1. PHASE
TRANSFORMATI
ON

2. Based on
Mass
transport
PHASE TRANSFORMATIONS
Diffusional
transformation
Diffusion less military
transformation
Based on
Order
PHASE TRANSFORMATIONS
Ist order nucleation
and growth
2nd order entire
volume transforms
No change in
composition
Change in
composition

3. Polymorphic Transformations:
Typically exhibited by single
component systems where different crystal
structures are stable over different temperature
ranges.
• E.g. bcc-fcc transformation in Fe
Major phase transformations that occur in solid phase are due to thermally
activated atomic movements.
The different types of phase transformation that is possible can be divided
into 5 groups:
► Polymorphic changes
► Precipitation Transformation
► Eutectoid transformation
► Ordering reactions
► Massive transformation
• .

4. • Precipitation Transformations:
• Generally expressed as α’→ α + β
• where α’ is a metastable supersaturated
solid solution
• β is a stable or metastable precipitate
• α is a more stable solid solution with the
same crystal structure as α’ but
composition closer to equilibrium

5. • Eutectoid Transformations:
• Generally expressed as γ→ α+β
• Metastable phase (γ) replaced by a more
stable mixture of α + β
• Precipitation and eutectoid transformations
require compositional changes in the
formation of the product phase and
consequently require long-range diffusion

6. • Massive Tranformations:
• Generally expressed as β→ α
• Original phase decomposes into one or more
new phases which have the same composition
as the parent phase but different crystal
structures
• Ordering Transformations:
• Generally expressed as α
(disordered) →α’ (ordered) . These
do not require long range diffusion

7. • PHASE transformation- Change in crystal structure+ Change in
composition.
• Surface creations always hinders the process of transformation. The
new phase always trys to create the surface, so energy needs to be
supplied. So volume free energy will try to decrease the energy but
surface free energy will try to increase the energy.
• ΔFv=VΔf
• V= Vol of the new crystal
• f=free energies of the new phase
• ΔFs = sν
• s = surface area of the new crystal
• ν = free energy per unit area

8. MECHANISM OF PHASE TRANSFORMATION :
• Changes of phase in the solid state involve a redistribution of the atoms in that solid and
the kinetics of the change necessarily depend upon the rate of atomic migration. The
transport of atoms through the crystal is more generally termed diffusion. This can occur
more easily with the aid of vacancies, since the basic act of diffusion is the movement of
an atom to an
empty adjacent atomic site.
• Let us consider that during a phase change an atom is moved from an α-phase lattice
site to a more favorable β-phase lattice site. The energy of the atom should vary with
distance as shown in Figure 1, where the potential barrier which has to be overcome
arises from the interatomic forces between the moving atom and the group of atoms
which adjoin it and the new site.
• Only those atoms (n) with an energy greater than Q are able to make the jump, where Qα
→ β = Hm-Hα
• barrier is given, from the Maxwell–Boltzmann distribution law, as proportional to exp
[−Q/kT ], where k is Boltzmann’s constant, T is the temperature and Q is usually
expressed as the energy per atom in electron volts.

9. FIG 1
• During the transformation it
is not necessary for the
entire system to go from α
to β in one jump and, in
fact, if this were necessary,
phase changes would
practically never occur.
Instead, most phase
changes occur by a process
of nucleation and growth.
• Chance thermal fluctuations
provide a small number of
atoms with sufficient
activation energy to break
away from the matrix (the
old structure) and form a
small nucleus of the new

10. • By this mechanism, the amount of material in the intermediate configuration of
higher free energy is kept to a minimum, as it is localized into atomically thin
layers at the interface between the phases. Because of this mechanism of
transformation, the factors which determine the rate of phase change are:
• (1) the rate of nucleation, N (i.e. the number
of nuclei formed in unit volume in unit time)
• (2) the rate of growth, G (i.e. the rate of increase in
radius with time).
• Both processes require activation energies, which in general are not equal, but the
values are much smaller than that needed to change the whole structure from α to
β in one operation.
• But Growth is more spontaneous process , because it already has surface to grow
over it. With more (surface/volume) ratio it tends to go faster with less
undercooling.
• Even with such an economical process as nucleation and growth transformation,
difficulties occur and it is common to find that the transformation temperature,
even under the best experimental conditions, is slightly higher on heating than on

11. • The combined effect of (a) and (b) is shown in the curve below:
• Where (a) is rate of crystal growth and (b) is rate of nucleation

12. • This sluggishness of the transformation is known as hysteresis,
and is attributed to the difficulties of nucleation, since diffusion,
which controls the growth process, is usually high at temperatures
near the transformation temperature and is therefore not rate
controlling. Perhaps the simplest phase change to indicate this is
the solidification
of a liquid metal.
•
• The transformation temperature, as
shown on the equilibrium diagram,
represents the point at which the free
energy of the solid phase is equal to that
of the liquid phase.Thus,we may
consider the transition, as given in a
phase diagram, to occur when bulk or
chemical free energy change,DGv is
infinitesimally small and negative, i.e.
when a small but positive driving force
exists.However such a definition ignores

13. • When the nucleus is formed the atoms which make up the interface
between the new and old phase occupy positions of compromise
between the old and new structures, and as a result these atoms have
rather higher energies than the other atoms. Thus, there will always be a
positive free energy term opposing the transformation as a result of the
energy required to create the surface
of interface. Consequently, the transformation will occur only when the
sum DGv + DGs becomes negative, where DGs arises from the surface
energy of solid/liquid interface.
• Normally, for the bulk phase change, the number of atoms which form
the interface is small and DGs compared with DGv can be
ignored.However, during nucleation DGv is small, since it is proportional
to the amount transformed, and DGs, the extra free energy of the
boundary atoms, becomes important due to the large surface area-to-
volume ratio of small nuclei. Therefore, before transformation can take
place the negative term DGv must be greater than the positive term DGs
D

14. • Undercooling: It is the gap between the temp predicted for the
transformation to occur and the temp at which the transformation
actually occurs. During the cooling of a liquid,
solidification (nucleation) will begin
only after the temperature has been
lowered below the equilibrium
solidification (or melting)
temperature Tm. This phenomenon
is termed supercooling or
(undercooling).
 The driving force to nucleate
increases as DT increases.
Small supercooling  slow
nucleation rate - few nuclei - large
crystals.

15. Effect of degree of undercooling on the rates of nucleation and growth
Tammann’s curve

16. • The transition from a highly disordered liquid to an ordered solid is
accompanied by a lowering in the energy state of the metal and the release
of thermal energy (latent heat of solidification), forming the arrest on the
cooling curve shown in the previous figure. This ordering has a marked and
immediate effect upon other structure-sensitive properties of the metal; for
instance, the volume typically decreases by 1–6%, the electrical conductivity
rises and the diffusivity, or ability of the atoms to migrate, falls.
• Solidification is a classic example of a nucleation and growth process. In the
general case of freezing
within the bulk of pure molten metal, minute crystalline nuclei form
independently at random points. After this homogeneous form of
nucleation, continued removal of thermal energy from the system causes
these small crystalline regions to grow independently at the expense of the
surrounding melt. Throughout the freezing process, there is a tendency for
bombardment by melt atoms to destroy embryonic crystals; only nuclei

17. • Rapid cooling of a pure molten metal reduces the time available for nuclei
formation and delays the onset of freezing by a temperature interval of dT.
This thermal undercooling (or super cooling), which is depicted in previous
figure , varies in extent, depending upon the metal and conditions, but can
be as much as 0.1–0.3Tm, where Tm is the absolute melting point.
• However, commercial melts usually contain suspended insoluble particles
of foreign matter (e.g. from the refractory crucible or hearth), which act as
seeding nuclei for so-called heterogeneous nucleation.
• Undercooling is much less likely under these conditions; in fact, very
pronounced undercooling is only obtainable when the melt is very pure
and extremely small in volume. Homogeneous nucleation is not
encountered in normal foundry practice.

19. • The growing crystals steadily consume the melt and eventually impinge
upon each other to form a structure of equiaxed (equal-sized) grains (in
upper 2 figures). Heterogeneous nucleation, by providing a larger
population of nuclei, produces a smaller final grain size than
homogeneous nucleation.
• The resultant grain (crystal) boundaries are several atomic diameters
wide. The angle of misorientation between adjacent grains is usually
greater than 10–15◦. Because of this misfit, such high-angle grain
boundaries have a higher energy content than the bulk grains and, on
reheating, will tend to melt first.
• (During a grain-contrast etch of diamond-polished polycrystalline metal,
the etchant attacks grain boundaries preferentially by an electrochemical
process, producing a broad ‘canyon’ which scatters vertically incident
light during normal microscopical examination. The boundary then
appears as a
black line.)

20. • During the freezing of many metals (and alloys), nucleated crystals
grow preferentially in certain directions, causing each growing crystal
to assume a distinctive, non-faceted1tree-like form, known as a
dendrite. In cubic crystals, the preferred axes of growth are <1 0 0>
directions.
• As each dendritic spike grows, latent heat is transferred into the
surrounding liquid, preventing theformation of other spikes in its
immediate vicinity. The spacing of primary dendrites and of dendritic
arms therefore tends to be regular. Ultimately, as the various crystals
impinge upon each other, it is necessary for the interstices of the
dendrites to be well fed with melt if interdendritic shrinkage cavities
are to be prevented from forming.
• Convection currents within the cooling melt are liable to disturb the
delicate dendritic branches and produce slight angular misalignments

21. • Gentle stirring of the melt encourages this process, which is
known as dendrite multiplication, and can be used to produce a
fine-grained and equiaxed structure (e.g. electromagnetic
stirring of molten steel). Dendrite multiplication is now
recognized as an important source of crystals in castings and
ingots.

22. NUCLEATION IN SOLIDS
• When the transformation takes place in the solid state, i.e. between two solid phases, a second factor
giving rise to hysteresis operates. The new phase usually has a different parameter and crystal
structure from the old so that the transformation is accompanied by dimensional changes.
• However, the changes in volume and shape cannot occur freely because of the rigidity of the
surrounding matrix, and elastic strains are induced. The strain energy and surface energy created by
the nuclei of the new phase are positive contributions to the free energy and so tend to oppose the
transition.
The total free energy change is :
DG = VDGS + VDGV + Aγ – (1)
• where A is the area of interface between the two phases and γ the interfacial energy per unit area, and
DGS is the misfit strain energy per unit volume of new phase. For a spherical nucleus of the second
phase:
• DG = (4/3)πr³(DGv − DGs) + 4πr²γ and the misfit strain energy reduces the effective driving force for
the transformation. Differentiation of equation (1) gives
• rc = −2γ/(DGv − DGs), and
• W = (16πγ³/3)(DGv − DGs)²

24. • The value of γ can vary widely from a few mJ m−2 to several hundred mJ
m−2 depending on the coherency of the interface. A coherent interface is
formed when the two crystals have a good ‘match’ and the two lattices
are continuous across the interface. This happens when the interfacial
plane has the same atomic configuration in both phases, e.g. {1 1 1} in
fcc and {0 0 0 1} in cph. When the ‘match’at the interface is not perfect it
is still possible to maintain coherency by straining one or both lattices, as
shown in Figure a.
• These coherency strains increase the energy and for large misfits it
becomes energetically more favorable to form a semi-coherent interface
(Figure b) in which the mismatch is periodically taken up by misfit
dislocations.
• The coherency strains can then be relieved by a cross-grid of dislocations
in the interface plane, the spacing of which depends on the Burgers
vector b of the dislocation and the misfit ε, i.e. b/ε. The interfacial energy

25. • The energy of a semi-coherent interface is 200–500 mJ m−2 and
increases with decreasing dislocation spacing until the dislocation
strain fields overlap. When this occurs, the discrete nature of the
dislocations is lost and the interface becomes incoherent.The
incoherent interface is somewhat similar to a high-angle grain
boundary with its energy of 0.5–1 J m−2 relatively independent of the
orientation.
• The surface and strain energy effects discussed above play an
important role in phase separation. When there is coherence in the
atomic structure across the interface between precipitate and matrix
the surface energy term is small, and it is the strain energy factor
which controls the shape of the particle.
• A plate-shaped particle is associated with the least strain energy,
while a spherical-shaped particle is associated with maximum strain

26. • On the other hand,surface energy determines the crystallographic
plane of the matrix on which a plate-like precipitateforms. Thus, the
habit plane is the one which allows the planes at the interface to fit
together with the minimum of disregistry.
• It is also observed that precipitation occurs most readily in regions
of the structure which are somewhat disarranged, e.g. at grain
boundaries, inclusions, dislocations or other positions of
highresidual stress caused by plastic deformation. Such regions have
an unusually high free energy and
necessarily are the first areas to become unstable during the
transformation. Also, new phases can form there with a minimum
increase in surface energy.

27. HOMOGENEOUS NUCLEATION
• Nuclei form uniformly throughout the parent phase; requires considerable
supercooling (typically 80-300°C).
• Quantitatively, since DGv depends on the volume of the nucleus and Gs is proportional
to its surface area, we can write for a spherical nucleus of radius r:
• DG = (4πr³DGv/3) + 4πr²γ,
• where DGv is the bulk free energy change involved in the formation of the nucleus of
unit volume and γ is the surface energy of unit area. When the nuclei are small the
positive surface energy term predominates, while when they are large the negative
volume term predominates, so that the change in free energy as a function of nucleus
size is as shown in next figure.
• This indicates that a critical nucleus size exists below which the free energy increases
as the nucleus grows, and above which further growth can proceed with a lowering of
free energy; DGmax may be considered as the energy or work of nucleation W . Both rc
and W may be calculated since dDG/dr = 4πr²DGv +8πrγ= 0,when r = rc and thus rc =
−2γ/DGv. Substituting for rc gives W = 16πγ³/3DGv² .

29. • The surface energy factor γ is not strongly dependent on temperature, but
the greater the degree of undercooling or supersaturation, the greater is
the release of chemical free energy and the smaller the critical nucleus size
and energy of nucleation. This can be shown analytically, since DGv = DH −
TDS, and at T = Te, DGv = 0,so that DH = TeDS.
• It therefore follows that DGv = (Te − T ) DS = DTDS and because DGv ∝ DT
, then W ∝ γ³/DT² .
HETEROGENOUS NUCLEATION
• It forms at structural inhomogeneities (container surfaces, impurities,
grain boundaries, dislocations) in liquid phase much easier since
stable “nucleating surface” is already present; requires slight
supercooling (0.1-10ºC ).

30. • This figure shows how this occurs at a
mold wall or pre-existing solid particle,
where the nucleus has the shape of a
spherical cap to minimize the energy and
the ‘wetting’angle θ is given by the balance
of the interfacial tensions in the plane of
the mold wall, i.e. cos θ = (γML −
γSM)/γSL.
• The formation of the nucleus is associated
with an excess free energy given by
DG = VDGv + ASLγSL + ASMγSM −
ASMγML= π/3(2− 3 cos θ + cos3 θ)r³DGv+
2π(1 − cos θ)r²γSL + πr²sin2θ(γSM − γLM).
• Differentiation of this expression for the
maximum, i.e. dDG/dr = 0, gives rc =
−2γSL/DGv and
W = (16πγ³/3DGv²)[(1 − cos θ) ²(2 + cos

31. • The shape factor S(θ) ≤ 1 is dependent on the value of θ and the work of
nucleation is therefore less for heterogeneous nucleation. When θ = 180◦,
no wetting occurs and there is no reduction in W; when θ → 0◦ there is
complete wetting and W → 0; and when 0< θ < 180◦ there is some wetting
and W is reduced.
• If rate kinetics of phase transformation is increased then the structure will
be finer and this is indicated by the Hall - Petch equation States that
decrease in grain size and with fineness in the structure the strength in
increased.
δo =δ + Ka (-1/2) → Hall-Petch Equation
• Where, δo = Friction stress
δ = in stress
a = grain size
K= locking parameter

32. SOLID STATE TRANSFORMATION
• During the solid state transformation still another factor acting
inhibiting the nucleation transformation nuclei.
• A new phase always differs from the initial one in its structure and
specific volume.
• Since the transformation develops an elastic crystalline medium,
change in specific volume should cause an development in elastic
strain energy in one or both the phases. This inhibits the
transformation and kinetics the free energy.
• Therefore, the certain elastic component ΔFel makes a +ve
contribution to the free energy change in the solid state
transformation
•
Fe
g
(Austenite)
Eutectoid
transformation
C FCC
Fe3C
(cementite)
a
(ferrite)
+
(BCC)

34. TTT CURVES
• To understand the character of transformation ot austenite to
the resulting phases, Davenport and Bain showed that by
studying the transformation isothermally (at constant
temperature) of an austenitised
a series of temperatures below A1, a characteristic, time-
temperature transformation (TTT) curved obtained.
• The diagram that illustrates the transformation of austenite as
a function of time at a temperature is a TTT, or isothermal
transformation ( I T) diagram. These curves have ‘C’ or ‘S’
shape in plain carbon and low alloy steels. Each steel
composition has its own different ‘S ’ curve for given grain
size (without inclusions). Fig. a illustrates ‘S’ curve for

35. • The structure produced when austenite is allowed to transform isothermally at
a given temperature can be conveniently represented by a diagram of the type
shown in Figure a, which plots the time necessary at a given temperature to
transform austenite of eutectoid composition to one of the
three structures: pearlite, bainite or martensite.
• Such a diagram, made up from the results of a series of isothermal
decomposition experiments, is called a TTT curve, since it relates the
transformation product to the time at a given temperature. It will be evident
from such a diagram that a wide variety of structures can be obtained from
the austenite decomposition of a particular steel; the structure may range
from 100% coarse pearlite, when the steel will be soft and ductile, to fully
martensitic, when the steel will be hard and brittle. It is because this wide
range of properties can be produced by the transformation of a steel that it
remains a major constructional material for engineering purposes.

36. • From the TTT curve it can be seen that, just below the critical temperature,
A1, the rate of transformation is slow even though the atomic mobility must
be high in this temperature range. This is
because any phase change involving nucleation and growth (e.g. the pearlite
transformation) is faced with nucleation difficulties, which arise from the
necessary surface and strain energy contributions to
the nucleus.
• Of course, as the transformation temperature approaches the temperature
corresponding to the knee of the curve, the transformation rate increases.
The slowness of the transformation below the knee of the TTT curve, when
bainite is formed, is also readily understood, since atomic migration is slow
at these lower temperatures and the bainite transformation depends on
diffusion. The lower part of the TTT curve below about 250–300◦C indicates,
however, that the transformation speeds up again and takes place
exceedingly fast, even though atomic mobility in this temperature range
must be very low.

37. • For this reason, it is concluded that the martensite transformation
does not depend on
the speed of migration of carbon atoms and, consequently, it is
often referred to as a diffusionless
transformation. The austenite only starts transforming to
martensite when the temperature falls below
a critical temperature, usually denoted by Ms. Below Ms the
percentage of austenite transformed to
martensite is indicated on the diagram by a series of horizontal
lines.

39. WHY TTT CURVE HAS A C- SHAPE
• The transformation of austenite doesnot start immediately on quenching the austenised sample to a
constant temperature bath
• Transformation of the austenite to its product occurs after a definite time interval – incubation
period
• Incubation period is that period in which transformation doesnot proceed because enough diffusion
has not taken placein austenite for the transformation to start. Thus the C shape shows that the
stability of austenite first decreases sharply to the minimum then increases again
• Thus the rate of austenite transformation is:
Nil at Ac1 temperature (free energy change is 0)
As temperature falls, it first increases and reaches maximum
(free energy change increases with increase in undercooling)
Nucleation rate increases as critical nucleus size decreases
Rate is maximum at nose
Below the nose the rate of increase in the transformation due to nucleation rate is ofset by in rate
of diffusion at low temperatures
The rate further decreases with the increase in undercooling (diffusion rate)
• Thus the TTT curve has a characteristic C shape.

40. POSSIBLE PHASES IN TTT DIAGRAM FOR EUTECTOID
STEEL

41. • As pointed out before one of the important utilities of the TTT diagrams comes from the
overlay of micro-constituents (microstructures) on the diagram.
• Depending on the T, the (γ+ Fe3C) phase field is labeled with micro-constituents like
Pearlite, Bainite.
• The time taken to 1% transformation to, say pearlite or bainite is considered as
transformation start time and for 99% transformation represents transformation finish.
• We had seen that TTT diagrams are drawn by instantaneous quench to a temperature
followed by isothermal hold.
• Suppose we quench below (~225°C, below the temperature marked Ms), then Austenite
Isothermal Transformation diagram for eutectoid steel
transforms via a diffusionless transformation (involving shear) to a (hard) phase known as
Martensite. Below a temperature marked Mf this transformation to Martensite is complete.
Once γ is exhausted it cannot transform to (γ + Fe3C).
• Hence, we have a new phase field for Martensite. The fraction of Martensite formed is not a
function of the time of hold, but the temperature to which we quench (between Ms and Mf).
• Strictly speaking cooling curves (including finite quenching rates) should not be overlaid on
TTT diagrams (remember that TTT diagrams are drawn for isothermal holds!).

42. • Isothermal hold at: (i) T1 gives us
Pearlite, (ii)T2 gives Pearlite+Bainite,
(iii) T3 gives Bainite.
• Note that Pearlite and Bainite
are both α+Fe 3C (but their
morphologies are different)
• To produce Martensite we should
quench at a rate such as to avoid the
nose of start of C curve called critical
cooling rate.
• if we quench between Ms and Mf we
will get a
mixture of Martensite and γ (called
retained
Austenite).

43. • Determination of TTT diagram for eutectoid steel
• For the determination of isothermal transformation (or) TTT diagrams, we
consider molten salt bath technique combined with metallography and hardness
measurements.
• In molten salt bath technique two salt baths and one water bath are used.
• Salt bath I is maintained at austenising temperature (780˚C for eutectoid steel).
• Salt bath II 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 austenite at A1+20-40°C for eutectoid, A3+20-
40°C for hypo-eutectoid steel and Acm +20-40°C for hyper-eutectoid 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 cold water.
• The microstructure of each sample is studied using metallographic techniques.
The type, as well as quantity of phases, is determined on each sample.
• Transformation of austenite to ferrite-cementite mixtures occurs after a definite
time (say t1) This time during which transformation does not proceed is known as

46. HOW TO DRAW TTT CURVE

58. FACTORS AFFECTING TTT
CURVES• EFFECT OF GRAIN SIZE ON THE TTT CURVES
• EFFECT OF ALLOYING ELEMENTS ON THE TTT CURVES
• EFFECT OF CARBON ON THE TTT CURVES
• 1.EFFECT OF GRAIN SIZE ON THE TTT CURVES
• All decomposition products of austenite nucleate heterogenously
at grain boundaries.
• Thus incubation period is reduced for fine grained steel
• S curve is more towards the left in fine grained steel
Fine grain
Larger grain
boundary area
More nucleation
sites

59. EFFECT OF ALLOYING ELEMENTS ON THE TTT
CURVES All alloying elements (except Co)
shift the S curve to the right
 Austenite stabilizers move the
curve to the right( Mn, Ni,etc)
 Carbide formers shift the S curve
further to the right because:
 Diffusion of alloying elements is
too slow(substitutional elements)
 Diffusion of carbon is slower as
carbide formers donot easily part
with the carbon
 Allotropic change γ -----> α is
reduced by solutes
 Bainitic transformation is lesser
affected ( no redistribution of
alloying elements)
nose
4340 Steel

60. EFFECT OF CARBON ON THE TTT
CURVES
HYPOEUTECTOID STEELS
 Ferrite is the nucleating phase on
decomposition of austenite
 As carbon increases from 0 to
0.77% :
Ferrite
content
decreases
Incubation
period
increases
Nose of S
curve move
more
towards the
right
EUTECTOID STEELS
 Have the maximum incubation
period

61. HYPEREUTECTOID STEELS
• Cementite is the nucleating phase
• As the carbon content increases
more than 0.77%:
Cementite
content
increases
Incubation
period
decreases
Nose of S
curve moves
more
towards the
left

62. TemperatureoC
Ms
Proeutectoid
phase starts to
form on this line
A +F
A
F + P
Pearlite reaction starts
Ac1
Ms
Ms Ms
A+P
P Fe3C +P
Fe3C +A
Proeutectoid
cementite starts
to form on this
line
BB
TTT curves for hypo , eutectoid and hyper-eutectoid steels
TTT curves for hypo , eutectoid and hyper-eutectoid steels

63. CONTINUOUS COOLING
CURVE

64. • The TTT diagrams are also called Isothermal Transformation
Diagrams, because the transformation times are
representative of isothermal hold treatment (following a
instantaneous quench).
• In practical situations we follow heat treatments (T-t
procedures/cycles) in which (typically)there are steps involving
cooling of the sample. The cooling rate may or may not be
constant.
• The rate of cooling may be slow (as in a furnace which has
been switch off) or rapid (like quenching in water).
• Hence, in terms of practical utility TTT curves have a
limitation and we need to draw separate diagrams called
Continuous Cooling Transformation diagrams (CCT), wherein
transformation times (also: products & microstructure) are
noted using constant rate cooling treatments.

65. • A diagram drawn for a given cooling rate (dT/dt) is typically used
for a range of cooling rates (thus avoiding the need for a separate
diagram for every cooling rate).
• However, often TTT diagrams are also used for constant cooling
rate experiments- keeping in view the assumptions &
approximations involved. Important difference between the CCT &
TTT transformations is that in the CCT case Bainite cannot form.
The CCT diagram for eutectoid steel is considered next.
• Determination of CCT diagram for eutectoid steel
• CCT diagrams are determined by measuring some physical
properties during continuous cooling. Normally these are
specific volume and magnetic permeability. However, the
majority of the work has been done through specific volume
change by dilatometric method. This method is supplemented
by metallography and hardness measurement.

66. • In dilatometry the test sample is austenitised in a specially
designed furnace and then controlled cooled. Sample dilation is
measured by dial gauge/sensor. Slowest cooling is controlled by
furnace cooling but higher cooling rate can be controlled by gas
quenching.
•

67. • Cooling data are plotted as temperature versus time (Fig. a). Dilation
is recorded against temperature (Fig. b). Any slope change indicates
phase transformation. Fraction of transformation roughly can be
calculated based on the dilation data as explained below
•

68. • The austenite-pearlite region (A-
--B) terminates just below the
nose. Continued cooling (below
Mstart) of austenite will form
martensite.

69. • For continuous cooling of a steel
alloy there exists a critical
quenching rate that represents
the minimum rate of quenching
that will produce a totally
martensitic structure.
• This curve will just miss the nose
where pearlite transformation
begins

70. Different cooling rates f
eutectoid steel

71. AUSTENITIZATIO
N

72. • The first step in the true heat treatment cycle of steel is the austenitisation
i.e. to get a homogeneous austenite by heating it to a predetermined
temperature in the austenite stability range.
Austenite can transform into various products depending on the composition and cooling rates.
Morphology of parent austenite(grain size) decides the morphology of products and thus its
properties.

73. FORMATION OF AUSTENITE
As the temperature is raised above the A1 temperature, it is the
pearlite which transforms to austenite first. When all the pearlite has
changed to austenite, this austenite grows consuming increasing
amount of free ferrite(in hypoeutectoid steels) or free cementite(in
hypereutectoid steels).

74. Experimentally, nucleation has been seen to occur at the interfaces of ferrite and
cementite lamellae within a pearlite colony but primarily at the intersections of
pearlite colonies.
Once the austenite has nucleated at the interface of ferrite and cementite, it
grows consuming both the ferrite and cementite of pearlite.
The rate of movement of austenite boundary into ferrite and cementite phases is
not equal.
This rate is inversely proportional to the concentration jump at the interface. As
the concentration jump at the austenite-cementite interface is higher(due to the
high concentration of carbon in cementite), austenite boundary moves much
faster into ferrite phase.
By the time when the whole of the pearlitic structure has transformed to

75. In hypoeutectoid steels, the size of proeutectoid ferrite grains is
much larger than the thickness of the ferrite lamellae in pearlite.
Thus the time for complete disappearance of free ferrite exceeds
the time needed for the disappearance of pearlite. The same is
true in case of proeutectoid cementite.
The austenite formed from cementite and ferrite is generally not
homogenous. Some heating is required to make it homogeneous.
Homogenization requires high temperature/time , or both
High temperatures are required if the rate of heating is high,
otherwise comparatively lower temperatures can achieve the
purpose.

76. KINETICS OF AUSTENITE FORMATION
The formation of austenite on heating occurs by nucleation
and growth
The kinetics depends on:
oTransformation temperature and holding time
oRate of heating
oInterface between ferrite and cementite
oGrain size
oNature of the alloying elements present

77. TRANSFORMATION TEMPERATURE
oThe rate of austenite formation increases with increase in temperature as it
increases the rate of carbon diffusion and the free energy is more negative
oTransformation takes a shorter time at higher temperatures of transformation
and vice versa
RATE OF HEATING
oFor higher rates of heating, transformation starts at higher temperatures
and for slower rates, at lower temperatures
oFor any rate of heating transformation occurs over a range of
temperature

78. INTERFACE BETWEEN FERRITE AND CEMENTITE
Higher the interfacial area faster is the transformation.
Interfacial area can be increased by:
Decreasing the inter-lamellar spacing between ferrite and
cementite:
The closer the ferrite – cementite lamellae, the higher is
the rate of nucleation.
Increasing the cementite or carbon content:
This will lead to more pearlite content in steels and thus
more interfaces.
Examples :
1. High carbon steels austenize faster than low carbon
steels
2. Tempered martensite structure austenizes faster than
coarse paerlite

79. GRAIN SIZE
The coarser the parent grain size the slower is the transformation rate.
This is because for a given volume of sample, the total grain boundary area
is less if the grain size is large.
NATURE OF ALLOYING ELEMENTS PRESENT
oAlloying elements in steel are present as alloyed cementite or as alloy
carbides.
oAlloy carbides dissolve much more slowly than alloyed cementite or
cementite.
oThe stronger the alloy carbide formed the slower is the rate of formation of
austenization.
oDiffusion of substitutional alloying elements is much slower than the
interstitial element, carbon.
oThus the rate of austenization depends on the amount and nature of
alloying element

80. IMPORTANCE OF AUSTENITIC GRAIN SIZE IN STEELS
The size of austenitic grains is the most important structural
characteristic of heated steel. The grain size strongly affects its own
transformation behaviour and the mechanical properties of the
microstructures formed from austenite.
Austenitic grain boundaries are preferred sites for the nucleation of
pro-eutectoid phases(pro-eutectoid ferrite in case of hypoeutectoid
steels and proeutectoid cementite in case of hypereutectoid steels) and
pearlite which are diffusion-controlled transformation products.
Coarse austenite grains have less grain boundary area for a given
volume of sample. Thus, fewer nucleation sites are available which leads
to the retardation of diffusion-controlled transformation of austenite
and paves way for the easy transformation to martensite.

81. EFFECT OF GRAN SIZE ON MECHANICAL PROPERTIES
The effect of grain size on different properties are given below:
YIELD STRESS
The dependence is given by Hall-Petch equation :
𝜎𝑜 = 𝜎𝑖 + 𝐾𝐷−1/2
where
𝜎0= yield stress
𝜎𝑖= friction stress opposing motion of dislocation
K is the extent to which dislocations are piled at barriers
D is the average grain diameter

82. Grain refinement improves the strength and ductility at the same time
IMPACT TRANSITION TEMPERATURE
Increase in grain size raises the impact transition temperature, so more
prone to failure by brittle fracture

83. CREEP STRENGTH
 Coarse grained steel has better creep strength above equicohesive temperature
 Below this fine grain structure have better creep strength
FATIGUE STRENGTH
 Fine grained steel have higher fatigue strength
HARDENABILITY
 Coarse grained steels have higher hardenability
 (smaller grain boundary area in coarse grained structure gives less sites for
effective diffusion, so martensite formation on cooling is favoured)
MACHINABILITY
 Coarse grain structure has better machinability due to ease in discontinuos chip
formation(low toughness)

84. PEARLITIC
TRANSFORMATIO
N

85. INTRODUCTION
It is a common micro constituent of a variety of steels where it increases the
strength of steel to a substantial extent.
It is formed when austenite in iron carbon alloys is transformed isothermally
at or below the eutectoid temperature (723K) .
The name ‘Pearlite’ is related to the fact that a polished and etched pearlitic
structure has the colourfulness of mother-of-pearl.

86. DEVELOPMENT OF
MICROSTRUCTURE
Schematic representations of the
microstructures for an iron–carbon
alloy of eutectoid composition
(0.76 wt % C) above and below the
eutectoid temperature.

87. DEVELOPMENT OF
MICROSTRUCTUREAs shown in the previous slide, an alloy of eutectoid composition (0.76 wt %
C) is cooled from a temperature within the phase region, say, 800°C.
Initially, the alloy is composed entirely of the austenite phase having a
composition of 0.76 wt % C and corresponding microstructure, also indicated
in the figure.
As the alloy is cooled, there will occur no changes until the eutectoid
temperature is reached.
 Upon crossing this temperature to point b, the austenite transforms according
to Equation discussed just a few slides before.
The microstructure for this eutectoid steel that is slowly cooled through the
eutectoid temperature consists of alternating layers or lamellae of the two
phases ( α and Fe3C) that form simultaneously during the transformation.

88. DEVELOPMENT OF
MICROSTRUCTUREIn this case, the relative layer thickness is approximately 8 to 1.
This microstructure, represented schematically in the previous figure, point b,
is called pearlite.
Below is a photomicrograph of a eutectoid steel showing the pearlite and
formation of pearlite from austenite.

89. MORPHOLOGY
s
Consider the isothermal
transformation diagram for a
eutectoid iron–carbon alloy,
with superimposed isothermal
heat treatment curve (ABCD).
Microstructures before, during,
and after the austenite-to-pearlite
transformation are shown

90. MORPHOLOGY
The thickness ratio of the ferrite and cementite layers in pearlite is
approximately 8 to 1. However, the absolute layer thickness depends on the
temperature at which the isothermal transformation is allowed to occur.
 At temperatures just below the eutectoid, relatively thick layers of both the
α-ferrite and 𝐹𝑒3 𝐶 phases are produced; this microstructure is called coarse
pearlite, and the region at which it forms is indicated to the right of the
completion curve in the previous figure.
At these temperatures, diffusion rates are relatively high, carbon atoms can
diffuse relatively long distances, which results in the formation of thick
lamellae.

91. MORPHOLOGY
With decreasing temperature, the carbon diffusion rate decreases, and the layers
become progressively thinner. The thin-layered structure produced in the vicinity of
540°C is termed fine pearlite.
Fig.(a) - Coarse Pearlite [ Formed at
higher temp and is relatively
soft ]
Fig.(b) - Fine Pearlite [ Formed at lower
temp and is relatively hard ]

92. MORPHOLOGY
It is a lamellar structure with cementite and ferrite.
The cementite and ferrite are present in a definite ratio of 8:1.
Each ferrite plate in the pearlitic lamellae is a single crystal and some
neighboring plates in a single colony have approximately the same
orientation of lattice. This holds for the cementite also.
In general, both sides of the line of discontinuity in a pearlite colony make a
small angle in lattice orientation with each other.
In the ferrite region near the boundary of pearlite colonies or grains, there are
net-works of dislocations or dislocation walls, at each node of which a
cementite rod is present.

93. MECHANISM
The growth of pearlite from austenite clearly involves two distinct processes:
• a redistribution of carbon (since the carbon concentrates in the cementite and
avoids the ferrite).
• a crystallographic change (since the structure of both ferrite and cementite
differs from that of austenite).
 Of these two processes it is generally agreed that the rate of growth is
governed by the diffusion of carbon atoms, and the crystallographic change
occurs as readily as the redistribution of carbon will allow.
The active nucleus of the pearlite nodule may be either a ferrite or cementite
platelet, depending on the conditions of temperature and composition which
prevail during the transformation, but usually it is assumed to be cementite.

94. MECHANISM
HULL-MEHL model easily explains the pearlite formation.

95. MECHANISM
 The nucleus may form at a grain boundary as shown in (Figure a) in
previous slide, and after its formation the surrounding matrix is depleted of
carbon, so that conditions favour the nucleation of ferrite plates adjacent to
the cementite nucleus (Figure b).
The ferrite plates in turn reject carbon atoms into the surrounding austenite
and this favours the formation of cementite nuclei, which then continue to
grow. At the same time as the pearlite nodule grows sideways, the ferrite and
cementite lamellae advance into the austenite, since the carbon atoms rejected
ahead of the advancing ferrite diffuse into the path of the growing cementite
(Figure c).
 Eventually, a cementite plate of different orientation forms and this acts as a
new nucleus as shown in (Figures d & e).

96. MECHANISM
Hull-Mehl mechanism for Pearlitic Transformation

97. MECHANISM
This process of formation of alternate plates of ferrite and cementite forms a
colony. A new cementite nucleus of different orientation may form at the
surface of colony forming another colony.
The point to be noted is if “austenite transforms to pearlite at a constant temp
then the interlamellar spacing is same in all the colonies”.
NATURE OF NUCLEUS
As pearlite is a 2 phase structure, it may be nucleated either by ferrite
or cementite in steels. In hyper-eutectoid steels, the pro-eutectoid cementite
nucleates pearlite, and in hypo-eutectoid steels, the pro-eutectoid ferrite
nucleates the pearlite. In eutectoid steel, the active nuclei (is defined as the first
one to form) could be either ferrite, or cementite, but may appear to be
cementite).

98. KINETICS
Kinetics of Pearlitic transformation is well explained by JOHNSON &
MEHL model.
JOHNSON & MEHL related the fraction of austenite transformed to pearlite
as a function of time by the equation:
where f(t) = fraction of austenite transformed to pearlite
.
N = Nucleation rate
.
G = Growth rate
t = Time

99. KINETICS
This equation makes the following assumptions:
1) The average nucleation rate is constant with time which actually isn’t true.
2) Nucleation occurs randomly, which isn’t truly correct.
3) The growth rate is constant with time, which can also change from one
nodule to other and with time.
4) Nodules maintain a spherical shape, but nodules may not be truly
spherical.

100. KINETICS
However, when f(t) is plotted against
the resulting sigmoidal curve illustrates that the basic kinetic behavior
of pearlite formation is a nucleation and growth process.

101. KINETICS
The time dependence of the nucleation rate in the early stages has been seen
to increase as the square of time as shown below.

102. KINETICS
The nucleation rate is not constant even at constant temp. If it is assumed to have an
average constant value, then the figure given below illustrates that the rate of
nucleation increases with decreasing temperature of transformation to become almost
maximum at around 550°C.
The nucleation rate is extremely
structure sensitive whereas growth rate is
structure insensitive.
Growth rate is significantly dependent on
temperature, specially on the degree of
undercooling.

103. KINETICS
• At lower critical temp, the free energy of austenite is equal to the free energy
of pearlite.
• Therefore at this temperature transformation of pearlite to austenite
transformation will be completed in infinite time.
• So the rate of transformation will be zero.
• So it is essential to undercool the austenite below the equilibrium (A1) temp.
• Below the lower critical temp, free energy of pearlite < free energy for
austenite and hence it is more thermodynamically stable.
• Lower the free energy more will be the stability of PEARLITE.

104. KINETICS
• Free energy of pearlite is less at lower temperature and so stability is
increased by increasing ΔT.
• The decomposition of austenite to pearlite proceeds by the redistribution of
carbon atoms of austenite into ferrite and cementite, and is essentially a
diffusion controlled process.
• The rate of diffusion decreases exponentially with decreasing temp
• This shows lower the transformation temp retards the rate of transformation.
• There is a transformation temp for which diffusion of C atoms is too small
resulting in diffusion controlled transformation
• Rate of diffusion of carbon atoms is negligible below 200 C

105. KINETICS
 This shows that undercooling affects the rate of transformation in 2 ways:
Undercooling
increased degree of
undercooling reduces
the transformation rate
by lowering the rate of
carbon diffusion
curve.
increased degree of
undercooling increases the
transformation rate by
providing greater difference in
free energies of austenite and
pearlite.

106. KINETICS
• The combined effect is shown in the curve below:
• Where (a) is rate of crystal growth and (b) is rate of nucleation

107. KINETICS
Effect of degree of Undercooling of the rates of nucleation and growth

108. KINETICS
• Hardness of pearlite increases as interlamellar spacing S0 decreases and also
same for strength.
• As S0 is inversely proportional to the degree of undercooling thus yield
strength and also UTS is linearly related to the interlamellar spacing or
degree of undercooling below eutectoid temp.
• As the pearlite content increases in C steels, impact transition temp is
substantially raised, decreasing ductility and toughness as the ferrite-
cementite interface provides sites for easy nucleation of cracks

109. KINETICS

110. EFFECT OF ALLOYING ELEMENTS
ADDITION ON
PEARLITIC TRANSFORMATION
 Almost alloying element except Co lower both the rate of nucleation and rate of growth.
 As compared to carbon other alloying element diffuse very slowly.
 As the diffusion rate for metallic atom is much slower than the
 carbon atom the formation of stable carbide during the transformation will be feasible only
at higher transformation temp.
 Partitioning of carbon gets delayed when Cr eats up C and forms carbide Cr23C6 when
alloyed with austenite.

111. BAINITIC
TRANSFORMATIO
N

112. INTRODUCTION:
• Bainite is an acicular microstructure (not a phase) that forms in steels at temperatures from
approximately 250-550°C (depending on alloy content).
• A fine non-lamellar structure, bainite commonly consists of cementite and dislocation-
rich Ferrite. The high concentration of dislocations in the ferrite present in bainite makes this
ferrite harder than it normally would be.
• Davenport and Bain originally described the microstructure as being similar in appearance to
tempered martensite.

113. EDGAR
BAIN

114. MECHANISM
• Diffusivity of carbon decreases rapidly with fall in temperature. This shows
along with diffusion some other mechanism is responsible for the
transformation to occur.
• Formation of bainite is accompanied by surface distortion so some shear
mechanism is responsible for its transformation.
• So it is a complex one and involves both diffusion less and diffusion
controlled phenomena .Hence, it is termed as a “Diffusion less diffusion
controlled transformation”.
• Two mechanisms are thought to be for the Bainite formation:
1. Diffusive theory
2. Displacive theory

115. DIFFUSIVE THEORY
• The diffusive theory of bainitic transformation process is based on short range diffusion at
the transformation front.
• Random and uncoordinated thermally activated atomic jumps control formation and the
interface is then rebuilt by reconstructive diffusion.
• When the austenite is undercooled below the Bs temp, C atoms redistribute in the
Austenite by diffusion. This redistribution leads to formation of regions with varying
carbon concentration in Austenite. Some of these regions are enriched in carbon while
others are deficient in C. Such a difference in C concentration will result in the
development of stresses.
• The theory is neither able to explain the shape nor surface relief caused by the bainite
transformation.

117. DISPLACIVE THEORY
• Diffusionless growth requires that transformation occurs at a temperature below T0 when
the free energy of bainite becomes less than that of austenite of the same composition.
• A locus of the T0 temperature as a function of the carbon concentration is called the T0
curve,an example of which is plotted on the Fe–C phase diagram. Growth without
diffusion can only occur if the carbon concentration of the austenite lies to the left of the
T0.
• When the plate of bainite forms without diffusion, any excess carbon is soon afterwards
rejected into the residual austenite. The next plate of bainite then has to grow from
carbon–enriched austenite. This process must cease when the austenite carbon
concentration reaches the T0 curve. The reaction is said to be incomplete, since the
austenite has not achieved its equilibrium composition (given by the Ae3 curve) at the
point the reaction stops.

118. Schematic Illustration Of The Origin Of The T0 Construction On
The Fe–c Phase Diagram.
Austenite With A Carbon Concentration To The Left Of The T0
Boundary Can In Principle Transform Without Any Diffusion.
Diffusionless Transformation Is Thermodynamically Impossible If
The Carbon Concentration Of The Austenite Exceeds The T0 Curve.

119. • It is found experimentally that the transformation to bainite does
indeed stop at the T0 boundary.
• The balance of the evidence is that the growth of bainite below the Bs
temperature involves the successive nucleation and martensitic
growth of sub–units, followed in upper bainite by the diffusion of
carbon into the surrounding austenite.
• The possibility that a small fraction of the carbon is nevertheless
partitioned during growth cannot entirely be ruled out.
• The carbon atoms partition into the residual austenite (or precipitate
as carbides),shortly after growth is arrested. The precipitation of
carbides is therefore a secondary event.

120. SHAPE DEFORMATION
• The formation of bainite causes a deformation which is an invariant–plane
strain with a shear component of about 0.26 and a dilatational strain normal
to the habit plane of about 0.03.
• Bainite forms at a relatively high temperature when compared with
martensite. The parent austenite is weaker at high temperatures and cannot
accommodate the large shape deformation elastically. It therefore relaxes by
plastic deformation in the region adjacent to the bainite.
• The effect of this plastic deformation is to stifle the growth of bainite plates
before they hit any obstacle. This is why each bainite plate grows to a size
which is often smaller than the austenite grain size and then comes to a
halt. Further transformation happens by the formation of a new plate and
this is why the sheaf morphology arises.

121. FIG: Atomic Force Microscope Image Of The Displacements Caused
On A Polished Surface Of Austenite By The Growth Of Bainite. Notice
The Shear Deformation (Dark Contrast) And Indeed The Plastic
Accommodation (Light Contrast Tapering From The Ridge Of The
Region Of Dark Contrast) Of The Shape Change In The Austenite
Adjacent To The Bainite Plates.

122. MORPHOLOGY
• On the basis of morphology bainite can be of two types:-
1) Upper bainite
2)Lower bainite

123. UPPER BAINITE
• Known as ‘feathery bainite’ as it resembles feather of a bird
• Forms in temperature range of 5500C-4000C.
• The structure consists of
i. Lath or needle-like ferrite which runs parallel to the longer axis and
ii. Carbide precipitates as fine plates, parallel to the direction of growth of bainite,
mainly at the lath boundaries
• Carbides are present as ‘discontinuous stringers’ when carbon content is low and
‘continuous stringers’ when carbon content is high.

124. • The ferrite laths have ‘sub laths’ with high dislocation density.
• Decrease in temperature produces finer and closely formed laths with smaller
spacing of carbide particles
• The ferrite and cementite in bainite have Kurdjumov–Sachs orientation
relationship with the parent austenite
• Diffusivity of carbon in this temperature range is high enough to cause partition
of carbon between ferrite and austenite.
• Structure is brittle and hard and the deposition of hard carbide stringers on the
soft ferrite makes it a completely useless structure.

125. Schematic growth mechanism of Upper
Bainite
Upper bainite in medium carbon
steel

127. LOWER BAINITE
• Known as ‘Plate bainite’.
• Forms in the temperature range of 4000C-2500C.
• The structure consists of
i.Lenticular plates of ferrite
ii.Fine rods or blades of carbide at an angle of 55 to 60o to the axis of bainite.
• Carbides can be cementite or ε-carbide, or a mixture depending on temperature of
transformation and composition of steel.

128. • Carbides precipitate within the ferrite plates
• Ferrite plates have smaller sub-plates with low angle boundaries between them
• Higher dislocation density than upper bainite
• Habit planes of ferrite plates are the same as martensite that forms at low
temperatures of the same alloy
• Alloying elements do not diffuse or form their carbides during bainite
transformation.

129. Lower Bainite structure in
medium carbon steel
Stages of formation of Lower Bainite
Schematic representation of lower bainite structure

132. INTRODUCTION
• Martensite is a product of a phase transformation that occurs by shear in various alloys like:
Cu-Al ; Au-Cd; Fe-Ni; Fe-C; some ceramics;etc.
• Martensite is a supersaturated solid solution of Carbon in Iron – named after German
metallurgist –Adolph Martens.
• In steels , the parent Austenite can transform to BCC(body-centred cubic), BCT(body-
centred tetragonal) or HCP(hexagonal closed packed) closed packed daughters.
• When rapid cooling occurs from Austenitic state-a very hard structure- Martensite ,forms
the basis of hardening of the steels.
• Morphologically ,Martensite can be found in steels in two forms:
->Plate Martensite
->Lath Martensite
• Martensite need not always be hard and brittle. For example Fe-Ni alloys have soft and
ductile Martensite.

134. MILITARY TRANSFORMATION:
• Most phase transformations studied in this course have been diffusional
transformations where long range diffusion is required for the (nucleation and) growth
of the new phase(s).
• There is a whole other class of military transformations which are diffusion less
transformations in which the atoms move only short distances in order to join the new
phase (on the order of the interatomic spacing).
• These transformations are also subject to the constraints of nucleation and growth.
• They are (almost invariably) associated with allotropic transformations.

135. AUSTENITE –MARTENSITE
TRANSFORMATION
• Martensite, the hardening constituent in quenched steels, is formed at
temperatures below about 200◦C.
• It is formed on quenching austenite, such that the diffusion of carbon is not
favored.
• The atoms move in an organized manner relative to their neighbours and
therefore they are known as a military transformations in contrast to diffusional
civilian transformations.
• Each atom moves by a distance less than one inter-atomic distance and also
retain its neighborhood undisturbed.
• But the total displacement increases as one moves away from the interphase
boundary which results in a macroscopic slip as can be observed as relief
structure on the surface of Martensite.

136. Plate Martensite Showing Coherency With Mother Grain
Structure

137. • At the beginning of the transformation Martensite takes the form of lens or
plates spanning the entire grain diameter
• The subsequent plates formed are limited by the grain boundaries and the
initial Martensite plates formed
• Where the plates intersect the polished surface they bring about a tilting of
the surface.
• But, macroscopically the transformed regions appear coherent to the
surrounding austenite.

138. Crystallography of Martensitic Transformation:
The martensite needles have been formed not with the aid of atomic diffusion but
by a shear process, since if atomic mobility were allowed the large strain energy
associated with the transformed volume would then be largely avoided.
 The lenticular shape of a Martensite needle is a direct consequence of the stresses
produced in the surrounding matrix by the shear mechanism of the transformation
and is exactly analogous to the similar effect found in mechanical twinning.
The strain energy associated with Martensite is tolerated because the growth of
such sheared regions does not depend on diffusion, and since the regions are
coherent with the matrix they are able to spread at great speed through the
crystal.
 The large free energy change associated with the rapid formation of the new phase
outweighs the strain energy, so that there is a net lowering of free energy.

140. CRYSTAL STRUCTURE OF MARTENSITE
A very significant aspect of austenite to martensite transformation is the very large difference in
solid solubility of carbon in gamma iron (0.77% of C at 727 ◦C) and in iron (0.02%C at 727 ◦C).
By rapid cooling of FCC austenite to room temperature the diffusion of carbon is suppressed and
carbon atoms are trapped in octahedral site of bcc structure to result in BCT Martensite.
Austenite, A◦ =3.548 + 0.044(%C)
Martensite, A◦=2.861 – 0.013(%C)
c= 2.861 + 0.16(%C)
Tetragonality is measured by the ratio between the axes, c/a increases with the carbon content as:
c/a=1+0.045 (%C)

141. • When the FCC γ- Fe transforms to bcc α-Fe, carbon is trapped in the octahedral sites of
body centered cubic structure to give body centered tetragonal (BCT) structure
• The trapped carbon atoms cause tetragonal distortion of BCC lattice.
• When carbon is more than 0.2%, BCT structure is formed.

142. IMPORTANT CHARACTERISTICS OF
MARTENSITE TRANSFORMATION:
1) Diffusionless/Military transformation
2) Athermal transformation.
3) Retained Austenite
4) Ms and Mf temperatures
5) Reversibility of Martensitic transformation
6) Habit planes
7) Bain distortion
8) Effect of applied stress on transformation
9) Hardness of Martensite
10) Stabilization of Martensite

143. DIFFUSIONLESS TRANSFORMATION:
• Martensite composition are exactly equal to its parent Austenitic phase.
• The Carbon atoms are present in the same Octahedral sites in
Martensite as that of these sites in FCC- Austenitic phase without
diffusion.
• Diffusionless behaviour can be understood by the fact that in other
alloy systems , the solid solutions remained ordered after this
transformation.

145. Diffusionless Shear Reaction:
• Shape deformation of plate martensite •
Lens formation

146. ATHERMAL TRANSFORMATION:
• Ms and Mf temperatures start from the y-axis of the TTT
curves, indicating the absence of incubation period for this
transformation.
• The first crystal of martensite forms at Ms temperature, and if
more martensite is to be formed, the steel must be cooled
continuously further within Ms-Mf range, but fully
transformation is not possible.

148. MS AND MF TEMPERATURE:
• For each steel, the Austenite to Martensite transformation
starts at a definite temperature called Ms temperature.
• This temperature can vary very widely over the range from
500C to room temperature.
• This variation depends upon the amount of austenite
stabilising elements in the steel (except Co & Al):
Ms (oC)=561 – 474(%C) – 33(%Mn) – 17(%Ni) – 17(%Cr)-21(%Mo).
• Carbon has a very strong effect on the Martensitic start
temperature.

149. • Over a wide range Ms temperature remains independent of
cooling rate , but at very high cooling rates it increases.
• Martensitic transformation can not be suppressed even at the
highest cooling rate attained ,i.e. Ms temperature is raised by
coarse grain of Austenite.

151. RETAINED AUSTENITE
• Martensitic transformation never goes to completion, so the Mf
temperature line is generally dotted .
• At Mf, less than 1% of Austenite is present in a highly stressed state,
along with 99% Martensite.
• Transformation thus is difficult due to unfavourable stress conditions
.
• But for all practical purposes the transformation is said to be
complete at Mf.
• Retained Austenite increases due to higher temperature and increase
in Carbon & alloying elements concentration.
• Steels with less than 0.4%C ,on quenching have very little Retained
Austenite.
• The substructure of Retained Austenite Is different from that of
Austenite due to higher density of dislocations, stacking faults, etc.

154. REVERSIBILITY OF MARTENSITIC
TRANSFORMATION
• With definite amount of superheating as the driving force,
Martensite to Austenite Diffusionless transformation may take
place
• This reverse transformation starts at temperature As
• This property can be seen in systems like:
1. Fe-Ni alloys
2. Al-Cu alloys
3. Ti alloys, etc.

156. • This reversibility has similar features as Ms
transformation like :
• Surface Relief
• As & Af Temperature
• Ad temperature,etc.
• In Fe-Fe3C system , before the reversal from Martensite
to Austenite, tempering reaction occurs.
• Tempering sets due to high (interstitial) diffusivity of C in
Supersaturated BCT Martensite.

157. HABIT PLANES
 The transformation is characterized by a well established relationship between
the orientation of parent austenite and the transformed martensite.
 Habit planes are those planes of the parent austenitic lattice on which
martensitic plates are formed and which lie parallel t the physical plane of the
martensitic plate.
 A habit plane is distorted by the martensite transformation though along it
shear displacement takes place during transformation.
 The habit planes for low, medium and high carbon steels are (111),(225), (259)

158. Martensite habit plane in various types of Steel

159. Martensitic Habit Planes and their conversion to BCT structure

160. An micrograph of austenite that was polished flat and then allowed to transform into martensite.
The different colours indicate the displacements caused when martensite forms.

162. BAIN DISTORTION MODEL:
• In 1924, Bain demonstrated how the BCT lattice could be obtained from
the FCC structure with the minimum of atomic movement, and the
minimum of strain in the parent lattice.
• We use the convention that x,y z and x', y'. z' represent the original and
final axes of the FCC and BCC unit cells.
• An elongated unit cell of the bcc structure can be drawn within two FCC
cells. Transformation to a BCC unit cell is achieved by contracting the
cell 20% in the z direction and expanding the cell by 12% along the x and
y axes.
• The volume expansion during this transformation is 4.3%.
• The Bain deformation results in the following correspondence of crystal
planes and directions:

163. Martensite
FCC
Austenite
FCC
Austenite
Alternate choice of
Cell
Tetragonal
Martensite
Austenite to Martensite → 4.3 % volume increase
Possible positions of
Carbon atoms
Only a fraction of
the sites occupied
20% contraction of c-axis
12% expansion of a-axis
In Pure Fe after
the Matensitic transformation
c = a
C along the c-axis
obstructs the contraction
C
BCT
C
FCC Quench
%8.0
)('
%8.0
)( ag
 

164. EFFECT OF APPLIED STRESS ON
TRANSFORMATION:
• Since the formation of martensite involves a homogeneous distortion of the parent
structure, it is expected that externally applied stresses will be of importance.
• Plastic deformation is effective in forming martensite above the Ms temperature,
provided the temperature does not exceed a critical value usually denoted by Md.
• However, cold work above Md may either accelerate or retard the transformation
on subsequent cooling.
• Even elastic stresses, when applied above the Ms temperature and maintained
during cooling, can affect the transformation; uniaxial compression or tensile
stresses raise the Ms temperature while hydrostatic stresses lower the Ms
temperature.

167. HARDNESS OF MARTENSITE:
• Martensite is the hardest phase found in Fe-C system.
• Reasons of hardness may be the following:
- The solid solution strengthening,
- The imperfections in structure,twins,
- The segregation of carbon to dislocations,
- Grain size of austenite,
- Some precipitated carbides,
-Volume expansion cause the shear and hydrostatic stresses in the lattice,
which lock the screw as well as edge dislocations which is the major cause
of increased hardness.

170. STABILIZATION OF MARTENSITE:
• When cooling is interrupted below Ms, stabilization of the remaining austenite often occurs.
• Thus, when cooling is resumed martensite forms only after an appreciable drop in
temperature.
• Such thermal stabilization has been attributed by some workers to an accumulation of carbon
atoms on those dislocations important to martensite formation. This may be regarded as a
direct analog of the yield phenomenon.
• The temperature interval before transformation is resumed increases with holding time and
is analogous to the increase in yield drop accompanying carbon build-up on strain ageing.
• Furthermore, when transformation in a stabilized steel does resume, it often starts with a
‘burst’, which in this case is analogous to the lower yield elongation.

171. • The transformation starts at a definite temperature –Ms ( Martensite start) temperature. The
transformation proceeds over a range of temperatures till Mf temperature
• The amount of martensite increases on decreasing transformation temperature between Ms and Mf.
• At Mf not all austenite is converted to martensite, but a certain amount is present as retained austenite
• Although the martensite transformation ends at Mf, some austenite still remains untransformed as
retained austenite
• Mf temperature depends on cooling rate . Slower cooling rates lower the Mf temperature
• Mf temperatures are also lowered by increase in carbon content
• Cooling below Mf does not change the amount of martensite.
• The velocity of the martensite transformation, in general, is independent of the transformation
temperature.
• The velocity of transformation is extremely fast almost 10-7 s. This is associated with a crying sound.
• Martensitic transformation is independent of holding time.
171
KINETICS OF MARTENSITIC TRANSFORMATION:

172. 172

173. • Martensite forms by three different modes:
Athermal (without thermal activation)
Burst
Isothermal (thermally activated diffusion-controlled)
1) ATHERMAL Martensite:
oThe amount of Martensite formed is a function of the temperature to which
the alloy is cooled.
oCooling to lower temperatures leads to formation of new plates.
oThis kinetics proceeds above Room temperature, so is dominant in
industrial practices.
oThe fraction of thermal Martensite formed is given by:
f=1- exp(- 1.10 × 10^-2 × ΔT)
where,
ΔT is the degree of undercooling below Ms temperature.
173

174. 2) BURST Kinetics (Jump-like Kinetics):
oFor some alloys like Fe-Ni and Fe-Ni-C , with sub-zero Ms temperatures,
the Burst phenomenon occurs.
oHere the plates of Martensite nucleate newer plates , known as auto-
catalysis .
oZigzag arrays of plates are formed.
oAll the plates form in a very small fraction of second accompanied with an
‘audible click’.
oThe amount of Martensite formed in a burst varies from a few percent to
even 70% of Austenite.
3) ISOTHERMAL Kinetics:
oOccurs in alloys like: Fe-Ni-Mn and Fe-Ni-Cr
oTransformation is a function of time at a constant temperature
oReaction starts slowly, then accelerates due to auto–catalysis, and then
decays

176. MORPHOLOGY OF MARTENSITE:
• Martensite transformation occur by combination by two shears. One of which called lattice
deformation(called Pure strain).
• Second shear is called inhomogeneous lattice deformation.
• Austenite lattice transforms to martensite lattice by it.
• This shear could be by slip or by twinning depending on composition of steel,
temperature of transformation and strain rate.
• Morphology of martensite means the shape of martensite particles. In steel two different
type of morphologies are observed:
- Lath Martensite
- Plate Martensite

178. LATH MARTENSITE
• A lath has the shape of a strip the length of which has largest dimension and is
limited by the grain boundary of austenite.
• Lathe has grouped together in parallel fashion.
• High dislocation density 10^15 – 10^16 /(m)^2.
• Lath Martensite is formed when Ms temperature is high.
• It is formed in low and medium carbon steel.
• The morphology of a lath with dimensions a > b >= c growing on a <111>
plane suggests a thickening mechanism involving the nucleation and glide of
transformation dislocations moving on discrete ledges behind the growing
front.
• It seems possible that due to the large misfit between the BCT and FCC lattices
dislocations could be self-nucleated at the lath interface.
• The criterion to be satisfied for dislocation nucleation in this case is that the
stress at the interface exceeds the theoretical strength of the material.

179. Growth of a Lath
Martensite

180. PLATE MARTENSITE
• The plate Martensite is acicular or lenticular martensite(Lens shaped) resembles
the shape of mechanical twins.
• It forms in steel having lower Ms temperature.
• It is formed in the steel having high percentage carbon.
• In medium and high carbon steels, or high nickel steels, the morphology of the
martensite appears to change from a lath to a roughly plate-like product.
• This is associated with lower Ms temperatures and more retained austenite.
• However, as mentioned earlier, there is also a transition from plates growing on
<225> planes to <259>, planes with increasing alloy content. The lower carbon or
nickel <225> martensite often consists of plates with a central twinned 'midrib',
the outer region of the plate being free of twins.
• It appears that the twinned midrib forms first and the outer (dislocation) region
which is less well defined than the midrib, grows afterwards. The high carbon or
nickel <259> martensite on the other hand is completely twinned and the habit
plane measurements have less scatter than the mixed structures.

181. Constraints
in the matrix
does not
allow parallel
plates but a
lens.
Growth of Plate Martensite:

182. AGE
HARDENING

183.  The strength and hardness of some metal alloys may be improved with
ageing time, by the formation of extremely small, uniformly dispersed
particles (precipitates) of a second phase within the original phase
matrix
 Hardness increases as function of Time
 Some alloys that can be Age-hardened or aged are:
 Copper-beryllium (Cu-Be)
 Copper-tin (Cu-Sn)
 Magnesium-aluminum (Mg-Al)
 Aluminum-copper (Al-Cu)
 High-strength aluminum alloys

184. PRECIPITATION HARDENING
• the strength and hardness of some metal alloys may be
enhanced by the formation of extremely small uniformly
dispersed particles of a second phase within the original phase
matrix.
• this is accomplished by appropriate heat treatments.
• the process is called precipitation hardening because the small
particles of the new phase are termed "precipitates”.

185. REQUISITE FEATURES ON PHASE DIAGRAM FOR AGE
HARDENING
1. Appreciable maximum solubility of component in the other.
2. Solubility limit that rapidly decreases with decrease in temperature
 Alloys can form Super-Saturated-Solid-Solution on cooling
 The SSSS can reject fine dispersion of precipitates on ageing.
3. The precipitates of 2nd phase should be coherent in nature
 “age hardening" is also used to designate this procedure
because the strength develops with time, or as the alloy
ages at designated temperatures below the “solvus”
temperature.
 alloys that are hardened by precipitation treatments
include Al-Cu, Cu-Be, Cu-Sn, and Mg-Al; and some

186. CONTINUED…
 The matrix should be relatively soft and ductile, and the
precipitate should be hard and brittle.
 The alloy must be quenchable.
Solvus
curve
Solvus
curve

187. STEPS IN AGE HARDENING HEAT
TREATMENT
1. SOLUTIONIZING
 first heat treatment where all solute atoms are dissolved to form a
single-phase solid solution.
( just above the solvus temperature)
 Heat to T0 and dissolve second phase
 Over heating is avoided as it may lead to:
 Melting
 Oxidation
 Grain growth
 Burning
 Decrease in ductility

188. STEPS IN AGE HARDENING HEAT
TREATMENT
2. QUENCHING
 Rapidly quench to very low temperature T1
 Metastable Super-Saturated–Solid-Solution i.e high temperature
state ( A phase solid solution supersaturated with B atoms)
formed
 Hot boiling water or air cooling or cold water used as required for
quenching

189. STEPS IN AGE HARDENING HEAT
TREATMENT
3. AGEING
 The supersaturated ‘a’ solid solution is usually heated to an
intermediate temperature T2 within the a+b region (diffusion rates
increase).
 The b precipitates begin to form as finely dispersed particles. This
process is referred to as aging.
 After aging at T2, the alloy is cooled to room temperature.
 Strength and hardness of the alloy depend on the precipitation
temperature (T2) and the aging time at this temperature.
 Ageing for a longer time results in coarsening of the precipitates-
overaging

190. 190
PRECIPITATION HARDENING
• The Process:
• Solution treatment, in which the
alloy is heated to a temperature
above the solvus line into the
alpha phase and held for a
period sufficient to dissolve the
beta phase.
• Quenching to room temperature
to create a supersaturated solid
solution
• Precipitation Treatment; alloy is
heated to a temperature below
Ts to cause precipitation of fine
particles of beta phase.

191. STEPS IN PRECIPITATION
HARDENING

192. 192
©2003 Brooks/Cole, a division of Thomson Learning, Inc. Thomson Learning™ is a trademark used herein under license.
The aluminum-copper phase diagram and the microstructures that
may develop curing cooling of an Al-4% Cu alloy.

193. 193
©2003 Brooks/Cole, a division of Thomson Learning, Inc. Thomson Learning™ is a trademark used herein under license.
The aluminum-rich end of the aluminum-copper phase
diagram showing the three steps in the age-hardening heat
treatment and the microstructures that are produced.

194. ARTIFICIAL AND NATURAL AGEING
• ARTIFICIAL AGEING
Ageing at a temperature higher than room temperature
Hardness peak comes in very short time
Growth is comparable to nucleation
Particles become large in short period and steel loses their hardness
• NATURAL AGEING
Ageing is done at room temperature
Requires long times- Several days to reach maximum hardness
Peak strength is higher than obtained in artificial ageing, no over
ageing occurs.
o
o
o
194

195. Hardness,VHN
Ageing time,
(change of scales at certain
intervals)

196. 196
AGEING

197. EFFECT OF AGEING TEMPERATURE ON
STRENGTH

198. EFFECT OF AGEING TEMPERATURE ON
DUCTILITY

199. QUENCHED IN VACANCIES
• On quenching from high temperature, high % of vacancies
get retained in steel
• These vacancies provide path for diffusion at lower
temperatures when diffusion rate is very slow
• Solute atoms move through few inter atomic distances with
the help of these vacancies to give very fine precipitation –
Ageing
• The fluctuation in solute concentration provide small
clusters in the crystal in solute which acts as nuclei for the
precipitation
• Size of precipitation becomes finer as temperature at which
precipitation occurs is lowered
199

200. PRECIPITATION SEQUENCE DURING AGEING OF
ALLOY (AL-4.5%CU)
• The precipitation occurs in steps involving several transition (metastable)
precipitates before equilibrium precipitate forms
• The equilibrium precipitate does not form instantly as nucleation barrier is
too high - incoherent
• The alloy is quenched from 550°C
• The sequence:
GP Zones
θ’’ (GP
Zone 2)
θ’ θ(CuAl2)
200

201. 201

202. • GP ZONES
Guinier- Preston Zones also called GP1 Zones
The first early stage of ageing
Fully coherent, same lattice structure as Alluminum with matrix thus
nucleation is favored
Plate-like clusters of Copper atoms segregated on {100} planes of
aluminum lattice
Diameter – 100Å , Thickness – 3-6Å
Density 1018 per cm3
Coherency or elastic strains develop
Occurs by diffusion of Cu atoms aided by Quenched-in vacancies over
short distances
Give first peak of hardness
PRECIPITATION SEQUENCE DURING AGEING OF
ALLOY (Al-4.5%Cu)
202

203. • θ’’ (GP2 ZONE)
Coherent intermediate precipitate
Composition is CuAl2
Plate like, Diameter- 1500Å, Thickness- 100Å
Tetragonal crystal Structure, a= 4.04Å, c =7.68Å
Have elastic coherency strains
Produce greater distortion than any other transition structure
PRECIPITATION SEQUENCE DURING AGEING OF
ALLOY (Al-4.5%Cu)
203

204. • θ
Equilibrium precipitate – CuAl2
Fully incoherent precipitate
Nucleates heterogeneously
Tetragonal crystal Structure, a= 6.07Å, c =4.87Å
Coherency strains are not present
Leads to Softening
Result of Overageing
PRECIPITATION SEQUENCE DURING AGEING OF
ALLOY (Al-4.5%Cu)
204

205. • With increasing time, the hardness increases, reaching
a maximum (peak), then decreasing in strength.
• The reduction in strength and hardness after long
periods is overaging (continued particle growth).
PRECIPITATION SEQUENCE DURING AGEING OF
ALLOY (Al-4.5%Cu)
205

206. • Maximum hardness is obtained when there is ‘ Critical-
dispersion’ of GP Zones , or any other intermediate
precipitates(θ’’or θ’) or both
• After peak hardness, further ageing tends to decrease hardness
– overaging
• During overageing, the particles coarsen at the cost of
neighboring particles
PRECIPITATION SEQUENCE DURING AGEING OF
ALLOY (Al-4.5%Cu)
206

207. KINETICS OF PRECIPITATION
• Rate of precipitation is faster at higher temperatures
• Rate of precipitation is faster in alloys of widely dissimilar
metals
• Rate of precipitation is increased with presence of impurities
• Rate of precipitation increases with application of plastic
deformation just before ageing
• Rate of precipitation at a ageing temperature is faster in a low
melting alloy
207

208. 208
Effects of Aging Temperature
and Time
The effect of aging
temperature and
time on the yield
strength of an Al-4%
Cu alloy.
©2003 Brooks/Cole, a division of Thomson Learning, Inc. Thomson Learning™ is a trademark used
herein under license.

209. 209
The operator of a furnace left for his hour lunch break without
removing the Al-4% Cu alloy from the furnace used for the aging
treatment. Compare the effect on the yield strength of the extra
hour of aging for the aging temperatures of 190o
C and 260o
C.
Effect of Aging Heat Treatment Time on the
Strength of Aluminum Alloys
©2003 Brooks/Cole, a division of Thomson Learning, Inc. Thomson
Learning™ is a trademark used herein under license.
Fig – 2 The effect
of aging
temperature and
time on the yield
strength of an Al-4%
Cu alloy.

210. 210
The magnesium-aluminum phase diagram is shown in Figure.
Suppose a Mg-8% Al alloy is responsive to an age-hardening heat
treatment. Design a heat treatment for the alloy.
Design of an Age-Hardening Treatment
©2003 Brooks/Cole, a division of Thomson Learning, Inc. Thomson Learning™ is a trademark used
herein under license.
Fig – 3 Portion of
the aluminum-
magnesium
phase diagram.

211. 211
©2003Brooks/Cole,adivisionofThomsonLearning,Inc.ThomsonLearning™isatrademarkusedhereinunderlicense.
Figure 11.14
Microstructural
changes that occur
in age-hardened
alloys during
fusion welding: (a)
microstructure in
the weld at the
peak temperature,
and (b)
microstructure in
the weld after
slowly cooling to
room temperature.

212. HARDENING MECHANISMS:
• According to dislocation theory, the strength of a material is
controlled by the generation and mobility of the dislocations.
The increased strength of an age-hardened alloy is due to the
interactions of the moving dislocations with the dispersed
precipitates.
Barriers to the motion of dislocations:
• 1) Coherency strains around the GP zones
• 2) GP zones or precipitates

213. REASONS OF HARDENING BY AGEING:
• 1) Coherency strain hardening
• 2) Chemical strain hardening
• 3) Dispersion strain hardening
DISLOCATION INTERACTION WITH:
• 1) Other dislocations-strain hardening
• 2) Grain boundaries-grain boundary strengthening
• 3) Solute atoms-solid solution strengthening
• 4) Precipitates-precipitation hardening
• 5) Dispersoids- dispersion strengthening

214. COHERENCY STRAIN HARDENING:
• This is used for hardening of materials that are not responsive to heat
treatment processes.
• Intensity of strain hardening can be gaged from the slope of the flow curve,
defined by the parameter strain hardening exponent, n.
• Higher the value of n, greater is the strain hardening.
• Increasing the temperature lowers the rate of strain hardening.
• Consequence of strain hardening is improved strength and hardness but
materials’ ductility will be reduced.
• Coherency strain acts as barriers to dislocation movements.
• If size difference between solute and solvent is high, then the strain energy
is also high.
• Higher stress can be applied to overcome the barrier.
• The internal stress increases on :
 increase in size difference between precipitate and matrix
 Increase in elastic modulus of matrix

215. COHERENT GRAINS INCOHERENT
GRAINS

216. DISPERSION HARDENING:
• Small second phase particles distributed in a ductile matrix can
hinder the dislocation motion and thus increase the strength of
the material.
• Second phase particles can be introduced by:
1. Mixing and consolidation(dispersion hardening)
2. Precipitated in solid state(precipitation hardening)
• In dispersion hardening, hard particles are mixed with matrix
powder and processed by powder metallurgy techniques.(here
2nd phase shall have very little solubility in the matrix even at
elevated temperatures)

218. DISLOCATION-CUT MECHANISM:
• Dislocations cut through the precipitate particles.
• Possible only when slip plane is continuous from the matrix through
the precipitate particle and when the stress to move a dislocation in
precipitate is comparable to that in matrix.
• Cutting of particles is easier for small particles.
• Properties that dictate the ease of shearing: coherency strains,
stacking-fault energy, ordered structure, modulus effect, interfacial
energy, morphology and lattice friction stress.
• Shearing disturbs the atomic arrangement along the slip plane.
• Greater is the disturbance, greater is the stress required to shear the
precipitate.
• Thus, the dislocations are pinned.
The dislocations move through the matrix according
to one of the following:

222. BY-PASS MECHANISM:
• Cutting of particles is not possible when there is an interface or an
abrupt change in orientation i.e. when precipitates are incoherent
and larger in size.
• Under such instances, dislocations have to bend around them and
bypass because stress required is too high.
• The dislocation bows around the precipitate and meets at ends X and
Y forming a loop.
• The nature of dislocation at X and Y are opposite and so annihilate.
• A loop of dislocations is left behind the precipitate.
• This is OROWAN MECHANISM, which is similar to the operation of a
Frank-Reed source.
• Stress required to bend a dislocation is inversely proportional to the
average interspacing (l) of particles.

224. • τ=Gb/l
Where:
G= is the shear modulus of the matrix
b= is the Burgers vector of the dislocation
l= is the distance between the dislocations
Every time a dislocation bypasses it leaves behind a loop of
dislocation the precipitate.
Thus l decreases and the stress needed for the next dislocation to
bypass increases
In over ageing precipitates, l increases so strength decreases.

226. RECOVERY,
RECRYSTALLIZT
ION AND
GRAIN
GROWTH

227. INTRODUCTION
• From the above statement , we get to know that Cold
Work leads to various kinds of defects and
dislocations and increase their density.
Cold work
↑ point defect density
↑ dislocation density
 Point defects and dislocations have strain energy associated with them
 (1 -10) % of the energy expended in plastic deformation is stored in the
form of strain energy.

228. EFFECT OF COLD WORK:
• When a metal is cold-worked, by any of the many industrial shaping
operations, changes occur in both its physical and mechanical
properties.
• While the increased hardness and strength which result from the
working treatment may be of importance in certain applications, it is
frequently necessary to return the metal to its original condition to
allow further forming operations (e.g. deep drawing) to be carried out
of for applications where optimum physical properties, such as
electrical conductivity, are essential.
• The treatment given to the metal to bring about a decrease of the
hardness and an increase in the ductility is known as annealing.

229. • This usually means keeping the deformed metal for a certain time at
a temperature higher than about one-third the absolute melting
point.
• Cold working produces an increase in dislocation density; for most
metals ρ increases from the value of 1010–1012 lines m-2 typical of the
annealed state, to 1012–1013 after a few per cent deformation, and up
to 1015–1016 lines m-2 in the heavily deformed state.
• Such an array of dislocations gives rise to a substantial strain energy
stored in the lattice, so that the cold-worked condition is
thermodynamically unstable relative to the undeformed one.
• Consequently, the deformed metal will try to return to a state of
lower free energy, i.e. a more perfect state.

230. • In general, this return to a more equilibrium structure cannot
occur spontaneously but only at elevated temperatures where
thermally activated processes such as diffusion, cross slip and
climb take place.
• Like all non-equilibrium processes the rate of approach to
equilibrium will be governed by an Arrhenius equation of the
form:
Rate = A exp [-Q/kT]
where the activation energy Q depends on impurity
content, strain, etc.

231. • The formation of atmospheres by strain-ageing is one method
whereby the metal reduces its excess lattice energy but this process
is unique in that it usually leads to a further increase in the structure
sensitive properties rather than a reduction to the value
characteristic of the annealed condition.
• It is necessary, therefore, to increase the temperature of the
deformed metal above the strain-ageing temperature before it
recovers its original softness and other properties.

233. • The removal of the cold-worked condition, or in other
words, the annealing process, may be divided into
three stages:
• Recovery
• Recrystallization
• Grain growth

235. FIGURE SHOWING EFFECT OF ANNEALING
PROCESSES ON THE VARIOUS PROPERTIES OF
MATERIAL.

237. • This process describes the changes in the distribution and density of defects with
associated changes in physical and mechanical properties which take place in
worked crystals before recrystallization or alteration of orientation occurs.
• It will be remembered that the structure of a cold-worked metal consists of dense
dislocation networks, formed by the glide and interaction of dislocations, and,
consequently, the recovery stage of annealing is chiefly concerned with the
rearrangement of these dislocations to reduce the lattice energy and does not
involve the migration of large-angle boundaries.
• This rearrangement of the dislocations is assisted by thermal activation.
• Mutual annihilation of dislocations is one process.
• When the two dislocations are on the same slip plane, it is possible that as they run
together and annihilate they will have to cut through intersecting dislocations on
other planes, i.e. ‘forest’ dislocations.
• This recovery process will therefore be aided by thermal fluctuations, since the
RECOVERY:

238. • When the two dislocations of opposite sign are not on the same slip plane, climb or
cross-slip must first occur, and both processes require thermal activation.
• One of the most important recovery processes which leads to a resultant lowering of
the lattice strain energy is rearrangement of the dislocations into cell walls.
• This process in its simplest form was originally termed Polygonization, whereby
dislocations all of one sign align themselves into walls to form small-angle or sub-
grain boundaries.
• During deformation a region of the lattice is curved, and the observed curvature can
be attributed to the formation of excess edge dislocations parallel to the axis of
bending.
• On heating, the dislocations form a sub-boundary by a process of annihilation and
rearrangement.
• As shown in Figure from which it can be seen that it is the excess dislocations of
one sign which remain after the annihilation process that align themselves into

241. oThe relaxation processes occurring during recovery are of two types:
► Annihilation of excess point defect, in particular vacancies.
► Rearrangement of dislocations, and in some process some
annihilation of them.
o The relaxation processes during recovery occur more or less simultaneously
throughout the deformed matrix.
• First relaxation processes starts at low temp during annealing
• Recovery is initially very rapid and more so when the annealing temp is
high.
• Random dislocations of opposite sign come together and annihilate each
other.

242. • Polygonization is a simple form of sub-boundary formation and the basic movement is
climb, whereby the edge dislocations change their arrangement from a horizontal to a
vertical grouping. This process involves the migration of vacancies to or from the edge of
the half-planes of the dislocations.
• The removal of vacancies from the lattice, together with the reduced strain energy of
dislocations which results, can account for the large change in both electrical resistivity
and stored energy observed during this stage, while the change in hardness can be
attributed to the rearrangement of dislocations and to the reduction in the density of
dislocations.
POLYGONIZATION:
POLYGONIZATIO
N

243. Polygonization shown in two Grains. Here it occurs as coalescence of two grains by
rotation of one of them.

244. POLYGONIZATION:
• The process of polygonization can be demonstrated using the Laue method of X-
ray diffraction.
• Diffraction from a bent single crystal of zinc takes the form of continuous radial
streaks. On annealing, these asterisms break up into spots, where each
diffraction spot originates from a perfect polygonized sub-grain, and the
distance between the spots represents the angular misorientation across the
sub-grain boundary.
• Direct evidence for this process is observed in the electron microscope, where, in
heavily deformed polycrystalline aggregates at least, recovery is associated with
the formation of sub-grains out of complex dislocation networks by a process of
dislocation annihilation and rearrangement. In some deformed metals and alloys
the dislocations are already partially arranged in sub-boundaries, forming diffuse
cell structures by dynamical recovery .
• The conventional recovery process is then one in which these cells sharpen and

245. Laue Photograph of Polygonized Zinc.

247. DYNAMIC RECOVERY:
• The low temperature recovery process is generally the reduction of
number of point defects to their equilibrium number.
• At high temperatures , the Recovery process in deformed poly-
crystalline material is the process of movement of dislocations to sub-
boundaries , i.e. the process of Polygonization and Annihilation of
excess dislocations.
• If this process takes place during cold working , then the recovery is
called Dynamic recovery.
• It can take place at low temperatures for pure metals , as it is being
simultaneously stressed ,but can be quite intensive at high
temperatures.
• As the CRSS decreases , when temperature rises, the Dynamic Recovery
decreases the rate of Work-hardening.

248. RECRYSTALLIZATIO
N

249. RECRYSTALLIZATION:
• Recrystallization is the process of formation of new strain free grains
from deformed grains in a solid body by the movement of high angle
boundaries.
• Unlike recovery, the process of recrystallization makes the mechanical
and physical properties of deformed metal to return to completely to
those of the annealed state.
• Mechanical properties like hardness, yield strength, tensile strength,
percentage elongation changes drastically over a very small temperature
range to become typical annealed material.
• Although physical properties like electrical resistivity undergo
appreciable decrease sharply during recrystallization.
• The most significant changes in the structure sensitive properties occur
during the primary crystallization stage. In this stage the deformed

250. • The orientation of the new grains differs considerably from that of
the crystals they consume, so that the growth process must be
regarded as incoherent, i.e. it takes place by the advance of large-
angle boundaries separating the new crystals from the strained
matrix.
•
Typical isothermal recrystallization curve
resembling phase transformation.(factors
kept constant are alloy composition ,
amount of cold work , grain size ,
annealing temperature )