CHARACTERIZATION OF STEEL MICROSTRUCTURES BY

MAGNETIC BARKHAUSEN NOISE TECHNIQUE

A THESIS SUBMITTED TO THE GRADUATE SCHOOL OF NATURAL AND APPLIED SCIENCES

OF MIDDLE EAST TECHNICAL UNIVERSITY

BY

KEMAL DAVUT

IN PARTIAL FULFILLMENT OF THE REQUIREMENTS FOR

THE DEGREE OF MASTER OF SCIENCE IN

METALLURGICAL AND MATERIALS ENGINEERING

DECEMBER 2006

Approval of the Graduate School of Natural and Applied Sciences Prof. Dr. Canan Özgen Director

I certify that this thesis satisfies all the requirements as a thesis for the degree of Master of Science. Prof. Dr. Tayfur Öztürk Head of Department

This is to certify that we have read this thesis and that in our opinion it is fully adequate, in scope and quality, as a thesis for the degree of Master of Science. Assoc. Prof. Dr. C. Hakan Gür Supervisor Examining Committee Members Prof. Dr. Tayfur Öztürk (METU, METE) Prof. Dr. Şakir Bor (METU, METE) Prof. Dr. Macit Özenbaş (METU, METE) Assoc. Prof. Dr. C. Hakan Gür (METU, METE) Dr. İbrahim Çam (METU, Central Lab.)

iii

I hereby declare that all information in this document has been obtained and presented in accordance with academic rules and ethical conduct. I also declare that, as required by these rules and conduct, I have fully cited and referenced all material and results that are not original to this work. Name, Last name : Kemal DAVUT

Signature :

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ABSTRACT

CHARACTERIZATION OF STEEL MICROSTRUCTURES BY

MAGNETIC BARKHAUSEN NOISE TECHNIQUE

DAVUT, Kemal

M.S., Department of Metallurgical and Materials Engineering

Supervisor: Assoc. Prof. Dr. C. Hakan Gür

December 2006, 80 pages

This aim of this thesis is to examine the possibility of using Magnetic

Barkhausen Noise (MBN) technique in characterizing the microstructures of

quenched and tempered low alloy steels as well as annealed low carbon

steels. To determine the average grain size by MBN, SAE 1010 steel

consisting of dominantly ferrite was used. The specimens were slowly cooled

in the furnace after austenitizing at different time and temperature variations.

By metallographic examination the average ferrite grain size of specimens was

determined. The magnetic parameters were measured by a commercial MBN

system. With increasing ferrite grain size, the magnetic Barkhausen jumps

caused by the microstructure were decreased due to the reduction in grain

boundary density per unit volume. A clear relationship has been observed

between average grain size and the magnetic Barkhausen noise signals. SAE

4140, 5140 and 1040 steels were used to characterize the microstructures of

quenched and tempered specimens. After austenitizing and quenching

identically, the specimens were tempered at various temperatures between

200oC and 600oC. Formation of the desired microstructures was ensured by

metallographic examinations and hardness measurements. The results show

that as tempering temperature increases the Barkhausen activity increases

due to the enhancement of domain wall displacement with softening of the

martensite. It has been shown that MBN is a powerful tool for evaluating the

microstructures of martensitic and annealed steels.

Keywords: Magnetic Barkhausen Noise, Microstructure, Grain Size,

Quenching, Tempering

v

ÖZ

ÇELİK İÇYAPILARININ MANYETİK BARKHAUSEN GÜRÜLTÜSÜ YÖNTEMİ

ile KARATERİZASYONU

Davut, Kemal

Yüksek Lisans, Metalurji ve Malzeme Mühendisliği Bölümü

Tez Yöneticisi: Doç. Dr. C. Hakan Gür

Aralık 2006, 80 Sayfa

Bu tezin amacı Manyetik Barkhausen Gürültüsü (MBG) tekniğinin su verilmiş

ve temperlenmiş düşük alaşımlı çeliklerin ve de tavlanmış düşük karbonlu

çeliklerin iç yapılarının karakterizasyonu için kullanılabilme olasılığını

incelemektir. Çeliklerde tane boyutunun MBG yöntemi ile tayini için ferrit içeriği

çok yüksek SAE 1010 çeliği kullanılmıştır. Numuneler farklı sıcaklık ve zaman

varyasyonlarında östenitlendikten sonra fırında soğutulmuştur. Metalografik

inceleme ile numunelerde ortalama ferrit tane boyutu ölçülmüştür. Numunelerin

manyetik parametreleri ticari bir MBN sistemi ile ölçülmüştür. İçyapıdaki ferrit

taneleri irileştikçe Barkhausen zıplamaları birim hacimdeki tane sınırı

yoğunluğuna bağlı olarak azalmakta ve zayıf manyetik gürültü sinyallerine

neden olmaktadır. Ortalama tane büyüklüğü ile MBN sinyalleri arasında

belirgin bir ilişki gözlemlenmiştir. Su verilmiş ve temperlenmiş yapıları

karakterize etmek için SAE 4140, 5140 ve 1040 çelikleri kullanılmıştır. Benzer

şekilde östenitlenip su verilen numuneler 200oC ve 600oC aralığında

temperlenmiştir. İç yapı metalografik inceleme ve sertlik ölçümleriyle

karakterize edilmiştir. Temperleme sıcaklığı arttıkça Barkhausen aktivitesi,

yumuşayan martensitik yapıda domen duvarlarının daha kolay hareket

edebilmesi sayesinde artmaktadır. Bu çalışma MBG tekniğinin martensitik ve

tavlanmış çeliklerin içyapılarını değerlendirmek için uygun bir yöntem olduğunu

göstermiştir.

Anahtar Kelimeler: Manyetik Barkhausen Gürültüsü, İçyapı, Tane büyüklüğü,

Su Verme, Temperleme

vi

To my family

vii

ACKNOWLEDGEMENTS The author wishes to express his deepest gratitude to his supervisor

Assoc. Prof. Dr. C. Hakan Gür for his guidance, understanding and

continuous support throughout the study.

The author would also like to thank METU-Central Laboratory for the

MBN measurements and would like to express his gratitude to Dr.

İbrahim Çam for introducing him with the concept of MBN and helping

with the measurements.

The author is deeply grateful to Cengiz Tan for his help with the SEM

appointments.

The author would also like to thank research assistant Caner Şimşir for

all his help throughout the study.

The technical assistance of Mr. Hüseyin Çolak and Mr. Özdemir Dinç in

the heat treatment and metallography laboratories are gratefully

acknowledged.

The author wishes to express his gratitude to his parents Lale and Haluk

Davut for always supporting him and encouraging him to continue.

The author would also like to thank Oya Araslı, Naciye Tezel, Serhat

Önol Şakar, Selen Gürbüz, Gül Çevik for giving him the strength to finish

his degree by sharing all the good and the bad times.

viii

TABLE OF CONTENTS PLAGIARISM.....................................................................................................iii ABSTRACT...................................................................................................... iv ÖZ.......................................................................................................................v ACKNOWLEDGMENTS...................................................................................vii TABLE OF CONTENTS..................................................................................viii LIST OF TABLES..............................................................................................xi LIST OF FIGURES.............................................................................................x Chapter 1 – Introduction ..................................................................................1

1.1 Theory of Magnetic Barkhausen Noise.....................................................1

1.2 Literature Survey.....................................................................................16

1.2.1 Determination of Grain Size of Steels by MBN ................................16

1.2.2 Characterization of Quenched and Tempered Steel Microstructures

by MBN .....................................................................................................17

1.3 The Aim of the study ...............................................................................18

Chapter 2 - Experimental................................................................................20 2.1 Experimental Procedure for Grain Size Determination ...........................20

2.2 Experimental Procedure for Investigation of Quenched and Tempered

Samples ........................................................................................................21

Chapter 3 Results ...........................................................................................22 3.1 Grain Size ...............................................................................................22

3.1.1 Microstructure ..................................................................................22

3.1.2 MBN Measurements ........................................................................29

3.2 Quenched and Tempered Microstructures .............................................34

3.2.1 Microstructure and hardness ...........................................................34

3.2.2 MBN measurements ........................................................................42

3.2.3 Detection of Faulty Quenching and Tempering Treatment by MBN 56

Chapter 4 Conclusions...................................................................................63 REFERENCES ...................................................................................................65

ix

LIST OF TABLES

TABLES

Table 2.1 Chemical composition of the SAE 1010 steel used (wt %)………..19

Table 2.2 Chemical composition of the steels used for quenching and

tempering (wt%)…………………………………………………………………….19

Table 3.1 Details of heat treatments applied to SAE 1010 specimens...........22

Table 3.2 Hardness values and MBN parameters of SAE 4140 specimens...47

Table 3.3 Hardness values and MBN parameters of SAE 4140 specimens...47

x

LIST OF FIGURES

FIGURES

Figure 1.1 Domains in a ferromagnetic material……….………………………...1

Figure 1.2 Structure of a domain wall...............................................................2

Figure 1.3 Colloid patterns collected over domain walls of a polycrystalline

silicon iron…………………………………………………………………………….4

Figure 1.4 Colloid patterns on cobalt with magnetic field applied normal to

surface (a) -130 oersteds; (b) +130 oersteds………………………………….....4

Figure 1.5 Domain structures on barium ferrite observed by the Kerr effect (a)

H = 2700 oersteds; (b) H = 3000 oersteds…………………………….………….5

Figure 1.6 Domains in gadolinium garnet with a field of varying strength

observed by the Faraday effect……………………………………….……………5

Figure 1.7 Series showing the process of magnetization reversal in a bloomed

film of Ni-Fe as observed by the Kerr effect......................................................6

Figure 1.8 Magnetization steps.........................................................................7

Figure 1.9 Barkhausen jumps during magnetization………..............................9

Figure 1.10 A typical MBN system configuration……………………………….9

Figure 1.11 Passage of a domain wall through an inclusion………………….10

xi

Figure 1.12 Variation of the probability of nucleation, annihilation, growth of

domains and domain wall density as a function of H…………………………11

Figure 1.13 MBN effect during hysteresis ……………………………………..12

Figure 1.14 Barkhausen event at various depths of specimen ………………13

Figure 1.15 The raw MBN data (a), MBN fingerprint (b), frequency spectrum

(c) and pulse height distribution (d).................................................................15

Figure 3.1 Optical micrographs and corresponding grain size distribution

histograms of group 1 specimens; (a) 1100A, (b) 1200A, (c) 1300A ………23

Figure 3.2 Optical micrographs and corresponding grain size distribution

histograms of group 2 specimens; (a) 1100B, (b) 1200B, (c) 1300B …...……24

Figure 3.3 Optical micrographs and corresponding grain size distribution

histograms of group 3 specimens; (a) 700A-1, (b) 700A-2, (c) 700A-3 ……...25

Figure 3.4 Effect of annealing temperature on grain configuration at various

stages of annealing…………………………………………………………………27

Figure 3.5 SEM micrograph (a) and corresponding EDX analysis taken from

(b) grain boundary and (c) grain interior of 700A-1 specimen…………………28

Figure 3.6 MBN profiles of group 1 specimens……………………………….30

Figure 3.7 MBN profiles of group 2 specimens……………………………….30

Figure 3.8 MBN profiles of group 3 specimens……………………………….30

Figure 3.9 Optical (a) and SEM micrographs (b); corresponding EDX analysis

taken from grain boundary (c) and grain interior (d) of 1300B specimen…….31

xii

Figure 3.10 The relation between RMS and AGS-0.5 …………………………. 33

Figure 3.11 SEM micrographs of quenched and tempered SAE 4140

specimens.......................................................................................................35

Figure 3.12 SEM micrographs of quenched and tempered SAE 5140

specimens.......................................................................................................36

Figure 3.13 Continuous cooling transformation (CCT) diagram and

hardenability curve of SAE 4140 steel……………………………………….…..37

Figure 3.14 Continuous cooling transformation (CCT) diagram and

hardenability curve of SAE 5140 steel……………………………………….…..38

Figure 3.15 MBN profiles of SAE 4140 specimens……………………………44

Figure 3.16 MBN profiles of SAE 5140 specimens……………………………44

Figure 3.17 The SEM micrographs of as quenched SAE 4140 (a) and SAE

5140 (b) steels and schematic representation of magnetic microstructure in

martensite........................................................................................................45

Figure 3.18 The SEM micrographs of SAE 4140 (a) and SAE 5140 (b) steels

tempered at 600oC and schematic representation of magnetic microstructure in

tempered martensite....................................................................................46

Figure 3.19 Frequency spectrums of SAE 4140 specimens…………….........48

Figure 3.20 Frequency spectrums of SAE 5140 specimens…………….........48

Figure 3.21 Pulse height distributions of SAE 4140 specimens………….…..50

Figure 3.22 Pulse height distributions of SAE 5140 specimens………….…..50

xiii

Figure 3.23 Coercivity (a), remanence (b), permeability (c) values of SAE 4140

specimens ………………………………………..…………………………..52

Figure 3.24 Coercivity (a), remanence (b), permeability (c) values of SAE 5140

specimens………………………………………….…………………………53

Figure 3.25 Hardness correlation for SAE 4140 specimens…………………..55

Figure 3.26 Hardness correlation for SAE 5140 specimens…………………..55

Figure 3.27 Continuous cooling transformation diagram of SAE 1040 steel..56

Figure 3.28 SEM micrographs taken from edge (a) and centre (b); MBN

profiles (c) and frequency spectrums (d) of as-quenched SAE 1040

specimen…………………………………………………………………………….58

Figure 3.29 SEM micrographs taken from edge (a) and centre (b); MBN

profiles (c) and frequency spectrums (d) SAE 1040 specimen tempered at

300oC……………………………………………………………..………………….59

Figure 3.30 SEM micrographs taken from edge (a) and centre (b); MBN

profiles (c) and frequency spectrums (d) SAE 1040 specimen tempered at

500oC……………………………………………………………..………………….60

Figure 3.31 SEM micrographs taken from edge (a) and centre (b); MBN

profiles (c) and frequency spectrums (d) SAE 1040 specimen tempered at

600oC……………………………………………………………..………………….61

Figure 3.32 Schematic drawing of edge and centre regions of SAE 1040

specimens………………………………………………………..………………….62

Figure 3.33 MBN profiles of edge regions of SAE 1040 specimens……….62

1

CHAPTER 1

INTRODUCTION

1.1 Theory of Magnetic Barkhausen Noise

An electrical charge in motion creates a magnetic field. Since electrons are electrical

charges their motion can produce magnetic moments in two ways.

i) Their orbital motion around nucleus generates a small magnetic field and they have

magnetic moment along their axis of rotation.

ii) The spinning of electrons also cause a magnetic moment directed along spin axis.

This moment can only be in up or down direction.

In materials having completely filled levels contribute zero magnetic moment, because pairs

of electrons in each level have opposite spin and cancel each other. Ferromagnetic materials

like iron, cobalt and nickel have occupied 4s levels leaving vacant orbits in 3d shell.

Therefore in ferromagnetic materials, instead of canceling each other, moments line up by

an exchange force of quantum mechanical nature [1]. Actually no single comprehensive

theory exists for explaining ferromagnetic behavior. Two distinct theories, band theory and

molecular field theory, try to explain the phenomenon on the basis of lowering exchange

energy. The exchange energy refers to the part of electrostatic energy of a system of

electrons which depends on the spin states of neighboring electrons. In case of

ferromagnetic materials this exchange energy is lowest when the spins of the 3d electrons in

adjacent atoms are aligned parallel [2]. Because of this ferromagnetic materials exhibit a

large spontaneous moment due to the cooperative alignment of unpaired electron spins

along a certain direction.

Figure 1. Domains in a ferromagnetic material [2]

2

Ferromagnetic materials have regions called domains where all magnetic moments are

aligned in the same direction as seen in Figure 1.1. The existence of domains is a

consequence of energy minimization. A single domain specimen would have large

magnetostatic energy due to Coulomb interaction between magnetic free poles. The breakup

of magnetization into domains provides flux closure at the ends of specimen and hence

reduces the magnetostatic energy which represents the total free energy of the domain

structure. In each domain the magnetization is equal to saturation magnetization and the net

magnetization of a material is the vector sum of the magnetization within all domains. The

direction of alignment varies randomly from domain to domain, although certain

crystallographic axes, easy directions of magnetization, are preferred especially in the

presence of a magnetic field. The internal magnetization is stable when pointing parallel to

one of these easy directions. There are transition layers in which the magnetic moments

realign between the domains and therefore belong to neither domain as seen in Figure 1.2.

These transition layers are called domain walls or sometimes referred as Bloch walls. The

interactions between the magnetic moments and the crystal lattice cause magnetization to lie

in easy directions of magnetization. The magnetocrystalline anisotropy energy is the energy

related, the dependence of internal energy on direction of inner magnetization, which is

minimum for domains located parallel to easy directions of magnetization [1, 3]. The

magnetocrystalline anisotropy energy tends to make the domain walls thinner in order to

increase the number of spins pointing in easy directions of magnetization. The exchange

energy tends to make the walls thicker since it is minimized when neighboring moments are

aligned parallel. As a result of this competition the domain wall has a certain finite width and

a certain structure [4, 5].

Figure 1.2 Structure of a domain wall [5]

3

Presence of domains has been proven by innumerable experimental observations, both

direct and indirect. The first conformation was the indirect detection of domains by the

Barkhausen effect in which the reorientation of domains caused discrete changes in

magnetic induction within a ferromagnet. Domains are normally so small that one must use a

microscope to see them directly [5]. What one sees then depends on the technique involved.

The most popular direct observation techniques fall in two groups:

1. Those which disclose domain walls (Bitter method, electron microscope). The

individual domains, whatever their direction of magnetization, look more or less the

same, but the domain walls are delineated. The Bitter or powder method involves

the application of an aqueous suspension of extremely fine (colloidal) particles of

magnetite Fe3O4 to the polished surface of the specimen and it can detect moving,

as well as stationary domain walls as shown in Figures 1.3 and 1.4. Transmission

electron microscopy, which is often called Lorentz microscopy, can disclose domain

walls in specimens thin enough to transmit electrons. The electrons passing through

the specimen will be deviated due to different orientations of local magnetization

vector across a domain wall. This method has the advantage of high resolution,

which allows the examination of the fine detail of domain structure. It also permits

the direct observation of interactions between domains, crystal imperfections and

grain boundaries.

2. Those which disclose domains (optical methods involving the Kerr or Faraday

effects). Here domains are magnetized in different directions appear as areas of

different color, and the domain wall separating them appears merely as a line of

demarcation where one color changes to the other. The Kerr effect is a rotation of

the plane of polarization of a light beam during reflection from a magnetized

specimen. The Kerr method is ideal for studies of domain walls in motion (Figures 1.5 and 1.7) and has largely supplanted the Bitter method for such studies. The

Faraday effect is a rotation of the plane of polarization of a light beam as it is

transmitted through a magnetized specimen (Figure 1.6). This method is similar to

Kerr effect but it is limited to specimens thin enough to transmit light [6].

Changes in domain structure can occur by two principal means. Either the magnetization

within each domain can coherently rotate to a direction parallel to the applied field or the

boundary between two domains can move causing the entire magnetization change to be

localized at the domain boundary. Thus magnetization changes as a result of both of domain

wall motion and domain rotation. At lower applied magnetic fields the domain walls are

stretching so that they return to the non-magnetized state on removal of the applied field. Up

4

Figure 1.3 Colloid patterns collected over domain walls of a polycrystalline silicon-iron [6]

Figure 1.4 Colloid patterns on cobalt with magnetic field applied normal to surface.

(a) -130 oersteds, (b) +130 oersteds [6]

5

Figure 1.5 Domain structures on barium ferrite observed by the Kerr effect

(a) H = 2700 oersteds, (b) H = 3000 oersteds [6]

Figure 1.6 Domains in gadolinium iron garnet with a field of varying strength observed by the

Faraday effect [5]

6

Figure 1.7 Series showing the process of magnetization reversal in a bloomed film of Ni-Fe

as observed by the Kerr effect [6]

7

to that point the wall motion is reversible. When the applied field is increased further, the

process occurring is irreversible wall motion, in other words the growth of domains which are

aligned closely parallel to applied magnetic field (H) grow at the expense of others until the

magnetic structure becomes a single domain pointing in one of easy directions of

magnetization. Still greater H makes domain rotation predominant and in this region work

must be done against anisotropy forces; a rather large increase in H is required to produce a

relatively small increase in magnetization [5, 7]. Figure 1.8 shows these steps of

magnetization.

Figure 1.8 Magnetization steps [7]

During the magnetization process imperfection like dislocations or impurity elements in the

metal cause an increase in the energy lost in the form of a kind of internal friction. These

imperfections give rise to hysteresis. Also higher magnetocrystalline anisotropy gives rise to

hysteresis. Many ferromagnetic materials can be characterized from parameters obtained

from hysteresis curves. These parameters are:

8

• Permeability (µ): Measure of the degree to which the material can be magnetized.

The relation between magnetic induction (B) and applied magnetic field strength (H)

is as follows:

HB .µ= (1.1)

• Remanence: When the field is reduced to zero after magnetizing a magnetic

material the remaining magnetic induction is called remanent induction (BR) and the

remaining magnetization is called the remanent magnetization (MR).These are

related by the following formula:

RoR MB .µ= (1.2)

where µ0 is the permeability of vacuum. The remanence is used to describe the

value of either remaining induction or magnetization when the field has been

removed after the material has been magnetized to saturation. The remanent

induction or magnetization is used to describe the remaining induction or

magnetization when the field has been removed after magnetizing to an arbitrary

level.

• Coercivity: By applying a reverse magnetic field strength of Hc the magnetic

induction declines to zero. The coercive field is the magnetic field needed to reduce

the magnetization to zero from an arbitrary level whereas coercivity is the magnetic

field required to reduce magnetization to zero from saturation.

• Saturation magnetization (Ms): As H is increased the magnetization finally reaches

saturation at value designated Ms. At saturation all the magnetic moments are

aligned in the direction of applied magnetic field.

All ferromagnets when heated to sufficiently high temperatures become paramagnetic due to

the increased randomness of atomic moment. This transition temperature, above which the

thermal energy overcomes the exchange forces, is called the Curie temperature. At this

temperature the permeability of the material drops suddenly and both coercivity and

remanence become zero. The Curie temperature of iron is 770oC.

When a ferromagnetic material is magnetized, changes in physical dimensions, in general

occur. These changes could be longitudinal, transverse, or volumetric and are known as

magnetostriction. The dimensional change occurring along the direction of induced magnetic

field is called Joule effect magnetostriction and the converse is known as the Villari effect. [7]

9

Figure 1.9 Barkhausen jumps during magnetization [4]

High resolution examination of magnetization cycles of ferromagnetic materials reveals

discontinuous flux changes as shown in Figure 1.9. This effect is discovered by Prof.

Barkhausen in 1919 and named Barkhausen effect. During magnetization if a search coil is

placed close to the surface of the specimen and connected to an oscilloscope or computer

as shown in Figure 1.10, voltage spikes can be observed. These voltage spikes are known

as Magnetic Barkhausen Noise (MBN). At first these discontinuities in magnetization was

attributed to sudden discontinuous rotation of domain but now it is known that discontinuous

domain wall motion is the most significant factor. In fact both of these mechanisms occur and

contribute to MBN [8].

Figure 1.10 A typical MBN system configuration [9]

High Resolution Examination

10

When a domain wall bisects an inclusion, magnetostatic energy can be reduced to zero, at

the cost of a little wall energy, if closure domains form as shown in figure 1.11(a).

Observation of moving domain walls in crystals has shown that domain walls are pinned by

interaction of the moving wall with the spike domains attached to inclusion rather than by interaction of the inclusions themselves. Figure 1.11(a-d) shows the passage of a domain

wall through an inclusion. In response to an upward applied field, the wall in (a) moves to

right, as seen in (b) dragging out the closure domains into the form of tubes and creating a

new domain just to the right of the inclusion. Further motion of the main wall lengthens the

tubular domains as in (c). The change from (a) to (b) to (c) is reversible and the domain

arrangement (a) can be regained if the field is reduced. But if the field is increased further

the tubular domains do not continue to lengthen infinitely because their increasing surface

area adds two much wall energy to the system. A point is reached when the wall suddenly

snaps of the tubular domains irreversibly and jumps a distance to the right, leaving two spike domains attached to the inclusion as in (d) [5]. This is a Barkhausen jump and detected in

the form of voltage pulses induced in a search coil positioned close to the specimen surface.

Not only inclusions but also other microstructural features such as dislocations and grain

boundaries can cause Barkhausen jumps by pinning domain walls.

Figure 1.11 Passage of a domain wall through an inclusion [5]

During MBN measurements representative magnetic hysteresis loop is induced in the small

volume due to the energy loss with the irreversible process of magnetization. This

irreversible process mentioned above is strongly related to the dynamic behavior of domains,

i.e., nucleation, annihilation and growth of domains. Grain/lathe boundaries, dislocations and

precipitates affect this dynamic behavior. A schematical illustration in Figure 1.12 shows the

variation of the probability of the nucleation, annihilation and growth of domains as well as

the domain wall density as a function of H. The given magnetization loop is divided into two

parts, namely Region 1 and 2 as seen in Figure 1.12(a).

11

Figure 1.12 Variation of the probability of nucleation, annihilation, growth of domains and

domain wall density as a function of H [10].

When existing field strength is reduced from saturation magnetization (Region 1), new

domains are nucleated close to precipitates and grain or lathe boundaries where magnetic poles are accumulated during the spin rotation. The domain nucleation can proceed when

the reduction in the magnetostatic energy associated with the poles during the formation of

new domains is greater than the work required to form domains. The probability of the

domain nucleation, Pn, increases with decreasing H in order to reduce the magnetostatic

energy associated with the poles as shown in Figure 1.12(b). Thus, the density of domain

walls, Fd, increases with decreasing H as Figure 1.12(c) shows. The driving force for the

growth of domains arises from the difference between the applied field energy and the

domain wall energy. Thus, the probability for the growth of domains, Pg, decreases with

decreasing H as the wall energy of the domain preferentially oriented to the H direction decrease. When a magnetic field in the reverse direction increases beyond Hc (Region 2),

domains formed in region 1 grow at the expense domains aligned unfavorably. This growth

causes annihilation of unfavorably aligned domains and results in a decrease of Fd. Thus,

the domain wall density, Fd, is highest around H=0. As the driving force for the domain wall

propagation increases with increasing H, Pg as well as the probability of domain annihilation,

Pa, increases as shown in Figure 1.12(b) [10].

12

The total angular displacement across a domain wall is often 180o or 90o, particularly in cubic

materials because of the anisotropy and the change in direction of moments take place

gradually over many atomic planes. MBN is principally caused by the motion of 180o walls.

At the same time non-180o domain wall motion will have to occur due to presence of closure

domains. However, contributions of the non-180o domain walls to MBN are smaller than

those of the 180o walls for two reasons:

i) The average velocity of 180o walls are larger than that of the non180o walls,

ii) The volume swept out by 180o walls is larger than that by non-180o walls [10]

The Barkhausen effect is strongest on the steepest part of the magnetization curve as shown

in Figure 1.13. MBN is sensitive to changes in mechanical stress, composition and

microstructural features such as grain size, inclusions, precipitates, dislocations. Because of

the number of influential variables, for material characterization this technique makes relative

comparisons between material states.

Figure 1.13 MBN effect during hysteresis [11]

Eddy currents arise in any conducting material in which magnetization is changing and flow

in such a direction as to produce magnetic fields opposing the change [12]. Eddy currents

become more significant as magnetizing frequency is increased. The MBN signal is

13

attenuated by this eddy current opposition (Figure 1.14). Thus we have a limited

measurement depth for MBN measurements given by the formula 1.3:

0....1

µµσπδ

rf= (1.3)

where

σ = conductivity

f = frequency content of MBN

µr = relative permeability

µo = permeability of vacuum

The MBN signal has frequency contents up to 2 MHz and measurement depths for practical

applications vary between 0.01 to 1.5 mm.

Figure 1.14 Barkhausen event at various depths of specimen [13]

For MBN method typical measurement parameters are as follows:

• Magnetizing frequency (Hz): Adjusts the frequency of magnetizing field applied to

the specimen. When magnetizing frequency is increased and magnetizing voltage

kept constant, magnetization level will decrease. On the other hand the magnetic

noise level, including peak heights and root-mean-square (RMS) voltages, will

increase due the fact that much more domain walls are participated during

magnetization at higher frequencies.

14

• Waveform: Different waveforms have been used for the alternating excitation field to

obtain MBN signals. Triangular and sinusoidal waves are used commonly whereas

square waves rarely. Similar MBN signals obtained when triangular or sinusoidal

waveforms are used.

• Magnetizing voltage (V): Adjusts the magnitude of magnetizing field applied to the

specimen. The actual level of magnetization depends on the sensor used and can

be checked. Increasing magnetizing voltage increases the level of magnetization.

• Magnetizing field strength (H): The field strength can be varied by changing the

magnetizing frequency and voltage simultaneously. In some sources the field

strength is preferred instead of magnetizing voltage for MBN signal analysis. The

signal level increases to a maximum at an intermediate field strength and then

decreased at higher fields. The increase was attributed to greater capacity for

overcoming pinning obstacles when the field is getting larger, and the decrease to

the predominance of domain rotation over domain wall motion at very high fields.

• Number of bursts: Determines how many magnetizing half cycles or Barkhausen

Noise bursts will be stored for signal analysis. Increasing the number of bursts

makes results more reliable whereas increases data analysis time.

• Sampling frequency: Determines how many samples per second are stored for

signal analysis. The sampling frequency is adjusted by regarding the magnetizing

and analyzing frequencies. The sampling frequency should be at least twice the

maximum analyzing frequency for consistency.

The raw Magnetic Barkhausen Noise data consists of series of voltage pulses and their

associated applied field values obtained as a function of time shown in Figure 1.15(a). The

raw MBN signal may be amplified and filtered for detailed analyses. By filtering the noise

with a band-pass filter the background noise can be eliminated and measurement depth can

be varied according to the formula 1.1. After being amplified and filtered the signal is ready

to be analyzed by the following methods and parameters:

• MBN fingerprints: Using a definite sampling frequency, a local root-mean-square

(RMS) value is calculated and plotted against applied field strength. This

instantaneous local value is averaged over several field cycles whose number is

determined by the number of bursts. Also a smoothing algorithm may be applied and

a fingerprint as shown schematically in Figure 1.15(b) is obtained. The fingerprints,

sometimes referred as MBN envelopes or profiles, are characterized by the

maximum noise amplitude referred as MBN peak height and the corresponding

magnetic field referred as peak position.

15

Figure 1.15 The raw MBN data (a), MBN fingerprint (b), frequency spectrum (c) and pulse

height distribution (d).

• Frequency spectrum: The frequency content of the noise can be obtained by using

Fourier analysis. In the literature a plot such as Figure 1.15(c) is referred as

frequency spectrum where y-axis is the amplitude or noise power (V2).

• Pulse height distribution: Another plot frequently encountered in literature is the

pulse height distribution which gives the size distribution of pulses as shown in

Figure 1.15(d) schematically.

• Representative B-H curves: The Barkhausen Noise is derived from magnetization

cycles. Thus the sum integral of rectified bursts gives a simulation of the hysteresis

loop. From that simulation coercivity, remanence and permeability can be calculated.

It should be stated that all these calculated parameters from the simulated

hysteresis can not be used as true values because the saturation magnetization of

the specimens are far beyond the actual level of magnetization reached locally

during MBN measurements.

In addition, single parameters such as total number of pulses, maximum pulse size and root

mean square (RMS) of all signal amplitudes can be used to characterize the noise signal.

16

1.2 Literature Survey Magnetic Barkhausen Noise (MBN) technique has proved its viability for characterization of

microstructures and it is considered as a valuable non-destructive evaluation (NDE)

technique for microstructural characterization of ferromagnetic materials. The dual sensitivity

of the phenomenon to stresses and to microstructure on which this study is focused, gives a

wide range of potential applications to the technique including determination of grain size of

steels.

1.2.1 Determination of Grain Size of Steels by MBN

S. Titto et al. studied non-destructive measurement of grain sizes of 500 samples of 0.17 –

0.37mm low alloy low carbon steel in soft annealed and temper annealed condition. Grain

sizes in the range of 5 – 25 µm were studied and a linear relation was found between grain

size and the amplitude of the MBN signal [14].

In a study of R. Ranjan et al. the peak heights and RMS values of the Barkhausen signal

was found to be increased with increasing grain size. Specimens of decarburized steel were

undergone the following sequence of hot rolling, cold rolling, annealing, temper rolling and

decarburization annealing (%C < 0.0005 wt%) and grain sizes of 70 – 120 µm were obtained

[15].

Another study investigated the influence of applied tensile stress and grain size on MBN in

SAE 1005 steels. Specimens having ferrite grain sizes of 20 – 45 µm were prepared by

applying variations of furnace and air cooling. The MBN peak heights decreased as ferrite

grains get coarser [16].

C. Gatelier-Rothea et al. made a similar study using ultrahigh purity iron with less than 20ppm impurities. By annealing the samples at 450oC, 600oC and 750oC for 3 hours

equiaxial grains of 50 – 300 µm were obtained. The results were similar to [16]; peak heights

decreased when grain size increased [17].

The effects of grain size and grain boundary misorientation were studied in pure iron

prepared by vacuum melting. Samples containing 50 -180 µm grains were prepared by

annealing at various temperatures between 1193 and 1393 K. Hall-Petch like relationship

was found between the grain size and RMS values. Moreover the frequency spectrums of

the MBN signals analyzed and the grain size was related to the ratio of two definite

frequency components. The RMS values and the ratio between the definite frequency

components decrease as grain size increases [18].

17

R. Ranjan used pure nickel samples, whose impurity content is less than 75ppm, instead of

iron or steel to investigate the effect of grain size on MBN. Samples composed of 20 – 240

µm grains were obtained by annealing at different temperatures between 500oC – 800oC for

three hours. The MBN fingerprints showed two peak behavior and the ratio between the

second and first peak heights decreased as grains became coarser [10].

Sakamoto et al. analyzed MBN signals theoretically and relate the RMS values of MBN to

grain size as follows:

2/1. −= dCRMS g (1.4)

In this Petch like relation “d” is the average grain size; Cg is independent of the grain size

and given by:

( ) ( )2

4/1

4/5

max

3.9.

28/

⎟⎟⎠

⎞⎜⎜⎝

⎛∆Φ=

KItC

NHdtdHC

S

totalvg

γπ

(1.5)

where H is the magnetic field strength, N is the total number of Gaussian pulses in the cross-

sectional unit area of a specimen, ∆Φ is the quantity of magnetic flux change in a

microregion, Cv is a constant, ttotal is the total time of generation of Gaussian pulses, γ is the

wall energy per unit area, Is is the saturation magnetization and K is the magnetic anisotropy

constant. In this study voltage pulses in each micro-region, that MBN is composed of, are

approximated by Gaussian pulses in order to facilitate the mathematical treatment [19].

1.2.2 Characterization of Quenched and Tempered Steel Microstructures by MBN

Blaow et al. used Ovako 677 steel to study effects of tempering on MBN. The specimens

were austenitised at 950oC for 1hour, followed by air cooling. The martensite structure

produced was tempered at 180oC and 400oC. The profiles obtained from MBN

measurements were characterized by peak height, peak position and area under the profile.

The differences between profiles of tempered structures were not significant [11].

J. Kameda used AISI 4340 steel austenitised at 850oC for 1 hour, oil quenched and

tempered at various temperatures between 100oC and 500oC for 1 hour. The peak height of

the MBN signal was sensitive to microstructure change induced during tempering [20].

18

O. Saquet et al. studied tempering induced changes on MBN, using water quenched XC 55

steel austenitised at 875oC and tempered at various temperatures ranging from 100oC to

600oC. The MBN fingerprints showed that peak heights increased and the peak positions

shifted to lower fields as tempering temperature increased. In addition a simple model was

proposed for the source mechanism of MBN signals [21].

Moorthy et al. used MBN to characterize the microstructures of water quenched 0.2%C steel

solutionised at 950oC for 1 hour and tempered at 600oC for 0.5 – 100 hour. Peak heights and

positions of MBN fingerprints were used to differentiate various carbide size distributions. A

single peak was observed in samples tempered for 0.5 and 1 hour whereas after 5 hour

tempering MBN fingerprints showed a clear two peak behavior [22].

The effect of tempering was also studied for case carburized EN 36 steel in another study by

Moorthy et al. The specimens were case carburized to produce a surface carbon content of

0.85%C and a case depth of 1mm followed by oil quenching from 900oC. The specimens

were tempered at 192oC for 2 hour and at 250oC for 4 hour. Effect of tempering was studied

using a range of magnetizing frequencies and a number of analyzing frequency ranges for

characterizing the MBN signal. A correlation between hardness depth profile and peak height

of the MBN fingerprint was found [23].

In another study 12% Cr Mo V stainless steel, which was solutionised in the range of 950oC

– 1150oC for 1 hour, was used to investigate the effects of tempering on MBN. The

solutionised samples were then air cooled and tempered in the range of 200oC -800oC for 1

hour. The noise energy and number of MBN pulses were the parameters used for signal

analysis. No significant change was observed in MBN signal due to tempering up to 500oC;

whereas increasing tempering temperature further caused a rapid increase in both of the

parameters [24].

1.3 The Aim of the study

Steel is the most important and widely used industrial commodity. Microstructure affects the

properties, especially mechanical properties, of steels. New designs require steels providing

longer service life with higher performance which makes quality control essential. Increasing

demand brings a growing need for non-destructive inspection of steel components. Magnetic

Barkhausen Noise (MBN) measurement provides a good alternative to traditional methods in

terms of fastness and accuracy. Nevertheless the MBN method has not yet been fully exploited

when compared with some other non-destructive methods ultrasound. The aim of this study is

non-destructive characterization of steel microstructures by Magnetic Barkhausen Noise method.

19

The aim of the first part of this study is determination of grain size of steels by MBN method.

Average ferrite grain size is one of the most important microstructural features that affect the

properties of steel components. Annealing is a widely used industrial process in order to

refine the grain, induce softness, improve electrical and magnetic properties, and improve

machinability as well. Annealing consists in heating the steel to the proper temperature and

then cooling slowly through the transformation range in the furnace up to low temperatures.

Annealing may be divided into 3 stages: recovery, recrystallization and grain growth. Various

grain sizes and distributions may be obtained by annealing due to differences in degree of

prior deformation, impurity content of steel, annealing time and temperature. Traditional

metallographic and mechanic methods, that involve taking representative specimens, can

not allow 100% inspection of steel work pieces. Besides being destructive and time

consuming makes these traditional methods too slow for present production rates. Instead of

traditional methods, MBN measurements may be used to determine grain size quickly, easily

and without damaging the material.

The aim of the second part of this study is characterizing the microstructures of quenched

and tempered steels by MBN method and correlating the hardness of these steels with MBN

findings. Steels are widely utilised in different industries, usually in the form of quenched and

tempered components. Any temperature up to the lower critical may be used for tempering;

thus an extremely wide variation in properties and microstructure ranging from those of as

quenched martensite to spheroidized carbides in ferrite can be produced by tempering.

Ultimately it is the balance of hardness (or strength) and toughness required in service that

determines the conditions of tempering for a given application [25]. If the principal desired

property is hardness or wear resistance, the part is tempered at about 200oC; if the primary

requirement is toughness, the part is then tempered above 400oC. Residual stresses are

relieved almost completely when the tempering temperature reaches 500oC. For consistency

and less dependence on time, quenched steel components generally tempered for 1 to 2

hours [26]. The industry has been searching for methods capable of characterizing material

properties accurately, quickly and without damaging the material. The MBN technique is

considered here a candidate method for such mechanical and microstructural

characterizations.

20

CHAPTER 2

EXPERIMENTAL

2.1 Experimental Procedure for Grain Size Determination 8 mm specimens were cut from cold drawn 50 mm diameter SAE 1010 steel. Table 2.1

shows the composition of the steel used. All the machining operations were done before the

heat treatments in order to avoid surface residual stresses.

Table 2.1 Chemical composition of the SAE 1010 steel used (wt%)

Steel Type C Cr Mo Mn Si P S Fe

SAE 1010 0.113 0.073 0.039 0.503 0.216 0.006 <0.001 Bal.

One group of specimens was annealed at 700oC for 2, 6, 16 and 24 hours. Another group

was austenitized at 900oC, 1000oC, 1100oC, 1200oC and 1300oC for 30 and 90 minutes

followed by furnace cooling.

For metallographic investigation the samples were finely ground and polished with diamond

paste. In order to reveal ferrite grains with enhanced contrast, color etching was used. After

polishing, the specimens were etched first by 5% Nital followed by bisulfate. The surfaces of

specimens were examined under optical microscope and scanning electron microscope.

Average ferrite grain size and grain size distributions for each specimen were obtained by

analyzing the photographs of examined surfaces using Clemex Image Analyzer software.

MBN measurements were performed using a commercial system (Rollscan, µscan 500-2). The

sensor S1-138-13-01 was used for the MBN measurements. A sinusoidal cyclic magnetic field

with an excitation frequency of 10 Hz was induced in a small volume of the specimen via a

ferrite core C-coil. The Barkhausen signals were filtered with a wide band-pass filter (1-200

kHz), amplified with a gain of 50 dB, and then, analyzed using the Rollscan-software. The peak

magnetizing voltage was 15V and sampling frequency was 2 MHz.

21

Before MBN measurements, in order to eliminate the effects of remanent magnetization on

results, all the specimens were passed through demagnetization tunnel. Zero remanent

magnetization for all specimens was ensured by Gauss-meter measurements.

2.2 Experimental Procedure for Investigation of Quenched

and Tempered Samples The specimens of 7 mm-thick and 22 mm diameter were prepared from the hot rolled SAE

4140 bar and 15 x 15 x 7 mm specimens from SAE 5140 bar. In order to investigate the

effects of low hardenability, specimens of 15 x 15 x 7 mm were prepared from SAE 1040

bar. Table 2.2 gives the chemical composition of the steels used. All the cutting and grinding

operations were done prior to the heat treatments in order to avoid surface machining

residual stresses. Austenitization was done under controlled atmosphere to avoid oxidation

and decarburization. All specimens were quenched in water after austenitization at 860oC for

30 minutes. Then, specimens were separately tempered at 200oC, 300oC, 400oC, 500oC and

600oC for 90 minutes. One specimen from each type of steel was left as-quenched.

Table 2.2 Chemical compositions of the steels used for quenching and tempering (wt%)

Steel Type C Cr Mo Mn Si P S Fe

SAE 1040 0.416 0.233 0.047 0.800 0.423 0.020 <0.001 Bal.

SAE 4140 0.475 0.942 0.224 0.840 0.202 0.023 0.015 Bal.

SAE 5140 0.491 1.143 0.053 0.730 0.312 0.016 0.047 Bal. Before metallographic investigation, the samples were finely ground, polished with diamond

paste and etched with 2% Nital. The through-thickness sections of the specimens were

examined using optical microscope and scanning electron microscope. For each specimen

an average hardness value was determined by measuring Vickers hardness at different

locations.

During MBN measurements a sinusoidal cyclic magnetic field with an excitation frequency of

125 Hz was induced in a small volume of the specimen via a ferrite core C-coil. The

Barkhausen signals were filtered with a wide band-pass filter (0.1-1000 kHz), amplified with

a gain of 20 dB, and then, analyzed using the Rollscan-software. The peak magnetizing

voltage was 10V and sampling frequency was 2 MHz. During the analyses average of 154

bursts were used to obtain Barkhausen parameters for each specimen where each burst

represents one half of the magnetization cycle.

22

CHAPTER 3

RESULTS

3.1 Grain Size

3.1.1 Microstructure

Figure 3.1, 3.2 and 3.3 show the optical micrographs and related grain size distribution

histograms of annealed SAE 1010 specimens. In order to enhance the contrast between

ferrite grains color metallographic techniques were used. Average grain sizes (AGS) and

size distribution histograms were obtained from direct measurement of about 100 grains per

specimen. The histograms indicate that the maximum distribution percentages are around

the average grain size value. Table 3.1 shows the details of heat treatments applied and the

corresponding average grain sizes (AGS) obtained for every SAE 1010 specimen. By

changing the time or temperature of annealing treatment, specimens composed of various

grain sizes are obtained.

The term annealing has been used in its broadest sense to refer to any heat treatment that

has as its objective the development of a nonmartensitic microstructure of low hardness and

high ductility. This understanding of annealing is much too broad, however, and a number of

Table 3.1 Details of heat treatments applied to SAE 1010 specimens

Specimen Code

Annealing Temperature

Annealing time

AGS (µm)

1100A 1100oC 30 min 40,2 ± 22

1200A 1200oC 30 min 54,1 ± 27

Gro

up 1

1300A 1300oC 30 min 58,2 ± 32

1100B 1100oC 90 min 40 ± 20

1200B 1200oC 90 min 47,5 ± 26

Gro

up 2

1300B 1300oC 90 min 63 ± 32

700A-1 700oC 2 hour 22 ± 10

700A-2 700oC 6 hour 24,1 ± 10

Gro

up 3

700A-3 700oC 24 hour 26,7 ± 11

23

Group 1 Specimens

0

5

10

15

20

25

0 10 20 30 40 50 60 70 80 90 100 110 120 130 140 150 160 170 180

Grain Size (mm)

Perc

enta

ge (%

)

0

5

10

15

20

25

0 10 20 30 40 50 60 70 80 90 100 110 120 130 140 150 160 170 180

Grain Size (mm)

Perc

enta

ge (%

)

0

5

10

15

20

25

0 10 20 30 40 50 60 70 80 90 100

110

120

130

140

150

160

170

180

Grain Size (mm)

Perc

enta

ge (%

)

Figure 3.1 Optical micrographs and corresponding grain size distribution histograms of group 1 specimens, (a) 1100A, (b) 1200A, (c) 1300A.

c) 1300A; AGS = 58,2 µm

b) 1200A; AGS = 54,1 µm

a) 1100A; AGS = 40,2 µm

24

Group 2 Specimens

0

5

10

15

20

25

0 10 20 30 40 50 60 70 80 90 100 110 120 130 140

Grain Size (mm)

Perc

enta

ge (%

)

0

5

10

15

20

25

0 10 20 30 40 50 60 70 80 90 100 110 120 130 140

Grain Size (mm)

Perc

enta

ge (%

)

0

5

10

15

20

25

0 10 20 30 40 50 60 70 80 90 100 110 120 130 140

Grain Size

Perc

enta

ge (%

)

Figure 3.2 Optical micrographs and corresponding grain size distribution histograms of group 2 specimens, (a) 1100B, (b) 1200B, (c) 1300B.

c) 1300B; AGS = 63 µm

b) 1200B; AGS = 47,5 µm

a) 1100B; AGS = 40 µm

25

Group 3 Specimens

0

5

10

15

20

25

30

0 5 10 15 20 25 30 35 40 45 50 55 60

Grain Size (mm)

Perc

enta

ge (%

)

0

5

10

15

20

25

30

0 5 10 15 20 25 30 35 40 45 50 55 60

Grain Size (mm)

Perc

enta

ge (%

)

0

5

10

15

20

25

30

0 5 10 15 20 25 30 35 40 45 50 55 60Grain Size (mm)

Perc

enta

ge (%

)

Figure 3.3 Optical micrographs and corresponding grain size distribution histograms of group 3 specimens, (a) 700A-1, (b) 700A-2, (c) 700A-3.

c) 700A-3; AGS = 26,7 µm

b) 700A-2; AGS = 24,1 µm

a) 700A-1; AGS = 22 µm

26

more specific annealing heat treatments have been developed and defined. Full annealing is

a heat treatment accomplished by heating steels into a single-phase austenite field and

slowly cooling, usually in a furnace, through the critical transformation ranges [25]. Group 1

and 2 specimens were full annealed at various temperatures as shown in Table 3.1.

The optical micrographs of Group 1 (Figure 3.1) and Group 2 (Figure 3.2) indicates that

specimens are composed of equiaxed ferrite grains with various AGS. The formation of pro

eutectoid products like proeutectoid ferrite, has all the manifestations of a process of

nucleation and growth; the nuclei appear at the austenite grain boundaries and grow until a

continuous layer is formed along the grain boundary. Further growth merely thickens the

grain boundary and at very low contents, this continues until all the austenite is transformed.

The transformation of austenite into proeutectic ferrite is a good example of heterogeneous

nucleation and occurs preferentially at grain boundaries. Any steel with fine austenite grains

has greater grain boundary surface area, consequently favors nucleation and reduces

incubation period [27]. Thus, with increasing austenite grain size parallel to the annealing

temperature, average grain size of proeutectic ferrite increases. The increase in austenite

grain size, which in turn affects ferrite grain size, is controlled by a time and temperature

dependent growth mechanism. The comparison of Group 1 and 2 specimens show that the

effect of temperature on growth of austenite is more dominant as expected.

Group 3 specimens were annealed below the eutectoid temperature; a treatment named as

process and recrystallization annealing. Prior to heat treatments the samples were in cold

deformed condition and this type of annealing is usually applied to soften and restore

ductility of cold worked steel products. Most of the energy expended in cold work is released

as heat during deformation. However, a small percent is stored by dislocations and some

other crystal imperfections introduced during deformation. The stored energy is the driving

force for the changes during annealing. On heating, high strain energy of the deformed

ferrite first drives recovery, a mechanism by which some of the crystal imperfections are

eliminated or rearranged into new configurations. Then ferrite grains having high dislocation

density are replaced by new grains having much lower dislocation density. The decrease of

energy associated with dislocations is the driving force for recrystallization. Eventually

recrystallized grains grow at the expense of other recrystallized grains. The driving force for

this grain growth is the reduction of energy associated with grain boundaries [28]. Various

grain sizes and distributions may be obtained by annealing due to differences in degree of

prior deformation, impurity content of steel, annealing time and temperature. The effects of

prior deformation and impurities were eliminated by using the samples from the same cold

deformed SAE 1010 steel bar. Figure 3.4 shows the effect of temperature on the grain

configuration at different stages of annealing for constant specified time period. The

comparison of Group 3 optical micrographs with Figure 3.4 indicates that the temperature

27

chosen was enough to recrystallize and grow all the grains. By changing the annealing time

differences in growth of domain and hence samples composed of various AGS were obtained.

However, impurity segregation at grain boundaries pinned the growing grains. Figure 3.5 shows the SEM micrograph and corresponding EDX analysis. The white regions in the SEM

micrograph are impurities and corresponding EDX analysis shows that sulphur (S)

concentration at grain boundary is much higher than grain interior. Thus, the difference

between AGS of specimens was not significant.

Figure 3.4 Effect of annealing temperature on grain configuration at various stages of annealing [26] In order to lessen the effects of pearlite, SAE 1010 low carbon steel was used. Pearlite

amount calculated from equilibrium iron-cementite phase diagram is about 12%. All the heat

treatments were done in a decarburizing atmosphere; thus much less pearlite is present in

the resulting microstructures. All the specimens are composed of nearly 100% ferrite so the

effects of pearlite on MBN signals are neglected.

The hardness measurements indicate that all specimens have close hardness values about

100 HV. Previous studies and second part of this study report that MBN signals are sensitive to

hardness changes. On the other hand the hardness difference between SAE 1010 steels in

this study is small enough to be negligible. In addition all specimens were cooled in furnace

which eliminates the effects of residual stresses. Regarding these, the differences between

MBN signals is due to the variations in AGS of ferrite.

28

b) Grain Boundary EDX-Analysis = 1.44% S + 98.56% Fe (wt.%)

c) Grain interior EDX-Analysis = 100% Fe (wt.%)

Figure 3.5. SEM micrograph (a) and corresponding EDX analysis taken from (b) grain boundary and (c) grain interior of 700A-1 specimen

a)

29

3.1.2 MBN Measurements Figure 3.6, 3.7 and 3.8 show the MBN profiles of group 1, 2 and 3 specimens respectively.

The peak heights of group 1 profiles show a clear difference between the specimens.

However such clear difference does not exist for the group 2 specimens. Group 3 specimens

have overlapping profiles since their grain size difference is not so significant. For all groups

of specimens differences in grain size does not affect the peak positions of profiles.

In the introduction part it was mentioned that the MBN signal is generated due to sudden

changes in magnetization and the irreversible motion of 180o domain walls is the main

contribution. Thus the MBN signal can be written as:

dHBHVHcMBN oo .).().(. 180180

∆= ∫ ρ (3.1)

where “c” is a constant which depends on the search coil, the permeability and conductivity

of the sample; “ )(180

Hhρ ” is the density of 180o domain walls at a field H; “ )(180 HV o ” is the

average critical velocity of a 180o domain wall when it is released from pinning sites; and

“ B∆ ” is the average change in the local magnetism due to unit displacement per unit area of

domain walls. As the grain size increases in a ferromagnetic material, the domains get larger

[29]; thus )(180

Hhρ decreases. Moreover, with increasing grain size, the mean free path of

domain wall motion increases. Consequently, the incremental field (∆H) required for bulging

the domain walls before they are unpinned also increases. Thus )(180 HV o decreases with

increasing grain size. As a result of this it is expected that the amplitude of MBN signals

decrease with increasing grain size [10].

The theory above may not be true in the presence of grain boundary precipitates which

increase the net magnetostatic energy and alter the domain structure. The precipitates may

affect the MBN signals in two ways:

i) They can act as nucleation sites when domains are nucleated during magnetization. As

the density of precipitates increases )(180

Hhρ will increase too.

ii) When the density of precipitates increases, the spacing between neighboring

precipitates will decrease, and hence )(180 HV o will increase.

With the exception of 1300B specimen of group 2, the MBN response of specimens is

consistent with theory described above. For all specimens peak positions of profiles are

30

0

2

4

6

8

10

12

14

-100 -50 0 50 100

Magnetic Field Strength (% of max.)

Avg

. MB

N le

vel (

% o

f 5V)

1100 A1200A1300A

Figure 3.6 MBN profiles of group 1 specimens

0

2

4

6

8

10

12

14

-100 -50 0 50 100

Magnetic Field Strength (% of max.)

Avg

. MB

N le

vel (

% o

f 5V)

1100B1200B1300B

Figure 3.7 MBN profiles of group 2 specimens

0

2

4

6

8

10

12

14

-100 -50 0 50 100

Magnetic Field Strength (% of max.)

Avg

. MB

N le

vel (

% o

f 5V)

700A-1700A-2700A-3

Figure 3.8 MBN profiles of group 3 specimens

31

c) Grain Boundary EDX-Analysis = 1.37% S + 98.63% Fe (wt.%)

d) Grain interior EDX-Analysis = 100% Fe (wt.%)

Figure 3.9. Optical (a) and SEM micrographs (b); corresponding EDX analysis taken from grain boundary (c) and grain interior (d) of 1300B specimen

a) Optical Microscope b) SEM

32

nearly same and are around zero magnetic field strength indicating MBN activity occurs at

very early stages of magnetization. During these very early stages domain nucleation is the

dominant mechanism. Grain boundaries have high internal energy due to unsatisfied atomic

bonding, which makes them preferential sites for domain nucleation. Grain boundary area

per unit volume decreases as grain size increases, which in turn reduces the potential

nucleation sites and causes difficulty in formation of new grains. During later stages of

magnetization the domains grow and eventually rotate as the material reaches magnetic

saturation.

The specimens contain practically 100% ferrite and in such ferritic structures domains can

move freely in ferrite grains until they face a grain boundary. These pinned domain walls at

grain boundary can continue their motion by Barkhausen jumps which in turn generates

magnetic noise signals. As grains coarsen Barkhausen jumps lessen due to reduced grain

boundary density. As a result, difficulties in domain nucleation, reduced number of

Barkhausen jumps and domain density cause low MBN activity in coarse grained structures.

The specimen 1300B of group 2, although coarse grained, exhibits unexpectedly high MBN

activity. Careful examination of optical and SEM micrographs revealed segregation of

impurities at grain boundaries (Figure 3.9). These impurities alter the magnetic structure and

cause an increase in domain wall density and in critical velocity of a domain wall as

explained above. Hence the MBN activity of this sample is unexpectedly high.

Another parameter that reflects the MBN behavior and influenced in the same way as peak

height of profile, is the root-mean-square (RMS) of noise signals. A theoretical study found a

Petch like relation between RMS and AGS:

5.0).( −= AGSkRMS (3.2)

where “k” is a constant and does not depend on AGS [19]. The details of this equation can

be found in “Literature Survey part”. Figure 3.10 compares the theoretical and experimental

results of group 1. The experimental results are consistent with the theoretical result which

suggests a linear relation between RMS and square root of AGS. The correlations of group 2

and 3 specimens are not shown since the magnetic structures of those specimens are

altered by grain boundary segregations as shown in Figures 3.5 and 3.9. In literature good

correlations between MBN and AGS was found in purer steels. Better correlations may be

obtained by using steels with lower impurity concentrations and hence preventing grain

boundary segregations.

33

0

0.5

1

1.5

2

2.5

3

3.5

4

4.5

5

0.1 0.11 0.12 0.13 0.14 0.15 0.16 0.17 0.18 0.19 0.2

AGS-0.5

RM

S (V

)

Theoretical [19]Experimental

Figure 3.10 The relation between RMS and AGS-0.5

34

3.2 Quenched and Tempered Microstructures

3.2.1 Microstructure and hardness Representative SEM micrographs (Figure 3.11and 3.12) and hardness values (Table 3.2

and Table 3.3) of the SAE 4140 and SAE 5140 show that typical martensitic and tempered

martensitic structures were successfully obtained.

The as-quenched specimens of both steels are the hardest of all specimens due to the fact

that their martensitic structure has a tetragonal lattice with interstitial carbon in solid solution

and high dislocation density formed by shear. Unlike ferrite or pearlite martensite forms by a

sudden shear process in the austenite lattice which is not accompanied by atomic diffusion.

Ideally the martensite reaction is a diffusionless shear transformation, highly crystallographic

in character, which leads to a characteristic lath or lenticular microstructure. The shear

involved in martensitic transformation cause severe elastic strains in both the martensite and

surrounding austenite matrix. The volume of elastically strained material is greatly reduced if

the shape of the martensite is lath or lenticular [28]. Also should be stated that the terms lath

and plate refer to the three dimensional shapes of individual martensite crystals. During

metallographic examinations the cross sections will appear to be needlelike or acicular [25].

Martensite is a supersaturated solid solution of carbon in ferritic iron. The carbon atoms tend

to order in such a way that the crystal structure changes from body-centered cubic to body

centered tetragonal. The tetragonality is measured by the ratio between the axes, c/a, and

increases with carbon content. The tetragonality of martensite arises as a direct result of

interstitial solution of carbon atoms in the bcc lattice, together with the preference for a

particular type of octahedral site imposed by the diffusionless character of the reaction [30]. In short the combined effects of solid solution and dislocation strengthening, lattice distortion

due to internal strains, fine particle size of martensite crystals make as quenched specimens

hardest [27]. The as quenched hardness values of both steels are expected to be high and equal since

they contain same amount of carbon. Figure 3.13 and 3.14 show the continuous cooling

transformation diagrams and hardenability curves of both steels. These steels have close

critical cooling rates; the critical cooling rate of SAE 4140 is slightly slower. Both steels are

austenitized and quenched identically so any ferrite or pearlite product would probably seen

35

Figure 3.11 SEM micrographs of quenched and tempered SAE 4140 specimens

a) As quenched b) Tempered at 200oC

c) Tempered at 300oC d) Tempered at 400oC

e) Tempered at 500oC f) Tempered at 600oC

36

Figure 3.12 SEM micrographs of quenched and tempered SAE 5140 specimens

a) As quenched b) Tempered at 200oC

c) Tempered at 300oC d) Tempered at 400oC

e) Tempered at 500oC f) Tempered at 600oC

37

Figure 3.13 Continuous cooling transformation (CCT) diagram and hardenability curve of

SAE 4140 steel [31]

38

Figure 3.14 Continuous cooling transformation (CCT) diagram and hardenability curve of

SAE 5140 steel [31]

39

in SAE 5140 whose critical cooling rate is higher. Another condition of full hardening is

cooling below Mf temperature. Austenite stabilizing elements like carbon, manganese, nickel,

chromium and molybdenum will lower both the Ms and Mf temperatures. The alloying

additions, especially manganese and molybdenum, of SAE 4140 are slightly higher than

SAE 5140 (Table 2.2). Since quenching mediums were same the lower hardness of SAE

4140 may be due to its lower Ms and Mf temperatures.

A fully martensitic component cannot be put to engineering use since it lacks toughness and

ductility. The brittleness of martensitic microstructures is due to a number of factors that may

include lattice distortion caused by carbon atoms trapped in the octahedral sites of the

martensite, impurity atom segregation at austenite grain boundaries, carbide formation

during quenching and residual stresses produced during quenching. Hence tempering is

necessary and it aims to:

(i) relieve internal stresses associated with the lattice shear due to the formation of martensite,

(ii) restore ductility and toughness at sacrifice of strength and hardness,

(iii) improve the dimensional stability through the breakdown of any retained austenite.

It is generally considered that tempered structures possess superior mechanical properties

than lamellar aggregates of equivalent hardness.

The structure of a steel quenched to form martensite is highly unstable. Reasons for the

instability include the supersaturation of carbon atoms in the body-centered tetragonal

crystal lattice of martensite, the strain energy associated with the fine dislocation or twin

structure of martensite, the interfacial energy associated with the high density of lath or plate

boundaries, and the retained austenite that is invariably present even in low-carbon steels.

The supersaturation of carbon atoms provides the driving force for carbide formation; the

high strain energy the driving force for recovery; the high interfacial energy the driving force

for grain growth or coarsening of the ferrite matrix; and the unstable austenite the driving

force for transformation to mixtures of ferrite and cementite on tempering.

In order to differentiate different stages of tempering the quenched specimens of both steel

types are tempered at various temperatures between 200oC - 600oC. The SEM micrographs

and hardness values indicate expected structural changes occur upon tempering of both

steel types.

Tempering up to 250oC does not change the microstructure and hence the hardness

significantly. In this first stage ε-carbides precipitate and martensite partially loses its

tetragonality due to the fact that carbon atoms can diffuse in the tetragonal lattice.

40

Experimental observations indicate that during this stage of tempering, apart from the initial bct

martensite of original carbon content, another bct martensite with c/a = 1,012 – 1,013

corresponding to ~0,3% carbon in solution appears in the structure. Carbon is precipitated

as ε-carbide which is approximately Fe2,4C and has hexagonal closed packed crystal

structure. This carbide precipitates as narrow laths or rodlets on cube planes of the matrix

and the carbon atoms are at octahedral interstities being as far away from each other as

possible. The epsilon carbide is usually nucleated first not because it is more stable but

because it has a better lattice matching with the matrix and hence greater probability for

nucleation as coherent nucleation can occur without much strain energy [27].

During tempering between 300oC and 400oC cementite replaces ε-carbides and martensite

loses tetragonality. Cementite first appears in the microstructure as plate like particles of

200nm long and ~15nm thickness. Once formed, the cementite particles agglomerate and

grow until a typical spheroidized structure is obtained. During this second tempering stage

the most likely sites for the nucleation of the cementite are the ε-carbide interfaces with

matrix and as the Fe3C particles grow, the ε-carbide particles gradually disappear. Other

possible nucleation sites for cementite are the interlath boundaries of the martensite and the

original austenite grain boundaries. The tetragonality of the matrix disappears and it is then,

essentially, ferrite, not supersaturated with respect to carbon [30]. The dislocation density is

effectively lowered not only by the reduction of dislocations within the laths but also by the

elimination of the low-angle lath boundaries [25].

Cementite to coarsens and spheroidizes; ferrite recrystallizes when tempering temperature is

increased further, up to 500oC and 600oC. At this third stage of tempering cementite particles

undergo a coarsening process and essentially become spheroidized. The spheroidization of

the cementite is encouraged by the resulting decrease in surface energy. The particles which

preferentially grow and spheroidize are located mainly at interlath boundaries and prior

austenite boundaries, although some particles remain in the matrix. The boundary sites are

preferred because of the greater ease of diffusion in these regions. At higher end of this

stage which could be up to 700oC, the martensite lath boundaries are replaced by more

equiaxed ferrite boundaries by a process which is best described as recrystallization [30]. In

addition residual stresses are relieved almost completely when tempering temperature

reaches 500oC [26]. The micro-stress relief is suggested to be due to the diffusion of carbon

from regions of compression to regions of tension [27].

As tempering temperature increases, by the mechanisms explained above, martensite softens

as indicated by hardness values of the steels. Although same tempering treatments are

applied, SAE 4140 steels soften more than SAE 5140. This behavior can be attributed to the

influence of alloying elements on tempering.

41

Tempering is a softening reaction which can be retarded by judicious choice of alloying

elements. The most effective elements in this regard are strong carbide formers such as

chromium, molybdenum, and vanadium. Without these elements, iron-carbon alloys and low

carbon steels soften rapidly with increasing tempering temperature. Alloying elements have

little influence on the first stage of tempering but may raise the second by as much as 100 or

200oC. The second stage, during which cementite precipitates, requires the diffusion of

carbon and the effect of alloying elements, in the absence of the formation of any alloy

carbides, can be appreciated in terms of their effect on diffusion of carbon. Alloying elements

retard softening through retarding the agglomeration of cementite by decreasing the rate of

carbon diffusion or by increasing the stability of cementite by dissolving in it. In carbon steels

the tetragonality of the lattice is observable up to 300oC but in alloy steels containing

chromium, molybdenum, tungsten, vanadium, cobalt and silicon, tetragonality may be

maintained after tempering at 400-500oC. The alloying elements also increase the stability of

low carbon martensite. In contrast manganese and nickel decrease stability.

Several alloying elements, notably silicon, chromium, molybdenum and tungsten, cause the

cementite to retain its fine structure to higher temperatures, either by entering into the

cementite structure or by segregating at the carbide-ferrite interfaces. They significantly

delay the softening process during 2nd and 3rd stages of tempering. This influence on the

cementite dispersion has other effects, in so far as the carbide particles, by remaining finer,

slow down the reorganization of the dislocations inherited from the martensite, with the result

that the dislocation substructures refine more slowly. The cementite particles are also found

on ferrite grain boundaries, where they control the rate at which the ferrite grains grow.

The alloying additions are more effective at first stage of tempering in SAE 5140 whose

chromium concentration is slightly higher. Chromium and manganese replace some iron

atoms from the ε-carbide, thereby increasing its stability and retarding the formation of

cementite [27]. On the other hand, chromium and molybdenum, two strong carbide formers

co-exist in SAE 4140. These elements increase the stability of cementite by dissolving in it

and significantly delay the softening process during higher temperatures. The molybdenum

content of SAE 4140 is 4 times higher than SAE 5140. The coefficient of diffusion of

molybdenum is extremely low as compared with carbon and its retardation effect appears

most prominently after tempering at about 500oC. [32]. The hardness changes reflect the

tempering behavior as well. The softening rate of SAE 5140 is slower at first stages whereas

at higher tempering temperatures like 300oC and 400oC its softening rate becomes faster

than SAE 4140 as attributed to alloying additions.

42

3.2.2 MBN measurements

MBN Profiles

Under the effect of an alternating magnetic field, a representative magnetic hysteresis loop

was induced in the small volume measured due to the energy loss with the irreversible

process of magnetization. This irreversible process is strongly related to the dynamic

behavior of domains, i.e., nucleation, annihilation and growth of domains. Grain/lathe

boundaries, dislocations and precipitates affect this dynamic behavior. Consequently, the

number of domain walls moving at a given instant and the mean free path of the domain wall

displacement decide the MBN peak height.

In a simple model proposed by Saquet et al. [21] change of the local magnetic moment

causing MBN was given by:

).( lSmrrrr δβδ = (3.3)

where β is a coefficient related to the atomic magnetic moment and type of domain wall, S is

the surface of the moving domain wall and δl is the wall displacement between pinning

obstacles. Previously it was mentioned that the principle source mechanism of MBN is the

motion of 180o domain walls. Hence an elementary Barkhausen event δm appears to be

mainly concerned by δl and S. δl is linked to microstructure which provides the pinning

obstacles and S is determined by the magnetic microstructure morphology which is also

related to microstructure. In a quenched steel small martensite needles/lathes determine the

size of domain and hence S. In addition domain wall displacements (δl) are short, so the

expected Barkhausen activity in quenched specimens is low. As tempering temperature

increases, the microstructure coarsens, thus larger domain wall displacements are possible.

Also the surface of a moving domain wall (S) will increase as domains get larger due the

coarsening of microstructure [29]. As a result the MBN activity is expected to increase as

tempering temperature increases.

Figures 3.15 and 3.16 show the graph for MBN signal versus applied field strength (MBN

profiles) for the quenched and tempered SAE 4140 and SAE 5140 specimens. As MBN is

symmetrical with respect to zero magnetic field, only the curves for the increasing applied

magnetic field are plotted. The peak positions were obtained by fitting a parabola to the 15%

of the top of the MBN curve in the fingerprints. The maximum point of this parabola is the

peak position. The as-quenched sample has the weakest MBN peak positioned at the

highest field linked with the high coercivity of martensite. Moreover, the peak position of the

43

signal shifts to the lower values of magnetic field due to tempering. As the tempering

temperature increases, the low amplitude broad peak of as-quenched martensite transforms

into a high amplitude peak situated at low magnetic field. The results show that magnetic

Barkhausen noise is influenced by the tempering which, as a function of the temperature,

causes changes in dislocation density, lattice straining (i.e., micro residual stresses) and the

morphology and size of cementite, and corresponding variations in hardness. The results are

in agreement with those of the previous studies.

In the as-quenched specimens, the body-centered tetragonal structure of martensite

determines the domain structure. Since the magnetic structure consists of very small

domains due to small needles\lathes, relative volume occupied by a domain wall is larger

(Figure 3.17). Also the easy direction of magnetization is restricted to the c-axis of tetragonal

structure only. Besides, high dislocation density in the martensite laths acts as a barrier to

the movement of domain walls. A strong field is required for the reversal of magnetization

because of low domain wall displacements and difficulty in nucleating domain walls.

Presence of micro residual stresses in the martensite needles has an additional effect on

reduction of the MBN response.

Tempering at 200oC changes the microstructure very slightly. Although ε-carbides form, the

microstructure is still needle shaped. Therefore, the height and position of the MBN peak do

not change significantly. When tempering temperature reaches 300oC and 400oC, cementite

replaces ε-carbides, the crystal structure of martensite loses its tetragonality, and dislocation

density reduces further. Corresponding magnetization orientation is no longer favored and

reverse domain nucleation and subsequent domain wall motions take place at lower

magnetic fields. All these factors make the domain wall motion easier, and therefore, the

amplitude of the MBN peak increases.

In tempering at 500oC and 600oC, carbides start spheroidizing and recrystallization of ferrite

begins. In parallel to the progressive coarsening of the microstructure, the average size of

the domain walls increases (Figure 3.18) These morphological changes and almost

complete relaxation of residual stresses result in a drastic increase in the MBN peak and a

clear shift to lower external magnetic field in the peak position by reducing the resistance to

the nucleation and movement of domains.

44

0

5

10

15

20

25

30

35

40

45

50

-100 -80 -60 -40 -20 0 20 40 60 80 100

Magnetic Field Strength (% of max.)

Avg

. MB

N le

vel (

% o

f 2V)

As QT 200T 300T 400T 500T 600

Figure 3.15 MBN Profiles of SAE 4140 specimens

0

2

4

6

8

10

12

14

16

-100 -80 -60 -40 -20 0 20 40 60 80 100

Magnetic Field Strength (% of max.)

Avg

. MB

N le

vel (

% o

f 2V)

As - QT 200T 300T 400T 500T 600

Figure 3.16 MBN profiles of SAE 5140 specimens

45

Figure 3.17 The SEM micrographs of as quenched SAE 4140(a); SAE 5140(b) steels and schematic representation of magnetic microstructure [33] (c) in martensite.

(c) Schematic representation

(b) As-quenched SAE 5140

(a) As-quenched SAE 4140

46

Figure 3.18 The SEM micrographs of SAE 4140 (a); SAE 5140 (b) steels tempered at 600oC and schematic representation of magnetic microstructure [33] (c) in tempered martensite.

(c) Schematic representation

(b) SAE 5140 tempered at 600oC

(a) SAE 4140 tempered at 600oC

47

The effects of alloying elements on tempering, retardation of softening during early stages

can easily be seen from the MBN profiles of the SAE 5140 steel. The profiles of as

quenched, and 200-300oC tempered specimens are very close and a clear separation exists

between the profiles of 300oC and 400oC tempered specimens, which indicates that the

structure continues to resist domain wall motion up to 300oC tempering. The profiles of SAE

4140 steel do not show such a clear separation between 300oC and 400oC; also between

500oC and 600oC. Thus, the differences between MBN profiles of SAE 4140 and 5140 are

consistent with microstructure and hardness variations upon tempering.

Table 3.2 Hardness values and MBN parameters of SAE 4140 specimens

Specimen Hardness(HV)

RMS (mV)

Peak Height(% of 2V)

Peak Position (%of max. field)

As Quenched 556 1.781 3.541 37.25 T 200 504 3.702 8.671 35.87 T 300 488 6.289 14.900 29.50 T 400 467 8.850 21.260 30.20 T 500 299 15.050 36.310 20.72 T 600 206 16.680 40.370 14.15

Table 3.3 Hardness values and MBN parameters of SAE 5140 specimens

Specimen Hardness(HV)

RMS (mV)

Peak Height(% of 2V)

Peak Position (%of max. field)

As Quenched 648 1.165 1.705 34.98 T 200 604 1.173 1.9 34.42 T 300 564 1.437 2.605 34.55 T 400 500 3.654 8.01 31.62 T 500 411 4.828 10.61 25.83 T 600 319 6.009 13.31 24.63

Frequency spectrum The raw MBN data is obtained with respect to time and signals in time domain can be

transformed to the frequency domain by application of Fourier transform in order to

determine the spectral content of time domain signals [34]. The frequency spectrum shows

the intensity of the noise given by the square of voltage (V2) as a function of the frequency.

The frequency spectrum is calculated in the given analyzing frequency range (0,1 – 1000

kHz) using “Fast Fourier Transform” (FFT) algorithm. Figure 3.19 and 3.20 show the

frequency spectrums of SAE 4140 and SAE 5140 specimens.

48

0

0.1

0.2

0.3

0.4

0.5

0.6

0.1 1 10 100 1000

Frequency (kHz)

Ampl

itude

(a.u

.)

As QT 200T 300T 400T 500T 600

Figure 3.19 Frequency spectrums of SAE 4140 specimens

0

0,05

0,1

0,15

0,2

0,25

0,1 1 10 100 1000

Frequency (kHz)

Am

plitu

de (a

.u.)

As QT 200 T 300T 400T 500T 600

Figure 3.20 Frequency spectrums of SAE 5140 specimens

49

The intensity of low frequency MBN signals increases as the tempering temperature

increases (Figure 3.19 and 3.20). Since frequency is inversely proportional to time, low

frequency content of MBN signal indicates larger domain wall displacements if the average

wall velocity is constant.

During MBN measurements the specimens are magnetized locally by irreversible

movements of domain boundaries. The applied field in this case produces reorientation of

the electron spins only within the width of the boundary walls as these pass across the

domains. Again the spin axes can reorient themselves only by some mechanism operating at

a finite speed [35]. Considering this, the average wall velocity can be assumed constant and

frequency content is directly related to distance between domain pinning sites.

In the as quenched martensite domain wall displacements are short due to small needles,

which is consistent with the absence of low frequency MBN signals. The frequency content

does not change significantly for tempering up to 200oC due to the fact that the

microstructure is still fine and needle shaped. As tempering temperature increases further

the amplitude of low frequency MBN signals increases, indicating larger domain wall

displacements. Such large displacements are expected due to progressive coarsening of the

microstructure.

Alloying elements that retard softening also affect frequency spectrums like in the case of

MBN profiles. This effect is clearly seen in the spectrums of SAE 5140 steel which was

explained in the previous section.

Pulse Height Distributions Pulse height distribution is a rarely used parameter and can be used to understand the

nature of MBN. This distribution, which is the number of events (pulses) against pulse

amplitude, depends on the number density and nature of pinning sites within the material [4].

The number of pulses detected in the MBN signal depends on the number of pinning sites

[36]. The pulse height distributions of the as-quenched and 600oC tempered specimens, the

most distinct cases, are given in Figure 3.21 and 3.22. As quenched specimens have about

180-200 thousand pulses detected whereas tempered ones have about 120000. In the as

quenched specimen each martensite needle/plate will act as a pinning site so the number of

pulses detected is very high. The quenched specimen has a pulse height up to only 0.4 V

indicating very small domain wall displacements.

50

020000400006000080000

100000120000140000160000180000200000

0 0,5 1 1,5

Pulse Height (V)

Num

ber o

f Pul

ses

As-QT 600

Figure 3.21 Pulse height distributions of SAE 4140 specimens

020000400006000080000

100000120000140000160000180000200000

0 0.5 1 1.5

Pulse Height (V)

Num

ber o

f Pul

ses

As QT 600

Figure 3.22 Pulse height distributions of SAE 5140 specimens

51

For the specimen tempered at 600oC the number of pulses decreases whereas height of the

pulses increases. The decrease in number of pulses is a consequence of softening of

microstructure which decreases the number of pinning sites respectively. Reduced dislocation

density and recrystallized ferrite enhance domain wall movement, thus allowing the domain walls to move longer distances or give larger jumps. These changes upon tempering are

reflected as higher amplitude pulses with heights more than 1.5 V, which are not present in the

as quenched specimen.

Hysteresis Loop Parameters

During MBN measurements a representative magnetic hysteresis loop is induced in the

small volume of specimen due to the energy lost in the form of a kind of internal friction. The noise signals obtained are proportional to rate of change in internal magnetization. Thus the

integration of noise signals along the whole applied magnetic field gives a representative

hysteresis loop from which parameters such as coercivity, remanence and permeability

could be obtained. These parameters may be used then in order to characterize samples.

Figure 3.23 (a) and 3.24 (a) show the coercivity values of SAE 4140 and 5140 steels. The

differences between quenched and tempered samples are not significant and the differences

between these values are small. It would be reasonable to assume that as quenched

samples have high coercivities due to the pinning of domain walls in the presence of small

martensite laths, needles and high dislocation density [37]. However during MBN measurements the actual level of magnetization reached locally is too small when compared

with the saturation magnetization of specimens. Also all specimens experience the same

applied field strength which makes it very difficult to distinguish coercivities. More reliable

results could be obtained from these specimens by obtaining true hysteresis curves form

magnetic pendulums and magnetometers.

Figure 3.23 (b) and 3.24 (b) show the remanence values of SAE 4140 and 5140 steels. The

remanences increase with increasing tempering temperature for both types of steels. The

quenched structures contain large amounts of crystal defects and it is possible that the elimination of part of these defects promoting an atomic rearrangement during the tempering

leads to an increase in the remanence of the steel [38].

Figure 3.23 (c) and 3.24 (c) show the permeability values of SAE 4140 and 5140 steels. For

both types of steels the permeabilities increase with increasing tempering temperature. As

tempering temperature increases magnetic softening of structure occurs parallel to

mechanical softening of the microstructure. The less efficient pinning of domain walls causes

magnetic softening, which also results in an easier domain wall motion [39]. It is expected

that enhanced domain wall motion increases the permeability of the tempered specimens.

52

As QT 200

T 300 T 400

T 500T 600

-

0.050

0.100

0.150

0.200

0.250

0.300

0.350

0.400

Frac

tion

of M

ax. A

pplie

d Fi

eld

Stre

ngth

(a)

As Q

T 200

T 300

T 400

T 500 T 600

0

10

20

30

40

50

60

Perc

ent o

f 2V

(b)

As Q

T 200

T 300

T 400

T 500 T 600

0

50

100

150

200

250

300

Slop

e at

Coe

rciv

e Po

int

(c)

Figure 3.23 Coercivity (a), remanence (b), permeability (c) values of SAE 4140

53

As QT 200

T 300

T 400T 500 T 600

0

0.05

0.1

0.15

0.2

0.25

0.3

0.35

0.4

Frac

tion

of M

ax. A

pplie

d Fi

eld

Stre

ngth

(a)

As Q T 200T 300

T 400

T 500

T 600

0

5

10

15

20

25

Perc

ent o

f 2V

(b)

As Q T 200 T 300

T 400

T 500

T 600

0

20

40

60

80

100

120

Slop

e at

Coe

rciv

e Po

int

(c)

Figure 3.24 Coercivity (a), remanence (b), permeability (c) values of SAE 5140

54

Hardness – MBN Correlation Generally a linear relation is found between mechanical hardness and the peak position of

the MBN. The linear relation between magnetic coercive force and hardness explains such

behavior. The microstructural features that impedes dislocation movement, also makes

domain wall movement harder [5]. Thus harder the material higher the position and lower the

height of the peak would be. Residual stresses have an additional effect on both hardness

and MBN in a similar manner. MBN is sensitive to both hardness and stresses which also

influence hardness. Regarding this a good correlation between MBN and hardness is

expected. Generally correlation between the peak height of MBN profile and hardness was

studied [40]. Another parameter that reflects the MBN behavior and is influenced in the same

way as MBN peak height upon tempering, is the root mean square of the Barkhausen signal

(RMS) as seen on Table 3.2 and 3.3.

The raw magnetic noise data consists of a series of voltage pulses and associated magnetic

field values. RMS of all signal amplitudes were sampled for the specified analyzing

frequency range according to the formula:

∑−

=

=1

0

21 n

iixn

RMS (3.4)

It is seen in Table 3.2 and 3.3 that the as-quenched specimens (the hardest ones) have the

lowest RMS values. As tempering temperature increases, in contrast to the decrease in

hardness, RMS value increases. Pinned domain walls due to high dislocation density and

small martensite needles cause lower RMS values. As tempering temperature increases

dislocation density decreases, micro residual stresses diminish and the magnetic structure

comes close to those of a ferrite. Thus, RMS value increases due to the enhancement of

domain wall displacement with softening of martensite. Figure 3.25 and 3.26 show the

correlation graph between the RMS values and peak heights of the MBN signal with the

hardness of specimens. The regression analysis shows an excellent correlation between the

RMS values and peak heights with hardness. The correlation between RMS and hardness is

better than the correlation with peak height. Thus RMS values can also be used for hardness

correlations.

55

Peak Height = -0.1057(HV) + 65.232R2 = 0.9396

RMS = -0.043(HV) + 26.796R2 = 0.9406

05

101520253035404550

0 100 200 300 400 500 600 700Hardness (HV)

Peak

Hei

ght o

r RM

S

Peak HeightRMS

Figure 3.25 Hardness correlations for SAE 4140 specimens

Peak Height = -0.0393(HV) + 26.29R2 = 0.9534

RMS = -0.0165(HV) + 11.408R2 = 0.9535

0

2

4

6

8

10

12

14

16

0 100 200 300 400 500 600 700Hardness (HV)

Peak

Hei

ght o

r RM

S

Peak HeightRMS

Figure 3.26 Hardness correlations for SAE 5140 specimens

56

3.2.3 Detection of Faulty Quenching and Tempering Treatment by MBN

When cooling rate shows significant variations as going from surface to interior, the phase

content and the residual stress state along the thickness of the specimen may differ. In such

cases, the hardness and the microstructure of the surface may not represent the whole

structure. Microstructural investigations showed that the thickness of the samples of SAE

4140 and SAE 5140 used in this study allowed the formation of desired microstructure

uniformly along a penetration depth of the MBN activity, which usually varies between 0.01

and 1.5 mm depending on the analyzing frequency. To demonstrate the effect of low

hardenability on microstructure and MBN, SAE 1040 steel was used. This steel type

practically has very low hardenability when compared with SAE 4140 and SAE 5140 as can

be seen from its continuous cooling transformation given in Figure 3.27.

Figure 3.27 Continuous cooling transformation diagram of SAE 1040 steel [31]

57

Figure 3.28, 3.29, 3.30, 3.31 show the SEM micrographs, MBN profiles and corresponding

frequency spectrums of quenched and tempered SAE 1040 specimens. In all specimens at

the edge regions where cooling rate is higher, fully martensitic or tempered martensitic

structures are observed. However, when going from edge to interior regions cooling rate is

lowered. The effect of low cooling rate combined with low hardenability is best observed at

centre portions of specimens. At the centre sections of all SAE 1040 specimens proeutectoid

ferrite is observed which influences the MBN signals as well.

In as-quenched specimen high dislocation density, small martensite needles/plates and

presence of micro residual stresses cause low MBN activity. As tempering temperature

increases MBN activity increases due to the softening of martensitic microstructure. The

edge sections of specimens where fully martensitic structure is present, the MBN activity is

as expected. On the other hand MBN activities of centre sections of specimens are

unexpectedly high. This high activity can be attributed to the presence of ferrite which has

practically no resistance to domain wall motion. Enhanced domain wall motion in ferrite

increases the MBN activity which can be directly observed from MBN profiles and frequency

spectrums. In the presence of ferrite higher amplitude profiles and spectrums are observed

in all specimens.

The difference between edge and centre portions is very significant for the as quenched

specimen whereas for specimens tempered at high temperatures, 500oC and 600oC, MBN

activities of fully martensitic and ferrite containing regions come closer. Such a behavior is

expected since at high tempering temperatures the magnetic microstructure developed is

very close to that of ferrite. It should also be mentioned that the volume fractions of

martensite and ferrite may also influence the MBN activity.

Figure 3.33 shows the MBN profiles of edge regions of SAE 1040 where fully martensitic

structures are present. The amplitudes of tempered specimens are far beyond the as

quenched specimen. This indicates that SAE 1040 shows practically no resistance to temper

softening. Since it contains no alloying elements upon tempering SAE 1040 softens rapidly,

where as in case of low alloy steels of SAE 4140 and especially 5140 the softening is

retarded.

58

As Quenched

05

1015202530354045

-100 -80 -60 -40 -20 0 20 40 60 80 100

Magnetic Field Strength (% of max.)

Avg

. MB

N le

vel (

% o

f 2V)

EdgeCentre

0

0.1

0.2

0.3

0.4

0.5

0.6

1 10 100 1000

Frequency (kHz)

Am

plitu

de (a

.u.)

EdgeCentre

Figure 3.28 SEM micrographs taken from edge(a) and centre(b); MBN profiles(c); frequency spectrums(d) of as-quenched SAE 1040 specimen.

(b) Centre (a) Edge

(d)

(c)

59

Tempered at 300oC

05

1015202530354045

-100 -80 -60 -40 -20 0 20 40 60 80 100

Magnetic Field Strength (% of max.)

Avg

. MB

N le

vel (

% o

f 2V) Edge

Centre

0

0.1

0.2

0.3

0.4

0.5

0.6

1 10 100 1000

Frequency (kHz)

Am

plitu

de (a

.u.)

EdgeCentre

Figure 3.29 SEM micrographs taken from edge(a) and centre(b); MBN profiles(c); frequency spectrums(d) of SAE 1040 specimen tempered at 300oC

(b) Centre

(d)

(a) Edge

(c)

60

Tempered at 500oC

05

1015202530354045

-100 -80 -60 -40 -20 0 20 40 60 80 100

Magnetic Field Strength (% of max.)

Avg

. MB

N le

vel (

% o

f 2V) Edge

Centre

0

0.1

0.2

0.3

0.4

0.5

0.6

1 10 100 1000

Frequency (kHz)

Am

plitu

de (a

.u.)

EdgeCentre

Figure 3.30 SEM micrographs taken from edge(a) and centre(b); MBN profiles(c); frequency spectrums(d) of SAE 1040 specimen tempered at 500oC

(b) Centre

(d)

(a) Edge

(c)

61

Tempered at 600oC

05

1015202530354045

-100 -80 -60 -40 -20 0 20 40 60 80 100

Magnetic Field Strength (% of max.)

Avg

. MB

N le

vel (

% o

f 2V) Edge

Centre

0

0.1

0.2

0.3

0.4

0.5

0.6

1 10 100 1000

Frequency (kHz)

Am

plitu

de (a

.u.)

EdgeCentre

Figure 3.31 SEM micrographs taken from edge(a) and centre(b); MBN profiles(c); frequency spectrums(d) of SAE 1040 specimen tempered at 600oC

(b) Centre

(d)

(a) Edge

(c)

62

0

5

10

15

20

25

30

35

-100 -80 -60 -40 -20 0 20 40 60 80 100Magnetic Field Strength (% of max.)

Avg

. MB

N le

vel (

% o

f 2V)

As - QT 300T 400T 500T 600

Figure 3.33 MBN profiles of edge regions of SAE 1040 specimens

Edge region

Centre region

Figure 3.32 Schematic drawing of edge and centre regions of SAE 1040 specimens

63

CHAPTER 4

CONCLUSIONS For the purpose of characterizing steel microstructures non-destructively the steel

specimens were heat treated using two common procedures. In the first part of this study the

effect of average ferrite grain size was investigated in the annealed low carbon steels. The

influence of tempering induced changes on MBN was studied in the second part.

Various average ferrite grain sizes were obtained by annealing the SAE 1010 specimens at

different time and temperature variations. The effects of pearlite and hardness on MBN were

neglected due to the fact that ferrite volume fraction in all specimens was about 90% and the

hardness differences were so small. Effect of residual stresses was eliminated by cooling the

specimens very slowly in the furnace. Thus only the differences in average ferrite grain size

influenced the MBN response of the specimens.

As grains become coarser, domain size increases and domain density decreases.

Consequently MBN peak heights and RMS values decrease. The peak positions of MBN

signals indicate that the MBN activity in the annealed specimens occurs during the early

stages of magnetization. At this stage domain nucleation is the predominant mechanism and

grain boundaries are the preferential sites for nucleation. As grain size increases, grain

boundary area per unit volume decreases, causing difficulty in nucleation of new domains.

An increase in average ferrite grain size cause reduced Barkhausen jumps and difficulty in

creating new domains; which decrease MBN activity. However grain boundary segregation in

some specimens alters the magnetic structure and increase the MBN activity. The results

obtained are consistent with the previous studies and theoretical expectations. It could be

concluded that MBN is sensitive to grain size differences.

Martensitic microstructures were obtained by quenching and tempering the low alloy steels

at different temperatures. MBN method is a powerful tool for evaluating different stages of

tempering. In the as quenched samples, pinned domain walls due to high dislocation density

and small martensite needles cause low MBN activity, and MBN peak is at higher magnetic

64

fields due to small domain wall displacements and difficulty in domain nucleation. As

tempering temperature increases dislocation density decreases, micro residual stresses

diminish and the magnetic structure comes close to those of a ferrite. Thus, MBN activity

gets higher due to the enhancement of domain wall displacement with softening of

martensite. Change in the number density of pinning sites and relative domain wall

displacements upon tempering influences the frequency spectrums, pulse height

distributions as well. RMS values can also be used instead of peak height, considering the

better correlation between RMS values and hardness, for hardness correlation. Via

establishing the quantitative relationships between MBN parameters and the microstructural

parameters, this method can be utilized efficiently and effectively for evaluating the hardness

and the microstructure of the steel components.

The alloying additions influence tempering behavior of steels. Alloying elements retard

softening of martensite structure either by stabilizing the carbides or by decreasing the rate

of carbon diffusion. MBN is sensitive to this softening retardation. In the absence of alloying

elements steels soften more rapidly upon tempering and hardenability decreases. Lower

hardenability causes pro-eutectic phases to form in some portions of specimens. Steel with

lower hardenability develop such structures and MBN can be used to detect such

treatments.

The results show that the MBN parameters are sensitive to tempering induced changes,

faulty treatments and grain size differences. The sensitivity of this phenomenon to

microstructural changes gives a wide range of potential applications to the technique.

For further investigations, ferromagnetic materials having nano-sized grains which should

have single domain magnetic structures may be a good candidate for MBN studies. Besides

components produced by powder metallurgy techniques can be investigated by MBN. As a

final suggestion domains and domain wall motion can be observed using Kerr effect and

transmission electron microscopy (TEM) in order to understand the domain mechanisms like

domain nucleation, domain annihilation and irreversible domain wall motion.

65

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