a=40mm

The maximum value occurred in the middle region of the specimen due to the existence of a triaxial

stress state at that region. The contrary was found near the surface, where a stress plane state took

place since, at the surface, the material is less constrained, which led to an inferior SIF value. This

previous evolution is similar to the presented in Nakamura’s [17] and Garcia-Manrique’s [18] works.

For the T specimen, travelling along the weld seam, the SIF evolution tends to vary significantly. Two

points were analysed, a=13 and 18 (Figure 6.27), in order to evaluate the different evolutions. The

respective evolutions are shown in Figure 6.28 and 6.29.

Figure 6.27: Location of the crack tip in both cases

Figure 6.28: SIF evolution along the thickness for T case a=13mm

200

205

210

215

220

225

230

235

240

245

250

0 1 2 3 4 5 6 7 8

SIF


)

z coordinate (mm)

5

10

15

20

25

30

35

0 1 2 3 4 5 6 7 8 9

SIF


)

z coordinate (mm)

75

Figure 6.29: SIF evolution along the thickness for T case a=18mm

Considering the crack length equal to 13mm, where the crack tip was located at the weld, SIF presented

a clear minimum value, on the side where the weld was located, which was the reason for T_case

presented the smallest average value (comparing with the other specimens) when smaller cracks length

were considered. This fact is due to the local increase of thickness which highlighted the presence of a

plane stress state. As the crack length increases, the difference between the two SIF values for the two

surfaces decreases, as the influence of the weld also decreases (Figure 6.19) and in both sides the SIF

value was fairly equal.

Moving further along the weld and considering a crack length in which the crack tip is near the weld

(a=23mm, Figure 6.30), a different evolution was obtained, (Figure 6.31).

Figure 6.30: Crack tip location for a=23mm

33

34

35

36

37

38

39

40

0 1 2 3 4 5 6 7 8 9

SIF


)

z coordinate (mm)

76

Figure 6.31: SIF evolution along the thickness T case for a=23mm

In this case, the minimum value occurs at the surface where the weld was not present. The weld material

tended to restrain the nearest material, at the surface, leading to a higher SIF on that side of the weld.

The triaxial effect was extended due to the local increase in thickness and propagated to the vicinity

areas (crack tip). A similar evolution was detected for the HAZ case for the several crack lengths

analysed, as shown in Figure 6.32. The reason behind this evolution is identical to the case previously

mentioned.

As it is possible to observe, the lowest SIF value occurred at the surface of the side where a weld was

not present. A similar evolution was observed on the T specimen when the crack tip was localized below

the weld, however, it was positioned at a relatively small distance of the weld.

Figure 6.32: SIF evolution along the thickness HAZ case for a=12mm

43

44

45

46

47

48

49

50

51

52

0 1 2 3 4 5 6 7 8

SIF


)

z coordinate (mm)

25

26

27

28

29

30

31

0.00 1.00 2.00 3.00 4.00 5.00 6.00 7.00 8.00

SIF


)

z coordinate (mm)

77

7. Conclusions and Recommendations

The analysis of the welding effect with four distinct cooling times (5, 10, 15 and 20 seconds) on

hardness and Yield Strength was performed. In addition, an evaluation on the distinct microstructures

due to different welding parameters at defined positions was carried out. The DIC method was applied

to measure the several crack lengths for the two fatigue tests. Finally, the SIF evolution with the crack

length value was evaluated for 3 different geometries, BM, T and HAZ. The SIF evolution was also

performed along the specimen thickness for the different specimens.

7.1. Conclusions

Regarding the fatigue results, it is only possible to mention that for both tests the Paris laws were well

addressed since the correlation coefficient was fairly close to the unit and the constant values were

within the typical values for steel. In the other hand, for the DC method several conclusion can be

expressed. It was detected that the FFT, which was applied to remove noise from the displacement field

derived from the DaVis software, did not reproduced the original vertical displacement. This problem

compromised the crack length measurements.

The combined results of the microstructure analysis and the hardness results revealed that the increase

in the heat input promoted the formation of ferrite. This led to a hardness reduction, especially near the

weld, since ferrite is a softer phase than bainite and martensite. This result is in agreement with previous

works, such as Prasad [57].

Regarding the hardness results, it was found that softest points were locates closer to the weld centre

in contrast with the hardest points. This means that the peak temperatures were higher for the softest

points than for the hardest, which it was confirmed by the temperature evolution on the face side and

root side.

Moreover, it was seen that the peak temperatures for the hardest and softest points ranged 700-750oC

and 900-9500C, respectively, with the exceptions of the hardest point in sample 5A and softest point in

sample 8A. In the first case, the low cooling time promoted the formation of harder structures at the

vicinity of the weld zone, while in the second one the high heat input value promoted the grain

coarsening, which allied with the formation of softer structures, led to a minimum value of hardness. In

fact, the grains at the weld tended to present a larger size than the grains of other zones, which, with

the grain coarsening due to the high heat input values and the formation of softer phases, promoted the

existence of a softer point at the weld. Another important conclusion is that the highest variations of the

hardness values occurred in lines located at the first weld run. As consequence the microstructure were

affected by the second weld run as well.

The main conclusion that can be taken from the yield profiles is that the sample 6A presented better

yield strength values than the remain samples and for that reason the set of welding parameters to

choose for the studied steel is the 6A set.

Finally, regarding the numerical analysis, the comparison between the average SIF values as a function

of crack length showed that the average SIF evolution was almost the same for the different weld

geometries, since it was hard to distinguish the different evolutions for the different geometries. So it is

78

possible to conclude that geometric effect introduced by the weld does not affected the SIF value. Only

in the T case and for lower crack length values, the average SIF values showed an inferior value when

compared to the remaining cases. This fact was justified by the SIF evolution for a crack length of 13

mm, which showed a non-symmetrical evolution, presenting a minimum value at the surface where the

weld was located. This value was considerably inferior to the remaining SIF values due to the local

increase in thickness. As the crack tip moved away from the weld, the evolution became symmetric,

which was shown for a 40 mm crack length and presented a similar evolution to the observed for the

BM case. For the HAZ case, the evolution was found non-symmetric as well, however, it was less evident

than in the previous case (T case with a crack length equal to 13 mm), since the difference between the

two SIF values at both surfaces was lower than observed previously. In this case, and comparing both

surface SIF values, the highest value was witnessed at the side which the weld was located,

consequence of the local increase in thickness. Although, in this case, the crack tip was not located at

the weld, it was located near it. In contrast to the evolution noticed for the T case, in this case the non-

symmetric evolution remained for all the crack lengths analysed.

7.2. Future works

From this work, there were some aspects identified it was identified which may be important to evaluate

in further researches, such as:

Evaluate the hardness profiles for other welding processes, such as Friction stir welding (FSW)

and submerged arc welding (SAW), and also study the influence of these on fatigue strength;

Analyse the Young’s modulus variation along the weld zone and heat affected zone, which can

be performed using the method proposed by Oliver [79];

DIC measurements for the other welding geometries beyond the geometry evaluated, even

using a different noise filter, instead of the FFT, to avoid issues on acquiring the crack length

data;

Compare the displacement field provided by the DaVis software with the analytical equation for

the displacement presented in Chapter 2;

A detailed analysis of the microstructure at the weld, near the surface, for the different weld

parameters, in order to understand the effect of the different microstructures on mechanical

properties at the weld;

Study the influence of different environmental condition, such as oxygen environment and high

temperatures, on fatigue crack propagation for this steel;

79

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85

Annexs

86

A1 Fatigue data

The following information indicates the crack length and the respective number of cycles necessary to

promote such crack length in the both cases.

Table A.1: Case 1

N (cycles

number) a (m)

715000 0,016

720000 0,016113

725000 0,01668

750000 0,0169

765000 0,01713

785000 0,01775

790000 0,017895

825000 0,018

835000 0,018455

855000 0,018471

865000 0,0188

870000 0,019

880000 0,019571

890000 0,020228

895000 0,020229

900000 0,0207

901000 0,020819

914000 0,021241

917000 0,021349

920000 0,021678

923000 0,021917

925000 0,022146

926000 0,022771

935000 0,023452

936000 0,023707

939000 0,023917

941000 0,02392

942000 0,024491

945000 0,024542

946000 0,025367

87

Table A.2: Case 2

A2 Temperature profiles and Modelled molten and austenittized zones

This information was obtained from the previous work [78]. It represents the temperature profile that the

steel was subjected at the welding stage. The last figure (Figure A5) indicates the areas where the steel

was completely molten and the austenitization process occurred at the welding stage, respectively.

948000 0,025444

949000 0,025583

950000 0,026

952000 0,026272

953000 0,026438

954000 0,026688

N (cycles

number) a (m)

705000 0,015915

710000 0,016182

720000 0,016857

725000 0,017765

735000 0,018703

745000 0,019875

750000 0,020944

752000 0,021013

753000 0,021262

754000 0,021505

755000 0,021636

756000 0,022058

757000 0,022263

758000 0,022552

759000 0,022807

760000 0,023136

761000 0,023202

762000 0,023777

763000 0,023942

764000 0,024273

765000 0,024603

766000 0,024961

767000 0,02561

768000 0,025794

769000 0,026434

770000 0,026891

771000 0,027395

773000 0,028518

774000 0,029389

775000 0,030169

776000 0,031431

955000 0,026875

956000 0,027245

957000 0,027438

958000 0,027931

959000 0,02826

960000 0,028667

961000 0,029117

962000 0,029359

963000 0,029768

964000 0,030263

965000 0,030719

966000 0,031978

967000 0,032391

968000 0,033081

969000 0,034143

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Figure A.1: Temperature profiles sample 5A - Face side (top image) Root side (bottom image)

Figure A.2: Temperatures profile sample 6A - Face side (top image) Root side (bottom image)

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Figure A.5: Modelled molten and austenitized zones overlapped with actual macrographs for root pass and for all second passes

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A3 Microstructure

Microstructures of the several positions established in Figure 4.9

Figure A.6: Position A, Top images 2000X magnification; Bottom images 5000X magnification

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Figure A.7: Position B, Top images 2000X magnification; Bottom images 5000X magnification

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Figure A.8: Position C, Top images 2000X magnification; Bottom images 5000X magnification

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Figure A.9: Position D, Top images 2000X magnification; Bottom images 5000X magnification

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Figure A.10: Position E, Top images 2000X magnification; Bottom images 5000X magnification

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Figure A.11: Position F, Top images 2000X magnification; Bottom images 5000X magnification

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Figure A.12: Position G, Top images 2000X magnification; Bottom images 5000X magnification

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Figure A.13: Position H, Top images 2000X magnification; Bottom images 5000X magnification

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Figure A.14: Position I, Top images 2000X magnification; Bottom images 5000X magnification

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