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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.
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