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

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 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

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 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

structure.

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

Lead 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

Variation in free energy during grain growth.

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

temperature.

The classical interpretation of these

observations is that nuclei form spontaneously in

the matrix after a so-called nucleation time t0,

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

t0.

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

equation k=∏NG3/3 and both N and G depend on T.

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

sub-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

(2) by a process of secondary recrystallization.

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

figure

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+ cementite]

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

sulphide inclusions, or bands in steels.

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

grain-oriented Si steels.

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 A1, and cool very slowly through

A1 to transform this inhomogeneous austenite

at a temperature not more than 50oC below

A1 temperature

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 20-40oC above Ac3, is cooled quickly

to the temp of isothermal holding( which is

below A1 temp 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

longer is the time of isothermal transformation.

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 processes.

• 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

• No change of dimensions in service life

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

• Ac3 + (40-60oC)

For hypereutectoid steels

• Acm + (30-50oC)

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

hardening or case hardening

EFFECT OF SOAKING TIME ON THE

MICROSTRUCTURE:

MICROSTRUCTURE AT THE STRIP SURFACE NORMALIZED AT 860oC

MICROSTRUCTURE AT THE STRIP SURFACE NORMALIZED AT 900oC

MICROSTRUCTURE AT THE STRIP SURFACE NORMALIZED AT 940oC

MICROSTRUCTURE AT THE STRIP SURFACE

NORMALIZED AT 960oC

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 microstructures than does normalizing.

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

is a necessity in complex and intricate parts

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 steel:Ac1 + (20 — 40°C)

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 achieve high hardness, if that is the objective.

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

hardness.

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 steels

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 operations.

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 requited.

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 cracks.

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

steel clean.

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 result of unequal cooling are called thermal stresses.

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 till Ms temperature is attained .

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 increase in stress levels.

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.

This untransformed austenite is retained austenite.

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

produce little overall change in size of the tools.

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 magnetic properties of the steels.

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 may also result in such defect.

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 in dimensions; oxidation and decarburisation

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 as-

quenched martensite.

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 plates, are observed (Fig. 9.1).

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 spheroidizes; recrystallization of ferrite.

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 orientation relationship .

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 reducing the mobility of C atoms.

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 Widmanstatten habit.

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 addition of alloying elements.

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

within the grains.

This process occurs most readily in carbon

steels.

At higher carbon content, the increased

density of cementite particles is much more

effective in pinning the ferrite boundaries, so

recrystallisation is much more sluggish.

The final process is the continued coarsening of

the cementite particles and gradual ferrite grain

growth.

Firstly, the hardness of the as-quenched martensite is largely influenced by the carbon content, as is the morphology of the martensite laths which have a up to 0.3 wt% C, changing at higher carbon contents.

The Ms temperature is reduced as the carbon content increases, arid thus the probability of the occurrence of auto-tempering is less.

During fast quenching in alloys with less than 0.2 wt % C, the majority (up to 90%) of the carbon segregates to dislocations and lath boundaries, but with slower quenching some precipitation of cementite occurs.

On subsequent tempering of low carbon steels up to 200°C further segregation of carbon takes place. but no precipitation has been observed.

Under normal circumstances it is difficult to detect any tetragonality in martensitic in steels with less than 0.2 wt% C, a fact which can explained by the rapid segregation of carbon during quenching.

The hardness change; during tempering are also very dependent on carbon content, as shown in figure for steels up to 0.4 wt% C.

Above this concentration, an increase in hardness has been observed in temperature range 50—150°C, as €-carbide precipitation strengthens the martensite.

However, the general trend is an overall softening, as the tempering temperature is raised.

The absence of other alloying elements means that the hardenability of the steels is low, so a fully martensitic structure is only possible in thin sections.

However, this may not be a disadvantage where shallow hardened surface layers are all that is required.

Secondly, at lower carbon levels, the Ms temperature is rather high, so autotempering is likely to take place.

Thirdly, at the higher carbon levels the presence of retained austenite will influence the results.

Added to these factors, plain carbon steels can exhibit quench cracking which makes it difficult to obtain reliable test results. This is particularly the case at higher carbon levels, i.e. above 0.5 wt% carbon.

The addition of allying elements to a steel has a substantial effect on the kinetics of the y →α transformation, and also of the pearlite reaction.

Most common alloying elements move the TTT curves to longer times, with the result that it is much easier to miss the nose of the curve during quenching.

This essentially gives higher hardenability, since martensite structures can be achieved at slower cooling rates and, in practical terms, thicker specimens can be made fully martensitic.

Alloying elements have also been shown to have a substantial effect in depressing the Ms temperature.

It is clear that certain elements, notably silicon

can stabilize the €-iron carbide to such an extent

that it is still present in the ‘microstructure

after tempering at 400°C in steels with 1-2 wt%

Si, and at even higher temperatures if the silicon

is further

increased.

The evidence suggests that both the nucleation

and growth of the carbide is slowed down and

that silicon enters into the €-carbide structure.

It is also clear that the transformation of €-iron

carbide to cementite is delayed considerably.

While the tetragonality of martensite

disappears by 300°C in plain carbon steels, in

steels containing some alloying elements,

e.g. Cr, Mo, W V, Ti, Si, the tetragonal lattice

is still observed after tempering at 450°C

and even as high as 500°C .

It is clear that these alloying elements

increase the stability of the supersaturated

iron-carbide solid solution.

In contrast manganese and nickel decrease

the stability.

Alloying elements also greatly influence the

proportion of austenite retained on

quenching.

Typically a steel with 4% molybdenum,

0.2%C, in the martensitic state contains less

than 2% austenite, and about 5% is detected

in a steel with 1% vanadium and 0.2%C.

The austenite can be revealed as a fine

network around the martensite laths, by

using dark field electron microscopy.

On tempering each of the above steels at 300°C, the austenite decomposes to give thin grain boundary films of cementite which, in the case of the higher concentrations of retained austenite, can be fairly continuous along the lath boundaries.

It is likely that this interlath cementite is responsible for tempered embrittlement frequently encountered as a toughness minimum in the range 300—350°C, by leading to easy nucleation of cracks, which then propagate across the tempered martensite laths.

Alloying elements can also restrain the coarsening of cementite in the range 400-700°C, a basic process during the fourth stage of tempering.

Several alloying elements, notably silicon, chromium. molybdenum and tungsten, cause the cementite to retain its fine Widmanstatten structure to higher temperatures, either by entering into the cementite structure or by segregating at the carbide-ferrite interfaces.

Whatever the basic cause may be, the effect is to delay significantly the softening process during tempering.

This influence on the cementite dispersion has other effects, in so far as the carbide particles, by remaining finer, slow down the reorganization of the dislocations inherited from the martensite, with the result that the dislocation substructures refine more slowly.

In plain-carbon Steel cementite particle begin to coarsen in the temperature range 350 -400°C and addition of chromium, silicon, molybdenum or tungsten delays the coarsening to the range 500-550°C.

It should be emphasized that up to 500°C the only carbides to form are those of iron.

A number of the familiar alloying elements in steels form carbides, nitrides and borides which are thermodynamically more stable than cementite.

It would therefore be expected that when strong carbide elements are present in a steel in sufficient concentration, their carbides would be formed in preference to cementite .

Nevertheless during the tempering of all ahoy steels, alloy carbides do not form until the temperature range 500-600°C , because below this the metallic alloying elements cannot diffuse sufficiently rapidly to allow alloy carbides to nucleate.

The metallic elements diffuse substitutionally, in contrast to carbon and nitrogen which move through the iron interstitially.

With the result that the diffusivities of carbon and nitrogen are of several orders of magnitude greater in iron than those of the metallic alloying elements.

Consequently higher temperatures are needed for the necessary diffusion of the alloying elements.

It is this ability of certain alloying elements to form fine alloy carbide dispersions in the range 500—600°C, which remain very fine even after prolonged tempering, that allows the development of high strength levels in many alloy steels.

Indeed, the formation of alloy carbides between

500 and 600°C is accompanied by a marked

increase in strength, often in excess to that of

the as-quenched martensite.

This phenomenon, which is referred to as

secondary hardening, is best shown in steels

containing molybdenum, vanadium, tungsten,

titanium, and also in chromium steels at higher

alloy concentrations.

This secondary hardening process is a type of

age-hardening reaction, in which a relatively

coarse cementite dispersion is replaced by a new

and much finer alloy carbide dispersion.

On attaining a critical dispersion parameter, the

strength of the steel reaches a maximum, and

as the carbide dispersion slowly coarsens, the

strength drops.

The process is both time and temperature

dependent, so both variables are often

combined in a parameter:

P = T(k + log t)

where:

T is the absolute temperature and

t the tempering time in hours,

while k is a constant which is about 20 for

alloy steels, usually referred to as the

Holloman-Jaffe parameter.

PHYSICAL PROPERTIES: Annealed metals are relatively soft

and can be cut and shaped more easily. They bend easily

when pressure is applied. As a rule they are heated and

allowed to cool slowly.

The animation above shows that an annealed metal is

usually softer and can be deformed more easily than metals

that are not annealed.

PHYSICAL PROPERTIES: Hardened metals are difficult to cut

and shape. They are very difficult if not impossible to

bend. As a rule they are heated and cooled very quickly by

quenching in clean, cold water.

The animation above shows that metals that have not been

annealed are very difficult to deform.

DIVYA -109MM0001

ATUL RANJAN -109MM0125

MANINDRA RANJAN -109MM0556

ZEBA KAMAL -109MM0575