WARM WORKING BEHAVIOUR OF ALPHA-IRON, Fe-Si, Fe-Co AND Fe-Ni ALLOYS :

A STUDY USING PROCESSING MAPS

A Thesis Submitted for the Degree of

f l a e t ~ r L of sc ience (Fxqineeriq) in the Faculty of Engineering

BY G . S. AVADHANI

DEPARTMENT OF METALLURGY INDIAN INSTI'MJTE OF SCIENCE

BANGALORE - 560 012 (INDIA) SEPTEMBER 1996

ACKNOWLEDGEMENT8

It is a great privilege to record my deep sense of

gratitude to Prof, Y . V . R . K . Prasad for introduc ing me to the

innovative science of Deformation Processing. His stimulating t guidance and constant encouragement has benefited me to learn the 4

best of research culture and the subject. In him I find a true friend, philosopher and guide with a kind heart.

I wish to thank Prof. D.H. Sastry, Chairman, Department of

Metallurgy for extending the facilities of the Department and his

support. I thank Prof.S.Ranganathan and Prof. K.S. Raman for

their encouragement.

I express my gratitude to Dr. M.Srinivas and Dr. G.

Malakondaiah of Defence Metallurgical Research

Laboratory(DMRL),Hyderabad, for providing all the materials used

in this . investigations.

I sincerely thank Mr. S. Sashidhara for his kind help in

conducting the experiments and for the encouragement he has

provided. I wish to thank Mr. G. Sivaram, Mr. R. Ravi and Mr. T.

Seshacharyulu for their help in computational work and for their

friendship. I thank Mr. Mohan Raj for his excellent sample preparation work. I also thank Mr. S. Sreenivas Murthy and Mr. D.

Molliah for their assistance in drafting and metallorgraphy

respectively. I thank Mr. Rajashekar for the excellent typing work of the thesis.

Finally, I express my gratitude to my parents and family

members. It is their affection, encouragement and inspiration

that made this achievement possible.

G.S. AVADWNI

In recent years, warm working of ferrous materials in the

temperature range 0.4-0.6Tm has assumed considerable importance

in manufacturing components with a combination of good strength &

ductility. The warm forged components are used in automobiles and

require close dimensional tolerence and good surface finish.In

the warm working temperature range (400-900~~), ferrous materials

exhibit the process of dynamic recovery which produces a stable

subgrain structure enhancing the strength.

The aim of the present investigation is to study the warm

working characteristics of alpha iron(bcc) , with a view to

understand the mechanism of deformation and to optimise the

process parameters, namely, temperature and strain rate. A

further aim of the investigation is to examine the effect of

alloying additions on the warm working behaviour of alpha iron

with a view to understand the response of alloy steels, during

warm working. Three binary alloys of iron, namely Fe-Si, Fe-Co

and Fe-Ni are chosen for this investigation. It is well known

that Si additions stabilize alpha phase while Ni additions

stabilize gamma phase. Co on the other hand, does not

significantly affect the transformation temperature. Furthermore,

the Curie temperature is also influenced differently by these

alloying additions.

The approach adopted in this investigation involves

developing of Processing Maps, using the concepts of Dynamic

Materials Modelling. Processing Maps consist of the variation of

the efficiency of power dissipation ( 9 = 2x1 / m+l , where m is the strain rate sensitivity of flow stress), as compared to an ideal linear dissipater, with temperature and strain rate. The

input required for developing the processing map is the variation

of flow stress with strain rate in a wide range at different

temperatures in the warm working range.

Compression tests were performed on alpha iron, Fe-5Si, Fe-

0.5C0, Fe-5Co, Fe-O.5Ni and Fe-5Ni alloys in the temperature

range 400-900~~ and strain rate range 0.001s-~ - 100s-~. From the

load - stroke curves, true stress-true strain data were

evaluated. The strain rate sensitivity of flow stress and the

efficiency of power dissipation 2m/(m+l) were obtained on the basis of the flow stress data as a function of temperature and

strain rate. Using a computer programme, processing maps,

consisting of power dissipation and instability maps were

developed for the various materials. Microstructures

corresponding to the various domains were examined using standard

techniques.

The processing map for alpha iron exhibited two domains: one

at 400~~/0.001s-~ with a peak efficiency 279 and another at 800c/ 0.001s-~ with a peak efficiency of 35%. The former domain represents dynamic recovery while the latter represents dynamic

recrystallization. The ductility reaches a peak value at 800c/ 0.001~'~ where efficiency also reaches a peak. Alpha iron

exhibits instabilities at strain rates greater than IS-' and

temperatures below 700'~ and these manifest as adiabatic shear

bands.

The addition of Si does not change the map significantly

except that the peak efficiency in the DRX domain (800~/0.001s~' ) has considerably increased. This is attributed to the decrease in the Curie temperature caused by Si addition which would

enhance the grain boundary migration in this alloy. The additions

of Co on the other hand, have shifted the DRX domain to higher

temperatures and higher strain rates since Co increases the Curie

temperature. Ni additions cause the formation of the two-phase

region(a1phatgamma) in which both phases deform in a compatible fashion. However, Fe-Ni alloys exhibit extensive instabilities in

a wide range of temperature and strain rate, making them

unsuitable for warm working. The influence of the concentration

of Co and Ni on the warm working characteristics is not very

significant.

In conclusion, the study shows that alpha iron exhibits

dynamic recrystallization at 800~/0. 001s-l. DRX is favoured by Si additions. Co shifts the DRX temperature to higher values

while Ni additions are unfavourable for warm working.

CONTENTS

Acknowledgements Synopsis Contents

C h a p t e r I : INTRODUCTION

1.1 Workability

1.2 Workability Test Techniques

1.3 Intrinsic Workability and Modelling Techniques

1.4 Literature Survey on High Temperature Deformation of Ferrous Materials

1.5 Aim of the Investigation and Approach

C h a p t e r I1 : EXPERIMENTAL

11.1 Materials and Specimen Preparation

11.2 Compression Testing

11.3 Development of Processing Maps

11.4 Tensile Testing

11.5 Metallography and Grain Size Analysis

C h a p t e r 1x1 : RESULTS AND DISCUSSION

111.1 Alpha Iron

111.2 Fe-SSi Alloy

111.3 Fe-Co Alloys

111.4 Fe-Ni Alloys

C h a p t e r IV : SUMMAIZY 24ND COHCLUSIONS

i ii iii

References

CHAPTER - I

INTRODUCTION

1.1 2 WORKABILITY

Mechanical working is an important step in the engineering

component manufacture and is done not only to give the required

shape to the material but also to obtain the specified

properties1. Constitutive behaviour of the workpiece under

processing conditions has to be understood for defining this goal

on repeatable basis. A parameter called 18workabilityt* which is

the ability of the workpiece to take different shapes in a metal

forming process without the onset of fracture or flow

instability, has to be optirnised for this purpose.

It is convenient to consider workability to consist of two

independent parts : (i) Intrinsic workability and (ii) state-of-

stress (SOS) dependent workability. Intrinsic workability depends

upon constitutive flow behaviour of the rnetal/alloy which

includes microstructure (chemical composition, prior processing)

and its response to the applied temperature (T), strain rate ( 2 ) and strain ( f ) in processing.

Normally, cold working is carried out at temperatures below

0.25 T,, hot working between 0.6 to 0.8Tm or more and warm

working between 0.4 to 0.6Tm, where T, is melting point

temperature. The recrystallization temperature is about 0.6 to

0.7 T,. In cold working, work piece acts as a storage of energy

and so the intrinsic workability also depends on strain. However,

for hot working, it acts as a dissipater of energy and therefore

strain effects are not significant from hot workability point of

view15.

Warm working has advantages of both the hot and cold working

giving good ductility and strength combinations to the product19.

~lso, we get good surface finish and better dimensional tolerance

in warm forged components making them useful in automobile

Industry in the manufacture of gears, ax\es and such critical

components.

In the warm range, the mechanical properties are changed by

the production of a low density, well-ordered sub grain structure

arising from the annihilation and polygonization of the

dislocations. This is called dynamic recovery and is associated

with elongated grains. In the hotter domain, in addition to the

above, new grains nucleate, grow and are deformed in a process of

dynamic recrystalization (DRX); the grains are equiaxed and also

contain a recovered substructure. As the temperature rises

further and strain rate falls, the strength declines in

association with an increase in the size of both subgrains and 17 grains . Therefore, the DRX domain is ideal for high

temperature deformation with optimum intrinsic workability.

State-of-stress (SOS) Workability depends upon the geometry

of the deformation zone (die design) and the applied stress state

in a mechanical process. For good SOS workability, the

hydrostatic components should be essentially compressive. Both

aspects of workability have to be optimised separately for

achieving defect free final product on repeatable basis without

trial and error.

1.2 : WORKABILITY TEST TECHNIQUES :

Some important experimental methods to evaluate deformation

behaviour are given below.

1. Uniaxial com~ression Test L In this a cylindrical specimen of -

standard dimension is compressed between flat dies with suitable

lubricants under isothermal conditions at a varying or constant

strain rate to obtain flow curves. Important workability

parameters like strain to initiate cracks can be estimated from

the microstructural investigation of the samples deformed at

different strains. This techniques is prefered over others since

accurate data Can be obtained at a constant true strain rate and

under conditions where the adiabatic temperature rise may be

measured.

2. Torsion Test L The large strains and strain rates attainable -

without lubricant breakdown or bulging problems make torsion test

suitable for large plastic strain investigations. Both solid or

tubular specimens are used. The number of revolutions taken for

failure is taken as an index of workability. The outer radius is

generally employed to characterise the material with associated

values of surface strain rate and shear stress.

3. Uniaxial ensi ion Test L The flow stress data, elongation to failure and the reduction in the cross-sectional area at fracture

are the important data that are derived from this test. It is

usually performed on round bars or thin sheet specimens.

While tensile, compression or torsion techniques may be

used, hot compression tests have decisive advantages. First of

all, in a compression test, it is easy to obtain a constant true

strain rate using an exponential decay of the actuator speed.

Secondly, in view of a simple geometry (cylindrical) of the

specimen, it is convenient to measure the adiabatic temperature

rise in the specimen so that temperature correction may be

incorporated. The test system must have the facility to achieve a

wide range of strain rates (10'~ to 10~s") and for an isothermal

heating of the specimen.

One or many of these tests may be employed to validate

workability data.

I. 3 t INTRINSIC WORKABILITY AND MODELLING TECHNIQUES FOR ITS OPTIMI2ATION:

(i) Intrinsic Workability The constitutive behaviour of the material under processing

conditions is the key to the understanding of intrinsic

workability since this decides the response of the material to

the applied temperature, strain rate and strain1. The

constitutive behaviour sensitively depends on the microstructure

which itself depends upon the chemistry of the alloy and the

processing history. As a part of the response of the material to

the applied process parameters, certain microstructural changes

(mechanisms) occur within the material. Frost and ~ s h b ~ ~ were the

first to represent this response in the form of deformation

mechanism maps of normalized stress vs homologous temperature

showing the area of dominance of each flow mechanism. The

kinetics of several flow mechanisms have been considered using

equations relating the flow stress to temperature, strain rate

and structure and the boundaries for each mechanism have been

established. Emphasis in all these maps was essentially placed on

the creep mechanisms and so on the lower strain rate regimes.

However, mechanical processing involves several mechanisms that

occur at higher strain rates. Considering strain rate as one of

the direct variables, ~ a j ~ extended the concept of Ashby to construct a processing map that represents the nucleation of

damage as a function of temperature and strain rate.

(ii) Rishi R a j maps ~ a j ~ considered two important damage mechanisms that are

relevant to processing. One is the cavity formation at hard

particles which do not deforn themselves but the matrix around it

deforms more than average, producing work hardening and high

stress concentration near the particles. When the stresses get

large enough, the interface may separate or the particle itself

may crack which may lead to the creation of damage due to cavity

formation ultimately contributing to ductile fracture. At

elevated temperatures, the rate of void formation will be reduced

because of lower work hardening rates due to recovery. Also,

lower strain rates will help in relieving the stress

concentration at the particle interfaces by the process of

diffusional transport of matter from regions of compression

around the particles to the regions of tension and by other creep

processes. On the basis of these temperature and strain rate

sf fects, ~a j4 calculated the lower bound condition for avoiding

cavitation at hard particles. The second damage mechanism is the

formation of wedge cracks at the grain boundary triple junctions to relieve the stress concentrations caused by the grain boundary

sliding occuring at high temperatures and lower strain rates. If

the strain rate is so high that the matrix deforms at a rate

faster than the boundaries can slide, then sliding effects will

be negligible and wedge cracking will not occur. If the strain

rate is very slow then there will be enough time to relax the

high stresses at the triple junctions. The upper bound condition for avoiding the wedge cracking at elevated temperature was

calculated by ~ a j ~ . Fig.1 shows the Raj map in which the lower bound for cavitation and upper bound for wedge cracking are

represented. In principle there is always a region which may be

termed as I1safeI1 for processing where these two damaging

mechanisms do not occur. The safe region has also a limit at very

large strain rates where flow localization due to the adiabatic

shear may occur. The 'lsafell region will consist of dynamic

recovery and dynamic recrystalization processes which impart

workability to the material.

(iii) Dynamic ater rials Model A recent development in the understanding of the

constitutive behaviour of the workpiece is the dynamic materials

model6, recently reviewed by Gegel et a17. This model is based

on systems engineering concepts8. Processing is modeled as a

system, an example of which is shown in Fig,, awith reference to

the extrusion process. The system consists of a source of power (a hydraulic powerpack), a store of power (tools like the container,

HIRAL MAPS

CRACKHG

' 0 4 0 6 0 8 T l ~ m

Fig. 1 Deformation processing map for aluminium. After ~ a ] 4 .

PROCESSING SYSTEM

"if SUBSYSTEMS ELEMENTS HYORAULIC ORWE SOURCE OF POWER EXTR . EW1PMENr STORE OF POWER WORKPIECE OlSSlPATOR OF KWER LUBRICA NT INTERFACE

~ig. 2Processing system with extrusion as an example.

the ram, the die holder and the die) and a dissipator of power

(the work piece). Energy is generated by the source, transmitted

to the tools which store the power and transfer it to the

workpiece through an interface (lubricant). The workpiece itself

dissipates the energy while it undergoes plastic flow in the

deformation zone. The response of each of the above system

elements depends upon their individual constitutive equations

which are to be known for modeling their behaviour. While the

above three elements are important parts of the system, the

transmission and interface elements of the system are less

understood and not modeled adequately. If the constitutive

behaviour of the system elements could be modeled accurately,

they can be linked together such that process control for energy

optimization may be achieved. In this system, it is important to

note that the power or energy per second is to be considered and

not energy per se since the response of the system depends on how

fast or slow the energy is input into the system and thus time

becomes an independent variable which makes the system "dynamicw.

While enough work has not been done in integrating the system

elements, the characteristics of the dissipator element

(workpiece) have been analysed using systems concepts. The

constitutive equation of the workpiece is its intrinsic

characteristic describing the manner in which energy is converted

at any instant into a form usually thermal or microstructural

which is not recoverable by the system.

On the basis of the above description of the dissipative

characteristics, the instantaneous response (d) of workpiece material to the applied strain rate ( k ) for creating a given

large plastic strain is represented in Fig.3(a). The dynamic

constitutive equation is assumed to follow a power law equation.

6 = K. im . . . . . (I) where K and m are constants.

The instantaneous total power dissipated will be given by a

rectangle of area (da). The constitutive equation decides the

dlkpath taken by the system to reach the applied strain rate

(limiting condition). If a largely different strain rate is

applied, a different 6.k path may be chosen by the material and

hence the values of K and m will change. For example, if high

strain rates are applied the material may choose a path leading

to internal fracture while the same material may deform by a safe

process like dynamic recovery at lower strain rates. The 6 - i path

so chosen by the system may be termed as lcdynamicn and is

dependent on the limiting conditions.

The total power dissipated, 4 . & , consists of two parts. In

systems modeling terminology8 these are given by the sum of two

integrals : i d P = t j . j = J ( * d + J 2 dd

0 0

The first integral is called G-content and the second J co-

content which is a complimentary function of G-content. The area

under the constitutive equation curve is represented by G-content

while the area above the curve is given by J co-content.

In order to understand the physical interpretation of G and

J of the dissipater, the atomistic processes of plastic

deformation in simple shear may be consideredlo. Plastic flow

occurs by slip along glide planes when a shear stress is

applied and is facilitated by the presence of dislocations. The

work done by 7 increases the potential and kinetic energies and

the potential energy is maximum when the atoms in the two rows

are opposite each other which makes the configuration unstable.

Thus the row or part of it moves fast without further assistance

from T into the next stable equilibrium configuration. A

considerable part of the potential energy is almost

instantaneously converted into kinetic energy. The total kinetic

energy created by plastic flow is converted into a temperature

rise. The major portion of the input is dissipated through this temperature rise and is represented by 6-content, the area under

the curve representing the dynamic constitutive equation.

The dislocations generated by plastic deformation will move

with certain velocity which is responsible for the strain rate

sensitivity of flow stress. The moving dislocations may group

themselves after some annihilation by mechanical or thermal

recovery and may also form interfaces which at sufficiently high

temperature may migrate to cause large scale annihilation of

dislocations. At lower temperatures, where the recovery processes

are slow, the dislocation groups may create internal cracks the

free surfaces of which form the sinks for the dislocation

annihilation. There can be several atomistic processes (e.g.

diffusion-aided flow, stress induced phase transformation) that

annihilate dislocations and dissipate energy. All these

metallurgical processes contribute to the dissipation of power to

a Smaller extent than G-content and represent power dissipation

through a complementary function J-co-content.

The power partitioning between G and J in a viscoplastic

material is controlled by the strain rate sensitivity (m) of flow

stress, since.

and thus m is a power partitioning factor. At one extreme, J can

only be as high as G since dislocations cannot be annihilated at

rates faster than they are generated (or the temperature rise).

This is the ideal case of a linear dissipator [Fig.z(b)] for which rn = 1 and J = Jmax = Gmin = 0.5P. At the other extreme, m =

0 and J = 0 and the material does not dissipate power through

metallurgical processes and will act as a llstorew of energy by

dislocation generation. Stable viscoplastic flow occurs between

the two extremes m = 0 and rn = 1.

J co-content may be explicitly evaluated from the integral

where Kf = (l/K)l/m is another constant. By combining with Eq. (1) one gets the J co-content as :

r n a d - i J = ......( 5 ) m + l Considering that the maximum possible rate of dislocation

annihilation can only be as fast as the dislocations are

generated, the power dissipation through J co-content may be

compared with that in a linear dissipater [m = 1 ; J,,, = ( 6 4 2 ) to define a dimensionless parameter called efficiency of power

dissipation (9 through metallurgical processes.

This parameter helps in mapping the dissipative

microstructural characteristics of the workpiece in a wide range

of strain rate and temperature. A schematic three-dimensional map

of the efficiency of power dissipation with temperature and

strain rate is shown in Fig.4 (a). In view of the non-linear

variation of the flow stress with strain rate, the map will have

hills and valleys. A better representation will be in the form of

an isoefficiency contour map obtained by sectioning the 3-D map

at constant efficiency levels. A schematic contour map is shown

in ~ig.4(b) .

(iv) Interpretation of Power dissipation maps The power dissipation maps are based on continuum approach

and will have to be interpreted in terms of the atomistic

mechanisms that are responsible for the power dissipation. Raj maps are the basis for this interpretation and the following

guidelines are available.

(i) In the low temperature (Ts 0.25 Tm), high strain rate regime

(10-100 s'l) , void formation occurs at hard particles and

that leads to ductile fracture. In the dissipation maps

these regions are characterized by a very high efficiency

and a rapid increase in efficiency with a decrease in

temperature and an increase in strain rate.

(ii) In high temperature (Tz 0.75 T ) , low strain rates

S ) , regime, wedge cracking caused by grain boundary

sliding occurs (except in superplastic materials in which

wedge cracking is at a minimum). In this region, the

t / h ( b )

Fig. 4 : (a) Schematic map of the variation of the power dissipation with temperature and strain rate;

(b) Contour map showing iso-efficiency contours.

efficiency of power dissipation is very high and increases

with a decrease in strain rate until a peak is reached.

(iii) In high temperature (T = 0.75 T,) and high strain rate regime (10" to 10 s") , dynamic recrystalization

dominates. This domain has a medium efficiency of power

dissipation (30-50%) .

(iv) At intermediate temperatures and strain rates dynamic

recovery process occurs and has a low efficiency ( ~ 2 0 % ) .

(V) At very high strain rates (2 10 s'l) there is a possibility for the occu ence of adiabatic shear bands r' and these lead to flow localization. Efficiency of power

dissipation is very low.

In complex alloys, there could be many more metallurgical

processes that contribute to power dissipation. Prior knowledge

of these processes is required to identify their characteristics

in a power dissipation nap. Also, after the domains are

identified, they have to be confirmed by microstructural studies

in each of the domains.

For optimizing intrinsic workability and for bulk working,

the domain of dynamic recrystalization (DRX) is of special

significance. Of the 'lsafelf metallurgical processes of power

dissipation, DRX has the highest efficiency. The process

reconstitutes the microstructure through the formation and

migration of grain boundaries.

(V) Instability Maps:

Some extremum principles in irreversible thermodynamics as

applied to continuum mechanics of large plastic flow are

considered by 2ieglerl0. It has been shown that any quasistatic

process involving a change of the strain (xk)* and not purely reversible implies a flux in phase space equivalent to an entropy

production d(i) S O . The entropy production may depend on the

present state of the system and on its history or otherwise

completely determined by the increment in strain (dxk). It can be

expressed in terms of the temperature (8) and the elementary

dissipation work ( d ~ ( ~ ) ) . C i) ~ i ) c;> dw = Xk * d z k = 0 . 4 5 + O . . . . . ( 7 )

where ~L'is the irreversible force*. The rate of dissipation work

is related to the rate of entropy production as. ( i )

-

c3 d L p = xk kk = e - dt d t . . . . . ( 8 )

where *k is the strain rate (velocity)* and P ( ~ ) power

dissipated, At any given stage of the process, the rate of

dissipation work is a function of the velocities Gk and is called a dissipation function D(&~) of the system

D c i l r ) >/ 0 ..... (9 The dissipation function represents the constitutive behaviour of

the material and is defined by

C L 3 without the aid of the irreversible forces Xk which occur only at

I r> the macroscopic level. Xk is related to D(&) as

* Different notations for strain, strain rate (velocity) and irriversible force are used to represent them as tensors. Effective strain, strain rate and flow stress are earlier denoted b y , b n d d.

ziegler10 noted that in a stage of the deformation process, the

entropy production inside the system, its motion and the

dissipation function are functions of strain rates alone,

provided the initial condition and strain rates are prescribed.

The strain itself defines the frame of the microsystem. On the

basis of the dissipation function and the principle of least

irreversible force, 2ieglerl0 proved that:

(i) There is a correspondence between velocity and force space. (if) The surfaces represented by the dissipation function are

convex.

( iii) The dissipation function increases monotonically with velocity outside the domain of D-0

(iv) Stable flow will occur if the differential quotient satisfies the inequality

where R=(dd? This is equivalent to saying that the irreversible force ~i'should increase with velocity xk.

(v) The principles of maximum rate of dissipation work and the

maximum rate of entropy production apply. This means that

the system should approach its final state on the shortest

possible way.

The above continuum principles have been used to develop

criteria for predicting metallurgical instabilities in

processing11. Since J determines the dissipation through

metallurgical processes, the dissipation function related to

metallurgical stability is given by J. By putting D=Jin Eq.(12),

one gets the condition for metallurgical stability at constant

temperature and strain: d 5 -

3 d > -

since J - (n/m+1)6.k the above equation for stability becomes

-

he left hand side term is called f (.&)parameter which goes negative when there is a metallurgical instability during

processing. This parameter may be evaluated as a function of

temperature and strain rate and superimposed on the map to

determine the instabilities. While Eq.(14) is a continuum

criterion, the manifestations of the instability may be different

in various materials. The well-known manifestations are adiabatic

shear bands, localized shear bands, Luders bands, kinked flow

bands, and flow rotations. On the basis of this, one can generate

the instability map and it can be superimposed on the efficiency

map which helps avoidunstable regions during processing.

1.4 t LITERATURE SURVEY ON HIGH TEMPERATURE DEFORMATION OF FERROUS MATERIALS:

Scanty data is available on warm working of ferrous

materials although there has been growing interests of automotive

industry in warm precision forging components1g

Creep studies on Iron and some steels have been reported in

the literature2r5r9 but these are at very low strain rates

where as warm working is carried out at higher strain rates.

However, creep study on pure iron by Karashima etal? shows that

change in temperature dependence of steady state creep rates in

ferro and para magnetic temperature regions is quite similar to

that of diffusion rates. This is attributed to the magnetic

effect on diffusion rate. The activation energy for creep (72

kcal/mole) was comparable with that for self diffusion in para

magnetic region whereas it was 81 kcal/mole in ferromagnetic

region. Also, similar change in creep rates near curie

temperature has been observed in Fe-Mo, Fe-Si and Fe-Co alloys.

In ferromagnetic region, creep rate decreased with increase in

alloy contents. This could be important for high temperature

deformation even at higher strain rates.

Crowther and ~intzl~in their study on hot ductility of steels

in the temperature range 550-950c (if one considers the

transformation temperature from CC to J phase which is 912Oc, one may call the deformation in ferrite region also a hot

working) have shown that carbon content plays an important role

in these series of plain C - Mn steels. Above 0.28% C, the change

in fracture mode suggested increasing the activation energy and

critical strain for DRX favouring linking of cracks formed by

grain boundary sliding. Torsional ductility of pure irons have

been studied by Robbins et a1. l3 in the temperature range of 600-

1200 C and 4 of 0.5 sol. They have found that 4 -iron is more ductile at elevated temperatures than )'-Iron which is attributed

to its bcc structure with higher rate of diffusion and more

number of slip systems. Also the ductility peak was observed at

8 0 0 ~ ~ in

which is attributed to the DRX. Simon~sen and ~ossin" have

studied the mechanical properties of zone refined iron upto 950c

with a & range of 3.3 x to 3.3 x 10-I .-I. They have reported the elongation maximum at 830c with of 3.3 x s-I

. Hot compression of armco iron and silicon steel at range of

0.05 to 1 S-I and temperature range of 600-1000~~ has been

carried out by Uvira and 3onas16. The calculation of activation

energy indicates that hot compression is a diffusion controlled

thermally activated process and the strain is not affected by

changes in temperature at constant k . High temperature deformation of o( -Iron has also been studied by Glover and

~ e l l a r s ~ ~ , in the temperature range of 5 0 0 - 8 0 0 ~ ~ over range of

2 in torsion. Their study has revealed DRX in < -Iron at low stresses. The results have been discussed in terms of a model for

DRX along with detailed microstructural evidences. Above the

Curie temperature (h.7700~), the activation energy for creep of

pure iron is greater than that for self diffusion because of

dynamic recrystallization (DRX) 'I. ~ l s o , ~ a ~ e n $ r ~ ~ has suggested that DRX occurs in high k deformation of O.1C steel. Kenne et a1. 23 have concluded that above 600c, the restoration process in

ferrite range of pure iron was DRX. Torsional hot workability has

been studied by Matuszewski et a1.18 in 0.45% C steel, in the

temperature range of 650 to 870'~ and 2 of 0.17 to 4.85 S-I. Their study also shows DRX in this steel with a ductility peak

at 760c, favouring warm working which is confirmed by Reynolds

and ~a~lor'~. Other advantages of warm working are mentioned and

these are better utilization of material (less redundant work),

improved surface finish and dimensional accuracy compared to hot

working and reduced press loads compared to cold working. Also,

better microstructura~ Control in as forged condition eliminates

the need for hardening and tempering treatments if worked in

ferrite region.

srinivas et a1.20 in their study on effect of solutes on

resistance to fracture in Armco Iron have reported that 3.5% Si

additions decrease the fracture toughness by 70% while 5% Co

addition increases it by 35%. Secondly, Si increases the yield

strength two fold where as Co decreases it by half as compared

with Armco Iron.

The survey revealed that there is a wide scope for

optimization of process parameters( T, ) in warm working of ferrous materials in the ferrite region and a study with wide

range of strain rate would add to a better understanding of warm

working behaviour of these materials.

1.5 AIM OF THE INVESTIGATION AND APPROACH

The aim of the present investigation is to study the warm

working characteristics of alpha iron(bcc), with a view to

understand the mechanism of deformation and to optimise the

process parameters, namely, temperature and strain rate. A

further aim of the investigation is to examine the effect of

alloying additions on the warm working behaviour of alpha iron

with a view to understand the response of alloy steels, during

warm working. Three binary alloys of iron, namely Fe-Si, Fe-Co

and Fe-Ni are chosen for this investigation. It is well known Si

additions stabilize alpha phase while Ni additions stabilize

gamma phase. Co on the other hand, does not significantly a fect

the transformation temperature. Furthermore, the Curie

temperature is also influenced differently by these alloying

additions24.

The approach adopted in this investigation involves

developing of processing Maps, using the concepts of Dynamic

Materials Modelling. Processing Maps consist of the variation of

the efficiency of power dissipation [q = 2m / (rncl) 1 , where m is the strain rate sensitivity of flow stress, as compared to an

ideal linear dissipator, with temperature and strain rate. The

input required for developing the processing map is the variation

of flow stress with strain rate in a wide range at different

temperatures in the warm working range, using compression tests.

CHAPTER - I1

EXPERIMENTAL

11.1 Materials and Specimen Preparation

Armco Iron and Fe-5Si, Fe-O.SCo, Fe-5Co, Fe-O.5Ni and Fe-5Ni

alloys were selected in this investigation. The chemical

composition and initial grain sizes are listed in Table 11.1 for

all the materials. Alpha iron was forged at 9 0 0 ~ ~ and annealed

for 2 hrs at 7 5 0 ~ ~ followed by furnace cooling. The binary alloys

were hot forged at 700c, annealed at 8 7 5 O ~ and furnace cooled.

Rods of 12mm diameter were the starting material for

machining specimens of geometry shown in Fig.II.l for compression

testing. Cylindrical Specimens of lOmm diameter and 15mm height

(8mm diameter and 1 2 m height for Ni alloy) were machined such

that their faces were parallel. Concentric grooves of about 0.5mm

depth were engraved on the specimens faces to facilitate the

retention of the lubricant. A lrnm45O chamfer was given to the

edges of the face to avoid fold over in the initial stages of the

compression. A 0.8mm diameter hole was drilled to a depth of 5mm

at half the height of each specimen for inserting a thermocouple

for measurement of actual sample temperature.

11.2 Compression Testing

The compression tests were carried out on a computer

controlled servo hydraulic (DARTEC, UK) machine with 5100kN load capacity. The actuator speed of the machine can be varied from a

minimum of 0.0003~m/s to a maximum of 1250mm/s with a maximum

stroke length of 50mm. ~t can be controlled to follow the

programmed variations with time. All the compression tests were

Table 11.1

chemical composition (Wt%) and Initial grain size of the materials used :

Grain Material C Mn S P Si Co Ni Fe Dia.

rm

Armco Iron 0.007 (0.03 0,005 0.003 -- -- -- Bal 118

Fe-5Si I .I -, do -- ...- 5 O m Bal 140 --

Fe-5Co -- ,, do -0 - 0 -0 -0 5 Bal 125 Fe-0 . 5C0 -- -- do o w 0- 00 0.5 -- Bal 130

Fe-5Ni -- -- do -- - - -0 0 0 5 Bal 100 Fe-O.5Ni W e ,, do -- -- ow (.. - 0.5 Bal 100

GROOVES

carried out at a constant true strain rate (i) in the range of 0.001 to 100 s-'- Platens made of Mar M-ZOO superalloy material

were used.

A three zone furnace with proportional temperature

controllers for each zone was used in order to have large uniform

temperature zone. A maximum of upto 1 1 0 0 ~ ~ can be attained with

this furnace. The temperature control was within +2Oc at all the

testing conditions. The adiabatic temperature rise during

compression was recorded using an icolet storage oscilloscope,

with the help of the thermocouple embedded in the specimen. The

specimens were compressed to a true strain of about 0.5.

Armco Iron and the five alloys were tested in the

temperature range 4 0 0 - 9 0 0 ~ ~ at 5 0 / 1 0 0 ~ ~ intervals and strain rate

range 0.00l-l00s'~ in a decade intervals. Molybdinum disulphide

with graphite was used as the lubricant for all the specimens.

11.3 Development of Processing Maps

Power dissipation maps are generated on the basis of

experimental data of flow stress as a function of temperature and

strain rate in a wide range. The variation of flow stress with

strain rate is fitted using a spline fit and the strain rate

sensitivity (m) as a function of strain rate is calculated from

the slope. Such data were obtained at various temperatures and

were used for the calculation of the efficiency of power

dissipation ( q ) using Eq.(6). On the basis of efficiency data as a function of strain rate and temperature, contour maps were

obtained using computer graphics. The instability maps were

generated on the basis of the continnum criterion given by

~q.14 and were also plotted as contour maps for different

values.

11.4 ensile Testing

cylindrical tensile samples were used to measure the hot

ductility values for Armco Iron. The gauge length of the sample

was 30mm and the diameter at the cross-section was 3mm. The

tension tests were conducted at a strain rate of 0.001s-~. The %

elongation at failure was taken as a measure of ductility.

11.5 Metallography and Grain Size Analysis

The deformed specimens were secti~oned parallel to the

compression axis and microstructural examination and grain size

measurements were conducted using standard metallographic

techniques.

Microstructures of air cooled specimens deformed at

selected regions of the power dissipation maps were examined. The

Olympus optical microscope was used for documenting the

microstructures.

All the specimen were pre-polished (ground) on different

grit emery papers and then polished on wheel with a Billiards

cloth and alcumina plus water as a lubricant. Finally, diamond

paste was used with kerosene as a lubricant on Selvyet cloth

before etching it with 4% Nital solution for about half a minute

for grain size and other microstructural observations in the

optical microscope.

Heyn intercept method26 was used for the grain size

measurements. The eye-piece cross-wire length was calibrated and

the magnification was chosen such that the number of grain

boundaries intersecting the line can be counted accurately. The

combination of line length and magnification was chosen in such a

way as to produce at least 15 intersections per field in order to

get an accurate estimate of average grain size measured in

micrometers (pm). In each material, number of measurements were repeated several times to get correct data.

CHAPTER I11

RESULTS AND DISCUSSION

I 1 ALPHA IRON

~ypical true-stress-true strain curves recorded at 5 0 0 ~ ~ and

8 0 0 ~ ~ at different strain rates for alpha iron are shown in

~ig.111.1(a) and ig.III.l(b), respectively. At 500c, the

material showed strain hardening at all strain rates. At 800c,

the curves corresponding to strain rates of 0.001 and 0.01 s-'

showed stready-state behaviour while at higher strain rates,

strain hardening is observed.

The variation of flow stress ( 4 ) with temperature (T) , strain rate (1) and strain ( ) for alpha iron is shown in Table 111.1. The flow stress is corrected for the adiabatic

temperature rise using linear interpolation of log4 vs. (1/T)

data at constant strain and strain rate and this correction was

found to be significant at lower temperatures and higher strain

rates.

The power dissipation map for alpha iron is shown in

Fig.III.Z(a) as a contour map. The maps obtained at other strains

are similar indicating that strain effects are not

significant. Similar behaviour was observed on other materials

investigated.

The power dissipation map for alpha iron exhibits two

domains. One in the temperature range 6 0 0 - 8 5 0 ~ ~ and strain rate

range 0.001-0.1~-~ with a peak efficiency of 35% at 8 0 0 ~ ~ and

10'~s". Another small domain at 4 0 0 ~ ~ and a strain rate of

0.001s-' has an efficiency of 27%.

TRUE PLASTIC STRAIN

TRUE PLASTIC STRAIN

Fig.III.1 True Stress-True plastic Strain Curves for Alpha Iron obtained in compression at (a) 5 0 0 ~ ~ and (b) 800c, a t different Strain rates.

Table III.1: Flow stress values (in MPa) of alpha-iron at different stram rates and temperatures for vanous strams (corrected for ad~abatlc temperature rise)

Stram Strain rate, s"

Temperature, O C 400 500 600 700 800 900

TEMPERATURE :L

TEMPERATURE ,OC

Fig. 111.2 (a) The power dissipation map for alpha iron (numbers represent efficiency in per cent.)

(b) Instability map for alpha iron.

The domain with the peak at 8 0 0 ~ ~ represents the process of

dynamic recrystallization. This interpretation is confirmed by

the observations described below concerning the variation of

grain size in the domain and the ductility of the material.

(i) Grain Size Variations :

A typical microstructure in the DRX domain for the alpha

iron is shown in Fig.III.3 (a) which corresponds to the peak

efficiency conditions (800c, 1 0 ~ s ) . The variations of the

average grain diameter with temperature at the strain rate

corresponding to the peak in the DRX domain (10'~s"') is shown in

Fig.III.4. The data is for air cooled specimens. However, the

variation for water-quenched samples showed similar behaviour

although the grain size was slightly finer. The grain size

increases with temperature following a sigmoidal variation up to

peak efficiency DRX temperature ( 8 0 0 ~ ~ ) beyond which, there is a

abnormal grain growth. Similar grain size variations were

recorded in the DRX domain of several other materials25.

(ii) Ductility :

The variation of ductility in torsion with the temperature

was measured by Robbins et a1.13 at a strain rate of 0.5 s-' and

this is shown in Fig.III.4. Also, the measured tensile

ductility values at a strain rate of 0.001 s'l at different

temperatures from the present investigation are plotted in the

same figure. Both the profiles show a ductility peak at 8 0 0 ~ ~

which matches with the temperature for peak efficiency in the DRX

domain. As the grain size increases, ductility drops sharply

beyond 8 0 0 ~ ~ . All these observations confirm the correlation

between the ductility and efficiency of power dissipation in the

Oig.III.3 (a) Microstructure of alpha iron deformed at BOO~C/O. 001s-' (DRX domain)

(b) Microstructure of the sample deformed at 500~C/loo s"showing localized flow.

ROBIN S et al.

grain growth

- E f FICIENCY

Fig.II1.4 (a) Ductility variation with temperature in alpha iron, (b) Grain Size variation with temperature in DRX domain and (c) Efficiency of power dissipation vs. temperature.

DRX domain. (Fig.III.4)

he peak efficiency in the DRX domain (35%) is lower than

that expected (about 50%) on the basis of the BCC structure for alpha iron which exhibits easy dynamic recovery due to its high stacking fault energy. This may be attributed to the magnetic

domain structures in alpha iron. The grain boundary migration

which is essential for dynamic recrystallization is slowed down

by the presence of magnetic domains below the Curie

temperature (-77 OOC) . The migrating boundary has to overcome the strong electron spins in the direction of magnetization. The

efficiency of power dissipation is lower as some energy is spent

to reorient the magnetic domains by the migrating grains. This is

also the reason for higher activation energy observed9 for self

diffusion of Iron in ferrite region .

As the temperature is increased beyond the Curie

temperature, the grain boundaries can migrate uninhibited by

magnetic domain structure giving rise to abnormal grain growth,

lowering the ductility and strength.

The above discussion indicates clearly that the higher

temperature domain observed in the map of alpha iron represents

the DRX process. The steady state behaviour of stress-strain

curves, the ductility data, the efficiency of power dissipation

and the grain size variation are in support of this conclusion.

The DRX efficiency is low because of magnetic domains structure

of alpha iron. The occurrence of DRX in alpha iron was also

reported by Glover et a1.17 on the basis of microstructural study

and kinetic analysis of hot torsion data,

On the basis of the instability criterion given by eq. (14)'

the instability parameter KC&) is plotted as a function of temperature and strain rate to obtain the instability map

(Fig.III.2b). The material will exhibit flow instability when

\(qis negative. According to this criterion, alpha iron will

exhibit flow instability in the temperature range 4 0 0 - 7 0 0 ~ ~ when

the strain rate is above 10s-I and at lower strain rates when the

temperature is around 4 0 0 ~ ~ . Microstructural examination of the

specimen deformed at 5 0 0 ~ ~ and 100s-I (Fig. III.3b) showed that

alpha iron exhibits flow localization in this regime which should

be avoided in processing.

111.2. Fe-561 Alloy

Typical true-stress vs true-strain curves recorded at 6 0 0 ~ ~

and 8 0 0 ~ ~ and at different strain rates for Fe-5Si alloy are

shown in Fig.III.S(a) and Fig.III.5(b) respectively. The curves

at 6 0 0 ~ ~ as well as at 8 0 0 ~ ~ and at lower strain rates, showed

steady state behaviour. The flow stress values at different

temperatures and strain rates are given in Table 111.2. The flow

stress data was corrected for the adiabatic temperature rise.

The iso-efficiency countour map for Fe-5Si alloy is shownin

Fig. 111.6 (a) . A strain of 0.5 is selected for easy comparison

with the Alpha iron Map and the maps at lower strains are not

significantly different . The processing map for this alloy is

very similar to that for alpha iron. The map exhibits two domains

similar to those recorded in alpha iron. The domain occurring

with a peak ef f icienoy of 56% at 800c/0. 001s-I represents

dynamic recrystallization. In comparison with in alpha iron, the

peak DRX efficiency is higher by about 20%. The DRX efficiency in

I I f I I I 0.1 0.2 0.3 0.4 0.5 0.6

TRUE PLASTK: STRAIN

600.

(a) 600 *C ,s-

100 10

1 0

0 1

0 01

Fig. 111.5 True Stress-True plastic Strain Curves for ~Lai- ~ 1 1 0 ~ obtained in compression at (a) 6 0 0 ~ ~ and - 8 4 (b) 800c, at different Strain rates.

120

I om

-

O6 I I I I I

0.1 0.2 0.3 0.4 0 5 C TRUE PLASTIC STRAIN

Table III.2: Flow stress values (rn MPa) of Fe-5Si alloy at drfferent strain rates and temperatures for vanous stram (corrected for adiabatic temperature nse)

Strzun Stram rate, C'

Temperature, "C ' 400 500 600 700 800 900

TEMPERATURE C

TEMPERATURE f c

Fig.III.6 (a) The power dissipation map for Fe-5Si alloy (numbers represent efficiency i n per cent.)

(b) Instability map for Fe-5Si alloy.

Fe-5Si has increased because the Curie temperature decreases due

to Si additions. For a 10% addition of Si to iron, the Curie

temperature decreases from 7 7 0 ~ ~ to 6 0 0 ~ ~ . The decrease in the

magnetisation of the Fe-Si alloys will enhance the grain boundary

migration compared to alpha iron.

The variation of the average grain diameter with temperature

in the DRX domain at 0.001s-~ is shown in Fig. 111.7 and compared

with the efficiency variations. The variation is sigmoidal as is

observed in several other materials25 and is very similar to that

obtained in alpha iron (Fig.III.4b) except for the high

temperature grain growth. Typical microstructure obtained on Fe-

5Si specimen deformed at 8 0 0 ~ ~ and 0.001s-~ strain rate is shown

in Fig. 111.8 (a) . The grain boundaries have a wavy configuration,

typical of DRX microstructures.

The variation of the instability parameter ( 4 ) with temperature and strain rate is shown in Fig.III.G(b). The

instability map shows that the material will exhibit flow

instability in the temperature range 4 0 0 - 6 0 0 ~ ~ when the strain

rate is above 0.1s" and upto 4 5 0 ~ ~ at lower strain rates.

Microstructural examination of the specimen deformed at 4 0 0 ~ ~ and

100s-l[~i~. 111.8 (b) ] showed that adiabatic shear bands occur in the instability regime which should be avoided in warm working of

the material. The instability regime is wider in the Fe-5Si alloy

than alpha iron as seen from a comparison of Fig.III.2(b) and

Fig.III.6(b).

111.3. Pa-Co Alloys

111.3.1 Be-5Co Alloy:

The true-stress vs. true-strain curves for this alloy at 500

GRAIN SIZE -

EFFICIENCY

TEMPERATU RE ,* C

Pig.III.7 (a) Grain Size variation of Fe-5Si alloy with temperature in DRX domain and

(b) Efficiency of power dissipation vs. temperature.

Fig.III.8 {a) Microstructur of Fe-5Si alloy deformed at 800C/0. OOls-P(DRX domain)

(b) Microstruc,f.ure of the sample deformed at 400~~/100s showing localized shear bands.

and 9 0 0 ~ ~ are shown in Fig. 111.9 (a) and 111.9 (b) respectively.

The curves obtained at 5 0 0 ~ ~ exhibit work hardening at rates

decreasing with strain, typical of dynamic recovery. At strain

rates higher than IS", the curves exhibited work hardening while

at strain rates at and below 0.1~"~ steady- state flow is

observed.

The data in Table 111.3 show flow stress variation with

temperature strain rate and strain, corrected for adiabatic

temperature rise.

The processing map for Fe-5Co alloy is shown in

Fig.III.lO(a) and the corresponding instability map is shown in

Fig.III.lO(b). Both are obtained at 0.5 strain for the Fe-SCo

alloy and the maps at other strains are similar.

The processing map for Fe-5Co alloy exhibits the following

domains.

(A) The domain in the temperature range 600-900C and strain rate range 0.001-1s-~ has a maximum efficiency of 33% occurring at

9 0 0 ~ ~ and 0.01s-l.

(B) The domain in the temperature range 400-500~~ and strain rate range 0.001-0.01s-I has a maximum efficiency of 20% occurring

at 4 0 0 ~ ~ and 0.001s".

The microstructure obtained on specimen deformed at 9 0 0 ~ ~ and

o. 1s'' is shown in Fig. 111.11 (a) which shows wavy grain boundary

structure typical of dynamic recrystallization process. The

variation of grain size with temperature at a strain rate of

0.01s-I is shown in Fig.III.12 which shows that the grain size

increases with temperature in a fashion similar to that observed

in alpha iron and Fe-5Si alloys. It is interesting to note that

0 0.1 0.2 0 . 3 0 4 0.5 0 6 TRUE PLASTIC STRAIN

I 1 I I I I 0.1 0.2 0.3 0.4 0.5 O f

T R M PLASTIC STRAIN

Pig.IIX.9 True Stress-True plastic Strain Curves for Fe-5Co Alloy obtained in compression at (a) 500'~ and (b) 900c, at different Strain rates.

Table Il3.3: Flow stress values (m m a ) of Fe-5Co alloy at different strain rates and temperatures for vanous strains (corrected for adiabatic temperature nse)

Strain Strain rate, s-'

-- - - - - -

Temperature, "C 400 500 600 700 800 900

TEMPERATURE :C

Fig.III.10 (a) The power dissipation map for Fe-5Co alloy (numbers represent efficiency in per cent.)

(b) Instability map for Fe-5Co alloy.

Fiq.111.11 (a) Microstruct re of Fe-5Co alloy deformed at -Y 900c/ 0.1s (DRX domain)

(b) Microstruc ure of the sample deformed at 400C/tl.ls-r showing localized flow.

EFFlCl ENCY (b)

Pig.XII.12 (a) rain Slze variation of Fe-5Co alloy with temperature in DRX domaln and

(b) Efficiency of power dissipation vs. temperature.

the DRX temperature corresponding to the peak efficiency has

increased to 9 0 0 ~ ~ in comparison with that in alpha iron and Fe-

~ ~ i ( 8 0 0 ~ ~ ) and the strain rate increased from 0.001 to 0.01s-l.

Also the maximum efficiency (33%) is similar to that in alpha

iron(35%) but lower than in Fe-5Si alloy. These effects may be

attributed the effect of Cobalt additions on the Curie

temperature of alpha iron. The Curie temperature of alpha iron

increases from about 770 to about 9 0 0 ~ ~ with the addition of 10%

Co. Thus Cobalt additions strengthen the magnetic domains of

alpha iron, restrict the grain boundary migration and reduce the

efficiency of DRX. Higher temperatures are therefore required to

achieve dynamic recrystallization. The higher strain rates may be

attributed to the larger grain sizes of the Fe-5Co alloy(-100pm)

than in alpha iron (-lOpm) .

The domain occurring in the lower temperature range (400-

5 5 0 ~ ~ ) represents dynamic recovery process. Under these

conditions, the material exhibits work hardening type stress -

strain curves. The rate of hardening decreases with straln which

is typical of dynamic recovery. The microstructure indicates

dynamic recovery with

associated elongated fine grained structures [Fig.III.13.].

The instability map for Fe-5Co alloy is shown in

Fig.III.lO(b) which shows the variation of the instability

parameter $ ( E } with temperature and strain rate. The material

shows instability when ) is negative. Fe-5Co alloy exhibits

flow instabilities in the temperature range 5 0 0 - 9 0 0 ~ ~ when the

strain rate is higher than about 10s-l.~ho material has exhibited

flow localization under these conditions. At lower

Z'fg.111.13 Dynamic rec very microstructure of Fe-5Co alloy a t 500oC,.oCls'~.

temperatures (400-500~~) , the material exhibits cracking along adiabatic shear bands at strain rates higher than 1.0s" and flow

localization at lower strain rates [Fig.III.ll(b)]. All these

instability regimes should be avoided in processing.

with a view to evaluate the effect of Cobalt concentration on

the warm working characteristics of Fe-Co alloys, the processing

map for Fe-O.5Co has been developed. The map IS shown in

Fig. 111.14 (a) and the instability map in Fig. 111.14 (b) . The

processing map exhibits a single domain in the temperature range

5 0 0 - 9 0 0 ~ ~ and strain rate range 0.001-1.0s-~ with a maxrmum

efficiency of 36% occurring at 900c/0. 001s-'. This represents

DRX of this alloy. A comparison of this map with that for Fe-

5Co [Fig.III.lO(a) ] clearly shows that the DRX behaviour is not significantly different.

The Fe-0.5Co alloy exhibits flow instability in the

temperature range 400-900~~ when the strain rate is above 1s".

In comparison with that in Fe-SCo, the instability is less

intense and occurs in a narrower range of temperature and strain

rate.

111.4 Pe-li alloys

111.4.1 Fe-5Ni:

The true-stress vs. true-strain curves obtained on Fe-5Ni

alloy at 700 and 9 0 0 ~ ~ are shown in Fig.III.lS(a) and 111.15 (b) .

The curves do not show any significant work hardening. On the

other hand, flow softening is observed at higher temperatures and

lower strain rates, The flow stress data are given in Table 111.4.

Pig. 111.14 (a) The power dissipation map for Fe-O.5Co alloy (numbers represent efficiency in per cent.)

(b) Instability map for Fe-0.5Co alloy.

TRUE PLASTIC STRAIN

TRUE PLASTIC STRAJN

Pig.XII.15 True Stress-True plastic Strain Curves for Fe-5Ni Alloy obtained in compression at (a) 7 0 0 ' ~ and (b) 900%, at different Strain rates.

The processing map for Fe-5Ni alloy is shown in

Fig.III.lb(a) which exhibits a single domain in the temperature

range 5 5 0 - 9 0 0 ~ ~ and strain rate range 0.001-10s" with a maximum

efficiency of 38% occurring at 900~/0. 001s-~. The processing map

is continuous inspite of the occurrence of the dual phase

(alpha+gamma) field in the temperature range 580-780~~. This

domain represents dynamic recrystalization of the gamma phase

with a maximum efficiency at 9 0 0 ~ ~ and 0.001s-~ where a single

phase (gamma) field exists.Typica1 DRX microstructure obtained on

a specimen deformed at 9 0 0 ~ ~ and 0.001s-~ is shown in Fig. 111.17

(a). The grain size variation with temperature at a strain rate

of 0.001s-~ is shown in Fig.III.18 which exhibits a discontinuity

at 7 8 0 ~ ~ corresponding to the (alpha+gamma) to gamma

transformation.This curve indicates gamma grain refinement at

temperatures above 7 8 0 ~ ~ .

It may be noted that Nickel is a gamma stabiliser and hence

the magnetic effects that influenced the DRX of alpha iron are

absent.The instability map for ~e-5Ni alloy showing the variation

of \ C ) with temperature and strain rate is given in Fig. 111.16 (b) . The regimes where E ) is negative correspond to flow

instabilities. Fig.III.l6(b) shows that the Fe-5Ni alloy exhibits

intense flow instabilities in the temperature range 4 0 0 - 9 0 0 ~ ~

when the strain rate is above 0.1s". In the temperature range

400-600c, the material exhibits instabilities even at lower

strain rates(>0,001). Flow localization associated with cracking

occurs in these regimes [Fig.III.l7(b)].

A comparison of the instability maps for alpha

iron[Fig.111.2(b) 1, ~e-5co[Fig.IXI.lO(b)] and Fe-5Ni

TEMPERATURE : C

A -0.88 B -0% C -0 63

TEMPERATURE :C

F%g.111.16 (a) The power dissipation map for Fe-5Ni alloy (numbers represent efficiency in per cent.)

(b) Instability map for Fe-5Ni alloy.

Fig. 111.17 (a) Microstructure of Fe-5Ni alloy deformed at ~ O ~ ~ C / O . O O ~ S - ~ ( D R X domain)

(b) Microstructure of the sample deformed at 4 0 0 ~ ~ and d . 1 ~ - l showing localized flow and cracking.

EFFICIENCY (b)

TEMPERATURE PC

1 1 1 8 (a) Grain Size variation of Fe-5N1 alloy w i t h temperature in DRX domain and

(b) Efficiency of power dissipation vs. temperature.

alloy[Fig. 111.16 (b) ] indicates that F e - 5 N i alloy has a wide instability regime and hence is not very suitable for warm

working at temperatures lower than 600'~. Even at higher

temperatures, the processing has to be done at strain rates

lower than 0.1s".

111.4.2. Fe-0.SNi:

For the purpose of finding the influence of Nickel content on

the warm working characteristics of Fe-Ni alloys, the processing

map for Fe-0.5Ni alloy was established. The map is shown in

Fig.III.lg(a). The map exhibits a domain in the temperature range

5 5 0 - 9 0 0 ~ ~ and strain rate range 0.001-10s-~ with a maximum

efficiency of 31% occurring at 900~~/0.1s'~. This domain is

essentially similar to that observed1 in Fe-5Ni alloy except that

the peak has shifted to higher strain rates (0.001 in Fe-5Ni to

0.1 in Fe-O.5Ni). This domain represents gamma dynamic

recrystalization which occuvs at higher strain rate when the

Nickel content is less.

Similar to Fe-5Ni alloy, intense cracking occurs at strain

rates higher than 1.0s" in the temperature range 400-500~~.

The instability map for Fe-O.5Ni alloy is shown in

Fig.III.lg(b). Flow localization occurs in the temperature range

400-500~~ at all strain rates in the range 0.001 to 0.1s-'.A$ strain rates higher than 0. ls", the material exhibits cracking

along adiabatic shear bands. The instability behaviour is similar

to that observed in Fe-5Ni alloy.

The results show that nickel additions are not favourable to

conduct warm working of alpha iron.

TEMPERATURE :C

TEMPERATURE (: C

Pig.IzI.19 (aJ The power dissipation map for Fe-0.5Ni alloy (numbers represent ef f ic iency in per cent.)

(b) Instability map for Fe-O.5Ni alloy.

CHAPTER IV

SOKMARY AND CONCLUSIONS

The warm working characteristics of alpha iron, Fe-Si, Fe-Co

and Fe-Ni alloy were studied in the temperature range 400-900~~

and strain rate range 0.001-100s-~. on the basis of the flow

stress data obtained as a function of temperature and strain rate

in compresslon, power dissipation maps and instability maps were

developed. The following conclusions are drawn from thls

investigation.

1. Alpha iron undergoes dynamic recrystalization in the

temperature range 600-850'~ and strain rate range 0.001-0.1~-~

with a maximum efficiency of 35% occurring at 8 0 0 ~ ~ and 0.001s-~.

At these conditions, the ductility reaches a peak value.

2. The lower than expected efficiency value for DRX is

attributed to the restrictive effect of magnetic domains to the

migration of grain boundaries.

3. Alpha iron exhibits adiabatic shear bands in the temperature

range 4 0 0 - 7 0 0 ~ ~ when tho strain rate is above 10s-l. This

instability regime should be avoided in warm working of the

material.

4 . Addition of Silicon increases the efficiency of power

dissipation for DRX without changing the DRX of alpha

iron. This result is attributed to the lowering of Curie

temperature by Silicon additions.

5. Cobalt additions increase the DRX temperature of alpha iron

by about 1 0 0 ~ ~ and this effect is also caused by the

stabilization of magnetic domains by Cobalt due to an increase

in Curie temperature.

0, 6. Nickel additions stabilize gamma phase which dynamiclly

4 recrystalizes at 900c and 0.001s-~ for Fe-5Ni and at

~ O O ~ C / O . 1s-I for Fe-0.5Ni.

7. Fe-Ni alloys exhibit wide regimes of flow instability and

therefore are not suitable for warm working.

8. The effect of change of concentration from 0.5 to 5% of Co

and Ni to alpha iron does not significantly change the warm

working characteristics.

REFERENCES

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