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PHASE TRANSFORMATIO N

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Pearlitic Transformation

PHASE TRANSFORMATION

Based onMasstransportPHASE TRANSFORMATIONSDiffusionaltransformationDiffusion less military transformationBased onOrderPHASE TRANSFORMATIONSIst order nucleation and growth2nd order entire volume transformsNo change in compositionChange in composition

Polymorphic Transformations: Typically exhibited by single component systems where different crystal structures are stable over different temperature ranges. E.g. bcc-fcc transformation in Fe

Major phase transformations that occur in solid phase are due to thermally activated atomic movements.The different types of phase transformation that is possible can be divided into 5 groups: Polymorphic changes Precipitation Transformation Eutectoid transformation Ordering reactions Massive transformation .

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

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

Massive Tranformations: Generally expressed as Original phase decomposes into one or more new phases which have the same composition as the parent phase but different crystal structures

Ordering Transformations: Generally expressed as (disordered) (ordered) . These do not require long range diffusion

PHASE transformation- Change in crystal structure+ Change in composition.Surface creations always hinders the process of transformation. The new phase always trys to create the surface, so energy needs to be supplied. So volume free energy will try to decrease the energy but surface free energy will try to increase the energy.

Fv=VfV= Vol of the new crystalf=free energies of the new phase Fs = ss = surface area of the new crystal = free energy per unit area

Mechanism of phase transformation :Changes of phase in the solid state involve a redistribution of the atoms in that solid and the kinetics of the change necessarily depend upon the rate of atomic migration. The transport of atoms through the crystal is more generally termed diffusion. This can occur more easily with the aid of vacancies, since the basic act of diffusion is the movement of an atom to anempty adjacent atomic site.Let us consider that during a phase change an atom is moved from an -phase lattice site to a more favorable -phase lattice site. The energy of the atom should vary with distance as shown in Figure 1, where the potential barrier which has to be overcome arises from the interatomic forces between the moving atom and the group of atoms which adjoin it and the new site. Only those atoms (n) with an energy greater than Q are able to make the jump, where Q = Hm-H barrier is given, from the MaxwellBoltzmann distribution law, as proportional to exp [Q/kT ], where k is Boltzmanns constant, T is the temperature and Q is usually expressed as the energy per atom in electron volts.

fig 1

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

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

The combined effect of (a) and (b) is shown in the curve below:

Where (a) is rate of crystal growth and (b) is rate of nucleation

This sluggishness of the transformation is known as hysteresis, and is attributed to the difficulties of nucleation, since diffusion, which controls the growth process, is usually high at temperatures near the transformation temperature and is therefore not rate controlling. Perhaps the simplest phase change to indicate this is the solidificationof a liquid metal.

The transformation temperature, as shown on the equilibrium diagram, represents the point at which the free energy of the solid phase is equal to that of the liquid phase.Thus,we may consider the transition, as given in a phase diagram, to occur when bulk or chemical free energy change,DGv is infinitesimally small and negative, i.e. when a small but positive driving force exists.However such a definition ignores the process whereby bulk liquid is transformed to bulk solid i.e. nucleation & growth.

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

D

Undercooling: It is the gap between the temp predicted for the transformation to occur and the temp at which the transformation actually occurs.

During the cooling of a liquid, solidification (nucleation) will begin only after the temperature has been lowered below the equilibrium solidification (or melting) temperature Tm. This phenomenon is termed supercooling or (undercooling). The driving force to nucleate increases as T increases.Small supercooling slow nucleation rate - few nuclei - large crystals.Large supercooling rapid nucleation rate - many nuclei - small crystals.

Effect of degree of undercooling on the rates of nucleation and growthTammanns curve

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

Rapid cooling of a pure molten metal reduces the time available for nuclei formation and delays the onset of freezing by a temperature interval of dT. This thermal undercooling (or super cooling), which is depicted in previous figure , varies in extent, depending upon the metal and conditions, but can be as much as 0.10.3Tm, where Tm is the absolute melting point. However, commercial melts usually contain suspended insoluble particles of foreign matter (e.g. from the refractory crucible or hearth), which act as seeding nuclei for so-called heterogeneous nucleation. Undercooling is much less likely under these conditions; in fact, very pronounced undercooling is only obtainable when the melt is very pure and extremely small in volume. Homogeneous nucleation is not encountered in normal foundry practice.

The growing crystals steadily consume the melt and eventually impinge upon each other to form a structure of equiaxed (equal-sized) grains (in upper 2 figures). Heterogeneous nucleation, by providing a larger population of nuclei, produces a smaller final grain size than homogeneous nucleation.The resultant grain (crystal) boundaries are several atomic diameters wide. The angle of misorientation between adjacent grains is usually greater than 1015. Because of this misfit, such high-angle grainboundaries have a higher energy content than the bulk grains and, on reheating, will tend to melt first.(During a grain-contrast etch of diamond-polished polycrystalline metal, the etchant attacks grain boundaries preferentially by an electrochemical process, producing a broad canyon which scatters vertically incident light during normal microscopical examination. The boundary then appears as ablack line.)

During the freezing of many metals (and alloys), nucleated crystals grow preferentially in certain directions, causing each growing crystal to assume a distinctive, non-faceted1tree-like form, known as a dendrite. In cubic crystals, the preferred axes of growth are directions.As each dendritic spike grows, latent heat is transferred into the surrounding liquid, preventing theformation of other spikes in its immediate vicinity. The spacing of primary dendrites and of dendritic arms therefore tends to be regular. Ultimately, as the various crystals impinge upon each other, it is necessary for the interstices of the dendrites to be well fed with melt if interdendritic shrinkage cavities are to be prevented from forming. Convection currents within the cooling melt are liable to disturb the delicate dendritic branches and produce slight angular misalignments in the final solidified structure (e.g. 510).These low-angle boundaries form a lineage (macromosaic) structure within the final grain, each surface of misfit being equivalent to an array of edge dislocations.Convection currents can also provide thermal pulses which cause dendritic branch tips to melt off and enter the main body of the melt, where they act as kindred nuclei.

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

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

The value of can vary widely from a few mJ m2 to several hundred mJ m2 depending on the coherency of the interface. A coherent interface is formed when the two crystals have a good match and the two lattices are continuous across the interface. This happens when the interfacial plane has the same atomic configuration in both phases, e.g. {1 1 1} in fcc and {0 0 0 1} in cph. When the matchat the interface is not perfect it is still possible to maintain coherency by straining one or both lattices, as shown in Figure a. These coherency strains increase the energy and for large misfits it becomes energetically more favorable to form a semi-coherent interface (Figure b) in which the mismatch is periodically taken up by misfit dislocations.The coherency strains can then be relieved by a cross-grid of dislocations in the interface plane, the spacing of which depends on the Burgers vector b of the dislocation and the misfit , i.e. b/. The interfacial energy for semi-coherent interfaces arises from the change in composition across the interface or chemical contribution as for fully coherent interfaces, plus the energy of the dislocations .

The energy of a semi-coherent interface is 200500 mJ m2 and increases with decreasing dislocation spacing until the dislocation strain fields overlap. When this occurs, the discrete nature of the dislocations is lost and the interface becomes incoherent.The incoherent interface is somewhat similar to a high-angle grain boundary with its energy of 0.51 J m2 relatively independent of the orientation.The surface and strain energy effects discussed above play an important role in phase separation. When there is coherence in the atomic structure across the interface between precipitate and matrix the surface energy term is small, and it is the strain energy factor which controls the shape of the particle. A plate-shaped particle is associated with the least strain energy, while a spherical-shaped particle is associated with maximum strain energy but the minimum surface energy.

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

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

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

This figure shows how this occurs at a mold wall or pre-existing solid particle, where the nucleus has the shape of a spherical cap to minimize the energy and the wettingangle is given by the balance of the interfacial tensions in the plane of the mold wall, i.e. cos = (ML SM)/SL.The formation of the nucleus is associated with an excess free energy given byDG = VDGv + ASLSL + ASMSM ASMML= /3(2 3 cos + cos3 )rDGv+ 2(1 cos )rSL + rsin2(SM LM).Differentiation of this expression for the maximum, i.e. dDG/dr = 0, gives rc = 2SL/DGv andW = (16/3DGv)[(1 cos ) (2 + cos )/4] or W(heterogeneous) = W(homogeneous)[S()].

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

Solid state transformation

During the solid state transformation still another factor acting inhibiting the nucleation transformation nuclei. A new phase always differs from the initial one in its structure and specific volume.Since the transformation develops an elastic crystalline medium, change in specific volume should cause an development in elastic strain energy in one or both the phases. This inhibits the transformation and kinetics the free energy. Therefore, the certain elastic component Fel makes a +ve contribution to the free energy change in the solid state transformation

Fe

g(Austenite)

Eutectoid transformation

CFCCFe3C(cementite)a (ferrite)+

(BCC)

TIME-TEMPERATURE-TRANSFORMATION(TTT) CURVES

TTT CURVESTo understand the character of transformation ot austenite to the resulting phases, Davenport and Bain showed that by studying the transformation isothermally (at constant temperature) of an austenitiseda series of temperatures below A1, a characteristic, time-temperature transformation (TTT) curved obtained. The diagram that illustrates the transformation of austenite as a function of time at a temperature is a TTT, or isothermal transformation ( I T) diagram. These curves have C or S shape in plain carbon and low alloy steels. Each steel composition has its own different S curve for given grain size (without inclusions). Fig. a illustrates S curve for eutectoid composition of steel.

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

From the TTT curve it can be seen that, just below the critical temperature, A1, the rate of transformation is slow even though the atomic mobility must be high in this temperature range. This isbecause any phase change involving nucleation and growth (e.g. the pearlite transformation) is faced with nucleation difficulties, which arise from the necessary surface and strain energy contributions tothe nucleus. Of course, as the transformation temperature approaches the temperature corresponding to the knee of the curve, the transformation rate increases. The slowness of the transformation below the knee of the TTT curve, when bainite is formed, is also readily understood, since atomic migration is slow at these lower temperatures and the bainite transformation depends on diffusion. The lower part of the TTT curve below about 250300C indicates, however, that the transformation speeds up again and takes place exceedingly fast, even though atomic mobility in this temperature range must be very low.

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

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

Possible phases in TTT diagram for eutectoid steel

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

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

Determination of TTT diagram for eutectoid steelFor the determination of isothermal transformation (or) TTT diagrams, we consider molten salt bath technique combined with metallography and hardness measurements. In molten salt bath technique two salt baths and one water bath are used.Salt bath I is maintained at austenising temperature (780C for eutectoid steel).Salt bath II is maintained at specified temperature at which transformation is to bedetermined (below Ae1), typically 700-250C for eutectoid steel.Bath III which is a cold water bath is maintained at room temperature.In bath I number of samples are austenite at A1+20-40C for eutectoid, A3+20-40C for hypo-eutectoid steel and Acm +20-40C for hyper-eutectoid steels for about an hour.Then samples are removed from bath I and put in bath II and each one is kept for different specified period of time say t1, t2, t3, t4,..........,tn etc.After specified times, the samples are removed and quenched in cold water.The microstructure of each sample is studied using metallographic techniques. The type, as well as quantity of phases, is determined on each sample.Transformation of austenite to ferrite-cementite mixtures occurs after a definite time (say t1) This time during which transformation does not proceed is known as incubation period.The magnitude of incubation period provides a qualitative idea about the relative stability of supercooled austenite. Smaller incubation period corresponds to lesser stability of austenite.

HOW TO DRAW TTT CURVE

54

FACTORS AFFECTING TTT CURVESEFFECT OF GRAIN SIZE ON THE TTT CURVESEFFECT OF ALLOYING ELEMENTS ON THE TTT CURVESEFFECT OF CARBON ON THE TTT CURVES

1.EFFECT OF GRAIN SIZE ON THE TTT CURVESAll decomposition products of austenite nucleate heterogenously at grain boundaries.

Thus incubation period is reduced for fine grained steelS curve is more towards the left in fine grained steel

EFFECT OF ALLOYING ELEMENTS ON THE TTT CURVESAll alloying elements (except Co) shift the S curve to the rightAustenite stabilizers move the curve to the right( Mn, Ni,etc)Carbide formers shift the S curve further to the right because:Diffusion of alloying elements is too slow(substitutional elements)Diffusion of carbon is slower as carbide formers donot easily part with the carbonAllotropic change -----> is reduced by solutesBainitic transformation is lesser affected ( no redistribution of alloying elements)

nose4340 Steel

EFFECT OF CARBON ON THE TTT CURVESHYPOEUTECTOID STEELSFerrite is the nucleating phase on decomposition of austeniteAs carbon increases from 0 to 0.77% :

EUTECTOID STEELSHave the maximum incubation period

HYPEREUTECTOID STEELSCementite is the nucleating phaseAs the carbon content increases more than 0.77%:

Temperature oCMsProeutectoid phase starts to form on this lineA +FAF + PPearlite reaction startsAc1MsMsMsA+PPFe3C +PFe3C +AProeutectoid cementite starts to form on this lineBBTTT curves for hypo , eutectoid and hyper-eutectoid steelsTTT curves for hypo , eutectoid and hyper-eutectoid steels

CONTINUOUS COOLING CURVE

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

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

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

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

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

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

Different cooling rates for eutectoid steel

AUSTENITIZATION

The first step in the true heat treatment cycle of steel is the austenitisation i.e. to get a homogeneous austenite by heating it to a predetermined temperature in the austenite stability range.Austenite can transform into various products depending on the composition and cooling rates.

Morphology of parent austenite(grain size) decides the morphology of products and thus its properties.

AustenitePearliteBainiteMartensite

FORMATION OF AUSTENITE

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

Experimentally, nucleation has been seen to occur at the interfaces of ferrite and cementite lamellae within a pearlite colony but primarily at the intersections of pearlite colonies.

Once the austenite has nucleated at the interface of ferrite and cementite, it grows consuming both the ferrite and cementite of pearlite.

The rate of movement of austenite boundary into ferrite and cementite phases is not equal.

This rate is inversely proportional to the concentration jump at the interface. As the concentration jump at the austenite-cementite interface is higher(due to the high concentration of carbon in cementite), austenite boundary moves much faster into ferrite phase.

By the time when the whole of the pearlitic structure has transformed to austenite, there is non-uniformity of carbon in austenite i.e. it has higher carbon content at the sites formerly occupied by cementite. Additional time is needed to obtain a homogeneous austenite.

In hypoeutectoid steels, the size of proeutectoid ferrite grains is much larger than the thickness of the ferrite lamellae in pearlite. Thus the time for complete disappearance of free ferrite exceeds the time needed for the disappearance of pearlite. The same is true in case of proeutectoid cementite.

The austenite formed from cementite and ferrite is generally not homogenous. Some heating is required to make it homogeneous.

Homogenization requires high temperature/time , or both High temperatures are required if the rate of heating is high, otherwise comparatively lower temperatures can achieve the purpose.

KINETICS OF AUSTENITE FORMATIONThe formation of austenite on heating occurs by nucleation and growth

The kinetics depends on:Transformation temperature and holding timeRate of heatingInterface between ferrite and cementiteGrain sizeNature of the alloying elements present

TRANSFORMATION TEMPERATUREThe rate of austenite formation increases with increase in temperature as it increases the rate of carbon diffusion and the free energy is more negativeTransformation takes a shorter time at higher temperatures of transformation and vice versa

RATE OF HEATINGFor higher rates of heating, transformation starts at higher temperatures and for slower rates, at lower temperaturesFor any rate of heating transformation occurs over a range of temperature

INTERFACE BETWEEN FERRITE AND CEMENTITE Higher the interfacial area faster is the transformation. Interfacial area can be increased by:Decreasing the inter-lamellar spacing between ferrite and cementite: The closer the ferrite cementite lamellae, the higher is the rate of nucleation.Increasing the cementite or carbon content: This will lead to more pearlite content in steels and thus more interfaces.Examples : 1. High carbon steels austenize faster than low carbon steels 2. Tempered martensite structure austenizes faster than coarse paerlite 3. Spheroidal pearlite takes longer time to austenize due to very low interfacial area

GRAIN SIZE The coarser the parent grain size the slower is the transformation rate.This is because for a given volume of sample, the total grain boundary area is less if the grain size is large.

NATURE OF ALLOYING ELEMENTS PRESENTAlloying elements in steel are present as alloyed cementite or as alloy carbides.Alloy carbides dissolve much more slowly than alloyed cementite or cementite.The stronger the alloy carbide formed the slower is the rate of formation of austenization.Diffusion of substitutional alloying elements is much slower than the interstitial element, carbon.Thus the rate of austenization depends on the amount and nature of alloying element

IMPORTANCE OF AUSTENITIC GRAIN SIZE IN STEELSThe size of austenitic grains is the most important structural characteristic of heated steel. The grain size strongly affects its own transformation behaviour and the mechanical properties of the microstructures formed from austenite.

Austenitic grain boundaries are preferred sites for the nucleation of pro-eutectoid phases(pro-eutectoid ferrite in case of hypoeutectoid steels and proeutectoid cementite in case of hypereutectoid steels) and pearlite which are diffusion-controlled transformation products.

Coarse austenite grains have less grain boundary area for a given volume of sample. Thus, fewer nucleation sites are available which leads to the retardation of diffusion-controlled transformation of austenite and paves way for the easy transformation to martensite.

EFFECT OF GRAN SIZE ON MECHANICAL PROPERTIES

Grain refinement improves the strength and ductility at the same time

IMPACT TRANSITION TEMPERATURE Increase in grain size raises the impact transition temperature, so more prone to failure by brittle fracture

CREEP STRENGTHCoarse grained steel has better creep strength above equicohesive temperatureBelow this fine grain structure have better creep strength

FATIGUE STRENGTHFine grained steel have higher fatigue strength

HARDENABILITYCoarse grained steels have higher hardenability(smaller grain boundary area in coarse grained structure gives less sites for effective diffusion, so martensite formation on cooling is favoured)

MACHINABILITYCoarse grain structure has better machinability due to ease in discontinuos chip formation(low toughness)

Pearlitic Transformation

INTRODUCTIONIt is a common micro constituent of a variety of steels where it increases the strength of steel to a substantial extent.

It is formed when austenite in iron carbon alloys is transformed isothermally at or below the eutectoid temperature (723K) .The name Pearlite is related to the fact that a polished and etched pearlitic structure has the colourfulness of mother-of-pearl.

DeVElopment of Microstructure

Schematic representations of the microstructures for an ironcarbon alloy of eutectoid composition (0.76 wt % C) above and below the eutectoid temperature.

DeVElopment of Microstructure

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

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

Morphologys

Consider the isothermal transformation diagram for aeutectoid ironcarbon alloy, with superimposed isothermal heat treatment curve (ABCD). Microstructures before, during, and after the austenite-to-pearlitetransformation are shown

Morphology

MorphologyWith decreasing temperature, the carbon diffusion rate decreases, and the layers become progressively thinner. The thin-layered structure produced in the vicinity of 540C is termed fine pearlite.

Fig.(a) - Coarse Pearlite [ Formed at higher temp and is relatively soft ] Fig.(b) - Fine Pearlite [ Formed at lower temp and is relatively hard ]

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

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

Mechanism HULL-MEHL model easily explains the pearlite formation.

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

mechanismHull-Mehl mechanism for Pearlitic Transformation

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

kineticsKinetics of Pearlitic transformation is well explained by JOHNSON & MEHL model.JOHNSON & MEHL related the fraction of austenite transformed to pearlite as a function of time by the equation:

where f(t) = fraction of austenite transformed to pearlite

.N = Nucleation rate.G = Growth ratet = Time

kineticsThis equation makes the following assumptions:The average nucleation rate is constant with time which actually isnt true.Nucleation occurs randomly, which isnt truly correct.The growth rate is constant with time, which can also change from one nodule to other and with time.Nodules maintain a spherical shape, but nodules may not be truly spherical.

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

kineticsThe time dependence of the nucleation rate in the early stages has been seen to increase as the square of time as shown below.

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

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

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

KINETICSThis shows that undercooling affects the rate of transformation in 2 ways:

Undercoolingincreased degree of undercooling reduces the transformation rate by lowering the rate of carbon diffusion curve.increased degree of undercooling increases the transformation rate by providing greater difference in free energies of austenite and pearlite.

KINETICSThe combined effect is shown in the curve below:

Where (a) is rate of crystal growth and (b) is rate of nucleation

KINETICS

Effect of degree of Undercooling of the rates of nucleation and growth

KINETICSHardness of pearlite increases as interlamellar spacing S0 decreases and also same for strength.

As S0 is inversely proportional to the degree of undercooling thus yield strength and also UTS is linearly related to the interlamellar spacing or degree of undercooling below eutectoid temp.

As the pearlite content increases in C steels, impact transition temp is substantially raised, decreasing ductility and toughness as the ferrite-cementite interface provides sites for easy nucleation of cracks

KINETICS

Effect of alloying ELEMENTS addition on PEARLITIC TRANSFORMATION

Almost alloying element except Co lower both the rate of nucleation and rate of growth.

As compared to carbon other alloying element diffuse very slowly.

As the diffusion rate for metallic atom is much slower than the

carbon atom the formation of stable carbide during the transformation will be feasible only at higher transformation temp.

Partitioning of carbon gets delayed when Cr eats up C and forms carbide Cr23C6 when alloyed with austenite.

BAINITIC TRANSFORMATION

INTRODUCTION:Bainiteis anacicularmicrostructure (not a phase) that forms in steels at temperatures from approximately 250-550C (depending on alloy content).

A fine non-lamellar structure, bainite commonly consists of cementiteand dislocation-richFerrite. The high concentration of dislocations in the ferrite present in bainite makes this ferrite harder than it normally would be.

Davenport and Bain originally described the microstructure as being similar in appearance to temperedmartensite.

EDGAR BAIN

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

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

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

Schematic Illustration Of The Origin Of The T0 Construction On The Fec Phase Diagram. Austenite With A Carbon Concentration To The Left Of The T0 Boundary Can In Principle Transform Without Any Diusion.Diusionless Transformation Is Thermodynamically Impossible If The Carbon Concentration Of The Austenite Exceeds The T0 Curve.

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

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

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

MORPHOLOGYOn the basis of morphology bainite can be of two types:- 1) Upper bainite 2)Lower bainite

UPPER BAINITE

Known as feathery bainite as it resembles feather of a birdForms in temperature range of 5500C-4000C.The structure consists of Lath or needle-like ferrite which runs parallel to the longer axis and Carbide precipitates as fine plates, parallel to the direction of growth of bainite, mainly at the lath boundariesCarbides are present as discontinuous stringers when carbon content is low and continuous stringers when carbon content is high.

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

Schematic growth mechanism of Upper BainiteUpper bainite in medium carbon steel

LOWER BAINITEKnown as Plate bainite.Forms in the temperature range of 4000C-2500C.The structure consists ofLenticular plates of ferriteFine rods or blades of carbide at an angle of 55 to 60o to the axis of bainite.Carbides can be cementite or -carbide, or a mixture depending on temperature of transformation and composition of steel.

Carbides precipitate within the ferrite platesFerrite plates have smaller sub-plates with low angle boundaries between themHigher dislocation density than upper bainiteHabit planes of ferrite plates are the same as martensite that forms at low temperatures of the same alloyAlloying elements do not diffuse or form their carbides during bainite transformation.

Lower Bainite structure in medium carbon steelStages of formation of Lower BainiteSchematic representation of lower bainite structure

MARTENSITICTRANSFORMATION

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

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

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

Plate Martensite Showing Coherency With Mother Grain Structure

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

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

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

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

Important characteristics of Martensite Transformation:Diffusionless/Military transformationAthermal transformation.Retained AusteniteMs and Mf temperaturesReversibility of Martensitic transformationHabit planesBain distortionEffect of applied stress on transformationHardness of Martensite Stabilization of Martensite

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

Diffusionless Shear Reaction:

Shape deformation of plate martensite Lens formation

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

Ms and Mf Temperature:For each steel, the Austenite to Martensite transformation starts at a definite temperature called Ms temperature.This temperature can vary very widely over the range from 500C to room temperature.This variation depends upon the amount of austenite stabilising elements in the steel (except Co & Al):

Ms (oC)=561 474(%C) 33(%Mn) 17(%Ni) 17(%Cr)-21(%Mo).

Carbon has a very strong effect on the Martensitic start temperature.

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

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

Reversibility Of Martensitic TransformationWith definite amount of superheating as the driving force, Martensite to Austenite Diffusionless transformation may take placeThis reverse transformation starts at temperature As This property can be seen in systems like:Fe-Ni alloysAl-Cu alloysTi alloys, etc.

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

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

Martensite habit plane in various types of Steel

Martensitic Habit Planes and their conversion to BCT structure

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

Bain Distortion Model:

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

Martensite

FCCAusteniteFCCAusteniteAlternate choice of CellTetragonal MartensiteAustenite to Martensite 4.3 % volume increasePossible positions of Carbon atomsOnly a fraction ofthe sites occupied

20% contraction of c-axis12% expansion of a-axis

In Pure Fe after the Matensitic transformationc = a

C along the c-axis obstructs the contraction

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

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

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

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

172

Martensite forms by three different modes:Athermal (without thermal activation)BurstIsothermal (thermally activated diffusion-controlled)

ATHERMAL Martensite:The amount of Martensite formed is a function of the temperature to which the alloy is cooled.Cooling to lower temperatures leads to formation of new plates.This kinetics proceeds above Room temperature, so is dominant in industrial practices.The fraction of thermal Martensite formed is given by: f=1- exp(- 1.10 10^-2 T) where, T is the degree of undercooling below Ms temperature.173

BURST Kinetics (Jump-like Kinetics):For some alloys like Fe-Ni and Fe-Ni-C , with sub-zero Ms temperatures, the Burst phenomenon occurs.Here the plates of Martensite nucleate newer plates , known as auto-catalysis .Zigzag arrays of plates are formed.All the plates form in a very small fraction of second accompanied with an audible click.The amount of Martensite formed in a burst varies from a few percent to even 70% of Austenite.ISOTHERMAL Kinetics:Occurs in alloys like: Fe-Ni-Mn and Fe-Ni-CrTransformation is a function of time at a constant temperatureReaction starts slowly, then accelerates due to autocatalysis, and then decaysIt does not go to complete transformation.

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

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

Growth of a Lath Martensite

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

Constraints in the matrix does not allow parallel plates but a lens.

Growth of Plate Martensite:

AGE HARDENING

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

Precipitation Hardening

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

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

CONTINUEDThe matrix should be relatively soft and ductile, and the precipitate should be hard and brittle.The alloy must be quenchable.

Solvus curve

Solvus curve

STEPS IN AGE HARDENING HEAT TREATMENTSOLUTIONIZINGfirst heat treatment where all solute atoms are dissolved to form a single-phase solid solution. ( just above the solvus temperature)Heat to T0 and dissolve second phaseOver heating is avoided as it may lead to:MeltingOxidationGrain growthBurningDecrease in ductility

STEPS IN AGE HARDENING HEAT TREATMENT2. QUENCHINGRapidly quench to very low temperature T1Metastable Super-SaturatedSolid-Solution i.e high temperature state ( A phase solid solution supersaturated with B atoms) formedHot boiling water or air cooling or cold water used as required for quenching

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

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

Steps in Precipitation Hardening

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

193

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

ARTIFICIAL AND NATURAL AGEINGARTIFICIAL AGEINGAgeing at a temperature higher than room temperatureHardness peak comes in very short timeGrowth is comparable to nucleationParticles become large in short period and steel loses their hardnessNATURAL AGEINGAgeing is done at room temperatureRequires long times- Several days to reach maximum hardnessPeak strength is higher than obtained in artificial ageing, no over ageing occurs.ooo194

Hardness, VHNAgeing time, (change of scales at certain intervals)

196Ageing

Effect of Ageing Temperature on Strength

Effect of Ageing Temperature on Ductility

QUENCHED IN VACANCIESOn quenching from high temperature, high % of vacancies get retained in steelThese vacancies provide path for diffusion at lower temperatures when diffusion rate is very slowSolute atoms move through few inter atomic distances with the help of these vacancies to give very fine precipitation AgeingThe fluctuation in solute concentration provide small clusters in the crystal in solute which acts as nuclei for the precipitationSize of precipitation becomes finer as temperature at which precipitation occurs is lowered

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PRECIPITATION SEQUENCE DURING AGEING OF ALLOY (Al-4.5%Cu)The precipitation occurs in steps involving several transition (metastable) precipitates before equilibrium precipitate forms The equilibrium precipitate does not form instantly as nucleation barrier is too high - incoherentThe alloy is quenched from 550CThe sequence:

(CuAl2)200

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GP ZONESGuinier- Preston Zones also called GP1 ZonesThe first early stage of ageingFully coherent, same lattice structure as Alluminum with matrix thus nucleation is favoredPlate-like clusters of Copper atoms segregated on {100} planes of aluminum latticeDiameter 100 , Thickness 3-6Density 1018 per cm3Coherency or elastic strains developOccurs by diffusion of Cu atoms aided by Quenched-in vacancies over short distancesGive first peak of hardness

PRECIPITATION SEQUENCE DURING AGEING OF ALLOY (Al-4.5%Cu)202

(GP2 ZONE)Coherent intermediate precipitateComposition is CuAl2Plate like, Diameter- 1500, Thickness- 100Tetragonal crystal Structure, a= 4.04, c =7.68Have elastic coherency strainsProduce greater distortion than any other transition structure

PRECIPITATION SEQUENCE DURING AGEING OF ALLOY (Al-4.5%Cu)203

Equilibrium precipitate CuAl2Fully incoherent precipitateNucleates heterogeneouslyTetragonal crystal Structure, a= 6.07, c =4.87Coherency strains are not presentLeads to SofteningResult of Overageing

PRECIPITATION SEQUENCE DURING AGEING OF ALLOY (Al-4.5%Cu)204

With increasing time, the hardness increases, reaching a maximum (peak), then decreasing in strength.The reduction in strength and hardness after long periods is overaging (continued particle growth).

PRECIPITATION SEQUENCE DURING AGEING OF ALLOY (Al-4.5%Cu)

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Maximum hardness is obtained when there is Critical- dispersion of GP Zones , or any other intermediate precipitates(or ) or both After peak hardness, further ageing tends to decrease hardness overagingDuring overageing, the particles coarsen at the cost of neighboring particles

PRECIPITATION SEQUENCE DURING AGEING OF ALLOY (Al-4.5%Cu)206

KINETICS OF PRECIPITATIONRate of precipitation is faster at higher temperaturesRate of precipitation is faster in alloys of widely dissimilar metalsRate of precipitation is increased with presence of impuritiesRate of precipitation increases with application of plastic deformation just before ageingRate of precipitation at a ageing temperature is faster in a low melting alloy207

208Effects of Aging Temperature and TimeThe effect of aging temperature and time on the yield strength of an Al-4% Cu alloy.

2003 Brooks/Cole, a division of Thomson Learning, Inc. Thomson Learning is a trademark used herein under license.

209The operator of a furnace left for his hour lunch break without removing the Al-4% Cu alloy from the furnace used for the aging treatment. Compare the effect on the yield strength of the extra hour of aging for the aging temperatures of 190oC and 260oC.Effect of Aging Heat Treatment Time on the Strength of Aluminum Alloys

2003 Brooks/Cole, a division of Thomson Learning, Inc. Thomson Learning is a trademark used herein under license.Fig 2 The effect of aging temperature and time on the yield strength of an Al-4% Cu alloy.

210The magnesium-aluminum phase diagram is shown in Figure. Suppose a Mg-8% Al alloy is responsive to an age-hardening heat treatment. Design a heat treatment for the alloy.Design of an Age-Hardening Treatment

2003 Brooks/Cole, a division of Thomson Learning, Inc. Thomson Learning is a trademark used herein under license.Fig 3 Portion of the aluminum-magnesium phase diagram.

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2003 Brooks/Cole, a division of Thomson Learning, Inc. Thomson Learning is a trademark used herein under license.Figure 11.14 Microstructural changes that occur in age-hardened alloys during fusion welding: (a) microstructure in the weld at the peak temperature, and (b) microstructure in the weld after slowly cooling to room temperature.

HARDENING MECHANISMs:

According to dislocation theory, the strength of a material is controlled by the generation and mobility of the dislocations. The increased strength of an age-hardened alloy is due to the interactions of the moving dislocations with the dispersed precipitates.Barriers to the motion of dislocations:1) Coherency strains around the GP zones2) GP zones or precipitates

REASONS OF HARDENING BY AGEING:1) Coherency strain hardening2) Chemical strain hardening3) Dispersion strain hardeningDISLOCATION INTERACTION WITH:1) Other dislocations-strain hardening2) Grain boundaries-grain boundary strengthening3) Solute atoms-solid solution strengthening4) Precipitates-precipitation hardening5) Dispersoids- dispersion strengthening

COHERENCY STRAIN HARDENING:

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

COHERENT GRAINSINCOHERENT GRAINS

DISPERSION HARDENING:

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

DISLOCATION-CUT MECHANISM:

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

The dislocations move through the matrix according to one of the following:

BY-PASS MECHANISM:

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

=Gb/lWhere:G= is the shear modulus of the matrixb= is the Burgers vector of the dislocationl= is the distance between the dislocationsEvery time a dislocation bypasses it leaves behind a loop of dislocation the precipitate.Thus l decreases and the stress needed for the next dislocation to bypass increases In over ageing precipitates, l increases so strength decreases.

RECOVERY, RECRYSTALLIZTION AND Grain Growth

INtroduction

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

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

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 increases from the value of 10101012 lines m-2 typical of the annealed state, to 10121013 after a few per cent deformation, and up to 10151016 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:RecoveryRecrystallization Grain growth

Figure showing Effect of annealing processes on the various properties of material.

RECOVERY

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

RECOVERY:

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 P