time temperature transformation

TTT diagrams
TTT diagram stands for “time-temperature-transformation” diagram. It is
also called isothermal transformation diagram
Definition: TTT diagrams give the kinetics of isothermal
transformations.
1

Determination of TTT diagram for eutectoid steel
Davenport and Bain were the first to develop the TTT diagram
of eutectoid steel. They determined pearlite and bainite
portions whereas Cohen later modified and included MS and
MF temperatures for martensite. There are number of methods
used to determine TTT diagrams. These are salt bath (Figs. 1-
2) techniques combined with metallography and hardness
measurement, dilatometry (Fig. 3), electrical resistivity
method, magnetic permeability, in situ diffraction techniques
(X-ray, neutron), acoustic emission, thermal measurement
techniques, density measurement techniques and
thermodynamic predictions. Salt bath technique combined
with metallography and hardness measurements is the most
popular and accurate method to determine TTT diagram.
2

Fig. 2 : Bath II low-temperature
salt-bath for isothermal treatment.
Fig. 1 : Salt bath I -austenitisation
heat treatment.
3

Fig . 3(a): Sample and
fixtures for dilatometric
measurements
Fig. 3(b) : Dilatometer
equipment
4

In molten salt bath technique two salt baths and one water
bath are used. Salt bath I (Fig. 1) is maintained at austenetising
temperature (780˚C for eutectoid steel). Salt bath II (Fig. 2) is
maintained at specified temperature at which transformation is
to be determined (below Ae1), typically 700-250°C for
eutectoid steel. Bath III which is a cold water bath is
maintained at room temperature.
In bath I number of samples are austenitised at AC1+20-40°C
for eutectoid and hypereutectoid steel, AC3+20-40°C for
hypoeutectoid steels for about an hour. Then samples are
removed from bath I and put in bath II and each one is kept for
different specified period of time say t1, t2, t3, t4, tn etc. After
specified times, the samples are removed and quenched in
water. The microstructure of each sample is studied using
metallographic techniques. The type, as well as quantity of
phases, is determined on each sample.
5

The time taken to 1% transformation to, say pearlite or bainite
is considered as transformation start time and for 99%
transformation represents transformation finish. On quenching
in water austenite transforms to martensite.
But below 230°C it appears that transformation is time
independent, only function of temperature. Therefore after
keeping in bath II for a few seconds it is heated to above
230°C a few degrees so that initially transformed martensite
gets tempered and gives some dark appearance in an optical
microscope when etched with nital to distinguish from freshly
formed martensite (white appearance in optical microscope).
Followed by heating above 230°C samples are water
quenched. So initially transformed martensite becomes dark in
microstructure and remaining austenite transform to fresh
martensite (white).
6

Quantity of both dark and bright etching martensites are
determined. Here again the temperature of bath II at which 1%
dark martensite is formed upon heating a few degrees above
that temperature (230°C for plain carbon eutectoid steel) is
considered as the martensite start temperature (designated MS).
The temperature of bath II at which 99 % martensite is formed
is called martensite finish temperature ( MF).
Transformation of austenite is plotted against temperature vs
time on a logarithm scale to obtain the TTT diagram. The
shape of diagram looks like either S or like C.
Fig. 4 shows the schematic TTT diagram for eutectoid plain
carbon steel
7

Fig.4: Time temperature transformation (schematic) diagram for plain carbon
eutectoid steel
t1 t3
t2 t4 t5
MF, Martensite finish temperature
M50,50% Martensite
MS, Martensite start temperature
Metastable austenite +martensite
Martensite
%
of
Phase
0
100
Temperature
Log time
Hardness
Ae1
T2
T1
50%
T2
T1
Pearlite
Fine pearlite
Upper bainite
Lower bainite
50% very fine pearlite + 50% upper bainite
At T1, incubation
period for pearlite=t2,
Pearlite finish time
=t4
Minimum incubation
period t0 at the nose
of the TTT diagram,
t0
MS=Martensite
start temperature
M50=temperature
for 50%
martensite
formation
MF= martensite
finish temperature
8

At close to Ae1 temperature, coarse pearlite forms at close to
Ae1 temperature due to low driving force or nucleation rate.
At higher under coolings or lower temperature finer pearlite
forms.
At the nose of TTT diagram very fine pearlite forms
Close to the eutectoid temperature, the undercooling is low so
that the driving force for the transformation is small. However,
as the undercooling increases transformation accelerates until
the maximum rate is obtained at the “nose” of the curve.
Below this temperature the driving force for transformation
continues to increase but the reaction is now impeded by slow
diffusion. This is why TTT curve takes on a “C” shape with
most rapid overall transformation at some intermediate
temperature.
9

Pearlitic transformation is reconstructive. At a given temperature (say
T1) the transformation starts after an incubation period (t2, at T1).
Locus of t2 for different for different temperature is called
transformation start line. After 50% transformation locus of that time
(t3 at T1)for different temperatures is called 50% transformation line.
While transformation completes that time (t4 at T1) is called
transformation finish, locus of that is called transformation finish line.
Therefore TTT diagram consists of different isopercentage lines of
which 1%, 50% and 99% transformation lines are shown in the
diagram. At high temperature while underlooling is low form coarse
pearlite. At the nose temperature fine pearlite and upper bainite form
simultaneously though the mechanisms of their formation are entirely
different. The nose is the result of superimposition of two
transformation noses that can be shown schematically as below one
for pearlitic reaction other for bainitic reaction (Fig. 6).
Upper bainite forms at high temperature close to the nose of TTT
diagram while the lower bainite forms at lower temperature but above
MS temperature. 10

Fig. 5(a) : The appearance of a (coarse) pearlitic
microstructure under optical microscope.
11

Fig. 5(b): A cabbage filled with water analogy of the three-
dimensional structure of a single colony of pearlite, an
interpenetrating bi-crystal of ferrite and cementite.
12

Fig. 5(c): Optical micrograph showing colonies
of pearlite . Courtesy of S. S. Babu.
13

Fig. 5(d): Transmission electron micrograph
of extremely fine pearlite.
14

Fig. 5(e): Optical micrograph of extremely
fine pearlite from the same sample as used to
create Fig. 5(d). The individual lamellae
cannot now be resolved.
15

Fig. 6: Time Temperature Transformation (schematic) diagram for plain carbon
eutectoid steel
MF
M50
MS
Metastable γ + M
M
Temperature
Log time
Hardness
Ae1
P
FP
UB
LB
50% very FP + 50% UB
Metastable γ
γ
γ=austenite
α=ferrite
CP=coarse pearlite
P=pearlite
FP=fine pearlite
UB=upper bainite
LB=lower bainite
M=martensite
MS=Martensite start
temperature
M50=temperature for
50% martensite
formation
MF= martensite finish
temperature
16

On cooling of metastable austenite 1% martensite forms at
about 230°C. The transformation is athermal in nature. i.e.
amount of transformation is time independent (characteristic
amount of transformation completes in a very short time) but
function of temperature alone. This temperature is called the
martensite start temperature or MS.
Below Ms while metastable austenite is quenched at different
temperature amount of martensite increases with decreasing
temperature and does not change with time.
The temperature at which 99% martensite forms is called
martensite finish temperature or MF. Hardness values are
plotted on right Y-axis. Therefore a rough idea about
mechanical properties can be guessed about the phase mix.
17

TTT diagram gives
Nature of transformation-isothermal or athermal (time
independent) or mixed
Type of transformation-reconstructive, or displacive
Rate of transformation
Stability of phases under isothermal transformation conditions
Temperature or time required to start or finish transformation
Qualitative information about size scale of product
Hardness of transformed products
18

Factors affecting TTT diagram
Composition of steel-
(a) carbon wt%,
(b) alloying element wt%
Grain size of austenite
Heterogeneity of austenite
Carbon wt%-
As the carbon percentage increases A3 decreases, similar is the case
for Ar3, i.e. austenite stabilises. So the incubation period for the
austenite to pearlite increases i.e. the C curve moves to right. However
after 0.77 wt%C any increase in C, Acm line goes up, i.e. austenite
become less stable with respect to cementite precipitation. So
transformation to pearlite becomes faster. Therefore C curve moves
towards left after 0.77%C. The critical cooling rate required to prevent
diffusional transformation increases with increasing or decreasing
carbon percentage from 0.77%C and e for eutectoid steel is minimum.
Similar is the behaviour for transformation finish time.
19

Pearlite formation is preceeded by ferrite in case of
hypoeutectoid steel and by cementite in hypereutectoid steel.
Schematic TTT diagrams for eutectoid, hypoeutectoid and
hyper eutectoid steel are shown in Fig.4, Figs. 7(a)-(b) and all
of them together along with schematic Fe-Fe3C metastable
equilibrium are shown in Fig. 8.
20

Fig. 7(a) :Schematic TTT diagram for plain carbon hypoeutectoid steel
MF
M50
MS
Metastable γ + M
M
Temperature
Log time
Hardness
Ae1
α+CP
α+P
FP
UB
LB
FP + UB
γ=austenite
α=ferrite
CP=coarse pearlite
P=pearlite
FP=fine pearlite
UB=upper Bainite
LB=lower Bainite
M=martensite
MS=Martensite start
temperature
M50=temperature for
50% martensite
formation
temperature
Ae3
t0
Metastable γ
21

Fig. 7(b): Schematic TTT diagram for plain carbon hypereutectoid
steel
M50
MS
Metastable γ
Metastable γ + M
Temperature
Log time
Hardness
Ae1
Fe3C+CP
Fe3C+P
Fe3C+FP
UB
LB
very FP +UB
Aecm
t0
γ=austenite
CP=coarse pearlite
P=pearlite
FP=fine pearlite
UB=upper Bainite
LB=lower Bainite
M=martensite
MS=Martensite start
temperature
M50=temperature for
50% martensite
formation
22

Fig. 8: Schematic Fe-Fe3C metastable equilibrium diagram
and TTT diagrams for plain carbon hypoeutectoid, eutectoid
and hypereutectoid steels
MS
(a) Fe-Fe3C
metastable phase
diagram
(b) TTT diagram for
hypoeutectoid steel
(c ) TTT diagram
for eutectoid steel
(d) TTT diagram for
hypereutectoid steel
γ=austenite
α=ferrite
CP=coarse
pearlite
M=martensite
MS=Martensite start temperature
M50=temperature for 50% martensite
formation
MF= martensite finish temperature
P=pearlite
FP=fine pearlite
UB=upper bainite
LB=lower bainite
23

Under isothermal conditions for various compositions
proeutectoid tranformation has been summarised below
(Fig. 9). In hypoeutectoid steel the observable ferrite
morphologies are grain boundary allotriomorph (α)(Fig.11(a)-
(d)), Widmanstätten plate (αW)(Figs. 12-16), and massive (αM)
ferrite (Fig.11(f)).
Grain boundary allotriomorphs form at close to Ae3
temperature or extension of Aecm line at low undercooling.
Widmanstätten plates form at higher undercooling but mainly
bellow Ae1. There are overlap regions where both
allotriomorphs and Widmanstätten plates are observed.
Equiaxed ferrite forms at lower carbon composition less than
0.29 wt%C.
24

Weight % carbon
Temperature
Ae3
0.0218 0.77
Austenite
Pearlite
Ae1
Plate martensite
Mix martensite
Lath martensite
MS
MF
Volume % of retained
austenite at room
temperature
Volume
%
of
retained
austenite
Fig 9: Temperature versus composition in which various morphologies
are dominant at late reaction time under isothermal condition
αW
CmW
αM
Upper bainite
Lower bainite
W=Widmanstätten
plate
M=massive
P=pearlite
αub=upper bainite
αlb =lower bainite
25

There are overlapping regions where both equiaxed ferrite
and Widmanstätten plates were observed. However at very low
carbon percentage massive ferrite forms. The reconstructive
and displacive mechanisms of various phase formation is
shown in Fig. 10.
In hypereutectoid steel both grain boundary allotriomorph and
Widmanstatten plates were observed. Massive morphology
was not observed in hypereutectoid steel. Grain boundary
allotriomorphs were observed mainly close to Aecm or close to
extension of Ae3 line but Widmanstätten plates were observed
at wider temperature range than that of hypoeutectoid steel. In
hypereutectoid steel there are overlapping regions of grain
boundary allotrioorph and Widmanstätten cementite.
26

Fig. 10: The reconstructive and displacive mechanisms.
27

Fig. 11(a): schematic diagram of grain boundary allotriomoph
ferrite, and intragranular idiomorph ferrite.
28

Fig.11(b): An allotriomorph of ferrite in a sample which is partially
transformed into α and then quenched so that the remaining γ
undergoes martensitic transformation. The allotriomorph grows
rapidly along the austenite grain boundary (which is an easy diffusion
path) but thickens more slowly. 29

Fig. 11(c): Allotriomorphic ferrite in a Fe-0.4C steel which is
slowly cooled; the remaining dark-etching microstructure is fine
pearlite. Note that although some α-particles might be identified as
idiomorphs, they could represent sections of allotriomorphs.
Micrograph courtesy of the DoITPOMS project. 30

Fig. 11(d): The allotriomorphs have in this slowly cooled low-
carbon steel have consumed most of the austenite before the
remainder transforms into a small amount of pearlite.
Micrograph courtesy of the DoItPoms project. The shape of
the ferrite is now determined by the impingement of particles
which grow from different nucleation sites.
31

Fig. 11(e): An idiomorph of ferrite in a sample which is partially
transformed into α and then quenched so that the remaining γ
undergoes martensitic transformation. The idiomorph is
crystallographically facetted.
32

Fig. 11(f ): Massive ferrite (αm) in Fe-0.002 wt%C alloy
quenched into ice brine from 1000°C. Courtesy of T. B.
Massalski
33

Fig. 12(a): Schematic illustration of primary Widmanstätten
ferrite which originates directly from the austenite grain
surfaces, and secondary αw which grows from allotriomorphs.
34

Fig. 12(b): Optical micrographs showing white-etching (nital)
wedge-shaped Widmanstätten ferrite plates in a matrix quenched to
martensite. The plates are coarse (notice the scale) and etch cleanly
because they contain very little substructure.
35

Fig. 13: The simultaneous growth of two self-
accommodating plates and the consequential tent-like
surface relief.
36

Fig.14: Transmission electron micrograph of what optically appears
to be single plate, but is in fact two mutually accommodating plates
with a low-angle grain boundary separating them. Fe-0.41C alloy,
austenitised at 1200°C for 6 hrs, isothermally transformed at 700°C
for 2 min and water quenched. 37

Fig. 15: Mixture of allotriomorphic ferrite, Widmanstätten ferrite
and pearlite. Micrograph courtesy of DOITPOMS project.
38

Fig. 16 (a) Surface relief of Widmanstätten ferrite Fe-0.41C
alloy, austenitised at 1200°C for 6 hrs, isothermally
transformed at 700°C for 30 min and water quenced, (b) same
field after light polishing and etching with nital. 39

For eutectoid steel banitic transformation occurs at 550 to
250°C. At higher temperature it is upper bainite and at lower
temperature it is lower bainite. As C increases the austenite to
ferrite decomposition becomes increasingly difficult. As
bainitic transformation proceeds by the nucleation of ferrite,
therefore banitic transformation range moves to higher timing
and lower temperature. With increasing percentage of carbon
the amount of carbide in interlath region in upper bainite
increases and carbides become continuous phase. However at
lower percentage of carbon they are discrete particles and
amount of carbide will be less in both type of bainites. For
start and finish temperatures for both types of bainites go
down significantly with increasing amount of carbon (Figs. 8-
9). However increasing carbon makes it easier to form lower
bainite.
40

Fig 17: Summary of the mechanism of the bainite reaction.
41

Fig. 18: Upper bainite; the
phase between the platelets
of bainitic ferrite is usually
cementite.
42

Fig. 19: Transmission electron micrograph of a sheaf of upper bainite in
a partially transformed Fe-0.43C-2Si-3Mn wt% alloy (a) optical
micrograph, (b, c) bright field and corresponding dark field image of
retained austenite between the sub units, (d) montage showing the
structure of the sheaf. 43

Fig. 20 : Corresponding outline of the sub-units near the sheaf tip
region of Fig. 19 44

Fig. 21 : AFM image showing surface relief due to individual bainite
subunit which all belong to tip of sheaf. The surface relief is
associated with upper bainite (without any carbide ) formed at 350°C
for 2000 s in an Fe-0.24C-2.18Si-2.32Mn-1.05Ni (wt% ) alloy
austenitised at 1200°C for 120 s alloy. Both austenitisation and
isothermal transformation were performed in vacuum. The
microstructure contains only bainitic ferrite and retained austenite.
The measured shear strain is 0.26±0.02. 45

Fig. 22: Optical micrograph illustrating the sheaves of lower bainite in
a partially transformed (395C), Fe-0.3C-4Cr wt% ally. The light
etching matrix phase is martensite. (b) Corresponding transmission
electron micrograph illustrating subunits of lower bainite.
a b
46

Fig. 23 : (a) Optical micrograph showing thin and spiny lower
bainite formed at 190°C for 5 hours in an Fe-1.1 wt% C steel. (b)
Transmission electron micrograph showing lower bainite midrib in
same steel. Courtesy of M. Oka
47

Fig. 24 : Schematic illustration of
various other morphologies: (a)
Nodular bainite, (b) columnar bainite
along a prior austenite boundary, (c)
grain boundary allotriomorphic
bainite, (d) inverse bainite
a
b
c
d
48

Within the bainitic transformation temperature range, austenite of
large grain size with high inclusion density promotes acicular
ferrite formation under isothermal transformation condition. The
morphology is shown schematically (Figs. 25-27 )
Fig. 25 : shows the morphology and nucleation site of
acicular ferrite.
49

Fig. 27: Replica transmission electron micrograph of
acicular ferrire plates in steel weld. Courtesy of Barritte.
51

For eutectoid steel martensite forms at around 230°C. From
230°C to room temperature martensite and retained austenite are
seen. At room temperature about 6% retained austenite can be
there along with martensite in eutectoid steel. At lower carbon
percentage MS temperature goes up and at higher percentage MS
temperature goes down (Fig. 4, Figs. 7-8, Fig. 28). Below 0.4
%C there is no retained austenite at room but retained austenite
can go up to more than 30% if carbon percentage is more than
1.2%. Morphology of martensite also changes from lathe at low
percentage of carbon to plate at higher percentage of carbon.
Plate formation start at around 0.6 % C. Therefore below 0.6 %
carbon only lathe martensite can be seen, mixed morphologies
are observed between 0.6%C to 1%C and above 1% it is 100%
plate martensite (Figs. 29-39).
52

Weight % carbon
Temperature
Austenite +cementite
Ae3
0.0218 0.77
Pearlite
Austenite
Ferrite + pearlite
Pearlite+cementite
Ae1
Plate martensite
Mix martensite
Lath martensite
MS
MF
Ferrite + austenite
Volume % of retained
austenite at room
temperature
Volume
%
of
retained
austenite
Fig. 28: Effect of carbon on MS, MF temperatures and retained austenite in plain carbon
steel
53

Lath
(Fe-9%Ni-0.15%C)
Lenticular
(Fe-29%Ni-0.26%C)
Thin plate
(Fe-31%Ni-0.23%C)
Substructure Dislocation
Dislocation
Twin (midrib)
Twin
Habit plane
{111}A
{557}A
{259}A
{3 10 15}A
{3 10 15}A
O.R. K-S
N-W
G-T
G-T
Ms high low
Fig. 29: Morphology and crystallography of (bcc or bct) martensite in
ferrous alloys
Courtesy of
T. Maki
54

Fig. 30: Lath martensite
Courtesy of
T. Maki
55

(T. Maki，K. Tsuzaki, I. Tamura: Trans. ISIJ, 20(1980), 207.)
Packet: a group of laths
with the same habit plane
( ~{111}g )
Block : a group of laths
with the same orientation
(the same K-S variant)
Fig. 31: effect of carbon
on martensite lath size
56

Fig. 33: Fe-31%Ni-0.28%C
(Ms=192K)
Lenticular martensite
(Optical micrograph)
Fig.32: Fe-29%Ni-0.26%C
(Ms=203K)
Fig.34: schematic diagram for
lenticular martensite
Courtesy of
T. Maki
57

cooling
after polished and etched
Fig. 35: Growth behavior of lenticular martensite
in Fe-30.4%Ni-0.4%C alloy
surface relief surface relief
surface relief
(T. Kakeshita, K. Shimizu, T. Maki, I. Tamura, Scripta Metall., 14(1980)1067.)
Courtesy of
T. Maki
58

midrib twinned region
schematic illustration
Fig. 36: Lenticular martensite in Fe-33%Ni alloy
(Ms=171K)
Optical micrograph
Courtesy of
T. Maki
59

Fig. 37: Optical microstructure of lath martensite (Fe-C alloys)
0.0026%C 0.18%C
0.61%C
0.38%C
Courtesy of
T. Maki
60

Block structure in a single packet (Fe-0.18%C)
Fig.39 : Orientation
image map
Fig. 38: SEM image
Courtesy of
T. Maki

Alloying elements: Almost all alloying elements
(except, Al, Co, Si) increases the stability of supercooled
austenite and retard both proeutectoid and the pearlitic reaction
and then shift TTT curves of start to finish to right or higher
timing. This is due to i) low rate of diffusion of alloying
elements in austenite as they are substitutional elements, ii)
reduced rate of diffusion of carbon as carbide forming
elements strongly hold them. iii) Alloyed solute reduce the rate
of allotropic change, i.e. γ→α, by solute drag effect on γ→α
interface boundary. Additionally those elements (Ni, Mn, Ru,
Rh, Pd, Os, Ir, Pt, Cu, Zn, Au) that expand or stabilise
austenite, depress the position of TTT curves to lower
temperature. In contrast elements (Be, P, Ti, V, Mo, Cr, B, Ta,
Nb, Zr) that favour the ferrite phase can raise the eutectoid
temperature and TTT curves move upward to higher
temperature.
62

However Al, Co, and Si increase rate of nucleation and growth
of both ferrite or pearlite and therefore shift TTT diagram to
left. In addition under the complex diffusional effect of various
alloying element the simple C shape behaviour of TTT
diagram get modified and various regions of transformation
get clearly separated. There are separate pearlitic C curves,
ferritic and bainitic C curves and shape of each of them are
distinct and different.
63

The effect of alloying elements is less pronounced in bainitic
region as the diffusion of only carbon takes place (either to
neighbouring austenite or within ferrite) in a very short time
(within a few second) after supersaturated ferrite formation by
shear during bainitic transformation and there is no need for
redistribution of mostly substitutional alloying elements.
Therefore bainitic region moves less to higher timing in
comparison to proeutectoid/pearlitic region. Addition of
alloying elements lead to a greater separation of the reactions
and result separate C-curves for pearlitic and bainitic regions
(Fig. 40). Mo encourage bainitic reaction but addition of boron
retard the ferrite reaction. By addition of B in low carbon Mo
steel the bainitic region (almost unaffected by addition of B)
can be separated from the ferritic region.
64

Fig. 40: Effect of boron on TTT diagram of low carbon Mo steel
MS
Metastable austenite + martensite
Temperature
Log time
Ae1
Bainite
Metastable austenite
Ae3
Ferrite C curve in low
carbon Mo steel Ferrite C curve in low
carbon Mo-B steel
Pearlitic C curve in low
carbon Mo steel Pearlitic C curve in low
carbon Mo-B steel
Addition of boron
Addition of boron
Bainite start
65

However bainitic reaction is suppressed by the addition of some
alloying elements. BS temperature (empirical) has been given by
Steven & Haynes
BS(°C)=830-270(%C)-90(%Mn)-37(%Ni)-70(%Cr)-83(%Mo)
(elements by wt%)
According to Leslie, B50(°C)=BS-60
BF(°C)=BS-120
Most alloying elements which are soluble in austenite lower MS, MF
temperature except Al, Co.
Andrews gave best fit equation for MS:
MS(°C)=539-423(%C)-30.4Mn-17.7Ni-12.1Cr-7.5Mo+10Co-7.5Si
(concentration of elements are in wt%).
Effect of alloying elements on MF is similar to that of MS. Therefore,
subzero treatment is must for highly alloyed steels to transform
retained austenite to martensite.
66

Addition of significant amount of Ni and Mn can change the nature of
martensitic transformation from athermal to isothermal (Fig. 41).
Temperature
Log time
Fig. 41: kinetics of isothermal martensite in an Fe-Ni-Mn alloy 67

Effect of grain size of austenite: Fine grain size shifts S curve
towards left side because it helps for nucleation of ferrite,
cementite and bainite (Fig. 43). However Yang and Bhadeshia
et al. have shown that martensite start temperature (MS) is
lowered by reduction in austenite grain size (Fig. 42).
Fig. 42: Suppression of Martensite
start temperature as a function
austenite grain size Lγ. MO
S is the
highest temperature at which
martensite can form in large
austenite grain. MS is the observed
martensite start temperature (at 0.01
detectable fraction of martensite).
Circles represent from low alloy data
and crosses from high alloy data.
68

T= MS. a, b are fitting empirial constants,
m =average aspect ratio of martensite=0.05 assumed, Vγ
=average volume of austenite. f=detectable fraction of
martensite=0.01 (taken).
It is expected similar effect of grain size on MF as on MS.
Grain size of austenite affects the maximum plate or lath size.
i.e. larger the austenite size the greater the maximum plate size
or lath size
69

Fig. 43 : Effect of austenite grain size on TTT diagram of plain carbon
hypoeutectoid steel
MF
M50
MS
Metastable γ + M
M
Temperature,
T
Log(time, t)
Hardness
Ae1
α+CP
α+P
FP
UB
LB
50% FP + 50% UB
γ=austenite
α=ferrite
CP=coarse pearlite
P=pearlite
FP=fine pearlite
UB=upper Bainite
LB=lower Bainite
M=martensite
MS=Martensite start
temperature
M50=temperature at
which 50% martensite
is obtained
temperature
Ae3
Metastable γ
For finer austenite
70

Heterogeinity of austenite: Heterogenous austenite increases
transformation time range, start to finish of ferritic, pearlitic
and bainitic range as well as increases the transformation
temperature range in case of martensitic transformation and
bainitic transformation. Undissolved cementite, carbides act
as powerful inocculant for pearlite transformation. Therefeore
heterogeneity in austenite increases the transformation time
range in diffussional transformation and temperature range of
shear transformation products in TTT diagram.
71

Applications of TTT diagrams
• Martempering
• Austempering
• Isothermal Annealing
• Patenting
Martempering : This heat treatment is given to oil hardenable
and air hardenable steels and thin section of water hardenable
steel sample to produce martensite with minimal differential
thermal and transformation stress to avoid distortion and
cracking. The steel should have reasonable incubation period
at the nose of its TTT diagram and long bainitic bay. The
sample is quenched above MS temperature in a salt bath to
reduce thermal stress (instead of cooling below MF directly)
(Fig. 44)
72

Surface cooling rate is greater than at the centre. The cooling
schedule is such that the cooling curves pass behind without
touching the nose of the TTT diagram. The sample is
isothermally hold at bainitic bay such that differential cooling
rate at centre and surface become equalise after some time.
The sample is allowed to cool by air through MS-MF such
that martensite forms both at the surface and centre at the
same time due to not much temperature difference and thereby
avoid transformation stress because of volume expansion.
The sample is given tempering treatment at suitable
temperature.
73

Fig. 44: Martempering heatreatment superimposed on TTT diagram
for plain carbon hypoeutectoid steel
MF
M50
MS
Metastable γ + martensite
Martensite
Temperature
Log time
Ae1
α+CP
α+P
FP
UB
LB
50% FP + 50% UB
γ=austenite
α=ferrite
CP=coarse pearlite
P=pearlite
FP=fine pearlite
t0=minimum incubation
period
UB=upper bainite
LB=lower bainite
M=martensite
MS=Martensite start
temperature
M50=temperature at which
50% martensite is obtained
temperature
Ae3
t0
Metastable γ Tempering
Tempered martensite
74

Austempering
Austempering heat treatment is given to steel to produce lower
bainite in high carbon steel without any distortion or cracking to
the sample. The heat treatment is cooling of austenite rapidly in a
bath maintained at lower bainitic temperature (above Ms)
temperature (avoiding the nose of the TTT diagram) and holding
it here to equalise surface and centre temperature (Fig. 45) and .
till bainitic finish time. At the end of bainitic reaction sample is
air cooled. The microstructure contains fully lower bainite. This
heat treatment is given to 0.5-1.2 wt%C steel and low alloy steel.
The product hardness and strength are comparable to hardened
and tempered martensite with improved ductility and toughness
and uniform mechanical properties. Products donot required to
be tempered.
75

Fig. 45: Austempering heatreatment superimposed on TTT diagram
for plain carbon hypoeutectoid steel
MF
M50
MS
Martensite
Temperature
Log time
Ae1
α+CP
α+P
FP
UB
LB
50% FP + 50% UB
γ=austenite
α=ferrite
CP=coarse pearlite
P=pearlite
FP=fine pearlite
period
UB=upper bainite
LB=lower bainite
M=martensite
MS=Martensite start
temperature
temperature
Ae3
t0
Metastable γ Tempering
Lower bainite
76

Isothermal annealing
• Isothermal annealing is given to plain carbon and alloy steels
to produce uniform ferritic and pearlitic structures. The
product after austenising taken directly to the annealing
furnace maintained below lower critical temperature and hold
isothermally till the pearlitic reaction completes (Fig. 46). The
initial cooling of the products such that the temperature at the
centre and surface of the material reach the annealing
temperature before incubation period of ferrite. As the
products are hold at constant temperature i.e. constant
undercooling) the grain size of ferrite and interlamellar
spacing of pearlite are uniform. Control on cooling after the
end of pearlite reaction is not essential. The overall cycle time
is lower than that required by full annealing.
77

Fig. 46: Isothermal annealing heat treatment superimposed on TTT
diagram of plain carbon hypoeutectoid steel
MF
M50
MS
Martensite
Temperature
Log time
Ae1
α+CP
α+P
FP
UB
LB
50% FP + 50% UB
γ=austenite
α=ferrite
CP=coarse pearlite
P=pearlite
FP=fine pearlite
period
UB=upper bainite
LB=lower bainite
M=martensite
MS=Martensite start
temperature
temperature
Ae3
t0
Metastable γ
Ferrite and pearlite
78

Patenting
Patenting heat treatment is the isothermal annealing at the nose
temperature of TTT diagram (Fig. 47). Followed by this the
products are air cooled. This treatment is to produce fine
pearlitic and upper bainitic structure for strong rope, spring
products containing carbon percentage 0.45 %C to 1.0%C. The
coiled ropes move through an austenitising furnace and enters
the salt bath maintained at 550°C(nose temperature) at end of
salt bath it get recoiled again. The speed of wire and length of
furnace and salt bath such that the austenitisation get over
when the wire reaches to the end of the furnace and the
residency period in the bath is the time span at the nose of the
TTT diagram. At the end of salt bath wire is cleaned by water
jet and coiled.
79

Fig. 47: Patenting heat treatment superimposed on TTT diagram of
plain carbon hypoeutectoid steel
MF
M50
MS
Martensite
Temperature
Log time
Ae1
α+CP
α+P
FP
UB
LB
50% FP + 50% UB
γ=austenite
α=ferrite
CP=coarse pearlite
P=pearlite
FP=fine pearlite
period
UB=upper bainite
LB=lower bainite
M=martensite
MS=Martensite start
temperature
temperature
Ae3
t0
Metastable γ
fine pearlite and upper bainite
80

Prediction methods
TTT diagrams can be predicted based on thermodynamic
calculations.
MAP_STEEL_MUCG83 program [transformation start
curves for reconstructive and displacive transformations for
low alloy steels, Bhadeshia et al.], was used for the following
TTT curve of Fe-0.4 wt%C-2 wt% Mn alloy (Fig. 48)
Fig. 48: Calculated
transformation start curve
under isothermal
transformation condition
81

The basis of calculating TTT diagram for ferrous sytem
1. Calculation of Ae3 Temperature below which ferrite
formation become thermodynamically possible.
2. Bainite start temperature BS below which bainite
transformation occurs.
3. Martensite start temperature MS below which martensite
transformation occurs
4. A set of C-curves for reconstructive
transformation (allotriomorphic ferrite and pearlite).
A set of C-curves for displacive
transformations (Widmanstätten ferrite, bainite)
A set of C-curves for fractional transformation
5. Fraction of martensite as a function of temperature
82

1. Calculation of Ae3 temperature for multicomponent
system. [Method adopted by Bhadeshia et al.]
(This analysis is based on Kirkaldy and Barganis and is applicable for total alloying
elements of less than 6wt% and Si is less than 1 wt%)
Where Xi=mole fraction of component i, γi=activity coefficient of component i,
R=universal gas constant, assuming 0 for Fe, 1for C, i=2 to n for Si, Mn, Ni,
Cr, Mo, Cu,V, Nb, Co, W respectively.
for Ae3 temperature, low
temperature phase γ to be
substituted by α and high
temperature phase L to be
substituted by γ)
Where Xo is the mole fraction of iron then
Assume T is the phase boundary temperature at which high
temperature phase L is in equilibrium with low temperature phase γ.
In case of pure iron then T is given by
General procedure of determination of phase boundary
83

The Wagner Taylor expansion for the activity coefficients
are substituted in the above equations.
and 0GL= standard Gibbs free energy of pure high temperature
phase and 0Gγ= Standard Gibbs free energy of pure low
temperature phase
Similarly for carbon ( n=1) or component i
84

The Wagner-Taylor expansions for activity coefficients are
Where =0 (assumed)
are the Wagner interaction parameters i.e. interaction between
solutes are negligible. The substitution of Wagner-Taylor
expansions for activity coefficients gives temperature deviation
∆T for the phase boundary temperature (due to the addition of
substitutional alloying elements
k=1 to 11 in this case
85

In multicomponent system, the temperature deviations due to
individual alloying additions are additive as long as solute
solute interactions are negligible. Kirkaldy and co-workers
found that this interaction are negiligible as long as total
alloying additions are less than 6wt% and Si is less than
1wt%].
Eventually ∆T takes the following form
Where To is the phase boundary
temperature for pure Fe-C system and To
is given by .
86

And where
for which
and
where n=1 or i and ∆°Ho and ∆°H1 are standard molar
enthalpy changes corresponding to ∆°Go and ∆°G1.
87

If the relevant free energy changes ∆oG and the interaction
parameters ε are known then ∆T can be calculated for any
alloy.
Since all the thermodynamic functions used are dependent on
temperature, ∆T cannot be obtained from single application
of all values (used from various sources) but must be deduced
iteratively. Initially T can be set as To, ∆T is calculated. Then
T=T+ ∆T is used for T and ∆T is found. Iteration can be
repeated for a few times (typically five times) about till T
changes by less than 0.1K.
This method obtains Ae3 temperature with accuracy of ±10K.
88

2. Bainite start temperature BS from Steven and Haynes formula
BS(°C)=830-270(%C)-90(%Mn)-37(%Ni)-70(%Cr)-83(%Mo)
( % element by wt)
Both bainite and Widmanstätten ferrite nucleate by same
mechanism. The nucleus develops into Widmanstätten ferrite if
at the transformation temperature the driving force available
cannot sustain diffusionless transformation. By contrast bainite
form from the same nucleus if the transformation can occur
without diffusion. Therefore in principle BS=WS.
Bainite transformation does not reach completion if austenite
enriches with carbon. But in many steels carbide precipitation
from austenite eliminates the enrichment and allow the
austenite to transform completely. In those cases bainite finish
temperature is given (according to Leslie) by
89

3. MS Temperature:
At the MS temperature
90

In the above equation, T refers to MS temperature in absolute
scale, R is universal gas constant, x=mole fraction of carbon,Yi is
the atom fraction of the ith substitutional alloying element, ∆Tmagi
and ∆TNMi are the displacement in temperature at which the free
energy change accompanying the γ→α transformation in pure
iron (i.e. ∆FFe
γ→α) is evaluated in order to allow for the changes
(per at%) due to alloying effects on the magnetic and non-
magnetic components of ∆FFe
γ→α , respectively. These values were
taken from Aaronson, Zenner. ∆FFe
γ→α value was from Kaufmann.
91

The other parameters are as follows
(i) the partial molar heat of solution of carbon in ferrite,
∆¯Hα=111918 J mol-1 (from Lobo) and
∆¯Hα=35129+169105x J mol-1 (from Lobo)
ii) the excess partial molar non-configurational entropy of
solution of carbon in ferrite ∆Sα=51.44 J mol-1K-1 (from Lobo)
∆Sγ=7.639+120.4x J mol-1K-1 (from Lobo)
ωα=the C-C interaction energy in ferrite=48570 J mol-1(average
value) (from Bhadeshia)
ωγ=the corresponding C-C interaction energy in austenite values
were derived , as a function of the concentrations of various
alloying elements, using the procedure of Shiflet and Kingman
and optimised activity data of Uhrenius. These results were
plotted as a function of mole fraction of alloying elements and
average interaction parameter ω¯γ was calculated following
Kinsman and Aaronson.
92

∆f*=Zener ordering term was evaluated by Fisher.
The remaining term, ∆FFe
γ→α’=free energy change from
austenite to martensite as only function of carbon content.
and is identical for Fe-C and Fe-C-Y steels as structure for
both cases are identical (Calculated by Bhadeshia )=-900 to -
1400 J mol-1 (for C 0.01 to 0.06 mole fraction, changes are not
monotonic). However Lacher-Fowler-Guggenheim
extrapolation gives better result of -1100 to -1400 J mol-1
(Carbon mole fraction 0.01-0.06).
The equation was solved iteratatively until the both sides of
the equation balanced with a residual error of <0.01%. The
results underestimate the temperature of 10-20K. The error
may be due to critical driving force for transformation
calculation consider only a function of carbon content.
93

4. Transformation start and finish C curves
The incubation period (τ) can be calculated from the following
equation [Bhadeshia et al.]
Where T is the isothermal transformation temperature in absolute
scale, R is universal gas constant, Gmax is the maximum free energy
change available for nucleation, Q is activation enthalpy for
diffusion, C,p, z=20 are empirical constant obtained by fitting
experimental data of T, Gmax, τ for each type of transformation
(ferrite start, ferrite finish, bainite start and bainite finish) . By
systematically varying p and plotting ln(τ Gp
max/ Tz=20) against 1/RT
for each type of transformation (reconstructive and displacive) till
the linear regression coefficient R1 attains an optimum value. Once p
has been determined Q and C follow from respectively the slope
and intercept of the of plot. The same equation can be used to
predict transformation time. 94

Table-I: Chemical compositions, in wt% of the steel
chosen to test the model
95

The optimum values of p and corresponding values of C, Q for
different types of steels (Table-I) where concentrations are in wt% are
summarised below [Bhadeshia et.al.] (Table-II).
Table-II: The optimum values of fitting constants
FS=ferrite start, FF=ferrite finish, BS=bainite start and BF=bainite
finish
96

Based on Q, C and Gmax value it can be predicted that Mo
strongly retard the formation of ferrite through its large
influence on Q. however it can promote bainite via the small
negative coefficient that it has for the Q of the bainitic C-
curve. Cr retard both both bainite and ferrite but net effect is
to promote the formation of bainite since the influence on the
bainitic C-curve is relatively small. Ni has a slight retarding
effect on tranformation rate. Mn has also retarding effect on
ferrite as as well as bainitic transformation
The bainite finish C-curve of the experimental TTT diagram
not only shifts to longer time but also but is also shifts to
lower temperature by about 120°C. Therefore this is taken
care by plotting against in
order to determine p, Q and C for the bainitic finish curve.
97

Fractional transformation curves
Fractional transformation time can be estimated from the
following Johnson-Mehl-Avrami equation.
X=transformation volume fraction, K1 is rate constant which is
a function of temperature and austenite grain size d, n and m
are empirical constants. By selecting steels of similar grain
size, the austenite grain size can be neglected then the above
equation simplifies to
98

Assuming x=0.01 for transformation start and x=0.99 for
transformation finish. For a given temperature transformation
start time and finish time can be calculated then K1 and n can
be solved for each transformation product an a function of
temperature.
Then fractional transformation curves for arbitary values of x
can therefore be determined using
99

Representation of intermediate state of transformation
between 0% and 100% can be derived by fitting to the
experimental TTT diagram as follows:
Where x refers to the fraction of transformation.
In most of TTT diagrams of Russell has a plateau at its
highest temperature. Therefore a horizontal line can be drawn
at BS and joining it to a C-curve calculated for temperatures
below BS.
100

Relation between observed and predicted values for ferrite start
(FS), ferrite finish (FF), bainite start (BS) and bainite finish
(BF) are shown in Fig. 49. Predicted value closely matches the
observed values for selected low alloy steels. Predicted TTT
diagrams are projected on experimental diagrams (Figs. 50-52).
The model correctly predicts bainite bay region in low alloy as
well as in selected high alloy steels. The model reasonably
predicts the fractional C curves (Fig. 52). Mo strongly retard the
formation of ferrite through its large influence on Q. however it
can promote bainite via the small negative coefficient that it
has for the Q of the bainitic C-curve. Cr retard both both bainite
and ferrite but net effect is to promote the formation of bainite
since the influence on the bainitic C-curve is relatively small. Ni
has a slight retarding effect on transformation rate. Mn has also
retarding effect on ferrite as as well as bainitic transformation
The model is impirical in nature but it can nevertheless be
useful in procedure for the calculation of microstructure in steel.
101

Fig. 49: Relation between observed and predicted Q(Jmol-1)
value for: (a) FS-ferrite start, (b) FF-ferrite finish, (c) BS-bainite
start and (d) BF-bainite finish curves. 102

Fig. 50: Comparison of experimental and predicted TTT diagram
for BS steel:(a) En14, (b) En 16, (c) En 18 and (d) En 110.
103

Fig. 51: Comparison of experimental and predicted TTT diagram for
US steel:(a) US 4140, (b) US 4150, (c) US 4340 and (d) US 5150
104

Fig. 52: Comparison of the experimental and predicted TTT
diagrams including fractional transformation curves at 0.1, 0.5
and 0.9 transform fractions: (a)En 19 and (b) En24. 105

Limitations of model
The model tends to overestimate the transformation time at
temperature just below Ae3. This is because the driving force
term ∆Gmax is calculated on the basis of paraequilibrium and
becomes zero at some temperature less than Ae3 temperature.
The coefficients utilized in the calculations were derived by
fitting to experimental data, so that the model may not be
suitable for extrapolation outside of that data set. Thus the
calculation should be limited to the following concentration
ranges (in wt%):C 0.15-0.6, Si 0.15-0.35, Mn 0.5-2.0, Ni 0-
2.0, Mo 0-0.8 Cr 0-1.7.
106

5. Fraction of martensite as a function of temperature
Volume fraction of martensite formed at temperature T =f and
f=1-exp[BVpdΔGv)/dT(MS-T)]
Where, B=constant, Vp=volume of nucleus, ∆Gv=driving force
for nucleation, MS =martensite start temperature. Putting the
measured values
the above equation becomes
f=1-exp[-0.011(MS-T)] [Koistinen and Marburger equation].
The above equation can be used to calculate the fraction of
martensite at various temperature.
107

108
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Any questions?
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● hamzaahmed0696@mail.me
108

time temperature transformation

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