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ADVANCED REAL TIME OPTICAL IMAGING

High-Temperature Confocal Laser Scanning MicroscopyStudies of Ferrite Formation in Inclusion-Engineered Steels:A Review

WANGZHONG MU ,1,3 PETER HEDSTROM,1 HIROYUKI SHIBATA,2

PAR G. JONSSON,1 and KEIJI NAKAJIMA1

1.—Department of Materials Science and Engineering, KTH Royal Institute of Technology,Brinellvagen 23, 100 44 Stockholm, Sweden. 2.—Institute of Multidisciplinary Research forAdvanced Materials, Tohoku University, 2-1-1 Katahira, Aoba-ku, Sendai, Miyagi 980-8577,Japan. 3.—e-mail: [email protected]

The concepts of oxide metallurgy and inclusion engineering can be utilized toimprove the properties of low-alloy steels. These concepts aim at controlling theformation of intragranular ferrite (IGF), often a desirable microstructure pro-viding good mechanical properties without the need for expensive alloying ele-ments. IGF formation is stimulated to occur at non-metallic inclusions and forman arrangement of fine, interlocking ferrite grains. A method that has con-tributed significantly to investigations in this field lately is high-temperatureconfocal laser scanning microscopy (HT-CLSM). HT-CLSM is suited for in situstudies of inclusion behavior in liquid steel and phase transformations in solid-state steel, where in particular, displacive phase transformations can be stud-ied, since they provide sufficient topographic contrast. The purpose of the pre-sent report is to provide a brief review of the state of the art of HT-CLSM and itsapplication for in situ observations of ferrite formation in inclusion-engineeredsteels. The scientific literature in this field is surveyed and supplemented by newwork to reveal the capability of HT-CLSM as well as to discuss the effect offactors such as cooling rate and parent grain size on IGF formation and growthkinetics. The report concludes with an outlook on the opportunities and chal-lenges of HT-CLSM for applications in oxide metallurgy.

INTRODUCTION

Clean steels are in increasing demand to optimizemechanical properties such as fatigue properties,and the vision of process metallurgists is to producesteels without inclusions. However, there is limitedknowledge regarding how to remove and modifyinclusions during secondary refining in steelmak-ing, and formation of inclusions is inevitable usingcurrent industrial steelmaking practice.1–4 Accept-ing this, researchers have also tried to benefit frominclusions; For example, it has been shown thatcertain non-metallic inclusions, such as Ti oxide,TiN, and V(C,N), can serve as effective nucleationsites for formation of intragranular ferrite (IGF).The preferential formation of IGF, instead of grainboundary ferrite (GBF), improves the toughness ofsteel. This concept of utilizing inclusions within the

physical metallurgy discipline is known as oxidemetallurgy, introduced by Takamura and cowork-ers5–8 in 1990. Their research described utilizationof fine oxide inclusions to improve the quality of thefinal steel product. Another related concept isknown as inclusion engineering, emphasizing con-trol of the amount, morphology, size distribution,and chemical composition of the inclusions formedin liquid steel during the refining process.9 In thepresent report, we use the concept of oxide metal-lurgy, since the emphasis here is on the microstruc-tural aspects, as opposed to the inclusioncharacteristics. Oxide metallurgy is a relativelynew concept,10 but utilization of IGF microstruc-tures can be dated back to six decades ago.11 Hence,oxide metallurgy mainly stresses the correlationbetween inclusions and microstructure and requiresconsideration of both process and physical

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metallurgy. Oxide metallurgy has so far primarilybeen used to improve the toughness and strength ofthe coarse-grained heat-affected zone (CGHAZ) inweldments.12 Recently, reports on how to improvethe properties of various steel grades, includinghigh-strength low-alloy (HSLA) steels, line pipesteels, and off-shore steels,13,14 have also beenpresented.

One methodology that has become valuable forinvestigation of austenite decomposition in oxidemetallurgy is high-temperature confocal laser scan-ning microscopy (HT-CLSM). It provides a robust toolfor in situ real-time studies of phase transformationsduring heat treatment. The pioneering work onapplication of HT-CLSM dates back to the work ofEmi and coworkers.15–19 In addition to investigationson austenite decomposition in steels, HT-CLSM hasalso been applied to study, for example, crystalgrowth during solidification of a steel melt,15 inclu-sion agglomeration at the liquid steel-Ar gas inter-face,17,18,20,21 and engulfment and pushing ofinclusions at the melt–solid interface of steel.19

The aim of the present report is to survey theapplication of the HT-CLSM technique for investi-gation of IGF formation in steels utilizing oxidemetallurgy. This literature survey is supplementedby new HT-CLSM work to investigate the austenitedecomposition in a steel with TiO2 addition. Thereport starts with an introduction to HT-CLSM andthe state of the art of HT-CLSM instrumentation.Thereafter, we provide a literature review of HT-CLSM studies on austenite decomposition withinoxide metallurgy. Subsequently, we present a morein-depth summary of selected insights gained fromHT-CLSM that have contributed to understandingof oxide metallurgy in steels. We conclude thisreport with a discussion about the opportunities andchallenges in applying HT-CLSM for studyingaustenite decomposition in steels and relate this toopportunities in neighboring fields.

HIGH-TEMPERATURE CONFOCAL LASERSCANNING MICROSCOPY

Working Principle and Instrumentation

The design of a modern HT-CLSM microscope isdetailed in Fig. 1. The sample is positioned in acrucible, which is placed inside the furnace chamberon a Pt sample holder. A thermocouple wire isincorporated into the Pt sample holder to measurethe temperature from the bottom of the crucible. Ahalogen infrared (IR) heating lamp is situatedunderneath the sample holder to heat the specimen,which is placed at the focal point of the IR beam thatis reflected by the gold coating covering the inside ofthe ellipsoidal chamber. In the first HT-CLSM studyreported by Chikama et al.,15 a He-Ne laser beamwith power of 1.5 mW and wavelength of 632.8 nm15

was scanned across the specimen surface using anacoustic optic deflector (AOD) in the horizontaldirection and a Galvano mirror in the vertical

direction. For the latest version of HT-CLSM,VL2000DX-SVF17SP, manufactured by LasertecCorporation22 and Yonekura Manufacturing Corpo-ration,23 a wavelength of 405 nm is used. As shownin Fig. 1, the scanning beam is focused on thesample surface via a polarizing plate, a long-focusobjective lens, and a viewport covered with a quartzwindow. The reflection from the sample surface ispassed through a polarizing beam splitter to a lens,and subsequently focused onto a charge-coupleddevice (CCD) camera positioned behind a pinhole.HT-CLSM can scan the surface at different focaldistances and subsequently reconstruct an image ofthe sample surface in a three-dimensional plot withsmall Z direction depth. Using the simpler approachwith focusing at only one focal distance, the imageacquisitions are much faster, and this is usually theoperation mode applied for in situ studies at hightemperatures. The excellent depth (Z) resolution ofabout 0.01 lm provided by the pinhole enables thedistinction between flat areas and even shallowgrooves or dimples on the sample surface.22,23 Thelateral resolution of CLSM is limited by the wave-length of the laser and is about 0.15 lm. This meansthat the maximum useful magnification is on theorder of 1000 times. For the latest version HT-CLSM, VL2000DX-SVF17SP, the maximum magni-fication is 30009.22,23 For further information onthe working principles and instrumentation of HT-CLSM, the reader is referred to Refs. 15 and 22–25.

DETAILS OF OPERATION

The maximum temperatures and heating ratesdepend on the capability of the halogen infrared (IR)heating lamp. According to the manufacturer,22,23

the maximum working temperature can be set to1800�C, the maximum heating rate can approach1200�C/min, and rapid cooling at an average rate of3000�C/min at the higher temperatures can beobtained by forced cooling using He gas. To studysurface phenomena in metallic samples, it is neces-sary to use a sample surface polished to mirrorfinish. Frequently used sample sizes are 4 mm and8 mm in diameter to match the standard cruciblesize.22,23

The temperature of the furnace is controlledbased on the primary thermocouple (TC) located atthe bottom of the crucible. Typical temperaturedifferences between the sample surface and thebottom of the crucible are about 20�C to 40�C,depending on the sample material, sample thick-ness, and heating or cooling rate. The specificationgiven here was evaluated for a low-carbon steelsample with 1 mm thickness, cooled at a ratebetween 3.6�C/min and 678�C/min.26 Because ofthis difference in temperature, it may be importantto measure the temperature of the sample surfaceby welding a second TC on the sample surface.Another approach is to calibrate the temperature byperforming melting experiments of metal reference

Mu, Hedstrom, Shibata, Jonsson, and Nakajima2284

samples with known melting temperatures; Forexample, lead, aluminum, copper, nickel, and ironcan be melted in separate experiments under thesame heating and cooling conditions.

It is important to prevent oxidation of the samplesurface, since this would obscure the observations ofthe phenomenon of interest. Therefore, a vacuumsystem with rotary and/or diffusion pumps isinstalled. Purging of inert Ar gas with high purityis also available to decrease the oxygen content ofthe furnace atmosphere. Materials that are verysensitive to oxidation such as pure silicon27 and Alalloys28,29 require ultralow oxygen partial pressure(PO2), whereas steel samples are less sensitive andnon-metallic oxide samples are not sensitive to PO2.A mixed gas atmosphere with, e.g., 5% H2 andultrapure Ar gas, can be used to control PO2 to about10�15 Pa to 10�20 Pa. However, high concentrationsof H2 may damage the gold coating of the furnacechamber and the TC wire, thus H2 is only usedwhen necessary. The experimental conditions aredifferent in each study, hence the specific conditionsshould be referred to when discussing each work ifrelevant. For further information on the detailsof operating conditions, the reader is referred toRefs. 22 and 23.

HT-CLSM STUDIES WITHIN OXIDEMETALLURGY

Overview of In Situ Studies in OxideMetallurgy by HT-CLSM

In situ HT-CLSM observations of ferrite forma-tion in steel date back to the work by Hanamuraet al.30 in 1999. In their work, the starting temper-ature and morphology of polygonal ferrite (PF) andWidmanstatten ferrite (WF) nucleated from oxidesand sulfides as well as precipitation were success-fully measured. They reported that PF started toprecipitate at oxide particles at 773�C and that WFstarted to precipitate at MnS particles at 730�C. Theundercooling responsible for the difference in mor-phology between the ferrite formed at an oxide andat MnS was determined to be 43�C.30 Phelan andcoworkers investigated formation of WF in a low-carbon steel grade.31,32 WF formed from the prioraustenite grain boundary, but IGF formation frominclusion surfaces was not reported. Thereafter,Kikuchi et al.33–36 applied HT-CLSM to investigatethe postsolidification microstructure refinement inlow-carbon high-manganese steels through Ti deox-idation. Several issues, including austenite graingrowth and decomposition, solidification structure,

Fig. 1. Schematic illustration of high-temperature confocal laser scanning confocal microscopy (HT-CLSM).

High-Temperature Confocal Laser Scanning Microscopy Studies of Ferrite Formationin Inclusion-Engineered Steels: A Review

2285

Table

I.Summaryofin

situ

HT-C

LSM

studiesofferriteformation

Auth

or

[References]

Year

Inclu

sion

Ferrite

morphology

Quantita

tiveasp

ects

offer-

riteformation

Ferrite

formation

kin

etics

Specialasp

ects

Composition

Sizeand

no.

Fraction

Formation

temp.

Grain

size

effect

Han

am

ura

etal.

30

1999

Oxid

e(T

i,S

i),

Mn

SN

oY

esN

oN

oN

oY

esF

irst

wor

ku

sin

gH

T-C

LS

Mon

IGF

;co

olin

gra

tean

dn

ucl

eico

mp

osit

ion

sare

inves

tigate

dP

hel

an

etal.

31,3

22004–2005

No

No

Yes

No

Yes

No

Yes

Wid

man

statt

enfe

rrit

efo

rmati

onin

low

-carb

onst

eel

Kik

uch

iet

al.

33–36

2007–2011

Ti-

oxid

e,M

nO

-TiO

2,

Ti-

Al-

(Mg)-

O,

TiN

Yes

Yes

No

No

Yes

No

Sol

idif

icati

onan

dp

osts

olid

ifica

tion

mic

rost

ruct

ure

Ter

asa

ki,

Kom

izo

an

dco

-wor

ker

s37–44

2011–2013

Ti-X

(X=

O,

N),

Al-

Mn

-O,

BN

No

Yes

Yes

Yes

Yes

Yes

Fer

rite

nu

clea

tion

inw

eld

men

tan

dh

eat-

aff

ecte

dzo

ne;

B-t

reate

dst

eel;

com

bin

ati

onof

HT

-CL

SM

wit

hH

R-X

RD

Jia

ng

etal.

45,4

62011–2015

Ti-

Mn

-O,

Ti-

Al-

O,

Al-

Si-

Mn

-ON

oY

esY

esY

esY

esN

oE

ffec

tof

rati

oof

Ti

an

dA

lon

IGF

chara

cter

isti

csW

en&

Son

g47

2011-2

012

CeX

(X=

O,

S),

MgA

l 2O

4,

MgX

No

Yes

No

No

Yes

No

Fir

stw

ork

focu

sin

gon

IGF

form

ati

onfr

omra

re-e

art

hin

clu

sion

Wan

,W

uan

dco

wor

ker

s.48–50,5

82013–2017

TiN

No

Yes

No

No

Yes

Yes

Fer

rite

gro

wth

kin

etic

s;gra

inre

fin

emen

tk

inet

ics

Mu

etal.

26,5

1–53

2014–2016

Ti-

oxid

e,T

iNY

esY

esY

esY

esY

esY

esQ

uan

tita

tive

an

aly

ses

ofin

clu

sion

an

dfe

rrit

ech

ara

cter

isti

csW

uet

al.

54

2015

Al-

Ti-

Mg-M

n-O

No

Yes

No

No

Yes

No

Foc

us

onA

l-T

i-M

g-M

n-O

incl

usi

onL

oder

&M

ich

elic

55–57

2015–2017

(Ti,

Mn

) xO

yan

d(T

i,A

l,M

n) x

OySz,

Mn

S,

MgO

incl

usi

ons

Yes

Yes

Yes

No

Yes

No

Eff

ects

of(T

i,A

l)O

x-,

Mn

S-,

an

dM

gO

-con

tain

ing

incl

u-

sion

s,m

ech

an

ical

test


and secondary deoxidation particles, were studied.Terasaki, Komizo, and coworkers37–44 were the firstto provide comprehensive quantitative analysis ofthe effect of inclusions on factors such as ferritefraction, phase transition temperature, and grainsize. Thereafter, Jiang et al.45,46 and Wen et al.47

followed this methodology and performed in situobservations of ferrite formation kinetics in steelscontaining Ti-Mn-O and Al-Si-Mn-O complex oxidesas well as rare-earth inclusions. Wan, Wu, andcoworkers48–50 observed grain refinement in HSLAsteels containing TiN during heat treatments. Theirheating and cooling profiles were chosen to simulatethe coarse-grained heat-affected zone (CGHAZ) ofweldments. Mu et al.26,51–53 performed anothercomprehensive in situ study on ferrite formationkinetics. The effects of the cooling rate, grain size,and inclusion composition on the ferrite fraction andphase-transition temperature were quantitativelyinvestigated in inclusion-engineered steels contain-ing TiN and Ti-oxide additions. The steel grades hadcarbon content varying between 0.2 wt.% and0.3 wt.%. Note that the steel matrix and nucleiparticle compositions were quite different from theresults of previous studies.37–44 Moreover, a theo-retical model based on classical nucleation theorywas established to quantify the different potenciesof nucleation sites.52,53 Subsequently, Wu et al.54

and Loder et al.55–57 performed in situ studies in

steels after treatment with complex deoxidizers.Several complex oxides such as Al-Ti-Mg-Mn-O,(Ti,Mn)xOy, and (Ti,Al,Mn)xOySz were found to beeffective nucleation sites. In their studies, both theinclusion and steel matrix characteristics wereinvestigated. A summary of the available in situstudies of ferrite formation till 2017 is presented inTable I.

In literature, HT-CLSM has mainly been used tostudy the following parameters in situ during heattreatments: (i) ferrite fraction evolution, (ii) onsetand completion of phase transformations, i.e., startand finish temperatures or times, (iii) grain coars-ening, and (iv) IGF growth kinetics, such as two-dimensional (2D) lengthening and widening of, e.g.,ferrite. Examples of typical HT-CLSM images ofGBF and IGF formation in a steel are given inFig. 2. In this case, it is a steel with TiO2 additionswhich was heated to 1400�C then directly cooled atrate of 70�C/min. The next section uses the new datafrom this steel in combination with results fromliterature to discuss how HT-CLSM has contributedand can further contribute to understanding offerrite formation within oxide metallurgy and alsomore generally to austenite decomposition in steels.More specifically, we discuss evolution of IGF,starting temperature of ferrite formation, austenitegrain growth, and ferrite growth kinetics in contin-uous cooling experiments.

Fig. 2. Typical HT-CLSM images of IGF formation and growth in a steel with TiO2 addition. (a) GBF transformation starts at 741�C, (b) IGFtransformation starts at 676�C, (c) GBF and IGF growth at 666�C, and (d) transformations of GBF and IGF finish at 641�C.


2287

Starting and Finishing Temperaturesof Phase Transformations

The starting and finishing temperature of phasetransformations can be directly observed by in situHT-CLSM. Examples of observations on the startingtemperatures of GBF and IGF affected by the prioraustenite grain size (PAGS) can be seen in Fig. 3.There are a number of studies in literature on thestarting temperatures of IGF (TGBF,s) and GBF(TIGF,s). Some of these results from literature arepresented together with the new data for alloy A andB53 and KT4 and 538 in Fig. 3. In Fig. 3a to c, thegrain size was controlled by heating the samples todifferent maximum temperatures of 1200�C, 1250�C,1300�C, and 1400�C without holding, and by furtherholding the samples at 1400�C from 0 min to10 min.53 In Fig. 3d, the PAGS was increased byincreasing the holding time from 0 min to 5 min at1400�C.38 The alloys were subsequently cooled atdifferent rates. The overall tendency is for TIGF,s tobe independent of PAGS, whereas TGBF,s is about20�C to 50�C lower for coarse compared with finegrains. It is also seen that IGF forms at a temper-ature about 20�C to 50�C lower compared with GBF.Moreover, both TIGF,s and TGBF,s decrease withincreasing cooling rate. As discussed in ‘‘Overview

of In Situ Studies in Oxide Metallurgy by HT-CLSM’’section, the ferrite formation behavior for the steelswith Ti2O3 and TiO2 additions is similar. The reasonis that TiOx inclusions are formed in both steelgrades, x ranges from 1.5 to 1.8, and these act aspreferential nucleation sites.51 Thus, the TIGF,s ofthese steels are quite close to each other.

Since HT-CLSM can only be used to observe thephase-transition temperatureonthe samplesurface, ahybrid methodology combining HT-CLSM and differ-ential scanning calorimetry (DSC) has been estab-lished to study all the austenite decompositionproducts, i.e., IGF, GBF, and pearlite.26 This can becombined with investigations of the sample surfaceand bulk microstructures using, e.g., electronbackscatter diffraction (EBSD). An example of theapplication of this methodology is presented here forthe present alloys. DSC results using different coolingrates are shown in Fig. 4a and b. In Fig. 4, the firstinflection point is identified as the starting tempera-ture of ferrite formation, named TF,s; subsequently,the second and third inflection points are identified asthe starting and finishing temperatures of pearliteformation, named TP,s and TP,f. Note that DSC canonly detectTF,s, but by combining it with HT-CLSM, itwas possible to also identify TGBF,s and TIGF,s. The

0 100 200 300 400 500 600 700 800 900 1000

800

550

600

650

700

750

(°C/min)Aver.678 °C/min

3.6 C/min

GBFIGF

Cooling rate

67870 3.6

70 C/min

Tem

pera

ture

/ C

Prior austenite grain size / μm

Steel with TiO2 addition(a) current sample

0 100 200 300 400 500 600 700 800 900 1000550

600

650

700

750

800

Aver.678 °C/min

(b) Mu-Alloy A [53]

70 °C/min

Tem

pera

ture

of s

peci

men

/ C


(°C/min)Cooling rate

GBFIGF

678 70 3.6

3.6 °C/min

0 100 200 300 400 500 600 700 800 900 1000550

600

650

700

750

800

(°C/min)

Aver. 678 °C/min

Cooling rate

3.6 °C/min

GBFIGF

678 70 3.6

70 °C/min

Tem

pera

ture

/ °C


(c) Mu-Alloy B [53]

50 100 150 200 250 300 350 400 450550

600

650

700

750

800

KT4-GBF KT4-IGF KT5-GBF KT5-IGF

Tem

pera

ture

/ °C


(d) Zhang et al. [38]Cooling rate: 300 °C/min

°

°°

°

°

Fig. 3. Effect of prior austenite grain size and cooling rate on starting temperature of IGF and GBF formation in (a) the current steel with TiO2

addition, (b) alloy A and (c) alloy B taken from Ref. 53, and (d) data taken from Ref. 38.


final continuous cooling transformation (CCT) dia-gram is shown in Fig. 5. It can be seen from the EBSDmicrographs that the microstructure becomes finerwith increasing cooling rate. These types of CCTdiagrams can clearly aid selection of the heat treat-ment cycle in oxide metallurgy.

Evolution of Fraction of Intragranular Ferrite(IGF)

The evolution of the fraction of IGF in a specific steelgrade depends on two factors, according to literature.Firstly, the fraction of IGF depends on the parent

austenite grain size (PAGS).38,53 Secondly, the IGFfraction evolution clearly depends on the cooling rate incontinuous cooling experiments.

Considering the PAGS first, the fraction of IGFincreases with increasing grain size, disregardingother factors such as the composition of inclusionsand steel grade. The reason has been reported to bethat a larger PAGS provides a larger physical spacefor formation of IGF.59 Examples of this behaviorare given in Fig. 6, where results from our ownwork on a steel with TiO2 addition are showntogether with results from literature for low-alloysteels with Ti2O3 addition (alloy A),53 with TiN

5000 5250 5500 5750 6000 6250

(GBF+IGF)starts

finishesstarts

PearlitePearlite

Ferrite

550

600

650

700

750

(a)

Time [s]

Tem

pera

ture

[°C

]

Steel with TiO2 addition800

622 °C655°C

764 °C

-100

-80

-60

-40

-20

200 µm

DS

C, m

W

7000 8000 9000 10000 11000

(b)

Time [s]

Tem

pera

ture

[°C

]

Steel with TiO2 addition

(GBF+IGF)starts

Ferrite

758 °C startsPearlite

672 °C

finishesPearlite650 °C

600

650

700

750

800

-60

-50

-40

-30

DS

C, m

W

Fig. 4. DSC analysis of current steel with TiO2 addition, cooled at rate of (a) 18.3�C/min and (b) 3.6�C/min.


2289

addition (alloy B),53 with Ti (KT4) and with Ti-B-alloying (KT5) where TiX (TiN and TiOx) and Mn-Al-O exist.38

The effect of cooling rate on IGF formation is alsopresented in Fig. 6. It can be seen in Fig. 6a and bthat, when the cooling rate is varied between 3.6�C/min, 70�C/min, and 678�C/min, the area fraction ofIGF increases with increasing cooling rate from3.6�C/min to 70�C/min. However, the IGF fraction islowered when increasing the cooling rate further to678�C/min. The reason for this is competing phasetransformations, where low cooling rates promoteGBF, whereas high cooling rates promote marten-sitic transformation. Hence, both cooling rate andalloy composition must be considered to evaluatewhether a microstructure with high fraction of IGFwill be obtained. It can also be noted that the resultsfor the steel with TiO2 addition in Fig. 6a are closelyrelated to those of alloy A with Ti2O3 addition inFig. 6b. As mentioned above, this is because the

100 101 102 103 104 105

P

IGF700

750

800

IGF finishesPearlite starts

Tem

pera

ture

of s

peci

men

/ °C

Time / second

678 °C/min 70 °C/min3.6 °C/min

Steel with TiO2 addition

Pearlite finishes

600

650

500

550

400

450

GBF IGF startsGBF starts

Cooling from 1400 °C

Fig. 5. Schematic CCT diagram of the steel with TiO2 addition,measured by HT-CLSM, DSC, and EBSD.

0 100 200 300 400 500 600 700 800 900 10000

102030405060708090

100(a)

As-cast [51]

Aver.678 °C/min

3.6 °C/min

70 °C/min

Are

a fra

ctio

n of

IGF

(%)


Current steel with TiO2 addition

Fe-0.165C-0.23Si-0.82Mn (TiO2-added)

0 100 200 300 400 500 600 700 800 900 10000

102030405060708090

100

Alloy AAlloy B

B:Fe-0.227C-0.38Si-0.91Mn (TiN-added)A:Fe-0.223C-024Si-1.02Mn (Ti2O3-added)

70

Aver.678 °C/min

3.6 °C/min

(°C/min)678

Are

a fra

ctio

n of

IGF

(%)


(b) Mu et al. [53]Cooling rate

3.6

70 °C/min

50 100 150 200 250 300 350 4000

102030405060708090

100

300 °C/min

Are

a fra

ctio

n of

IGF

(%)


Cooing rate(c) Zhang et al. [38]

KT4 (Ti-killed) Fe-0.075C-0.23Si-1.45Mn-0.024Ti

KT5 (Ti-killed, B addition) Fe-0.075C-0.23Si-1.46Mn-0.024Ti-10ppm B

Fig. 6. Effect of prior austenite grain size and cooling rate on fraction of IGF in (a) current steel with TiO2 addition and in literature data publishedby (b) Mu et al.53 and (c) Zhang et al.38


inclusions that are acting as effective nucleationsites are the same in both steels and the slightdifferences are instead related to the minor dif-ference in matrix composition.51 It should also benoted that, when the cooling rate is 300�C/min, asin Fig. 6c, the increase of the IGF fraction withPAGS is steeper, thus the effect of PAGS isnegligible above 250 microns, whereas for theslower cooling rate of 70�C/min in Fig. 6a and b,there is a difference with PAGS up to about 500microns.

Austenite Grain Size Control

It is important for the design of an appropriatefinal microstructure in oxide metallurgy that theparent austenite grain size can be controlled. Thetransformation to austenite and the coarsening ofthe austenite grains can also be followed by in situHT-CLSM due to grain boundary grooving generat-ing a topographic contrast. Figure 7 shows typicalHT-CLSM images of the coarsened austenite grainsize in steel after heating to 1200�C and 1400�C atrate of 20�C/min. It is clear that coarsening isstimulated by high temperature and prolongedsoaking, but the details of specific alloys andheating and cooling cycles can be investigated. Theeffect of holding temperature on PAGS in thecurrent alloy and in literature data by Mu et al.53

is summarized in Fig. 8a. Moreover, the effect ofholding time at 1400�C on PAGS from a previousstudy by Zhang et al.38 is summarized in Fig. 8b.The chemical compositions of the steels presented inFig. 8 are summarized in Table II. Note that similarsteel compositions exhibit similar austenite grain-coarsening behavior (see the current alloy and alloy

Fig. 7. Typical HT-CLSM images of increase of prior austenite grain size by heating samples and (a) 1200�C and (b) 1400�C at a rate of 20�C/min.

1200 1250 1300 1350 14000

200

400

600

800

1000

1200 (a)steel

Current

(Ti2O3(Alloy A), TiO2)

Ti-oxides

TiN (Alloy B)

TiN

Ti-oxides

Alloy B[53]Alloy A[53]

10 min5 min

Prio

r aus

teni

te g

rain

siz

e /μ

m

Holding Temperature / °C

Holding time 0 min

0 50 100 150 200 250 300 3500

100

200

300

400

500

600Zhang et al. [38] KT3

KT4 KT5

Prio

r aus

teni

te g

rain

siz

e / μ

m

Holding time / second

(b)

Holding temperature: 1400 °C

Fig. 8. Effect of (a) temperature53 and (b) holding time38 on prioraustenite grain size.


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A in Ref. 53). However, the pinning effect of TiN isobvious in alloy B, and it therefore has a signifi-cantly smaller PAGS. Zhang et al.38 also investi-gated the pinning effect of BN, which was alsostrong, as evidenced by their KT5 sample withsmaller PAGS at each holding condition. As ageneral observation in oxide metallurgy, it shouldalso be mentioned that, in order to control thetoughness of the coarse-grained heat-affected zone

(CGHAZ) of weldments, the austenite grain sizeshould be controlled to a moderate value, some-where between 100 lm to 400 lm, according to HT-CLSM results published to date.38,53 If it is less than100 lm, IGF formation is not sufficient, whereas ifit is over 400 lm, the large grain size can lead to lowtoughness. The grain size in CGHAZ is affectedmainly by the heat input and must be controlledduring welding.

Table II. Chemical compositions of steels presented in Fig. 8

Steel C Si Mn Al Ti S N O Comments

Current sample 0.165 0.23 0.82 < 0.002 0.010 0.008 0.016 0.0066 TiO2 additionMu-Alloy A53 0.223 0.24 1.02 0.002 0.013 0.009 0.017 0.0072 Ti2O3 additionMu-Alloy B53 0.227 0.38 0.91 < 0.002 0.008 0.009 0.029 0.0071 TiN additionZhang-KT338 0.075 0.23 1.48 0.027 0.013 0.003 0.0047 0.001 Al-killed (4 ppm B)Zhang-KT438 0.075 0.23 1.45 < 0.001 0.018 0.003 0.0036 0.002 Ti-killed (1 ppm B)Zhang-KT538 0.075 0.23 1.46 < 0.001 0.024 0.003 0.0042 0.002 Ti-killed, B addition (10 ppm B)

Fig. 9. Typical HT-CLSM images of IGF growth in the steel with TiO2 addition for cooling rate of (a–d) 3.6�C/min, (e–h) 70�C/min, and (i, j)678�C/min.


Ferrite Growth Kinetics

Figure 9 shows typical images of the IGF growthkinetics in the steel with TiO2 addition for threeselected cooling rates. Based on these images, thelength of IGF lath at different cooling times can bemeasured, and subsequently the slope of the fit lineconnecting these data points corresponds to the IGFgrowth speed. These types of results have also beenreported by Wan et al.48,58 based on HT-CLSM

measurements. Note that the growth direction ofIGF may not be parallel to the specimen surface,and it may be growing beneath the surface. Aschematic illustration of the real and observedgrowth rate of IGF is presented in Fig. 10. In orderto know the actual growth speed of IGF, a plausiblemethodology using only HT-CLSM can be to collectseveral IGF growth rates from the same alloy withthe same cooling condition. The maximum value canbe considered as the actual growth speed, as seen inthe example in Fig. 11a. Note that the ‘‘maximumvalue’’ is currently decided in a somewhat ad hocmanner, but it is the intention of the authors topresent more sophisticated analysis of the maxi-mum value of the ferrite growth rate (actual growthrate) in future work. The growth rate of IGFreported by Wan et al.48,58 is also summarized inFig. 11b. It is seen that this speed ranges between33.3 lm/s to 190 lm/s, and that the maximum valuecan be considered as the actual one. In addition, thegrowth rate of IGF reported by Wan et al.48,58 isover 60 times faster than that in the currentsample, which is due to the much faster coolingrate (300�C/min) used in their work compared withthe current value of 3.6�C/min. This research topicwill be investigated in depth in future work by theauthors. Besides the current methodology to detectferrite growth kinetics based on only HT-CLSMmeasurements, further investigations in three-di-mensional characterization, in conjunction within situ observation, can be considered to understandthe difference in the IGF growth kinetics on thespecimen surface and in the bulk.

Opportunities and Challenges in HT-CLSMfor Oxide Metallurgy

The unique benefit of HT-CLSM is its capabilityfor in situ observations of surfaces up to about1800�C with relatively rapid heating and coolingrates. These types of conditions are difficult toachieve when using other in situ microscopy tech-niques such as in situ transmission electron micro-scopy (TEM) or in situ scanning electron microscopy(SEM). The fastest cooling rate for state-of-the-artHT-CLSM today is approximately 3000�C/min athigh temperatures, which fulfills the requirement tomimic the conditions in many metallurgical pro-cesses and for solidification of metals. However, thecooling rate drops significantly at lower tempera-tures that are relevant for studies of solid-statephase transformation in steels; For example, it isnot possible to study the martensitic phase trans-formation or to study low-temperature phase trans-formations such as bainite in isothermal conditionsin many low-alloy steels, since their hardenability istoo low. If a more powerful cooling system could beimplemented to increase the cooling rate to sayabout 100�C/s, it would open up new applicationareas for HT-CLSM.

Fig. 10. Schematic illustration of difference between measured andactual length of an IGF unit.

0 50 100 150 200 250 300 350 400 4500

50

100

150

200

250

300

350

Cooling rate: 3.6 °C/min

Growth speed μm/s]

0.56 2.46

0.33

3.17 (a) Current work

IGF

leng

th (L

IGF)

[μm

]

Relative time for IGF growth ( t) [s]

0

20

40

60

80

100

120

2.52.01.51.00.5

Ref.

190.5 [58] 33.1 [48]

(b) Wan et al. [48, 58]

Cooling rate: 300 °C/min

IGF

leng

th (L

IGF)

[ μm

]

Relative time for IGF growth ( t) [s]

Growth rate[μm/s]

0.0Δ

Δ

Fig. 11. Length of IGF unit against relative time in (a) current steelcooled at 3.6�C/min and (b) steel cooled at 300�C/min reported byWan et al.48,58


2293

HT-CLSM only provides images of the surface ofthe material investigated. This can sometimes be aproblem, since surface behavior may differ from thebulk behavior. An example was given in ‘‘FerriteGrowth Kinetics’’ section, considering the stereologyissue for the growth rate of IGF. Future workshould try to improve the methodologies to analyze,e.g., the growth rate of a three-dimensional (3D)object from a 2D image, and complementary char-acterization techniques may assist this work.Reconstructive phase transformations such as pear-lite and allotriomorphic ferrite formation are muchmore difficult to investigate than displacive trans-formations. The reason is that displacive transfor-mations in steels result in clear surface relief,whereas this is not the case for reconstructivetransformations. Therefore, for example, pearlitecan only be seen in some cases as a faint contrast,and in other cases, e.g., when the surface is oxidizedsomewhat, it will obscure the pearlite observations.

A great development for HT-CLSM would be toenable more analytical information in connectionwith surface observations; For example, it would bevery valuable to be able to study crystallographyand chemistry simultaneously. This could possiblybe achieved by, e.g., micro-focused x-ray diffractionand fluorescence, respectively. In large-scale facili-ties, i.e., synchrotrons, this type of hybrid method-ology has already been implemented where time-resolved synchrotron x-ray diffraction (TRXRD) hasbeen combined with HT-CLSM. This could mostlikely also be achieved on laboratory scale, since thehigh penetration of hard x-rays is not important oreven desirable when HT-CLSM is probing only thesurface.

Another issue that is more difficult to solve is theresolution limitation of HT-CLSM, which is inprinciple on the microscale and physically limitedby the wavelength of the laser light. This resolutionis good enough for the majority of process metallur-gical research, e.g., non-metallic inclusion motionbehavior in liquid metal and metal droplet inrefining reactions, but this issue still limits itsapplication in physical metallurgy, for example, forobserving growth of fine precipitates.

The current application of HT-CLSM in metal-lurgy mainly focuses on ferrous materials, and it isforeseen that also non-ferrous metallurgy couldbenefit significantly as well, for instance, observingthe solidification, metallic cluster dissolution, andmelting in aluminum and titanium alloys. Anotheralloy category where so far no HT-CLSM studieshave been reported is high-entropy alloys, e.g.,FeNiCoCrMn-x multicomponent alloys. Their high-temperature interfacial phenomena should also beof great interest. Another area of interest for HT-CLSM could be solidification in additive manufac-turing, since heating of the metal surface by thehalogen lamp and the solidification process of themetal surface in the HT-CLSM could potentiallymimic AM processing reasonably well.

CONCLUSION

HT-CLSM has proven to be a robust tool forin situ real-time observations of high-temperaturephenomena related to metallurgical processes suchas non-metallic inclusion motion behavior in liquidmetal and their dissolution behavior in liquid slag,crystal growth during solidification as well asaustenite decomposition, and ferrite growth inphysical metallurgy. Even if this methodology hasseveral limitations, numerous topics in metallurgi-cal research have been investigated effectivelyusing HT-CLSM. The present review provides asummary of selected insights gained from bothliterature and new results on austenite decomposi-tion within the field of oxide metallurgy or inclusionengineering using HT-CLSM. Opportunities fordeveloping HT-CLSM by establishing hybridmethodologies in combination with other character-ization methodologies as well as expanding itsapplications in related neighboring fields of manu-facturing are suggested.

ACKNOWLEDGEMENTS

W.M. and P.H gratefully acknowledge Jernkon-toret, the Swedish Steel Producers’ Association, forfinancing this work. W.M. also wants to thank TheSwedish Foundation for International Cooperationin Research and Higher Education (STINT) forfinancial support by granting a Postdoctoral Tran-sition Grants for Internationalisation (PT 2017–7330).

OPEN ACCESS

This article is distributed under the terms of theCreative Commons Attribution 4.0 InternationalLicense (http://creativecommons.org/licenses/by/4.0/), which permits unrestricted use, distribution, andreproduction in any medium, provided you giveappropriate credit to the original author(s) and thesource, provide a link to the Creative Commons li-cense, and indicate if changes were made.

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