1. 1
IIWIC 2008, Chennai 8-10, Jan. 2008 370-380
A Differential Scanning Calorimetry (DSC) study of welding
consumables for modified 9Cr-1Mo ferritic steel
B. Jeyaganesh1
, S. Raju1*
, S. Murugesan, E. Mohandas1
,
M. Vijayalakshmi1
, V. Ramasubbu2
, S. K. Albert2
and A. K. Bhaduri2
1
Physical Metallurgy Division,
2
Materials Technology Division
Indira Gandhi Centre for Atomic Research (IGCAR), Kalpakkam, 603 102, India
• Corresponding Author; E-mail : sraju@igcar.gov.in
• Tel : 91 44 27480306; Fax : 91 44 274 80 081
Abstract
A comprehensive characterisation of the thermal stability of four welding
consumables that are developed for welding mod. 9Cr-1Mo steel has been performed
using DSC. The various transformation arrest points such as Ac1, Ac3, the solidus and
liquidus temperatures and the enthalpy effects associated with the α→γ phase change
and melting have been determined precisely. It is found that both Ac1 and Ac3
temperatures exhibit a decrease with increasing (Mn+Ni) content. In addition, the
kinetics of austenite formation from a tempered martensitic microstructure has been
modelled in terms a simple isochronal version of the Kolmogorov-Johnson-Mehl-
Avrami (KJMA) formalism for the nucleation and growth phenomenon. It is found
that the apparent activation energy for austenite formation is sensitive to the heating
rate, in that the activation energy is more for slower rates of heating. An apparent
activation energy of about 260 kJ mol-1
is suggested for the α→γ transformation in
high chromium steels
(Key words : mod. 9Cr1Mo steel, welding electrode, phase transformation, DSC, activation energy)

2. 2
1.0. Introduction
A study on ferritic steels is of interest on both and applied grounds. On the
application front front, the use of Cr –Mo based ferritic steels in fossil fuel power
plants has a long history [1]. The combination of good thermal conductivity and low
thermal expansion [2] of ferritic alloys is particularly attractive from the point of view
of minimising the hazard from thermal stresses [1,3]. Besides, ferritic steels also
possess fairly good oxidation resistance at intermediate temperatures [3]. More over
the relatively low cost of ferritics as compared to 300- series of high nickel austenitic
stainless steels is an added economic incentive. The emerging pressure from both
economical and ecological quarters to enhance the thermal efficiency of conventional
fossil fuel fired power plants has only re-emphasised the need for developing
advanced grades of ferritic alloys that can function for longer durations at higher
temperatures, in the range 600 - 650 0
C [1,4]. The improved neutron irradiation
swelling resistance of ferritics over austenitic steels has also made them the preferred
choice as core structural candidates for realising an extended fuel burn up in liquid
metal cooled fast reactors, employing metallic fuels [5]. As a passing remark, we may
also note that ferritic steels have aroused significant interest among the fusion reactor
materials community as potential plasma facing first wall material [6]. On the basic
front, the various phase transformations that occur in ferritic steels as a function of
both temperature and composition and their influence on resulting microstructure
constitute an interesting play field for studying many physical metallurgical
phenomena.
It is a well-known fact that welding constitutes a major step in the fabrication of
power plant components. Therefore designing suitable welding process and welding
consumables in conjunction with appropriate pre and post weld heat treatment

3. 3
(PWHT) procedures is essential in ensuring quality welds. The advent of the newer
grades of ferritic steels in nuclear materials scenario [6-11] has catalysed an extensive
basic research programme on welding metallurgy, especially on those aspects that are
related to tailoring of composition of welding consumables with a view to minimise,
or avoid if possible, the formation of delta ferrite during the course of primary weld
solidification [12, 13]. Thus for example, a primary austenitic solidification mode
may be ensured by adjusting the combined concentration of austenite stabilising
elements such as, (Ni + Mn + Co + Cu + N + . .) in the overall composition of
welding consumables [13]. In some cases, the partial replacement of Mo by W, a
comparatively weak ferrite stabiliser [14], is also practised in addition so that
preferably nil or only a small delta ferrite is retained at the end welding process [15].
It must also be kept in mind that in tailoring the compositional limit of austenite
stabilisers, the attending penalty on the resulting low Ac1 transformation temperature
must also be given due consideration, for Ac1 decides the theoretical upper limit for
the post weld heat treatment (PWHT) temperature [13, 16, 17]. In addition to what has
been argued above, an appropriate balancing of the composition of welding
consumable is also called for in ensuring proper tempering of martensite during
PWHT, since the tempered microstructure resulting from the adoption of a tempering
temperature that is marginally higher than Ac1 often leads to reduced fracture
toughness and embrittlement during service [17, 18]. Thus, argued in any manner, a
comprehensive metallurgical characterisation of welding consumables is called for
from the standpoint of drawing appropriate welding specification.
In the construction of India’s first 500 MWe prototype fast breeder reactor
(PFBR), mod. 9Cr-1Mo steel has been identified as the candidate for the steam
generator tubes. Towards the cause of devising appropriate welding procedure, a few

4. 4
welding consumables with carefully tailored compositions have been chosen and
tested. As part of the broad based characterisation programme, we have undertaken in
this study a detailed differential scanning calorimetry (DSC) investigation of the
thermal stability and phase transformation characteristics of the welding consumables.
The results of this study are reported in this paper.
2.0. Experimental Procedure
2.1. Chemical composition, weld pad preparation and metallography
In table 1, the chemical composition of the four welding consumables, as
determined by direct reading optical emission spectroscopy is listed. These
consumables are : the E 9016 shielded metal arc welding (SMAW) electrodes
obtained from Kobe, Boehler and Midhani (E9016 – B9) and the filler wire for gas
tungsten arc welding (GTAW) obtained from M/S. Midhani, India. For comparison,
the composition of the base metal has also been presented in table 1.
A weld pad of dimension, 100 × 75 × 12 mm is made by the SMAW process for
the electrodes and by the GTAW method for the filler wire using the stringer bead
technique. A total of eight layers were built without any interpass treatment. The
following process parameters are adopted in making the weld pads: 90 A/25 V, heat
input of 900 – 1000 J for SMAW of electrodes; and 80 A/15 V, heat input of 300 J for
GTAW of the Midhani filler wire [19].
The samples for the present study were sliced from the portion that is located
approximately 2 mm below the top surface of the weld pad. For optical metallographic
characterisation, the specimens were prepared in the standard manner, starting from
coarse dry grinding to 1000 grade emery polishing and then to final finishing with
alumina and diamond polish. These were given a mild initial etching with a 2% nital
solution and then subsequently by the Villela’s reagent for about a minute. These

5. 5
polished samples were also used for x-ray diffraction (INEL, Cu Kα,) studies and
microhardness (Leitz, 100 g) measurements (table 2).
2.2. DSC experiments
The samples for DSC experiments were cut from the thin sections sliced originally
from the weld pad using diamond coated wire saw. These were further cleaned and
polished to regular and nearly identical cubes of mass varying from 50 to 70 ± 0.1 mg.
The DSC experiments were performed with Setaram Setsys 16®
heat-flux type high-
temperature differential scanning calorimeter, employing recrystallised alumina
crucibles of about 100µ L volume. The crucibles were ultrasonically cleaned in
methanol and rinsed with dilute HCl to remove any trace of contaminants arising from
previous use. The equipment details as well as the calibration procedure have been
discussed in our previous publication [20, 21]. Stated briefly, the experiments were
performed under a constant flow (50 ml per minute) of high purity argon. A range of
heating rates varying from 1 to of 100 K per minute is employed. While lower heating
rates (1-2 K min-1
) resulted in mild decarburisation of the sample due to the
appreciable resident time at high temperatures, high heating rates induce significant
thermal lag between the actual sample temperature and that sensed by the non-contact
probe. However, by calibrating the sample signal against the pure iron (Fe- Aldrich,
impurities ≤ 80 ppm) reference signal, especially around the region of phase
transitions, the temperature lag has been duly accounted for. It is generally found in
the case of 9Cr –1Mo ferritic steels, that for a heating rate of about 10 K min-1
, the
thermal lag is of the order of ≤ 2 K [22]. The sample mass and shape have been kept
nearly the same for both the reference and steel samples used in each set of
experiments. Fresh samples were used for each run, and multiple runs under identical
conditions are performed for calibrating the precision of the measured transformation

6. 6
temperature. The temperature calibration is performed using recommended high pure
melting point standards, namely, Sn, Al, Pb, In and Au. In addition, the measurement
of the enthalpy of α (bcc) → γ (fcc) allotropic transformation in pure iron under
identical experimental conditions is also employed for the heat flow calibration [22].
The heat flow rate or specific heat calibration is performed using the literature data on
the specific heat of pure iron [23, 24]. The measured transformation temperatures are
accurate to ± 1 K; and the transformation enthalpies even on a conservative estimate
are accurate to ± 5 % at 10 K min-1
. Since, a reliable and reproducible temperature and
heat flow calibration of DSC for cooling cycle is comparatively a difficult proposition
[25] due to the absence of reliable high temperature standards with known degrees of
undercooling, the enthalpy and the heat flow calibration of the signal were carried out
only for the heating cycle in the present study. It is presumed that the accuracy of
sample temperature measurement during cooling cycle is just the same as that of the
heating cycle [22].
3.0. Results
3.1. Optical microscopy, microhardness and x-ray diffraction studies
In figure 1(a-e), the optical micrographs of the three electrodes obtained from
Kobe, Boehler and Midhani and the filler wire supplied by Midhani are presented.
The characteristic martensite microstructure is generally seen in all the cases.
Although, no interpass treatment is allowed in the present study, the samples having
been sliced at about 2 mm below the top surface, might have suffered a mild auto
tempering due to the heat effects generated during subsequent welding passes. In the
case of Boehler electrode (fig 1a), the presence of a nicely outlined prior austenite
grain structure with some sporadic coarse carbide particles is readily apparent. The
microstructure of the Kobe electrode (fig 1b) on the other hand reveals a finer lath

7. 7
structure with no clear traces of prior austenite boundaries. The midhani electrode (fig
1c) reveals a coarser and somewhat patchy martensite while the microstructure of the
filler wire (fig. 1d) indicates considerable globularisation (fig. 1d). The base metal
microstructure (fig. 1e) is typical of tempered martensite with extensive decoration of
the ferrite subboundaries by numerous fine carbide particles. It should be mentioned
that this varying spectrum of microstructure witnessed at optical level is due to the
different extent of the residence time that each composition spends in different phase
fields during its thermal history [13]. In table 2 the microhardness values measured on
each sample under a load of 100 g are tabulated.
In figure 2, the room temperature x-ray diffraction profiles obtained using Cu-Kα
radiation for different samples are stacked together. The predominant peaks of α-
ferrite are clearly marked. Besides, a few very weak reflections from the (Fe,Cr)23C6
are also observed [26]. The estimated lattice parameter for the α-ferrite phase is found
to be 0. 2876 ± 0.0003 nm. It may be noted that this value is nearly the same for all
the welding consumables and is in agreement with that observed for 9Cr-1Mo low
carbon steel [20].
3.2. DSC studies : transformation temperatures and transformation enthalpy
In figure 3(a) a typical DSC profile obtained with the Midhani electrode is
illustrated. In 3(b), an exclusive, expanded view around the high temperature region is
illustrated for all the four welding consumables. In order to avoid clutter and needless
proliferation of figures, the thermogram of only one electrode composition is shown
in fig. 3(a). The details for others are tabulated (table 3). It must be mentioned that the
qualitative features of the DSC profiles are much the same for all compositions. Only
the onset and peak temperature values of various phase transformations are dependent
on the composition. A careful inspection of figure 3(a) reveals the following sequence

8. 8
of phase changes that take place upon slow heating (3 K min-1
) of a tempered
martensitic microstructure.
(i) ferromagnetic to paramagnetic transformation at TC, the curie point
of the α-ferrite phase. Depending on the composition, TC is found to
vary from 730 to 740 0
C. TC is more clearly attested during heating
than on cooling.
(ii) α-ferrite + M23C6 + MX → γ-austenite + M23C6 + MX
on crossing the Ac1 temperature. The transformation reaches a fair
degree of completion, depending on the heating rate, at Ac3. It has been
observed in high chromium steels that practically a little dissolution of
M23C6 alone is realised in this reaustenitisation step [27-30], although
the Thermo Calc®
based equilibrium simulations suggest small, yet
finite dissolution (5 %) by the time Ac3 is reached on very slow heating
[30-34].
(iii) γ + M23C6 + MX → γ + MX; (dissolution of M23C6 carbide).
Basically, the dissolution of carbide is a continuous reaction after
Ac3, and is generally seen as a shallow endothermic trough in an
otherwise smooth DSC profile [20]. This is observed better in slow
heating rate scans [22].
(iv) γ + MX → δ - ferrite + γ + MX; appearance of high temperature δ -
bcc phase.
(v) δ + γ + MX → δ + γ ; (dissolution of MX phase in γ )
(vi) δ + γ → liquid + δ + γ . (appearance of liquid)
(vii) liquid + δ + γ → liquid + δ (dissolution of γ in liquid)
(viii) liquid + δ → liquid (completion of melting)

9. 9
The region covered by the last three arrest points signifies the solidus – liquidus gap.
In the case of (9-12) Cr-Mo steels with a nominal carbon concentration of about 0.1
wt.%, the calculated equilibrium diagram(s) using Thermo Calc®
software and
associated steel database [34] suggests that one may encounter a very small existence
domain of the (liquid + δ + γ ) three phase field just prior to attaining the (δ + γ) two
phase region on slow continued heating [35]. The actual attainment is of course
dependent on kinetic factors [36]. This aspect has been kept in mind, while annotating
the high temperature region of the DSC profiles shown in figure 3(b). As can be seen,
the presence of a higher austenite stabilising (Mn + Ni + N) content in our weld
consumables as against the typical mod. 9Cr-1Mo base metal composition such as
P91 or E91, supports the possibility that a three phase (liquid + δ + γ ) equilibrium
field is in principle realised during slow (1-3 K min-1
) heating experiments. In figure
3(b), it is also interesting to note the subtle difference in the nature of the DSC
profiles for different consumables around (liquid + δ ) region. A somewhat flat
profile, such as the one witnessed for the midhani filler wire suggests that this
composition may be very close to the L + δ → L peritectic point, as compared to
other compositions. But this observation needs further scrutiny.
In the case of mod. 9Cr-1Mo steel, the M23C6 and MX phases refer
respectively to (Cr,Fe)23C6 and (V,Nb) (C,N) type mixed carbides and carbonitrides
[37, 38]. It must also be mentioned that for a given DSC instrument and a well
prescribed set of experimental conditions, the sharpness or the clearly defined nature
of various transformation arrest points is a direct function of the associated enthalpy
effect [25]. Unlike the case of α → γ structural transformation or melting, the other
phase changes such as the dissolution of M23C6 and MX phases, are accompanied by
only a meagre change in enthalpy [22]. Therefore, these changes manifest as mild yet

10. 10
distinct inflections in the base line compensated DSC profile [22]. In fact the various
transformation arrest temperatures are fixed by taking recourse to the derivative of the
actual DSC signal, in which these inflection points are more clearly marked.
The enthalpy effects associated with α → γ change and melting are given
directly in terms of the total area enclosed under the respective peak [39]. The
conversion factor required for getting enthalpy in terms of J g-1
is obtained from the
corresponding signal from pure iron [22]. The measured enthalpy values are listed in
the last two rows of table 3. As for the α → γ phase change is concerned, there is
reasonable agreement among the transformation enthalpy values recorded for the
midhani electrode, midhani filler wire and base metal. But the corresponding values
for both Kobe and Boehler electrodes are way of from this general trend. At this point,
it must be mentioned that for lower heating rates (1-10 K min-1
), the α → γ
transformation peak for the Kobe electrode is found to be very diffuse, contributing
thereby to a higher degree of uncertainty in the measured peak area and hence
enthalpy values. In addition, a critical inspection of the composition of Kobe electrode
suggests a relatively higher (Mn + Ni) content, which means the measured enthalpy
corresponds almost exclusively to the conversion of relatively destabilised ferrite into
austenite. In addition, the low carbon concentration of about 0.06 wt.% of Kobe
material supports the argument that it is exclusively the substitutionally alloyed ferrite
matrix that is getting transformed to austenite in the Ac1 – Ac3 intercritical domain;
any probable contribution from possible M23C6 dissolution is assumed to be nil in this
case. Both these factors can serve to decrease the enthalpy of α → γ phase change.
As for the melting enthalpy is concerned, the measured values are in the range 240 -
260 J g-1
. These values are of expected order and are in line with typical values
observed for low carbon alloy steels [40, 41].

11. 11
3.3. α-ferrite + carbides → γ-austenite transformation kinetics
Since α → γ phase change is a nucleation and growth controlled process, the
observed transformation onset and finish temperatures, namely Ac1 and Ac3 show
appreciable dependence on the rate of heating, β (Table 4). This point is graphically
illustrated in figure 4 for the midhani electrode. For other consumables, the relevant
data are listed in table 4. In figure 4(a), the change in the sharpness of the α → γ peak
profile with increasing heating rate is depicted. The methodology adopted for
obtaining Ac1 and Ac3 is also marked. In figure 4(b), the heating rate dependencies of
Ac1, Ac3 and Acα→γ peak temperatures are brought out. It is clear from this figure that
the transformation characteristics are strongly non linear in β, especially in the lower
range, of about 1-40 K min-1
. But this tendency reaches saturation for higher values of
β. With in the range of heating rates investigated in this study (100
to 102
K min-1
), the
variation of the transformation temperature with β can be satisfactorily expressed by
the following power law representation,
Tf (β) = To (β/βo)m
. (1)
In the above, Tf (β) stands for the general transformation temperature Tf measured as a
function of heating rate β, To is the transformation temperature for an arbitrarily
chosen reference heating rate (βo) and m is the power law exponent. For convenience,
βo may be taken as unity (βo = 1), and in which case, To becomes equal to the
observed transformation temperature at unit heating rate. From figure 4(b), it is
evident that Ac3 is more sensitive to the heating rate variation than Ac1, which
contributes to the kinetics induced widening of the (α + γ ) two phase field
(intercritical region) for higher heating rates. This is also inferred from the increasing
values of (Ac3-Ac1) with β (table 4). The reason for this behaviour is the fact that α

12. 12
→ γ transformation gets progressively difficult during its later stages, due possibly to
hard impingement effects which contributes to the delayed as well as incomplete
finishing of the transformation [42]. Although this is true for slow heating rates as
well, the steeply rising nature of Ac3 is however markedly evident for higher values of
β.
The fraction of austenite formed as a function of temperature γ(T) in the
intercritical zone (Ac1 ≤ T≤ Ac3) is estimated using the following expression
γ(T) = {Ac1∫T
ϕ dT) / (Ac1∫Ac3
ϕ dT)}. (2)
Here, Ac1∫T
ϕ dT, is the partial area under the peak in the temperature domain Ac1 – T.
The denominator Ac1∫Ac3
ϕ dT stands for the total peak area covering the entire
transformation temperature range. Eq. (2) assumes that austenite formation is
complete upon reaching Ac3, although this is not strictly true for higher heating rates
[42]. The transformation plots obtained using Eq. (2) are displayed in figure 5 for the
midhani electrode. It is clear that austenite formation from tempered martensite
supports the classical sigmoidal behaviour that is characteristic of a nucleation and
growth phenomenon.
In the usual methodology adopted for analysing the transformation kinetics on
continuous heating, the following functional form is often invoked to represent the
instantaneous reaction rate, (∂γ/∂T)β [43].
(∂γ/∂T)β = f(γ) k(T) (1/β). (3)
In this formalism, f(γ) is often taken to be an empirical, but suitable reaction model
that is consistent with the established kinetic features of the transformation under
consideration. The empirical rate constant k(T) is normally assumed to be the of the
Arrhenius form
k = k0 exp(-Qeff/RT), (4)

13. 13
with Qeff being the effective or apparent activation energy for the overall
transformation process. In a formal theoretical development of the overall
transformation kinetics for a nucleation and growth process, Qeff is actually identified
with a physically based model of nucleation and growth [44] and in such a case it can
be shown that it is actually a weighted sum of individual activation energies involved
in nucleation and growth processes [45-47]. For the case of α → γ transformation, the
activation energy for nucleating the γ phase, namely QN, is much smaller than that
involved in propagating the γ/α interface into ferrite QG [48-50]. In deference to the
overall scope of this study, the theoretical treatment of the transformation rate will not
be discussed here; nevertheless, it is sufficient to note that relative values of Qeff can
be given a consistent interpretation.
In the present study, the overall transformation kinetics is fitted to the following
nonisothermal version of the Kolmogorov-Johnson-Mehl-Avrami (KJMA) form for
the fraction of austenite γ(T) formed as a function of temperature at a constant heating
rate [46].
γ (T) = 1-exp{-kn
[R(T-TS)2
/β Qeff]n
}. (5)
It may be noted that in the above model, we have chosen T-TS, the temperature
increment with respect to the experimentally observed threshold or onset temperature
(TS) as the independent variable, since this corrects in an apparent manner for the
error incurred in not accounting precisely for the true start of the transformation
corresponding to zero transformed fraction. The experimental data on γ(T) for
different heating rates (β) are fitted using Eq. (5) by means of a standard nonlinear
optimisation routine and the resulting values for the kinetic quantities namely Qeff, k0,
and n are listed in table 5. As may be judged from this table, that notwithstanding the

14. 14
difference in composition, there is a nice agreement seen among the Qeff values
obtained for different consumables. The average value of n, the avrami exponent for
the overall transformation kinetics is about one. It also emerges from table 5, that the
kinetic parameters are somewhat sensitive to the heating rate; in particular the
apparent activation energy Q exhibits a gentle decrease with increasing heating rate
(β). While this point will be addressed in the discussion section, it is sufficient to note
that a standard KJMA model is able to account for the observed transformation
kinetics.
4.0. Discussion
4.1. Mn +Ni dependence of Ac1 temperature
One of the major outcomes of this investigation is the elucidation of the effect
of (Mn+Ni) content on the transformation temperatures. From the point of view of
welding physical metallurgy of high chromium ferritic-martensitic steels, two basic
considerations contribute to the proper choice of the welding consumable
composition. They are:
(i). the nature of the primary solidification mode, which needs to be austenitic
and
(ii). the on-heating Ae1 or Ac1 temperature, an appropriate value of which
decides the PWHT schedule. A standard manifestation of the balancing influence
between ferrite and austenite stabilising elements is the resulting Ac1 temperature.
Besides, the knowledge of Ac3 is also called for in rationalising the process related
heat treatments like homegenisation, normalising or any other hot working that is
done in the fully austenitic phase field. These metallurgical requirements are easily
satisfied by going in for a slightly higher quantum of austenite stabilising elements in
the overall composition. But an arbitrary enhancement of austenite stability by the

15. 15
way of increasing for example the combined (Mn+Ni) content is not recommended;
since this option interferes with the PWHT temperature and also results in
concomitant lowering of MS, the martensite start temperature and hence comes in the
way of realising a fully martensite structure in weld deposits [51]. The presence of a
duplex microstructure is inimical to toughness, creep ductility and so on [52].
The present study clearly reveals the role of combined (Mn+Ni) content in
lowering the Ac1 temperature in arc welding consumables as compared to the base
metal (figure 6). The present results on Ac1, can be satisfactorily fitted to the
following empirical relation
(Ac1/o
C) = 855 – 25.9 (Mn+Ni / wt.% ) – 4.0 (Mn+Ni/ wt.% )2
. (6)
Unfortunately, the number of data points catering to different (Mn+Ni) levels is too
small to support a statistical analysis; but nevertheless, the qualitative trend advocated
is in accord with the basic tenets of metallurgy. Further work on other compositions is
currently underway.
4.2. Kinetics of α → γ transformation on continuous heating
The second outcome of this study is concerned with the elucidation of
reaustenitisation reaction on heating. Admittedly, the formation of austenite from
tempered martensite microstructure is not a simple metallurgical phenomenon [53,
54]. It is clear that upon reaching Ac1, the lower critical temperature, both the
dissolution of carbide and the structural transformation of α-ferrite → austenite are
energetically feasible, but for kinetic factors, these concurrent reactions occur with
different velocities [55]. Of these, the structural transformation of (carbon lean) ferrite
to austenite is basically an interface diffusion-controlled process [55-57] while
carbide dissolution is a bulk or volume diffusion controlled phenomenon [27-29, 36].
The heterogeneous nucleation of austenite is catalysed by the presence of abundant

16. 16
internal interfaces to begin with, although at later stages, some other less potent
nucleation sites need to get activated [57]. Thus, in principle the nucleation rate of
austenite is one of continuously decreasing with time. However, assuming that
nucleation of austenite is adequately facilitated, the issue then is one of sustaining its
growth by enabling the transfer of atoms across the growing austenite/ferrite interface
[56]. From a mechanistic point of view, the austenite growth is possible even at a
higher heating rates, if it is mediated by the transfer of mobile species across the
growing γ/α interface [58].
In the case of carbide dissolution however, the steep concentration gradient
prevailing at the carbide/austenite or carbide/ferrite interfaces, both with respect to the
slow diffusing substitutional elements (Cr, Mo) and also to some extent with respect
to the relatively fast moving carbon atoms [36], dictates that volume diffusion in
austenite is the rate limiting one [36, 53, 54]. With this conceptual background the
role of heating rate in influencing the kinetics of carbide dissolution and hence
indirectly the austenitisation reaction is discussed next.
For very slow heating rates, the system can be assumed to evolve with the
possibility of realising local thermodynamic equilibrium with respect to mobile
carbon atoms at the interface. Under such circumstances, it is likely that the carbide
dissolution may get initiated even at relatively lower degree of superheating, that is
T>Ac1. The reaction proceeds to some extent, but certainly not to completion upon
reaching Ac3. The dissolution of carbides results in a gradual change in austenite
composition in the immediate vicinity and this affects subsequent austenite formation
kinetics in an indirect manner by effecting an increase in the activation energy for
carbon diffusion in austenite [40, 59].

17. 17
If on the contrary, the austenitisation is sought under relatively fast heating
conditions, the time scales provided are simply inadequate for the successful progress
of carbide dissolution [30]. Nevertheless, the ferrite to austenite structural change can
still be realised at these higher heating rates, with possibly a different mechanism, as
the work of Kaluba et al has demonstrated [58]. But even in such cases, it is highly
unlikely that transformation to austenite proceeds to 100 % completion upon reaching
Ac3. The presence of undissolved carbide particles may pin and retard the
advancement of γ/α interface. A slight superheating and or isothermal hold above Ac3
are often required to realise the complete formation of homogeneous austenite [20].
For intermediate heating rates, the actual picture is rather diffuse in the sense that both
carbide dissolution and α → γ phase change may proceed to different extents.
In terms of accounting for Qeff, it may be said that Qeff is reflective of the simple
interface diffusion controlled α → γ structural change under fast heating conditions.
However for slower heating rates, it is influenced certainly by the disintegration of the
carbide phase and its dissolution in the growing austenite phase. Judging the present
results in the light of the picture portrayed above, the gradual decrease in Q values
with increasing β can be rationalised. One final point is that the measured values of
Qeff is of the same order as that of the activation energy for chromium diffusion in the
austenite phase of a low alloy–low carbon steel (~ 260 kJ mol -1
)[59]. However, the
present estimate of Q is in direct contradiction to the normal value of about 110 – 120
kJ mol-1
quoted for austenitisation reaction in plain carbon or low alloy steels, wherein
the cementite or pearlite colony dissolution is more readily enabled. The higher values
of Qeff encountered in 9Cr-1Mo steels implies that simple carbon diffusion in austenite
model is not applicable for the α → γ reaction in highly alloyed power plant steels.
The experimental study of Lenel [27, 28] and the electron microscopy characterisation

18. 18
of austenite formation reaction in high chromium model alloys by Shtansky et al [29]
reinforce this viewpoint. It must be added that the understanding of the kinetics of
austenite formation and carbide dissolution in highly alloyed power plant and nuclear
grade steels is still far from complete and more studies need to be done to shed further
light onto this issue.
Conclusions
(i). A comprehensive differential scanning calorimetry characterisation of thermal
stability has been performed on four different varieties of mod. 9Cr-1Mo
welding consumables. Precise measurement of α → γ transformation
temperatures as a function of heating rate has also been made.
(ii). The transformation temperatures Ac1 and Ac3 are found to exhibit a strong
non-linear variation with heating rate, with the net effect of expanding the α+γ
two phase region to higher temperatures for higher rate of heating.
(iii). It is found that for nominally the same Cr/Mo content, the transformation
temperatures Ac1 and Ac3 exhibited a decrease with respect to increasing
(Mn+Ni) concentration. A high value of (Mn+Ni) content is also found to
record a lower enthalpy for the α-ferrite + carbide → γ -austenite, phase
change.
(iv). The kinetics of the α-ferrite + carbide → γ -austenite transformation has been
modelled after standard KJMA formalism and the apparent activation energy
for the overall phase change is estimated to be about 320-280 kJ mol-1
. The
activation energy is found to be sensitive to the heating rate; it decreased for
increasing rates of heating.
(v). For higher heating rates, it is suggested that the kinetics of α-ferrite + Carbide
→ γ -austenite transformation is found to be mainly interface controlled as no

19. 19
significant carbide dissolution is realised under such circumstance. For slow
heating, the accompanying M23C6 carbide dissolution in growing austenite
also makes an indirect contribution to overall activation energy.
Acknowledgements
The authors acknowledge sincerely the encouragement and sustained support
received from Dr. Baldev Raj, Dr. P. R. Vasudeva Rao and Dr. K. Bhanusankara Rao
during the course of this research. Shri. B. Jeya Ganesh acknowledges the award of
the junior research fellowship by DAE, India, as a part of which the present research
is carried out.

20. 20
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24. 24
Table 1
Listing of the compositional details of mod. 9 Cr-1Mo welding consumables used in
this study. The composition is determined by direct reading optical emission
spectroscopy.
Elements Kobe
E9016
Boehler
E9016
Midhani
(E9016-B9)
Midhani
Filler Wire
Midhani
Base metal
( weight percent )
Carbon 0.06 0.075 0.1 0.084 0.097
Manganese 1.37 1.20 1 0.60 0.37
Silicon 0.30 0.30 0.24 0.30 0.31
Nickel 0.95 0.65 0.7 0.75 0.38
Phosphorous 0.005 0.008 0.009 0.007 0.018
Sulphur 0.001 0.007 0.012 0.0034 0. 0047
Chromium 9.24 9.3 9.0 8.89 9.29
Molybdenum 1.05 1.0 0.99 0.93 0.92
Niobium 0.03 0.009 0.05 0.08
Vanadium 0.17 0.23 0.17 0.22 0.26
Nitrogen 0.03 0.04 0.055 0.0375 0.057
Copper - - < 0.05 - -
Aluminium - - < 0.010 - -
Iron balance balance balance balance balance
Mn+ Ni 2.32 1.85 1.7 1.35 0.75
V+Nb 0.20 0.23 0.179 0.27 0.34

25. 25
Table 2
Listing of the Vickers microhardness values (VHN) for mod. 9Cr-1Mo welding
consumables. Italicised values are not considered for the final analysis.
Sample
description
Midhani
electrode
Kobe
electrode
Boehler
electrode
Base metal Midhani
Filler
1 503 473 503 339 560
2 514 493 483 336 572
3 514 488 498 339 560
4 514 488 503 322 542
5 514 488 483 342 548
6 514 488 464 351 519
Average 514 485 487 341 556

26. 26
Table 3
The on-heating transformation temperatures for mod. 9Cr-1Mo welding consumables
as measured by DSC. The values quoted in brackets are in 0
C.
Description of phase change
Transformation temperature
K
(o
C)
Midhani
electrode
Kobe
electrode
Boehler
electrode
Filler
wire
Midhani
Base
metal
Midhani
α'
(martensite)→ α (ferrite)+ M23C6
930
(657)
947
(674)
937
(664)
944
(671)
903
(630)
TC, Curie temperature
1004
(731)
1003
(730)
1013
(740)
1004
(731)
1012
(739)
α + (V,Nb)(C,N)+ M23C6 → α+γ+
(V,Nb)(C,N)+ M23C6
(Ac1)
1072
(799)
1045
(772)
1080
(807)
1072
(799)
1102
(829)
α +γ+ (V,Nb)(C,N) + M23C6 → γ+
(V,Nb)(C,N)+ M23C6
(Ac3)
1120
(847)
1100
(827)
1117
(844)
1120
(847)
1138
(865)
γ+ (V,Nb)(C,N)+ M23C6→ γ+
(V,Nb)(C,N)
1513
(1240)
1522
(1249)
1520
(1247)
1543
(1270)
1522
(1249)
γ+ (V,Nb)(C,N) → δ + γ +
(V,Nb)(C,N)
1567
(1294)
1583
(1310)
1572
(1299)
1598
(1325)
1573
(1300)
δ + γ + (V,Nb)(C,N) → δ+ γ
1632
(1359)
1653
(1380)
1599
(1326)
1655
(1382)
1659
(1386)
δ+ γ → Liquid + γ+δ
1793
(1520)
1790
(1517)
1795
(1522)
1795
(1522)
1793
(1520)
Liquid + γ +δ → Liquid + δ
1796
(1523)
1793
(1520)
1796
(1523)
1798
(1525)
1794
(1521)
Liquid + δ → Liquid
1804
(1531)
1800
(1527)
1806
(1533)
1806
(1533)
1802
(1529)
∆o
Hf, Enthalpy of melting
(J g-1
)
242 240 266 241 239
∆o
H α → γ , Enthalpy of α → γ
transformation
(J g-1
)
17 8 26 16 13

27. 27
Table 4
Listing of measured transformation temperatures for the four mod. 9Cr-1Mo welding
consumables as a function of heating rate.
Midhani electrode Kobe electrode
Heating
rate
(K min-1
)
Ac1
(K)
Peak
(K) Ac3
(K)
Ac3-Ac1
(K)
Ac1
(K)
Peak
(K)
Ac3
(K)
Ac3-Ac1
(K)
2 1053 1081 1117 64
5 1064 1087 1107 43 1028 1064 1099 -
10 1070 1090 1115 44 1045 1077 1101 56
20 1073 1097 1120 48 1049 1085 1118 69
30 1079 1103 1130 51 1048 1085 1122 74
40 1081 1109 1137 57 1052 1093 1138 86
50 1083 1113 1147 64 1056 1093 1147 91
60 1083 1116 1159 76 1062 1099 1148 86
75 1091 1124 1169 78 1071 1113 1154 83
100 1095 1126 1180 85 1084 1111 1159 75
Boehler electrode Midhani filler wire
Heating
rate
(K min-1
)
Ac1
(K)
Peak
(K)
Ac3
(K)
Ac3-Ac1
(K)
Ac1
(K)
Peak
(K)
Ac3
(K)
Ac3-Ac1
(K)
3 1072 1087 1112 40 1094 1108 1125 30
5 1075 1094 1111 36 1098 1111 1129 31
10 1080 1099 1117 37 1104 1116 1137 33
20 1081 1107 1127 45 1105 1121 1145 40
30 1085 1110 1136 51 1108. 1125 1155 47
40 1088 1113 1143 55 1109 1129 1165 56
50 1091 1118 1153 62 1109 1133 1174 65
60 1093 1118 1153 60 1111 1136 1178 67
75 1095 1121 1162 68 1116 1140 1189 73
100 1103 1131 1177 73 1119 1144 1198 79

28. 28
Table 5
Listing of the KJMA fit parameters for the α-ferrite + M23C6 → γ -austenite phase
change upon continuous heating.
Midhani electrode Kobe electrode
Heating
rate
(K min-1
)
Q
(kJ mol-1
)
k0
(s-1
)
n
Heating
rate
(K min-1
)
Q
(kJ mol-1
) k0
(s-1
)
n
2 313 1.27 E15 1.26 3
5 308 3.02 E15 1.38 5 297 6.67 E14 0.70
10 298 1.32 E15 1.50 10 296 1.309 E15 0.91
20 296 1.47 E15 1.37 20 295 8.83 E14 1.18
30 294 1.15 E15 1.49 30 292 9.77 E14 1.14
40 291 8.08 E14 1.45 40 290 6.77 E14 0.99
50 288 5.33 E14 1.31 50 287 6.15 E14 1.13
60 60 284 4.81 E14 1.08
70 284 4.02 E14 1.28 75 282 2.17 E14 1.08
99 281 3.34 E14 1.17 99 281 5.58 E14 0.93
Boehler electrode Filler wire-Midhani
Heating
rate
(K min-1
)
Q
(kJ mol-1
) k0
(s-1
)
n
Heating
rate
(K min-1
)
Q
(kJ mol-1
) k0
(s-1
)
n
3 291 4.3 E14 1.26 3 320 9.286 E15 0.90
5 289 4.4 E14 1.44 5 312 7.964 E15 1.14
10 286 5.5 4E14 1.45 10 305 3.731 E15 1.13
20 283 3.79 E14 1.60 20 302 2.581 E15 1.17
30 283 4.47 E14 1.52 30 299 2.383 E15 1.16
40 280 3.77 E14 1.41 40 295 1.823 E15 1.17
50 279 2.80 E14 1.35 50 292 1.748 E15 1.09
60 276 3.06 E14 1.28 60 288 1.064 E15 1.18
70 274 1.99 E14 1.30 70 - - -
99 273 1.67 E14 1.22 99 286 1.085 E14 1.01

29. 29
a
b

30. 30
c
d

31. 31
Figure 1 (a-e) The optical micrographs of different electrode compositions : (a)
Boehler electrode, (b) Kobe electrode, (c) Midhani electrode, (d) Midhani filler, and
(e) base metal
e

32. 32
Figure 2. The room temperature x-ray diffraction profiles of different welding
consumables taken using Cu-Kα radiation.
20 40 60 80
(531)
(440)
(422)
a = 0.2875 nm
a = 0.2876 nm
a = 0.2876 nm
a = 0.2873 nm
M23
C6
carbide peak
Midhani : make E9016-B9
Kobe make : E 9016
Boehler make : E 9016
Midhani base metal
Mod. 9Cr1Mo Ferritic
Welding Consumables
αααα-(211)
αααα-(200)
αααα-(110)
Intensity(arbitraryunits)
2θθθθ (degrees)

33. 33
Figure 3. The DSC thermograms for the Midhani electrode and (b) the expanded
region around melting
200 400 600 800 1000 1200 1400 1600 1800
-10
-8
-6
-4
-2
0
2
α
+
γ
α
+
γ
α
+
γ
α
+
γ
γ+δ
γ+δ
γ+δ
γ+δ
αααα + γγγγ + M23
C6
+ MX
γγγγ+M23
C6
+M
X
γ+γ+γ+γ+M
X
αααα
γγγγ
+
δ
+
+
δ
+
+
δ
+
+
δ
+
M
X
δδδδ+γγγγ
Mod. 9Cr 1Mo Consumable
MIDHANI electrode
endo
Ac3
Ac1
Tc
Cooling
cycle
Heating Cycle
Cr (9);C(0.1);Mo(0.99);Mn(1);Si(0.24);Ni(0.7);(Mn+Ni)(1.70);
S(0.012);P(0.009);Nb(0.009);V(0.17);N(0.055);Cu(<0.05);Al(<0.010)
DSCOutput(µµµµV)
Temperature (K)
(a)
1780 1790 1800 1810
DSCoutput
(arbitraryunits)
endo
Mod. 9Cr 1Mo
Weld consumables
Midhani
Kobe
Boehler
filler wire
γ + δγ + δγ + δγ + δ liquid + δδδδ + γγγγ
liquid + δδδδ
liquid
Temperature (K)
(b)

34. 34
Figure 4: The variation of (a) α → γ transformation profile, and (b) transformation
temperatures, Ac1, Ac3 and Acα→γ peak with heating rate.
0 20 40 60 80 100
1050
1075
1100
1125
1150
1175
α - γ peak
Ac3
Ac1
Ac3
(K) = 1078K * (ββββ/K min
-1
)
0.0178
Acαααα-γγγγ peak
(K) = 1072K *(ββββ /K min
-1
)
0.01039
Ac1
(K) = 1051K * (ββββ/K min
-1
)
0.00835
Mod. 9Cr 1Mo welding consumable : MIDHANI electrode
TransformationTemperature(K)
Heating rate (K min
-1
)
(b)
1075 1100 1125 1150 1175
0
5
10
15
20
25
αααα - γγγγ peak
Ac3Ac1
2
30
10
20
3
5
50
70
99Mod. 9Cr 1Mo Weld consumable : MIDHANI electrode
DSCOutput(µµµµV)
Temperature (K)
(a
)

35. 35
Figure 5. The fraction of austenite formed as a function of temperature for different
heating rates.
1040 1060 1080 1100 1120 1140 1160 1180
0.0
0.2
0.4
0.6
0.8
1.0
heating rate
increasing
9970504030201053
Mod. 9Cr 1Mo :
Weld consumable
Midhani electrode
(α - γ transformation)
Fractionofaustenite
Temperature (K)

36. 36
Figure 6 : The variation of Ac1 and Ac3 transformation temperatures with respect to
combined (Mn+Ni) content
0.8 1.2 1.6 2.0 2.4
900
950
1000
1050
1100
1150
Filler
Ac1
Ac3
K
obe
Boehler
Mod. 9Cr 1Mo welding consumables
M
idhani
M
idhani
Basemetal
Ac1
,Ac3
(K)
Mn + Ni (wt%)