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EFFECT OF DEFORMATION TEMPERATURE ON MICROSTRUCTURE AND MECHANICAL
BEHAVIOUR OF WARM WORKING VANADIUM MICROALLOYED STEELS
C. García-Mateo1 , B. López2 and J.M. Rodriguez-Ibabe2
1 CENIM-CSIC. Dept. Physical Metallurgy. Avda. Gregorio del Amo 8, E-28039 Madrid, Spain
2 CEIT and Tecnun (University of Navarra), P.o M. Lardizabal 15, E-20018 San Sebastián, Basque
Country, Spain
* Corresponding author: Materalia Research Group, Department of Physical Metallurgy, Centro Nacional
de Investigaciones Metalúrgicas (CENIM), Consejo Superior de Investigaciones Científicas (CSIC),
Avda. Gregorio del Amo, 8. E-28040 Madrid, Spain. Tel: 0034 91 553 89 00 (Ext 373); Fax 0034 91 534
7425; E-mail: cgm@cenim.csic.es
Abstract
Plane strain compression tests of two V microalloyed steels and one plain C-Mn steel have been done to
analyse the influence of the deformation temperature, in the warm working range, on the final
microstructure and subsequent mechanical behaviour. In the case of V microalloyed steels, the reheating
temperature has an effect on the amount of vanadium in solution prior to deformation. This factor
influences the austenite evolution during warm deformation and the transformation during cooling. As a
consequence, in the microalloyed steels complex multiphase microstructures are obtained that lead to a
wide range of strength-toughness combinations. In contrast, in the case of the plain C-Mn steel minor
effects are observed in the deformation range from 800 to 870ºC.
Keywords: Warm forging; vanadium microalloyed steel; mechanical properties
INTRODUCTION
Nowadays there is a continuous demand, particularly from the automotive industry, for cheaper, lighter
and more reliable components and, at the same time, steel is facing strong competition from various other
groups of materials in a market that was traditionally its own. It is not surprising then that during the last
25 years, the steel industry has experienced some of the most important changes at two different fronts,
steel design and processes. Steel manufacturing processes face major competitions in the automotive
industry to produce lighter, cheaper and more efficient components that exhibit dimensions that are more
precise, need less machining and require less part processing. In that respect, mechanical and
metallurgical properties, as well as the manufacturing parameters, play decisive roles [1,2].
Hot forging grades have been traditionally plain carbon, alloyed and microalloyed steels. Their final
mechanical properties are achieved after applying quenching and tempering heat treatments or directly
cooling after forging. In all the cases machining operations are required to obtain the final shape and
dimensions. It is worth emphasising that the distribution of the costs in producing conventional forging
components shows that more than 50% is assigned to machining, besides of the cost of the base material,
the forging process and the heat treatment [3]. Taking into account this aspect, a great interest has
emerged in the near-net-shape technologies, as cold and warm forging processes. Concerning warm
forging (usually in the temperature range from 600 to 900ºC), this route has a number of advantages over
traditional forging procedures [4-7]. Among them, the following advantages need to be considered: better
utilisation of material, improved surface finish, dimensional accuracy compared with hot forging, reduced
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press loads compared with cold forging, and control of the microstructure such that the desired properties
can be obtained without further heat treatments. The process fills the niche between the closer tolerance,
but sometimes expensive, cold forging process and the somewhat lower precision of hot forging route.
The production of warm forged components was initially based on plain carbon steels with final ferrite-
pearlite microstructures at room temperature [6]. In the last years, warm forging techniques have also
been extended to components based on low alloy steel grades, requiring specific studies concerning flow
stress behaviour of these grades in the 600-900ºC temperature range [8,9]. In a similar way to what
happened with conventional forging, the introduction of microalloyed steel grades in warm forging has
been considered as an option to obtain better strength-toughness combinations in comparison to plain C-
Mn steels. Initially vanadium was considered as the main microalloyed option [5], while more recently
Nb microalloying has appeared as a very promising possibility [10,11].
In the case of hot forging, the introduction of vanadium assigned to this element its conventional role of
precipitation strengthening during cooling at the exit of the forging. In contrast, this situation changes
significantly in the case of warm forging [12]. It is necessary to take into account that, depending on the
reheating temperature prior to warm forming, a fraction of V(C,N) particles present in the as-rolled
condition will remain undissolved, while in conventional hot forming all the vanadium is in solution and,
as a consequence, available to precipitate during cooling. These undissolved particles can interact with
the microstructure in different scenarios. For example, it has been reported that undissolved V(C,N)
precipitates can reduce the austenite grain growth during reheating and similarly, they can delay austenite
static recrystallisation kinetics after deformation [13,14]. Both factors can have beneficial effects on final
room temperature properties by promoting additional microstructural refinement. In contrast, depending
on the reheating temperature the reduced amount of vanadium available to precipitate during cooling can
reduce its contribution to precipitation strengthening.
The aim of this work is to study the potential that warm forging in combination with V microalloying
may have on the final properties of steels. For this purpose two vanadium microalloyed steel grades and a
typical forging C-Mn steel, for comparison, have been used.
MATERIAL AND EXPERIMENTAL PROCEDURE
Two commercial vanadium microalloyed steels with different C and V contents and, for the sake of
comparison, a plain C-Mn steel were selected for this study. Their chemical compositions are listed in
Table I. content. Warm forging simulations were performed in a servo hydraulic press by testing plane
strain compression specimens of 25x60x10 mm3, machined from industrially hot rolled steel bars.
Specimens were heated in a resistance furnace and soaked at the deformation temperature for 10 min
before testing. In order to minimize the effect of friction and temperature gradient profiles, which may
lead to inhomogeneities in deformation distribution, samples were lubricated with boron nitride and both,
samples and die tools, were kept inside the furnace during all the experiment. Finally, specimens were
deformed to an strain of = 0.3 at an strain rate of 10 s-1. Following deformation, specimens were
subjected to two different cooling rates, i.e. air cooling (1ºC/s) and accelerated cooling ( 4ºC/s). Three
different testing temperatures were considered 800, 835 and 870ºC, all of them in the austenite range.
Prior austenite grain size (PAGS) before deformation was determined metallographically from quenched
specimens after soaking at the testing temperature. In order to reveal the PAGS it was necessary to temper
the specimens at 450-500ºC for periods in excess of 24 hours for V microalloyed steels and two hours in
the case of the plain C-Mn steel, respectively. After the treatment, the samples were polished and etched
in a solution of saturated picric acid. Prior austenite grain size was measured by the mean equivalent
diameter method.
Microstructural characterisation was carried out using optical and scanning electron microscopy on
metallographic samples cut longitudinally from the plane strain compression specimens. Round tensile
samples (4 mm diameter and 10 mm gauge length) and V-notch Charpy impact specimens (sub-size
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samples of 5x10 mm) were machined from plane strain compression specimens. In the case of the Charpy
tests, the absorbed energies were converted to standard specimen dimensions [15].
Carbon replicas were prepared by conventional methods, and observed in a transmission electron
microscope (TEM) fitted with an energy dispersive X-ray system. The size and distribution of V
precipitates were quantified for the different conditions, measuring in each case a minimum of 500 and a
maximum of 1200 particles.
MTDATA [16] fitted with NPL-plus database for steels was used to theoretically estimate the amount of
V in solid solution and the mass fraction of V precipitates as a function of temperature.
RESULTS
MICROSTRUCTURAL CHARACTERISATION
Figure 1 shows the mean austenite grain size as a function of the reheating temperature. Data
corresponding to reheating temperatures higher than initially considered for warm forging conditions (that
is, lower than 900ºC) have been included for comparison.
The microstructures of the vanadium microalloyed steels at room temperature consist of a mixture of
ferrite, pearlite/bainite and martensite/austenite constituent (MA), as shown in Figure 2. The pearlite
present in these steels has not its normal lamellar structure and can be considered as “degenerated
pearlite”. Besides this pearlite bainite is also identified although it is difficult, because of the morphology
of the pearlite, to distinguish between both constituents. In some regions, mainly in the case of V2 steel,
the MA constituents show a banded tendency. In the case of the plain carbon steel, the microstructure is
ferrite-pearlite with well-defined lamellar pearlite.
The metallographic measurements of the ferrite mean grain size and ferrite volume fraction are shown in
Figure 3 as a function of the deformation temperature. In both vanadium microalloyed steels, at 800 and
835ºC deformation temperatures the ferrite grain size remains very small with a mean size ranging
between 2.1 and 3.4 m after air cooling. The accelerated cooling does not lead to an additional
refinement, in contrast to what happens after deforming at 870ºC. In this later case, the grain sizes are
slightly coarser than those measured at the other two conditions. In relation to the plain C-Mn steel the
ferrite grain size is coarser that those corresponding to the V microalloyed steels, although it continues
being fine.
In relation to the ferrite volume fraction, the results obtained for the three steels are shown in Figure 3 as
a function of deformation temperature and cooling rate. While at 800ºC and 870ºC the ferrite fraction
tends to decrease, regardless of the cooling rate, with the carbon content present in the steel, at 835 some
changes in this behaviour have been identified. Thus, at 835ºC in the V2 steel the ferrite content is higher
than that measured in steel V1 for both cooling conditions,. In all the conditions, the accelerated cooling
reduces the ferrite fraction in relation to the air cooled samples. Finally, in the plain C-Mn steel the ferrite
fraction remains nearly unaffected by the deformation temperature.
The volume fraction of the MA constituent present in the vanadium microalloyed steels has been drawn
in Figure 4 as a function of the deformation temperature. Increasing the deformation temperature the MA
volume fraction increases, ranging from 5% at 800ºC to 18% at 870ºC in the air cooled samples. These
values are significantly higher when accelerated cooling is applied. On the other hand, the mean size of
the MA constituents slightly increases with the deformation temperature, ranging from 2.0 to 5.5 m.
Nevertheless, the effect of cooling rate is not always the same. For example, while coarser MA
constituents appear at 800ºC and accelerated cooling, the opposite occurs at 835 and 870ºC. Comparing
both V1 and V2 steels, similar or even coarser sizes have been measured in steel V1 at all the conditions.
Due to the different roles that precipitates can play in both microalloyed steels, an exhaustive analysis of
the precipitate size distribution and density formed under different conditions was performed. In the as
received condition, the spherical particles observed in both vanadium microalloyed steels were identified
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as V rich precipitates. The size distribution of these particles is shown in Figure 5. In the figure, it can be
seen that V1 steel exhibits a distribution of very fine precipitates with a mean diameter of 8 nm, 71% of
the precipitates being smaller than 9 nm, while in V2 steel the precipitates are slightly coarser with a
mean size of 14 nm. The number of particles per area unit () is also indicated in the figure.
Figure 6 shows the size distribution of V precipitates after reheating for 10 min. at 870 and 835ºC
followed by water quench. If these results are compared with those of Figure 5 (as received condition) it
is possible to see a displacement of the size distribution towards coarser particle sizes as heating
temperature increases. This behaviour denotes that there has been a dissolution process involving, above
all, the smallest precipitates, which is also confirmed by the significant decrease in the number of
particles per area unit. Finally, the precipitate size distributions after deformation followed by air cooling
are shown in Figure 7 for steel V1 at deformation temperatures of 800 and 870ºC and for steel V2 at a
deformation temperature of 870ºC, respectively. Results indicate that deformation at 870ºC induces a
refinement in the particle size distribution while lower deformation temperatures has the opposite effect,
further details will be given during discussion.
MECHANICAL PROPERTIES
In warm forging processes, the flow stress of the steel during deformation is an important factor to
evaluate as it affects steel workability and tools wear. The mean flow stress values (MFS) were calculated
from the stress-strain curves obtained in the plane strain compression tests. The MFS is defined as the
area under the stress-strain curve divided by the applied strain. The values obtained with the three steels
are drawn in Figure 8 as a function of the deformation temperature. As the temperature decreases from
870 to 800ºC there is an increase in the MFS value close to 20% in the three steels and, for a given
temperature, the microalloyed steels exhibit higher values than the C-Mn one.
Figure 9 summarizes the mechanical properties at room temperature as a function of the deformation
temperature (mean values of two samples). In most conditions, the highest yield stress and UTS
correspond to the V2 steel, followed by V1 steel. This tendency changes slightly in the samples deformed
at 800ºC and accelerated cooled. On the other hand, for a given steel the effect of the deformation
temperature does not show a well defined behaviour. While for the air cooled samples in the V
microalloyed steels a slight decrease occurs in the yield stress as the deformation temperature increases,
the opposite arises when accelerated cooling is applied. In the case of the UTS, for both cooling rates an
increasing tendency is observed with higher deformation temperatures. Concerning the C-Mn steel, the
influence of the deformation temperature seems to be smaller in both yield stress and UTS values,
although a discrepancy is observed in the samples deformed at 870ºC and air cooled.
Another aspect to take into account is that while in all the tensile tests done with the V microalloyed
steels over air cooled microstructures the yield curves have a well defined plateau, in the accelerated
cooled conditions this plateau reduces or completely disappears, showing a continuous yielding curve, as
the deformation temperature increases.
Ductility, measured as the reduction in area, increases as forging temperature decreases in the V
microalloyed steels, showing higher values for the air cooled than for accelerated cooled conditions.
Between both V steels, only small differences have been quantified. In contrast, the ductility of the C-Mn
steel remains nearly constant independently on the deformation temperature and cooling strategy.
The influence of the deformation temperature on the Charpy curves is shown in Figure 10. In the case of
the V microalloyed steels, the air cooled samples deformed at 870ºC show a remarkable shift of the
curves towards higher temperatures, almost 50ºC, compared to the other two deformation conditions. In
the case of accelerated cooled specimens, the associated increase in strength is accompanied by
impairment of toughness, more notorious in the microstructures corresponding to 830 and 870ºC
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conditions. Finally, in the case of the plain C-Mn steel the influence of the deformation temperature on
the Charpy curves is less important.
DISCUSSION
MICROSTRUCTURES BEFORE DEFORMATION
The austenite grain coarsening behaviour observed in Figure 1 may be directly related to the composition
of the steels. Grain growth can be influenced by the solute drag effect of any elements in solid solution
and by the pinning forces associated to precipitates. The amount of V in solid solution calculated with
MTDATA for both V microalloyed steels is shown in Figure 11 as a function of temperature. These
results indicate that an important fraction of precipitates present in the as received conditions will remain
undissolved. Similarly, the size distributions of precipitates measured in the samples quenched after
reheating at 835 and 870ºC confirm that they remain very fine and that, in consequence, are able to exert
a control in the austenite grain size. As the austenitisation temperature increases and more vanadium is
put in solution, grain growth starts over 900ºC but sizes continue being significantly smaller than those
measured in the C-Mn steel. This observation agrees with previous results obtained with vanadium
microalloyed steels that associate the inhibition of the austenite grain growth with a combined effect of
pinning by incompletely dissolved vanadium precipitates and a solute drag effect of vanadium [17].
The other microstructural aspect to take into account before deformation is applied is the situation of
V(C,N) precipitates. If the results of particle size distribution after reheating shown in Figure 6 are
compared with those of Figure 5 (as received condition) it is possible to appreciate that there has been a
dissolution process involving, above all, the smallest precipitates. This is supported by data in ref [18-20],
where it is estimated that the time needed at 870ºC to completely dissolve a vanadium precipitate changes
from few minutes for a 6 nm diameter particle to nearly one hour for a 18 nm diameter one. Although it is
not probable due to the moderate reheating temperatures, it could be argued if coarsening of the size
distribution is assisted or not by Ostwald ripening. Using Wagner formalism [21], the time required for
the coarsening of V precipitates by Ostwald ripening was calculated. The results indicate that a particle of
6 nm needs 1.8 hours and 7 hours to achieve 10 nm at 870ºC and 835ºC, respectively. This means that
coarsening by Ostwald ripening mechanism is very improbable to occur with the applied treatments, thus
supporting the idea that the larger size of precipitates after reheating compared to the as received
condition may exclusively arise from the dissolution of the smallest particles, as mentioned above.
Finally, and according to the equilibrium calculations shown in Figure 11, the precipitates present in the
as received condition in steel V2 will be more stable than those in steel V1, meaning that their dissolution
becomes more difficult. This observation is supported by the fact that in steel V2 an increase in the
reheating temperature from 835 to 870ºC implies a 2 nm enlargement in the mean diameter and a drop of
about 28 m- in the measured precipitate density, while the same figures for steel V1 are far bigger, 10
nm and 60 m-2 respectively. According with the latter discussion, it is reasonable to assume that at
800ºC, where small dissolution has taken place as predicted in Figure 11, the size distribution and density
of particles must be very close to that of the as received conditions, i.e. higher density of smaller particles
as compared to 870 and 835ºC reheating temperatures.
MICROSTRUCTURES AFTER DEFORMATION
Figures 3 and 4 indicate that some room temperature microstructural features are very dependent on the
deformation temperature in both microalloyed steels. This temperature can interact with the
transformation during cooling by modifying two factors: the austenite grain conditioning before
transformation and the vanadium dissolution/precipitation situation.
Concerning the austenite microstructure before transformation, in a previous work [14] it was quantified
that the presence of undissolved V(C,N) precipitates retards the static recrystallisation process occurring
after deformation in the warm temperature regime. This delay becomes longer for lower deformation
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temperatures. Taking into account the results of the aforementioned work, it can be concluded that in the
samples tested at 870ºC, before transformation begins, there will be sufficient time to achieve a
completely recrystallisation of austenite. In contrast, at 835 and 800ºC the recrystallisation will be
incomplete. This implies that the austenite grain boundary area per unit volume (Sv) before transformation
will increase in these last two conditions.
It is well known that higher Sv values favour the microstructure refinement during transformation [22].
This refinement effect is enhanced as the cooling rate is increased. However, it has been reported, in both
low and high carbon steels, that as Sv increases the effect of cooling rate to introduce additional
refinement becomes weaker [23,24]. In this context, in both V microalloyed steels the initial austenite
grain size prior to deformation is very small. As a consequence, the resulting Sv value in the partially
recrystallised microstructures deformed at 800 and 835ºC will be high enough to bring about a very fine
ferrite grain size, as shown in Figure 3, with a minor effect of the cooling strategy. This can explain the
behaviour of the ferrite grain size remaining nearly constant in both V microalloyed steels at 800 and
835ºC. In contrast, as at 870ºC the austenite microstructure is completely recrystallised before
transformation, Sv will be smaller than in the previous cases, thus leading to some refinement in the ferrite
mean grain size when accelerated cooling is applied. In the case of the plain C-Mn steel, a complete
recrystallisation of the austenite after deformation [25] and some grain growth could also happen for the
three deformation temperatures. As a result, smaller Sv values will be produced, which can explain the
slightly coarser ferrite grain size obtained in comparison to the V microalloyed steels.
The evolution of the ferrite volume fraction with deformation conditions shown in Figure 3 seems more
difficult to explain. Concerning the C-Mn steel, f remains nearly independent on the deformation
temperature. This can be explained assuming that the possible changes in the austenite grain size between
the different deformation conditions are small enough to not affect the CCT curve of the steel.. The
smaller ferrite fraction resulting at higher cooling rates follows also the expected tendency. In contrast to
this behaviour, in both V microalloyed steels some dependence on the deformation temperature is
observed.
The dependency of the ferrite fraction with deformation temperature can be considered together with the
tendency of increasing MA fraction shown in Figure 4. Different factors can be interacting
simultaneously. Firstly, the accumulated strain in the austenite (at low deformation temperatures) will
accelerate the ferrite transformation, i.e. ferrite will start forming at higher temperatures and shorter
times, thus favouring higher ferrite fraction contents [24]. Secondly, vanadium in solution delays slightly
the start of ferrite transformation and moves the end of pearlite and bainite transformations to longer
times [26]. The third factor to take into account is the amount of nitrogen in solution. Staiger et al.
identified an important influence of free nitrogen on the increase of the MA microconstituents formation
in low C high Mn steels [27]. These authors observed that nitrogen in solution increases the hardenability
affecting mainly to those regions where Mn was locally segregated.
These three factors can explain qualitatively the behaviour observed in steel V1, where a decrease in the
ferrite and an increase in MA volume fractions occur as deformation temperature increases. In the case of
steel V2, the MA fraction evolution follows a similar trend to that observed in steel V1. The lower
fractions at 835 and 870ºC, in comparison to those measured in steel V1, could be related to the more
stable V(C,N) precipitates (Figure 11) and, in consequence, to the smaller free nitrogen percentage. In
relation to the ferrite volume fraction, while at 800ºC it follows the general behaviour, the increase in f
measured in the samples deformed at 835ºC, with higher contents than those measured in steel V1 with
lower C content, deviate from the expected behaviour. Assuming that the difference in the austenite
conditioning in both microalloyed steels at 835ºC is small, the possible role of vanadium remains unclear.
The V(C,N) precipitate size distributions after deformation followed by air cooling shown in Figure 6
allow one to evaluate the evolution of the amount of vanadium available in solution after reheating. The
possibility of strain induced precipitation in austenite can be ruled out because of the absence of a plateau
denoting this effect in the recrystallised fraction versus time plots reported in ref. 14. An additional
argument for the absence of strain-induced precipitation is that there is no thermodynamic driving force
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for this to happen since the deformation temperature is the same as the austenitization temperature
Therefore, all the V in solution will be available for further precipitation during cooling, either as new
precipitates or on already existing ones.
A higher reheating temperature will lead to an increase in the precipitation rate of new particles through a
higher supersaturation during cooling but, on the other hand, lower dissolution temperatures will favour
accelerated precipitation of the solutes on pre-existing particles, which act as preferred sites [28,29]. This
is essentially the situation detected in both microalloyed steels. When comparing data of Figure 6 and
Figure 7 for steel V1, it is clear that a refinement in the particle size distribution is attained after
deformation applied at 870ºC. The mean size decreases slightly from 21 to 18 nm and there is an
important increase in the fraction of particles smaller than 9 nm. This implies that a fine fresh
precipitation is taking place during subsequent cooling after deformation, resulting in an important
increase in the density of precipitates detected i.e. from 11 to 91 m-2. This situation differs from that
observed in the samples deformed at 800ºC. Assuming that before deformation the precipitation state at
800ºC should be very similar to that of the as received condition, Figure 5, there is an increase in the
average precipitate size from 8 to 14 nm, in addition to very similar values of and an important increase
in the amount of particles bigger than 17 nm. This suggests that the small amount of solute vanadium
available is precipitating mainly on pre-existing particles. The behaviour just described for steel V1 is
applicable also to steel V2, being the only difference the formation of coarser precipitates for similar
deformation conditions.
MECHANICAL PROPERTIES
The mean flow stress values can be used for evaluating the forces associated with the warm forging.
Several expressions have been proposed to calculate MFS as a function of composition and deformation
parameters. One of these is the Misaka equation (Eq. 1) [30]:
13.021.0
2
2112029682851
594.075.1126.0exp8.9)(
TCC
CCMPaMFS
Eq. 1
where [C] is the carbon content in wt.%, T the temperature in K, the applied strain and
the strain rate.
As observed, this equation takes only into account the effect of carbon content. Its application to
microalloyed grades was done by Kirihata et al. through the introduction of the following multiplying
factor that considers the alloy additions [31], all in wt.%:
f = 0.835+0.098[Mn] +0.5[Nb] +0.128[Cr]0.8 +0.144[Mo]0.3 +0.175[V] +0.01[Ni] Eq. 2
In Figure 8 the experimental MFS values are compared with the predictions as a function of the absolute
inverse temperature, considering Eq 1 for the C-Mn steel and with the correction of Kirihata et al for the
microalloyed steels, including the vanadium in solution corresponding to each condition in agreement
with the predictions of Figure 11. The results show that there is a good correlation between measured and
predicted values for the case of the C-Mn steel, confirming the valuable use of Eq. 1. In contrast, in both
microalloyed steels the predictions significantly underestimate the MFS. These differences suggest that
other factors, no included in the empirical expressions of Misaka and Kirihata et al., are operating. In this
context, it might be suggested that the precipitates present during deformation at the warm reheating
temperature range analysed in this study are affecting the flow stress of the steel. It is worth emphasising
that Dutta et al. [32] observed that Nb(C,N) precipitates present in austenite lead to a strengthening
increment in the flow stress that could be associate directly to the particles. In order to confirm that, an
additional plane strain compression test was done at 1025ºC with steel V1. At this temperature it could
be considered that all the vanadium is in solution. The MFS value obtained from the flow stress curve
was 117 MPa, in very good agreement with the value of 114 MPa calculated from the Misaka-Kirihata
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equations. This confirms that in warm forging of V microalloyed steels there can be an additional increase
in the MFS and, as a consequence, in the forces required to deform the steel, associated to the precipitates
present prior to deformation application.
Concerning the room temperature mechanical behaviour, the application of the warm working to the V
microalloyed steels has led to multiphase microstructures that made difficult the quantification of the
contribution of each hardening mechanism to the overall strength.. Nevertheless, there are some general
rules that can be taken into account. For example, the influence of the ferrite on the overall strength
decreases as higher deformation temperatures are selected in the V microalloyed steels, while it remains
nearly constant in the C-Mn steel. This influence has been considered in Figure 12 through the fd-0.5
product. In the figure, a linear relationship with f has been chosen, as this type of dependences between
the soft phases and the overall strength have been proposed in the case of multiphase microstructures
[33], although for the case of ferrite-pearlite microstructures (that is, the case of the C-Mn steel) the f1/3
proposed by Gladman et al. is generally accepted [34]. Independently of selecting f or f1/3, the
decreasing tendency with deformation temperature shown in Figure 12 prevails.
The fact that the yield stress and the tensile strength decrease slightly or increase with deformation
temperature, depending on the steel and cooling condition, and that the ferrite contribution decreases,
implies a higher contribution of other hardening mechanisms, that is, hard constituents as MA (confirmed
by the data of Figure 4), precipitation strengthening and dislocation density (mainly in the accelerated
cooled samples). Precipitation hardening evidence has been presented in ref. [35], where nano-hardness
measurements of ferrite were made after different deformation conditions. The results thus obtained
showed that ferrite present at the microstructure becomes softer as the deformation temperature and
subsequent applied cooling rate decrease. The results are rationalized in terms of the effectiveness of the
precipitation strengthening contribution, i.e. precipitate size and distribution. All of these hardening
mechanisms will have a deleterious effect on the toughness, measured as ductile-brittle transition
temperature.
This toughness deterioration as deformation temperature and cooling rate are increased is confirmed by
Figure 10. As mentioned above, these results could be explained in terms of a higher fraction of harder
MA constituent (fM) as deformation temperature and cooling rate are increased [2,36,37]. This aspect was
also corroborated by further analysis of the fracture surfaces of brittle broken specimens, which showed
crack initiation in MA constituents. A similar dependence has been reported in ferrite-martensite
microstructures [38]. This deleterious effect of the amount of MA can be observed also in Figure 13
where the room temperature Charpy energy has been drawn as a function of MA volume fraction. The
figure indicates a rapid drop in CVN energy in a small interval of MA fraction: while for MA fractions
smaller than ~10% the behaviour corresponds to the complete ductile fracture in both microalloyed steels,
once this value is surmounted the absorbed energy abruptly decreases. The different contributions of the
microstructural features to the strength and toughness simultaneously can be observed in Figure 14, where
the tensile strength has been drawn against the room temperature Charpy energy. The figure indicates that
in the case of the plain C-Mn steel, the different deformation temperatures lead to a reduced strength-
toughness window. In contrast, the introduction of V microalloyed steels combined with deformation
temperatures between 800 and 835ºC produces an increase in toughness with strength levels similar or
higher than those obtained with the C-Mn steel, depending on the amount of V.
When deformation is carried out at 870ºC, the higher amount of V and nitrogen in solution has lead to an
increase in the fraction of hard phases that impairs toughness without a significant improvement in
strength. The application of higher cooling rates after deformation provides the possibility of obtain
higher strength, but at the expenses of a further decrease in toughness. At these conditions it will be
possible to achieve higher strength values than those obtained with the plain C-Mn steel, but in spite of a
loss in toughness.
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Summarising, because of the different amounts of V that are put in solution depending on the applied
reheating temperature before warm deformation, it is possible to have important differences in the room
temperature microstructure that will affect mechanical properties, showing a higher sensitivity to process
parameters than plain C-Mn steels. Taking it into account, in the range of 800-870ºC, the selection of low
deformation temperatures should lead to the best strength-toughness combinations. On the other hand,
Figure 8 indicates that microalloying requires higher force tools during forming in comparison to plain C-
Mn steels. This aspect should also be taken into account when selecting deformation conditions. Finally,
as the loss of toughness is associated with a higher amount of hard phases, similar approaches to those
applied to conventional forging, that is, an increase in the Si content of the steel [12,39] or a reduction in
Mn could be considered. This would permit to increase the deformation temperature with the consequent
reduction in the requirements of the tool forces.
CONCLUSIONS
The main conclusions of this work could be summarised as follows:
The microstructural and mechanical behaviour of V microalloyed steels show important
differences compared to conventional C-Mn steels in warm working processes. These differences
arise from the different amounts of V that is put in solution which depends on the applied
reheating temperature before warm deformation.
Given the low temperatures implied in warm forging procedures, i.e. low levels of dissolution and
limited coarsening by Ostwald ripening, the size and distribution of the vanadium precipitates in
the as received condition are of great importance. These undissolved particles can play different
roles depending on the processing temperature: control of the austenite grain growth during
reheating, increase of the required tool forces during forming and delay of austenite static
recrystallisation after deformation.
Microalloying with V leads to fine multiphase microstructures formed by ferrite, pearlite/bainite
and MA phases. Depending on the reheating temperature, the vanadium in solution will precipitate
during cooling as fine particles or precipitate over existing undissoved particles. The amount of
each constituent is affected by the reheating temperature.
At the deformation temperatures of 800-830 ºC followed by air cooling both V microalloyed
steels results in better strength-toughness combinations compared to the studied C-Mn steel.In
contrast, at 870ºC the increase of MA volume fraction significantly reduces the toughness. With
accelerated cooling a further increase in strength can be achieved but room temperature impact
toughness is markedly reduced.
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Tables
Table 1. Chemical composition of the steels examined (wt%)
Steel
C
Mn
Si
Cr
Ni
Mo
Cu
V
N
V1
0.24
1.56
0.28
0.10
0.09
0.04
0.24
0.18
0.010
V2
0.33
1.49
0.25
0.08
0.11
0.04
0.27
0.24
0.010
C-Mn
0.47
0.78
0.24
0.17
0.08
0.02
0.10
0.002
0.009
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27
28
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30
31
32
33
34
35
36
37
38
39
40
41
42
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Figures
0
5
10
15
20
25
30
750 850 950 1050
Temperature (ºC)
PAGS (m)
V1
V2
C-Mn
0
5
10
15
20
25
30
750 850 950 1050
Temperature (ºC)
PAGS (m)
V1
V2
C-Mn
Figure 1. Austenite grain size evolution with reheating temperature
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20 m
20 m
(a) (b)
20 m
20 m
(c) (d)
20 m
(e)
Figure 2. SEM micrographs of the microstructures generated in: V1 steel after deformation at 870ºC
followed by (a) air cooling and (b) accelerated cooling; V2 steel after deformation at 870ºC followed by
(c) air cooling and (d) accelerated cooling, and (e) C-Mn steel after deformation at 870ºC followed by air
cooling.
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V1 V2 V3
0
2
4
6
0 1 2 3 4 5 6 7
Ferrite grain size (m)
Steel
870ºC
0
2
4
6
0 2 4 6
Ferrite grain size (mm)
Steel
800ºC
20
40
60
0 1 2 3 4 5 6 7
Ferrite voolume fraction (%)
Steel
870ºC
20
30
40
50
60
0 1 2 3 4 5 6 7
Ferrite voolume fraction (%)
Steel
835ºC
20
30
40
50
60
0 1 2 3 4 5 6 7
Ferrite voolume fraction (%)
Steel
800ºC
V1 V2 V3
Steel Steel
Figure 3. Ferrite mean grain size and volume fraction as a function of deformation temperature. Black
symbols: air cooled and open symbols accelerated cooled microstructures. Error bars indicate the standard
deviation
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Figure 4. Evolution of MA volume fraction and mean size as a function of the reheating (deformation)
temperature. Error bars indicate the standard deviation.
0
10
20
30
40
50
1 5 9 13 17 21 25 29 33 37 41 45 49 53 57
Diámetro nm
%
Dm = 8 nm
0
5
10
15
20
25
30
1 5 9 13 17 21 25 29 33 37 41 45 49 53 57
Diámetro nm
%
Dm = 14 nm
%
= 171 m-2
71% < 9 nm
6.5% > 17 nm
= 128 m-2
53% < 14 nm
16% > 21 nm
%
Diameter/ nm
V1
V2
Figure 5. As received V(C,N) size distribution in V1 and V2 steels. stands for the density of V(C,N) per
area unit.
0
2
4
6
780 800 820 840 860 880
MA mean size (m)
Temperature (ºC)
V1(air)
V1(accel)
V2(air)
V2(accel)
0
10
20
30
780 800 820 840 860 880
MA fraction, fM(%)
Temperature (ºC)
V1(air)
V1(accel)
V2(air)
V2(accel)
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7
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0
5
10
15
20
25
30
1 5 9 13 17 21 25 29 33 37 41 45 49 53 57 61
Diámetro nm
%
Dm = 20 nm
0
5
10
15
20
25
30
1 5 9 13 17 21 25 29 33 37 41 45 49 53 57 61
Diameter nm
%
Dm = 18 nm
0
5
10
15
20
25
30
35
40
1 5 9 13 17 21 25 29 33 37 41 45 49 53 57
Diámetro nm
%
Dm = 21 nm
0
5
10
15
20
25
30
35
40
1 5 9 13 17 21 25 29 33 37 41 45 49 53 57
Diámetro nm
%
Dm = 11 nm
%
%
Diameter/ nm
V2 870ºC
Diameter/ nm
%%
= 71 m-2
45% < 9 nm
12% > 17 nm
= 11 m-2
8.4% < 9 nm
57% > 17 nm
= 10 m-2
36% < 14 nm
21% > 21 nm
= 38 m-2
37% < 14 nm
33% > 21 nm
V2 835ºC
V1 870ºC
V1 835ºC
Figure 6. V(C,N) size distribution in V1 and V2 steels after reheating for 10 min. at the indicated
temperatures and then water quench. stands for the density of V(C,N) per area unit.
0
10
20
30
40
50
1 5 9 13 17 21 25 29 33 37 41 45 49 53 57
Diámetro nm
%
Dm = 18 nm
0
10
20
30
40
50
1 5 9 13 17 21 25 29 33 37 41 45 49 53 57
Diámetro nm
%
Dm = 14 nm
0
5
10
15
20
25
30
1 5 9 13 17 21 25 29 33 37 41 45 49 53 57
Diameter nm
%
Dm = 15 nm
%
%
Diameter/ nm
V2 870ºC
Diameter/ nm
%
V1 870ºC
V1 800ºC
= 146 m-2
33% < 9 nm
30% > 17 nm
= 91 m-2
16% < 9 nm
48% > 17 nm
= 40 m-2
62% < 14 nm
23% > 21 nm
Figure 7. V(C,N) size distribution in V1 and V2 steels after deformation and air cooling at the specified
temperatures. stands for the density of V(C,N) per area unit.
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49
50
51
52
53
54
55
56
57
58
59
60
61
62
63
64
65
150
175
200
225
250
275
8.7 8.9 9.1 9.3
10000/T (1/K)
MFS (MPa)
V1
V2
C-Mn
Figure 8. Mean flow stress (MFS) obtained from plane strain compression tests as a function of
deformation temperature (close symbols). Experimental results are compared with those predicted by
Misaka’s modified equation (open symbols).
1
2
3
4
5
6
7
8
9
10
11
12
13
14
15
16
17
18
19
20
21
22
23
24
25
26
27
28
29
30
31
32
33
34
35
36
37
38
39
40
41
42
43
44
45
46
47
48
49
50
51
52
53
54
55
56
57
58
59
60
61
62
63
64
65
Air cooled
Accelerated cooled
600
700
800
900
1000
1100
775 800 825 850 875 900
UTS (MPa)
Temperature/ ºC
300
400
500
600
700
775 800 825 850 875 900
YS ( MPa)
Temperature (ºC)
300
400
500
600
700
775 800 825 850 875 900
YS ( MPa)
Temperature/ ºC
600
700
800
900
1000
1100
775 800 825 850 875 900
UTS (MPa)
Temperature/ ºC
30
40
50
60
70
775 800 825 850 875 900
AR (%)
Temperature (ºC)
V1
V2
C-Mn
30
40
50
60
70
775 800 825 850 875 900
AR (%)
Temperature (ºC)
V1
V2
C-Mn
Figure 9. Mechanical properties as a function of forging temperature.
1
2
3
4
5
6
7
8
9
10
11
12
13
14
15
16
17
18
19
20
21
22
23
24
25
26
27
28
29
30
31
32
33
34
35
36
37
38
39
40
41
42
43
44
45
46
47
48
49
50
51
52
53
54
55
56
57
58
59
60
61
62
63
64
65
0
20
40
60
80
100
120
140
-70 -30 10 50 90 130 170
Absorbed energy (J)
Temperature (ºC)
V1 acc
835ºC
870ºC
800ºC
0
20
40
60
80
100
120
140
-70 -30 10 50 90 130 170
Absorbed energy (J)
Temperature (ºC)
V2 air
835ºC
870ºC
800ºC
0
20
40
60
80
100
120
140
-70 -30 10 50 90 130 170
Absorbed energy (J)
Temperature (ºC)
V2 acc
835ºC
870ºC
800ºC
0
20
40
60
80
100
120
140
-70 -30 10 50 90 130 170
Absorbed energy (J)
Temperature (ºC)
C-Mn air
835ºC
870ºC
800ºC
0
20
40
60
80
100
120
140
-70 -30 10 50 90 130 170
Absorbed energy (J)
Temperature (ºC)
C-Mn acc
835ºC
870ºC
800ºC
Figure 10. Charpy test results.
1
2
3
4
5
6
7
8
9
10
11
12
13
14
15
16
17
18
19
20
21
22
23
24
25
26
27
28
29
30
31
32
33
34
35
36
37
38
39
40
41
42
43
44
45
46
47
48
49
50
51
52
53
54
55
56
57
58
59
60
61
62
63
64
65
0
0.2
0.4
0.6
0.8
1
750 800 850 900
Fraction of V in solution
Temperature (ºC)
V1
V2
Figure 11.Theoretical mass fraction of V in solid solution as a function of temperature.
0
2
4
6
8
10
12
780 800 820 840 860 880
Temperature (ºC)
fd-1/2 (mm-0.5)
V1(air) V1(acc) V2(air)
V2(acc) CMn(air) CMn(acc)
Figure 12. Evolution of the ferrite phase relevance in the strength, considered through the
f.d-0.5 factor, as a function of the deformation temperature.
1
2
3
4
5
6
7
8
9
10
11
12
13
14
15
16
17
18
19
20
21
22
23
24
25
26
27
28
29
30
31
32
33
34
35
36
37
38
39
40
41
42
43
44
45
46
47
48
49
50
51
52
53
54
55
56
57
58
59
60
61
62
63
64
65
0
20
40
60
80
100
120
010 20 30
fM (%)
CVN (J)
V1 (air)
V1 (acc)
V2 (air)
V2 (acc)
Figure 13. Dependence of room temperature Charpy
energy with hard phase (MA) volume fraction.
0
20
40
60
80
100
120
700 800 900 1000 1100
Tensile strength (MPa)
CVN (J)
800ºC
835ºC
870ºC
V1 V2
C-Mn
V1
V2
0
20
40
60
80
100
120
700 800 900 1000 1100
Tensile strength (MPa)
CVN (J)
800ºC
835ºC
870ºC
V1 V2
C-Mn
V1
V2
Figure 14. Relationship between room temperature Charpy energy and tensile strength in air
cooled (closed symbols) and accelerated cooled (open symbols) conditions.
