HEAT TREATMENT
Heat Treatment is the controlled heating and
cooling of metals to alter their physical and mechanical properties without changing the product shape.
Steels are particularly suitable for heat treatment,
since they respond well to heat treatment and the commercial use of steels exceeds that of any other material.
Generally, heat treatment uses phase
transformation during heating and cooling to change a microstructure in a solid state.
In heat treatment, the processing is most often
entirely thermal and modifies only structure. Thermomechanical treatments, which modify component shape and structure, and
thermochemical treatments which modify surface chemistry and structure, are also important
processing approaches which fall into the domain of heat treatment.
Iron carbon diagram form the basis of heat
treatments:
Various types of heat treatment processes
are used to modify the following
properties or conditions of the steel:
Improve the toughness Increase the hardness Increase the ductility
Improve the machinability Refine the grain structure Remove the residual stresses Improve the wear resistance
ANNEALING
Annealing is a heat process whereby a metal is
heated to a specific temperature, held at that
temperature and then allowed to cool slowly.
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.
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.
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.
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.
The removal of the cold-worked condition, or
in other words, the annealing process, may be
divided into three stages:
Recovery
Recrystallization
Grain growth
This is primarily a low temperature process, and
the property changes produced do not cause appreciable change in microstructure or the properties, such as tensile strength, yield strength, hardness and ductility.
The principal effect of recovery is the relief of
internal stresses due to cold working and this
prevents stress corrosion cracking and minimizes the distortion produced by residual stresses.
Electrical conductivity is also increased
appreciably during the recovery stage.
Commercially, this low temperature treatment in
the recovery range is known as stress relief annealing or process annealing.
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.
The recovery stage of annealing is chiefly
concerned with the rearrangement of the cold
worked dislocations to reduce the lattice energy
and does not involve the migration of
large-angle boundaries.
This rearrangement of the dislocations is
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 activation energy for such a cutting process is small.
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
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 subgrain boundaries.
During deformation a region of the lattice is curved,
as shown in Figure 7.50a, 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. This is shown in Figure 7.50b, 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 walls.
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.
Recrystallization takes place by a combination of
nucleation of strain free grains and the growth of these nuclei to absorb the entire cold worked material.
The term recrystallization temperature does not refer to a definite temperature below which recrystallization will not occur, but refers to the approximate temperature at which a highly cold worked material completely recrystallizes in one hour.
Pure metals have low recrystallization temperatures as compared with alloys.
Zinc, tin and lead have recrystallization temperatures below room temperature. This means that these metals cannot be cold worked at room temperature since they crystallize spontaneously, reforming a strain free
The recrystallization temperatures of several metals and alloys are listed as in the table below:
Material Recrystallization Temperature (oF)
Copper ( 99.99 %) 250
Copper ( 5 % Zinc) 600
Aluminum (99.99 %) 175
Aluminum alloys 600
Low carbon steel 1000
Zinc 50
Tin 25
In the primary recrystallization stage the deformed
lattice is completely replaced by a new unstrained one by means of a nucleation and growth process, in which practically stress free grains grow from nuclei formed in the deformed matrix.
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.
During the growth of grains, atoms get
transferred from one grain to another across the
boundary. Such a process is thermally activated
as shown in Figure
It is well known that the rate of recrystallization depends on several important factors, namely:
(1) the amount of prior deformation (the greater the degree of cold work, the lower the recrystallization temperature and the smaller the grain size),
(2) the temperature of the anneal (as the temperature is lowered the time to attain a constant grain size increases exponentially)
(3) the purity of the sample (e.g. zone refined aluminium recrystallizes below room temperature, whereas
aluminium of commercial purity must be heated several hundred degrees).
(4)Increasing the annealing time decreases the recrystallization temperature for the start of recrystallization.
Measurements, using the light microscope, of the increase in
diameter of a new grain as a function of time at any given temperature can be expressed as shown in the figure. The diameter increases linearly with time until the growing grains begin to impinge on one another, after which the rate
necessarily decreases.
Figure 7.52 Variation of grain diameter with time at a constant
The classical interpretation of these
observations is that nuclei form spontaneously in
the matrix after a so-called nucleation time t
0,
and these nuclei then proceed to grow steadily
as shown by the linear relationship.
The driving force for the process is provided by
the stored energy of cold work contained in the
strained grain on one side of the boundary
relative to that on the other side.
Such an interpretation would suggest that the
recrystallization process occurs in two distinct
stages, i.e. first nucleation and then growth.
During the linear growth period the radius of a
nucleus is R= G(t-to), where G, the growth rate, is dR/dt assuming the nucleus is spherical, the volume of the recrystallized nucleus is
If the number of nuclei that form in a time
increment dt is N dt per unit volume of
unrecrystallized matrix, and if the nuclei do not impinge on one another, then for unit total volume
This equation is valid in the initial stages when
f<<1. When the nuclei impinge on one another the
rate of recrystallization decreases and is related
to the amount untransformed (1 – f) by
where, for short times, equation reduces to last
equation.
This Johnson–Mehl equation is expected to apply
to any phase transformation where there is
random nucleation, constant N and G and small
t
0.
In practice, nucleation is not random and the rate
not constant so that equation will not strictly apply.
For the case where the nucleation rate decreases
exponentially, Avrami developed the equation
where k and n are constants, with n≈ 3 for a fast and n≈ 4 for a slow, decrease of nucleation rate.
Provided there is no change in the nucleation
mechanism, n is independent of temperature but k is very sensitive to temperature T; clearly from
The process of recrystallization may be pictured as
follows.
After deformation, polygonization of the bent lattice
regions on a fine scale occurs and this results in the formation of several regions in the lattice where the strain energy is lower than in the surrounding matrix; this is a necessary primary condition for nucleation.
During this initial period when the angles between the
grains are small and less than one degree, the sub-grains form and grow quite rapidly.
However, as the sub-grains grow to such a size that the
angles between them become of the order of a few degrees, the growth of any given sub-grain at the expense of the others is very slow.
Eventually one of the sub-grains will grow to such a
size that the boundary mobility begins to increase with increasing angle. A large angle boundary, θ≈(30– 40°), has a high mobility because of the large lattice irregularities or ‘gaps’ which exist in the boundary transition layer.
The atoms on such a boundary can easily transfer
their allegiance from one crystal to the other.
This sub-grain is then able to grow at a much faster
rate than the other subgrains which surround it and so acts as the nucleus of a recrystallized grain.
The further it grows, the greater will be the difference in orientation between the nucleus and the matrix it meets and consumes, until it finally becomes recognizable as a new strain-free crystal separated from its surroundings by a large-angle boundary.
Whether recrystallization nucleus it grows to become a strain-free grain depends on three factors:
(1) the stored energy of cold work must be sufficiently high to provide the required driving force,
(2) the potential nucleus should have a size advantage over its neighbours
(3) it must be capable of continued growth by existing in a region of high lattice curvature (e.g.transition band) so that the growing nucleus can quickly achieve a high- angle boundary.
Prior deformation, for example, will control the extent to which a region of the lattice is curved.
The larger the deformation, the more severely will
the lattice be curved and, consequently, the smaller will be the size of a growing sub-grain when it
acquires a large-angle boundary.
The importance of impurity content on recrystallization temperature is also evident from the effect impurities have on obstructing dislocation sub-boundary and grain boundary mobility.
The intragranular nucleation of strain-free grains, as discussed above, is considered as abnormal subgrain growth, in which it is necessary to specify that some sub-grains acquire a size advantage and are able to grow at the expense of the normal subgrains.
It has been suggested that nuclei may also be formed by a process involving the rotation of individual cells so that they coalesce with neighbouring cells to produce larger cells by volume diffusion and dislocation rearrangement Segregation of solute atoms to, and precipitation on, the
grain boundary tends to inhibit intergranular nucleation and gives an advantage to intragranular nucleation,
provided the dispersion is not too fine.
Small, finely dispersed particles retard recrystallization by reducing both the nucleation rate and the grain
boundary mobility, whereas large coarsely dispersed particles enhance recrystallization by increasing the nucleation rate.
When primary recrystallization is complete (i.e.
when the growing crystals have consumed all the strained material) the material can lower its energy further by reducing its total area of grain surface.
In this stage the tensile strength and hardness
continue to decrease but at a much less rate than the recrystallization stage.
With extensive annealing it is often found that grain
boundaries straighten, small grains shrink and larger ones grow.
The general phenomenon is known as grain growth;
the major change observed during this stage is the growth of the grain boundaries and reaching the original grain size .
The most important factor governing the process is the surface tension of the grain boundaries.
A grain boundary has a surface tension, T (= surface-free energy per unit area) because its atoms have a higher free energy than those within the grains.
Consequently, to reduce this energy a polycrystal will tend to minimize the area of its grain boundaries.
Second-phase particles have a major inhibiting effect on boundary migration and are particularly effective in the control of grain size.
The pinning process arises from surface tension forces exerted by the particle–matrix interface on the grain boundary as it migrates past the particle.
Nevertheless, even after growth has finished the grain
size in a specimen which was previously severely cold-worked remains relatively small, because of the large number of nuclei produced by the working treatment.
Exaggerated grain growth can often be induced,
however, in one of two ways, namely:
(1) by subjecting the specimen to a critical strain- anneal treatment, or
By applying a critical deformation (usually a few
per cent strain) to the specimen the number of
nuclei will be kept to a minimum, and if this strain is followed by a high-temperature anneal in a
thermal gradient some of these nuclei will be made more favourable for rapid growth than others.
The only driving force for secondary
recrystallization is the reduction of grain
boundary-free energy, as in normal grain growth. Special conditions are necessary. In some regions the grain boundaries become free (e.g. if the
inclusions slowly dissolve or the boundary tears
away) and as a result the grain size in such regions becomes appreciably larger than the average.
EFFECT OF ANNEALING ON TENSILE
STRENTH, HARDNESS, DUCTILITY AND
GRAIN SIZE
Annealing results in modification of the properties of
steel. The figure in the next slide shows the change in some of these properties on annealing.
During the recovery stage the decrease in stored energy
and electrical resistivity is accompanied by only a slight lowering of hardness.
The greatest simultaneous change in properties occurs
during the primary recrystallization stage.
During recrystallization stage, there is a significant drop
in tensile strength, hardness and a large increase in the ductility of the material which is illustrated in the
EFFECT OF ANNEALING ON TENSILE STRENTH,
HARDNESS, DUCTILITY AND GRAIN SIZE
The benefits of annealing are:
Improved ductility
Removal of residual stresses that result
from cold-working or machining
Improved machinability Grain refinement
At the same time, annealing also has a
few disadvantages as it
reduces the
hardness, yield strength and tensile
strength of the steel
Full annealing is the process by which the
distorted cold worked lattice structure is
changed back to one which is strain free through the application of heat. This process is carried out entirely in the solid state and is usually
followed by slow cooling in the furnace from the desired temperature.
The austenitising temp is a function of carbon
content of the steel and can be generalised as:
For hypoeutectoid steels and eutectoid steel
Ac3+(20-40oC) [to obtain single phase
austenite]
For hypereutectoid steels
Ac1+(20-40oC) [to obtain austenite+
To refine the grain size of steel castings, or
of hot worked steels to improve the
mechanical properties.
To soften the steel
To relieve internal stresses
To improve machinability
It also reduces some defects like aligned
The three important parts of full annealing are:
Proper austenitising temperature
Soaking time
Very slow cooling through A1(critical
temperature)
The formation of austenite destroys all structures
that have existed before heating. Slow cooling
yields the original phases of ferrite and pearlite in
hypoeuetectoid steels and that of cementite and
pearlite in hypereutectoid steels.
The subsequent heating, soaking and hot working
homogenises the structure to a large extent
since diffusion of C is very fast at high temp and the simultaneous plastic deformation breaks the dendrites with different portions moving in
relation to each other which also facilitate diffusion.
The main aim of homogenising annealing is to
make the composition uniform, i.e to remove chemical heterogeneity
The alloy steel ingots are homogenised at higher
temp like 1150-1200oC for 10-20 hours as the diffusion of substitutional solid solution forming elements is very low
The impact energy and ductility of the steel
increase as the homogenizing temperature
increases and the hardness, yield strength
and tensile strength decrease with an
increase in the homogenizing temperature.
Homogenising annealing has a few
shortcomings as well. It results in:
Grain coarsening of austenite, thereby impairing the
properties
Thick scales on the surface of steels It is an expensive process
Also known as process annealing or sub-critical
annealing.
Recrystallization annealing consists of heating the cold
worked steel above its recrystallization temp, soaking at this temp, and then cooling thereafter.
Process Annealing is used to treat work-hardened
parts made out of low-Carbon steels (< 0.25%
Carbon). This allows the parts to be soft enough to undergo further cold working without fracturing.
The main aims of recrystallization annealing are:
To restore ductility
To refine coarse grains
To improve electrical and magnetic properties in
No phase change takes place and the final
structure consists of strain-free, equiaxed grains of fine ferrite produced at the expense of deformed elongated ferrite grains.
Recrystallization temp(Tr) is given by:
Tr= (0.3-0.5)Tm.p
As little scaling and decarburisation occurs in
recrystallization annealing, it is preferred over full annealing.
However It would produce very coarse grains if the
steel has undergone critical amount of
deformation. In such cases, full annealing is preferred.
o Spheroidisation annealing consists of heating, soaking and
cooling, invariably very slowly to produce spheroidal pearlite or globular form of carbides in steels.
o Hypereutectoid steels consist of pearlite and cementite.
The cementite forms a brittle network around the pearlite. This presents difficulty in machining the hypereutectoid steels.
o To improve the machinability of the annealed
hypereutectoid steel spheroidize annealing is applied.
o This process will produce a spheroidal or globular form of a
carbide in a ferritic matrix which makes the machining easy.
o Prolonged time at the elevated temperature will completely
break up the pearlitic structure and cementite network. The structure is called spheroidite.
SPHEROIDISING METHODS:
Heating the steel to just below Ac1 temp,
holding at this temperature for a very long period followed by slow cooling.
Austenitise the steel at a temp not more than
50oC above A
1, and cool very slowly through
A1 to transform this inhomogeneous austenite at a temperature not more than 50oC below
Spheroidizing process applied at a temperature
below the LCT.
Spheroidizing process applied at a temperature below and above the LCT.
ADVANTAGES OF SPHEROIDISATION
ANNEALING:
minimum hardness
maximum ductility
maximum machinability
maximum softness
Commonly used for alloy steels to soften the
steel
In isothermal annealing, steel austenitised at
a temp 20-40
oC above Ac
3
, is cooled quickly
to the temp of isothermal holding( which is
below A
1temp in the pearlitic range), held
there for the required time so that complete
transformation of austenite takes place and
then normally cooled in air
The closer the temp of isothermal holding is to A1, coarser is the pearlite, softer is the steel, but
ADVANTAGES OF ISOTHERMAL
ANNEALING:
As cooling can be done in air after the
transformation is complete, total time of
heat treatment and the cost of annealing is
less.
The productivity of the furnace is high
The microstructure obtained is more uniform
and thus better control over hardness can be
obtained as transformation takes place at a
constant temperature.
Improved machinability with good surface
finish
Reduced warping in subsequent hardening
• Stress-Relief Annealing is useful in removing
residual stresses due to heavy machining or other cold-working processes.
• It is usually carried out at temperatures below
the LCT, which is usually selected around 1000oF.
• The main aims of stress-relief annealing are:
• To relieve the internal stresses, and thus, allow
higher external loads to be applied
• Increase fatigue life and prevent intercrystalline
corrosion.
• To reduce chance of warpage or cracking, or risk
of distortion in cracking
• To increase impact resistance and lower
susceptibility to brittle fracture
NORMALIZING
The normalizing of steel is carried out by heating
above the UCT (Upper Critical Temperature) to single phase austenitic region to get
homogeneous austenite, soaking there for some time and then cooling it in air to room
temperature.
The austenitising temperature range are:
For hypoeutectoid steels and eutectoid
steel
•
A
c3+ (40-60
oC)
For hypereutectoid steels
During normalising we use grain refinement
which is associated with allotropic
transformation upon heating γ→α
Parts that require maximum toughness and
those subjected to impact are often
normalized.
When large cross sections are normalized,
they are also tempered to further reduce
stress and more closely control mechanical
properties.
The microstructure obtained by normalizing
depends on the composition of the castings
(which dictates its hardenability) and the
cooling rate.
Figure below shows the normalizing temperatures for hypoeutectoid and hypereutectoid steels
To produce a harder and stronger steel than
full annealing
To improve machinability
To modify and/or refine the grain structure
To obtain a relatively good ductility without
reducing the hardness and strength
Improve dimensional stability
Produce a homogeneous microstructure
Reduce banding
Provide a more consistent response when
EFFECT OF SOAKING TIME ON THE
MICROSTRUCTURE:
MICROSTRUCTURE AT THE STRIP SURFACE
NORMALIZED AT 860
oC
MICROSTRUCTURE AT THE STRIP SURFACE
NORMALIZED AT 900
oC
MICROSTRUCTURE AT THE STRIP SURFACE
NORMALIZED AT 940
oC
MICROSTRUCTURE AT THE STRIP SURFACE
NORMALIZED AT 960
oC
COMPARISON OF ANNEALING AND
NORMALIZING
The metal is heated to a higher temperature and then
removed from the furnace for air cooling in normalizing rather than furnace cooling.
In normalizing, the cooling rate is slower than that of a
quench-and-temper operation but faster than that used in annealing.
As a result of this intermediate cooling rate, the parts will
possess a hardness and strength somewhat greater than if annealed.
Fully annealed parts are uniform in softness (and
machinability) throughout the entire part; since the entire part is exposed to the controlled furnace cooling. In the case of the normalized part, depending on the part geometry, the cooling is non-uniform resulting in non-uniform material
properties across the part.
Internal stresses are more in normalizing as compared to
annealing.
Grain size obtained in normalizing is finer than in annealing. Normalizing is a cheaper and less time-consuming process.
The slower cooling of annealing results in higher temperature transformation to ferrite and pearlite and coarser
Annealing and normalizing do not present a significant difference on the ductility of low carbon steels. As the carbon content increases, annealing maintains the %
elongation around 20%. On the other hand, the ductility of the normalized high carbon steels drop to 1 to 2 % level.
The tensile strength and the yield point of the
normalized steels are higher than the annealed steels.
Normalizing and annealing do not show a
significant difference on the tensile strength and yield point of the low carbon steels.
However, normalized high carbon steels
present much higher tensile strength and yield point than those that are annealed. This can be illustrated from the figures.
Low and medium carbon steels can maintain similar
hardness levels when normalized or annealed. However, when high carbon steels are normalized they maintain higher levels of hardness than those that are annealed.
ADVANTAGES OF NORMALIZING OVER
ANNEALING
Better mechanical properties
Lesser time-consuming
Lower cost of fuel and operation
ADVANTAGES OF ANNEALING OVER
NORMALIZING
Greater softness
Complete absence of internal stresses which
It is the process of heating the steel to
proper austenitizing temperature , soaking at
this temperature to get a fine grained and
homogeneous austenite , and then cooling
the steel at a rate faster than its critical
cooling rate.
The aims of hardening are:
1. Main aim of hardening is to induce high
hardness. The cutting ability of a tool is proportional to its hardness.
2. Many machine parts and all tools are hardened
to induce high wear resistance higher is the hardness , higher is the wear and the abrasion resistance .For example ,gears, shaft.
3. The main objective of hardening machine
components made of structural steel sis to develop high yield strength with good
toughness and ductility to bear high working stresses.
The steel is first heated to proper austenising
temperature to obtain a homogeneous and
fine grained austenite. This temperature
depends on the composition(carbon as well
as alloying elements).
The austenitising temperature of plain
carbon steels depends on the carbon content
of the steel and is generalised as :
For hypo-eutectoid steels :Ac3 + (20 — 40°C)
For hyper-eutectoid steels and eutectoid
Hypereutectoid steels, when heated in above
temperature range, to obtain homogeneous and fine-grained austenite which on quenching
transforms to fine-grained (very fine
needles/plates), and hard martensite as is desired to be obtained.
Heating these steels only up to critical range
(between Ac3 and Ac1) is avoided in practice.
Steel then has austenitic and ferrite.
On quenching, only austenite transforms to
martensite, and ferrite remains as it is, i.e., incomplete hardening occurs .
The presence of soft ferrite does not permit to
If the aim is to get high strength by the process
of tempering ferrite does not permit this as it has low tensile and yield strengths .
In fact, ferrite forms the easy path to fracture. Quenching of hypoeutectoid steels from
temperatures much above the proper
temperatures , when austenite has become coarse, results in coarse acicular form of martensite.
Coarse martensite is more brittle, and a unit or
two lower in hardness. It lowers the impact strength even after tempering, and is more prone to quench-cracking.
Hypereutectoid steels, when heated in the above
range, i.e., just above Ac1 have fine grains of austenite and proeutectoid cementite.
On quenching austenite transforms to fine
martensite and cememtite remains unchanged.
As the hardness of cementite (≈ 800 BHN) is
more than that of martensite (650-750 BHN), its presence increases the hardness, wear and
abrasion resistance as compared to only martensitic structure.
If temperature of austenitisation is much higher
than Ac1 but still below Acm temperature, a part of proeutectoid cementite gets dissolved to
increase the carbon content of austenhlc(> 0.77%)
On quenching as-quenched hardness is less,
because :
1.
Lesser amount of proeutectoid cementite is
present.
2.
Larger amount of soft retained austenite is
obtained as the dissolved carbon of
cementite has lowered the Ms and Mf
temperature.
3.
A bit coarser martensite has lesser
Heating hypereutectoid steels to a temperature
higher than Acm results in 100% austenite . It is very coarse austenite as very rapid grain-growth occurs due to dissolution of restraining
proeutectoid cementite .
The as-quenched hardness is low because of:
1) Absence of harder cementite.
2) As more carbon has dissolved in austenite, more
retained austenite is obtained.
3) Coarser martensite is a bit less hard and more
brittle.
Thus, these temperatures are avoided in carbon
When a heated steel object (say at 840°C) is plunged into a stationary bath of cold it has three stages as: Stage A -vapour-blanket stage:
Immediately on quenching, coolant gets vapourized as the steel part is at high temperature, and thus, a continuous vapour- blanket envelopes the steel part. Heat escapes from the hot surface very slowly by
radiation and conduction through the blanket of water vapour.
Since the vapour-film is a poor heat conductor, the cooling rate is relatively low (stage A in fig ). This long stage is undesirable in most quenching
Stage B-Intermittent contact stage (Liquid-boiling stage):
Heat is removed in the form of heat of
vaporization in this stage as is indicated by the steep slope of the cooling curve.
During this stage, the vapour-blanket is broken
intermittently allowing the coolant to come in contact with the hot surface at one instant, but soon being pushed away by violent boiling action of vapour bubble.
The rapid cooling in this stage soon brings the
metal surface below the boiling point of the coolant.
The vaporization then stops. Second stage
corresponds to temperature range of 500◦ to 100◦c , and this refers to nose of the CCT curve of the steel , when the steel transforms very rapidly ( to non martensite product ).
Thus, the rate of cooling in this stage is of great
importance in hardening of steels.
Stage C-Direct-Contact stage (Liquid-cooling stage):
This stage begins when the temperature of steel
surface Is below the boiling point of coolant.
Vapours do not form. The cooling is due to
convection and conduction through the liquid. Cooling is slowest here.
As the aim is to get martensite, the coolant should have quenching power to cool austenite to let it
transform to martensite. The following factors effect the quenching power of the coolant :
The cooling rate decreases as the temperature of
water and brine increases, i.e., it increases stage ‘A’, i.e., helps in persistence of the vapour blanket stage. The increased temperature brings it closer to its
boiling point, and thus, requires less heat to form vapour, specially above 60°C.
Good range of temperature for water as coolant is 20-40°C.
Oils in general, show increased cooling rates with the rise of temperature, with optimum cooling rates in range 50°—80°C.
In oils, the increase of temperature increases the
persistence of vapour-blanket, but this resulting decrease in the cooling rate is more than compensated by the
decrease of viscosity (with the rise in temperature) to
result in increase of rate of heat removal through the oil.
If the boiling point of a coolant is low, vapours form easily
to increase the ‘A’ stage of cooling. ¡t is better to use a coolant with higher boiling point. A coolant with low specific heat gets heated up at a faster rate to form vapours easily.
A coolant with low latent heat of vapourisation changes
into vapour easily to promote ‘A’ stage, i.e., decreases the cooling rate.
A coolant with high thermal conductivity increases the
cooling rate. Coolants with low viscoity provide faster cooling rates and decrease the ‘A’ stage.
A coolant should be able to Provide rate of cooling fast enough to avoid transformation of austenite to pearlite and bainite . Plain carbon steel invariably require çooling in water or brine. Whereby alloy steels are quenched normally in oils.
But milder the cooling medium , lesser the internal stresses developed , and thus lesser the danger of distortion , or cracks . An ideal quenching medium is one which is able to provide very fast cooling rate near the nose of the curve ( 650 -550°C)and at the same time it should provide very considerable slower rate if cooling within the range of martensitic
transformation( 300 - 200°C) to minimize internal stresses .
The oldest and still the most popular quenching medium,
water meets the requirements of low cost ,general easy availability, easy handling and safety.
The cooling characteristics change more than oil with the
rise of temperature, specially there is a rapid fall in cooling capacity as the temperature rises above 60°C, because of easy formation of vapour-blanket.
The optimum cooling pover is when water is 2O-4O°C. Thc cooling power of water is between brine and oils. Water provides high cooling power to avoid the
transformation of austenite to pearlite/bainite, but the major draw back is that it also provides high cooling rate in the the temperature range of martensite formation.
At this stage, the steel is simultaneously under the
influence of structural stresses (non-uniform change in structure) and thermal stresses which increase the risk of crack formation.
Sodium chloride aqueous solutions of about 10% by weight
are widely used and are called brines.
The cooling power is between 10% NaOH aqueous solution
and water.
These are corrosive to appliances.
The greater cooling efficiency of brines, or other aqueous
solutions is based as :
In brine heating of the solution at the steel surface causes
the deposition of crystal of the salt on hot steel surface .
This layer of solid crystals disrupts with mild explosive
violence, und throws off a cloud of crystals. This action destroys the vapour-film from the surface, and thus permits direct contact of the coolant with the steel surface with an accompanying rapid removal of heat.
Brines are used where cooling rates faster than water arc
Oils have cooling power between water at 40°C to water at
90°C.
In oil-quench, considerable variation can be obtained by the use
of animal, vegetable, or mineral oil, or their blends.
Oils should be used at 50- 80°C when these are more fluid, i.e
less VISCOUS, which increases the cooling power.
As the oils used generally have high boiling points, moderate
increase of temperature of oil does not very much increase the vapour blanket stage. However, oils in contrast to water, or brine, have much lower quenching power .
Its this relatively slow cooling rate in the range of martensitic
formation is atlvantageous as it helps in minimsing the danger to crack formation.
Oils with high viscosity are less volatile, and thus have decreased
vapour-blanket stage (increase thecooling rate). As lesser volatile matter is lost, their cooling power is not affected much with use.
polymer quenchants cool rapidly the heated
steel to Ms temperature, and then rather slowly when martensite is forming .
Polymer quenchants are water-soluble organic
chemicals of high ,molecular weights, and are
generally polyalkylene glycol-based, or polyvinyl pyrolidene-based.
Widely different cooling rates can be obtained
by varying the concentration of Organic additives in water; higher the additions, slower is the
cooling rate of solution.
There are little dangers of distortions and
It is an ideal quenching medium for a steel of
not very large section but with good
hardenabilty.
Addition of O.3-O.5’% water almost doubles
the cooling capacity. Normally holding time is
2-4 minutes/cm of section thickness.
Salt baths used for austenitising keep the
internal stresses are produced due to non-uniform plastic deformation. In quenching of steels ,this may be caused by thermal stresses, structural stresses, or both, or even premature failure of part in service. Cooling during quenching lakes place non-uniformly,
i.e., causes temperature gradient across the section. Surface layers contract more than the central
portion.
Contraction of surface is resisted by the central
portion, and this puts the central portion under the compressive stresses, and the surface layers in
tension .
If the magnitude of stress becomes more than the yield stress of steel (at that deformation occurs. These stresses that develop in a quenched part as a
Structural stresses are the stresses which develop due to
due to phase change (mainly austenite to martensite), and at different times.
Structural stresses are developed due to two main reasons:
1. Austenite and its transformation products have unequal specific volume i.e. change in volume occurs when
transformation occurs.
2. Phase changes occur at different times in the surface and in centre.
Under right conditions, both types of stresses get
superimposed to become larger than the yield strength to cause warping. but when the tensile internal stresses
become larger than the tensile strength cracks appear.
If an austenitised steel is quenched, it contracts thermally
figure(a) illustrates this in stage 1 As surface cools faster than centre, i.e., contracts more than centre distribution of stresses across the section is
illustrated in fig (b), i.e, the surface is under tensile nature of stress, while centre is under compressive stresses.
Only thermal Stresses are produced in stage 2 ,
surface having attained Ms temperature, transforms to martenSite, and thus expands, while the centre is still contracting as it is getting cooled.
In stage Il, centre may get plastically deformed ,as it is still ductile austenite.
In stage 3, martensite of surface and austenite of centre continue contracting leading to slight
In stage IV, centre has attained M5 temperature,
and begins to expand as it forms martensite, while surface is still contracting.
The centre, as it expands, puts the surface in
higher stress levels .
The surface has little deformation as it consists
of brittle martensie.
It is during this stage, the greatest danger of
cracking exists.
Thus, stress levels are highest not in the
beginning of the quench, but when the centre is transforming to martensite.
However, higher is the Ms temperature of the
steel, lesser is the expansion, there is reduced danger of quench-cracking.
Increase of carbon and alloying elements lower
the Ms temperature making the steel more prone to quench cracking.
Martensitic transfomiation is essentially an athermal
transformation.
Austenite begins to transform to martensite at Ms, and the
amount of martensite formed increases as the temperature decreases to complete at Mf temperature.
Less than 1 % of austenite may not transform because of
unfavourable stress conditions.
The Ms and Mf temperatures are lowered as the amounts of
carbon content and alloying elements(except cobalt and aluminium) increase in the steel.
In a quenched steel, the amount of martensite formed depends
on the location of Ms and Mf and the temperature of the coolant (which is normally room temperature. As long as room
temperature lies between Ms and Mf temperatures, austenite does not transform completely to martensite as it has not been cooled below Mf temperature.
10% retained austenite is normally desirable as
its ductility relieves some internal stresses developed during hardening. This reduces the danger to distortion or cracks.
The presence of 30-40% retained austenite
makes straightening of components possible after hardening.
Non distorting tools owe their existence to
retained austenite . It tries to balance
transformational volume changes during heating as well as cooling cycles of heat treatment to
1) The presence of soft austenite decreases the
hardness of hardened steels.
2) As retained austenite may transform to lower
bainite, or martensite later in service ,increase in dimensions of the part occurs.
3) This creates problems in precision gauges or
dies.
4) Stresses may develop in the part itself as well
as in adjacent pans. Grinding-cracks are mainly due to retained austenite transforming to
martensite.
5) Austenite is non-magnetic, decreases the
Retained austenite is generally undesirable. It is eliminated by one of the methods:
1. Sub-Zero treatment (cold treatment):
It Consists of cooling the hardened steel (having
retained austenite) to a temperature below 0°C or its Mf temperature.
There is no reason to cool a steel much below its Mf temperature.
Sub-zero treatment is more effective if it is done
immediately after the quenching operation (normally done to room temperature).
Sub-zero treatment is done in a low
temperature-cooling unit, which consists of double-walled vessel. The interior is made of copper in which the parts to
be deep-frozen are kept, and the exterior is made of steel provided with good heat-insulation.
The space in between the vessels is filled with
some coolant.
The sub-zero coolants could be, dry ice (Solid
CO2) + acetone (— 78°C); Ice + NaCl (—23°C); liquid air (—183°C); liquid N2 (— 196°C); Freon (— 111°C).
Total time of cooling in this unit is 1/2 to 1 hour. As this treatment transforms austenite to
martensite, steels after sub-zero treatment have high hardness, wear and abrasion resistance, and have no danger of grinding-cracks.
The stresses are increased further and thus,
tempering should be done immediately after sub-zero treatment.
Carburised steels, ball-bearing steels, highly
alloy tool steels, are normally given cold treatment.
2. Tempering:
The second stage of tempering eliminates the
retained austenite in most steels.
In high alloy steels, multiple tempering is able
to eliminate the retained austenite during cooling from the tempering temperature.
The main defects produced during hardening
are:
1. Mechanical properties not up to specifications:
The common defect in hardened tools, or
component is too low a hardness.
One or more of the followings could be the
cause of such a defect.
Insufficient fast cooling due to overheated or
even polluted coolant could be responsible for a defect.
The presence of scale, or oil, etc. on the surface
also decreases the cooling rate.
Circulation of coolant, or agitation of component
A shorter austenitising time can also cause such a defect. Lower austenitising temperature can also result in such a defect.
Decarburisation can also result in low surface
hardness. If too high temperature had been used, which produces larger amount of retained austenite can result in low surface hardness.
2. Soft Spots:
Soft areas on the hardened surface are called ‘soft spots’.
The adhering scale, or decarburisation of some areas or prolonged vapour-blanket stage due to overheated coolant or insufficient agitation or circulation of
coolant, or rough surface could cause presence of soft spots surface.
3. Quench Cracks:
Quench cracks form as a result of internal stresses
developed of tensile nature exceeding the tensile strength of the steel.
Steel with lower Ms temperature due to higher contents of
alloying elements are more prone to quench cracks. Higher carbon also results in more brittle martensite.
Quench cracks can form if there is more time lag between
the process of quenching and tempering.
Overheating of steel or a wrong coolant which gave much
faster rate of cooling, or there is faulty design of the component with sharp corners and sharp transition between sections, or a wrong steel has been chosen.
Presence of large amounts of retained austenite causes
grinding cracks.
The other defects could be distortion and warpage; change
Martensite is a very strong phase, but it is
normally very brittle so it is necessary to modify the mechanical properties by heat, treatment in the range 150—700°C.
Essentially, martensite is a highly Supersaturated
solid solution of carbon in iron which, during tempering, rejects carbon in the form of finely divided carbide phases.
The end result of tempering is a fine dispersion
of carbides in an α-iron matrix which often bears little structural similarity to the original
In the as-quenched martensite structure,the
laths or plates are heavily dislocated to an
extent that individual dislocations are very
difficult to observe in thin-foil electron
micrographs.
A typical dislocation density for a 0.2 wt%
carbon steel is between 0.3 and 1.0 x 1012
cm cm-3. As the carbon content rises above
about 0.3 wt%, fine twins about 5—10 nm
wide are also observed.
Often carbide particles, usually rods or small
These occur in the first-formed martensite,
i.e. the martensite formed near Ms, which
has the opportunity of tempering during the
remainder of the quench.
This phenomenon, which is referred to as
autó-tempering, is clearly more likely to
occur in steels with a high Ms.
On reheating as-quenched martensite, the
tempering takes place in four distinct but overlapping stages:
Stage 1, up to 250°C — precipitation of E-iron
carbide; partial loss of tetragonality in martensite.
Stage 2, between 200 and 300°C —
decomposition of retained austenite .
Stage 3, between 200 and 350°C — replacement
of &iron carbide by cementite; martensite loses tetragonality.
Stage 4, above 350°C — cementite coarsens and
Martensite formed in medium and high carbon
steels (0.3—1.5 wt% C) is not stable at room temperature because interstitial carbon atoms can diffuse in the tetragonal martensite lattice at this temperature.
This instability increases between room
temperature and 250°C, when €-iron carbide precipitates in the martensite (Fig. 9.2).
This carbide has a close-packed hexagonal
structure, and precipitates as narrow laths or rodlets on cube planes of the matrix with a well-defined
At the end of stage 1 the martensite still possesses a
tetragonality, indicating a carbon content of around 0.25 wt%.
It follows that steels with lower carbon contents are
unlikely to precipitate €-carhide.
This stage of tempering possess an activation energy of
between 60 and 80 kJ mo1, which is in the right range for diffusion of carbon in martensite. The activation energy has been shown to increase linearly with the carbon
concentration between 0.2 and 1.5 wt% C.
This would be expected as increasing the carbon
concentration also increases the occupancy of the preferred
interstitial sites, i.e. the octahedral interstices at the mid-points of cell edges, and centres of cell faces, thus
During stage 2. austenite retained during
quenching is decomposed usually in the temperature range 230-300°C.
In martensitiC plain carbon steels below 0.5
carbon. the retained austenite is often below 2%, rising to around 6 % at 0.8 wt C and over 30 % at 1.25 wt C.
The little available evidence suggests that in the
range 230-300°C, retained austenite decomposes to bainitic ferrte and cementite, but no
detailed comparison between this phase and lower bainite has yet been made.
During the third stage of tempering, cementite
flrst appears in the microstructure as a
Widmanstatten distribution of plates which have a well-defined orientation relationship with the matrix which has now lost its tetragonality and become ferrite.
This reaction commences as low as 100°C and is
fully developed at 300°C, with particles up to 200 nm long and 15 nm in thickness.
Similar structures are often observed in lower
carbon steels as quenched, as a result of the formation of Fe3C during the quench.
During tempering, the most likely sites for the
nucleation of the cementite are the €-iron
carbide irterfaces with the matrix (Fig 9.2) and as the Fe3C particles grow, the €-iron carbide particles gradually
disappear.
The twins occurring in the higher carbon
martensites are also site for the nucleation and growth of cementite which tends to grow along the twin boundaries forming colonies of similarly oriented lath shaped particles (Fig. 9.3) which can be readily ditinguished from the normal
A third site for the nucleation of cementite is
the grain boundary regions (Fig, 9.4)of both the interlath boundaries of martensite and
the original austenite grain b0unjaries.
The cementite can form as very thin films
which are difficult to detect but which gradually sp1eroidise to give rise to welI-defined particles of Fe3C in the grain boundary regions.
There is some evidence to show that these.
boundary cementite films can adversely affect ductility. However it can be modified by
During the third stage of tempering , the
tetragonality of thc matrix disappears and it
is then, essentially, ferrite, not
supersaturated with
respect to carbon.
Subsequent changes in the morpriology of
cementite particles occur by process where
the smaller particles dissolve in the matrix
providing carbon for the selective growth of
the larger particles.
During the third stage of tempering , the
tetragonality of thc matrix disappears and it
is then, essentially, ferrite, not
supersaturated with
respect to carbon.
Subsequent changes in the morpriology of
cementite particles occur by process where
the smaller particles dissolve in the matrix
providing carbon for the selective growth of
the larger particles.
It is useful to define a fourth stage of tempering in which the cementite particles undergo a coarsening process and essentially lose their crystallographic morphology, becoming spheroidized.
It commences between 300 and 400◦C, while spheroidizatiun takes place increasingly up to 700◦C.
At the higher end of this range of tempera.
ture the martensite lath boundaries are replaced by more equi-axid fèrrite grain boundaries by a process which is best described as recrystallization.
The final result is an equi-axed array of ferrite grains with coarse spheroidized particles of Fe3C (Fig. 9.5), partly, but not exclusively, by the grain boundaries.
The spherodisation of the Fe3C is encouraged by the
resulting decrease in surface energy.
The particles which preferentially grew and spheroidize
are located mainly at interlath boundaries and prior
austenite boundaries, although some particles remain in the matrix.
The boundary sites are preferred because of the greater
ease of diffusion in these regions. Also, the growth of cementite into ferrite is associated with a decrease in density so vacancies are required to accommodate the growing cementite.
Vacancies will diffuse away from cementite particles which
are redissolving in the ferrite and towards cementite particles which are growing, so that the rate controlling process is likely to be the diffusion of vacancies.
The original martensite lath boundaries remain stable up
to about 600°C, but in the range 350—600°C. there is
considerable rearrangement of the dislocations within the laths and at those lath boundaries which are essentially low angle boundaries.
This leads to a marked reduction in the dislocation density
and to lath-shaped ferritic grains closely related to the packets of similarly oriented laths in the original
martensite.
This process, which is essentially one of recovery, is
replaced between 600 and 700°C by recrystallization
which results in the formation of equi-axed ferrite grains with spheroidal Fe3C particles in the boundaries and
