Memorie - Analisi e Calcolo

Transcript

Metallurgia
La
Organo ufficiale
dell’Associazione Italiana di Metallurgia
Rivista fondata nel 1909
Poste Italiane spa - Spedizione in abbonamento postale - DL 353/2003 (conv. in L. 27/02/04 n. 46) art. 1., comma1 DCB UD.
Anno 107, n. 1 - Gennaio 2015 - Periodico mensile
Italiana
N. 1
Gennaio 2015
Anno 107
CONSEDIT - Viale Europa Unita, 29 - 34073 Grado (GO) - Tel. 0431 876070 - Fax 0431 886507 - www.consedit.com - [email protected]
Giornate Nazionali
sulla
Corrosione
e Protezione
XI EDIZIONE
Ferrara - 15-17 giugno 2015
www.aimnet.it/gncorr2015.htm
Organizzate da
ASSOCIAZIONE
ITALIANA
DI METALLURGIA
Con il patrocinio di
Provincia di Ferrara
Coordinatore delle Giornate
Prof. Cecilia Monticelli
Presentazione
Le Giornate Nazionali sulla Corrosione e Protezione si terranno a Ferrara dal 15 al 17 giugno 2015. Nel corso
delle loro edizioni, le Giornate Nazionali sulla Corrosione e Protezione si sono affermate su scala nazionale
come punto di incontro di ricercatori e aziende interessati a discutere questioni scientifiche, tecnologiche e
produttive, nell’ambito della corrosione e protezione dei materiali. Il Convegno prevede la presentazione dei
risultati raggiunti da vari gruppi di ricerca pubblici e privati. Anche in questa undicesima edizione sono stati
istituiti dei premi, destinati a giovani ricercatori che si distingueranno, nell’ambito della manifestazione, per
l’importanza e l’attualità dei temi proposti nelle loro letture.
Il Convegno sarà l’occasione ideale per commemorare la figura del Prof. Giordano Trabanelli, eminente studioso dei
fenomeni di corrosione e in particolare dei fenomeni di inibizione della corrosione di metalli e leghe, direttore per
lunghi anni del Centro di Studi sulla Corrosione “A. Daccò” di Ferrara e recentemente scomparso. A tale scopo, il prof.
Yuri Kuznetsov dell’Istituto Frumkin di Chimica Fisica ed Elettrochimica dell’Accademia Russa delle Scienze a Mosca,
studioso noto a livello internazionale per lo studio dei fenomeni di inibizione della corrosione, parteciperà alle Convegno con una sessione plenaria. Inoltre, alla memoria del Prof. Trabanelli sarà dedicata la sessione “Inibitori di corrosione”.
Aree tematiche principali
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Spazio aziende
È previsto uno spazio per l’esposizione di apparecchiature, per la presentazione dei servizi e per la distribuzione di materiale promozionale. Informazioni più dettagliate possono essere richieste alla Segreteria organizzativa del convegno ([email protected]).
Presentazione di memorie
Gli interessati a presentare memorie scientifiche dovranno inviare entro il 27 febbraio 2015, il titolo
della memoria, i nomi degli autori della memoria e la loro affiliazione ed un sommario di circa 500 parole
mediante il modulo online presente sul sito www.aimnet.it/gncorr2015.htm.
Date importanti:
Invio titolo e riassunti
Notifica accettazione
27 febbraio 2015
20 marzo 2015
Apertura iscrizioni
27 marzo 2015
Invio dei testi completi 30 aprile 2015
Atti
Gli atti del Convegno saranno predisposti sotto forma di CD-Rom e distribuiti agli iscritti all’inizio dei lavori.
Sede
La manifestazione si terrà dal 15 al 17 giugno presso le sale Imbarcadero del Castello Estense di Ferrara in Largo Castello
'FSSBSB1FSNBHHJPSJJOGPSNB[JPOJDPOTVMUBSFJMTJUPXXXDBTUFMMPFTUFOTFJU
Segreteria organizzativa
AIM - Associazione Italiana di Metallurgia
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EFC
Event no. 389
La Metallurgia Italiana
Metallurgia
La
Italiana
International Journal of the Italian Association for Metallurgy
Organo ufficiale dell’Associazione Italiana di Metallurgia.
Poste Italiane spa - Spedizione in abbonamento postale - DL 353/2003 (conv. in L. 27/02/04 n. 46) art. 1., comma1 DCB UD.
Anno 107, n. 1 - Gennaio 2015 - Periodico mensile
Mensile dell’Associazione Italiana di Metallurgia fondata nel 1946
N. 1
Gennaio 2015
Anno 107
Organo ufficiale
dell’Associazione Italiana di Metallurgia
Direttore Responsabile:
Gianangelo Camona
Comitato scientifico - Editorial Panel:
Livio Battezzati, Riccardo Carli, Mario Conserva, Augusto Di
Gianfrancesco, Franco Dinucci, Carla Gambaro, Gian Luca
Garagnani, Bevis Hutchinson, Chong Soo Lee, Alberto Molinari,
Roberto Montanari, Elena Pereloma, Giorgio Poli, Emilio
Ramous, Claudia Rinaldi, Roberto Roberti, Hans J. Roven, Dieter
Senk, Piotr R. Scheller, Pierre Soulignac, Jean-Marc Steiler,
Stefano Trasatti, George F. Vander Voort, Maurizio Vedani
CONSEDIT - Viale Europa Unita, 29 - 34073 Grado (GO) - Tel. 0431 876070 - Fax 0431 886507 - www.consedit.com - [email protected]
N. 1/Gennaio 2015
Anno 107 - ISSN 0026-0843
Segreteria di redazione:
Antonella Donzelli
Comitato di redazione:
Federica Bassani, Gianangelo Camona,
Antonella Donzelli, Ottavio Lecis, Carlo Mapelli
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Acciaio
The effects of composition and thermal path on hot ductility of forging steels
B. M. Connolly, J. Paules, A. DeArdo................................. 3
Evaluating freckle tendency for electro-slag remelting
30CrMnSiNi2A ingots by experiments and simulation
Y. Zhang, W. Q. Chen, L. Chen , Q. Z. Yan, C. W. Liz............11
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Vantaggi del riscaldamento ad induzione, nuove possibilità
di efficienza e flessibilità per i laminatoi per prodotti lunghi
A. Lainati............................................................................. 35
Processi e impianti
Riduzione delle perdite al camino e delle emissioni nocive
grazie ai nuovi bruciatori recuperativi a canali suddivisi e
rigenerativi
J. G. Wünning, A. Lazzaretto, A. Milani................................21
Alluminio e leghe
Effect of Grain Refinement on Microstructure and Wear
Behavior of Cast Al-7Si Alloys
P.C.Meena, A. Sharma, S. Singh.........................................25
Colata continua
High-precision numerical simulation for effect of casting
speed on solidification of 40Cr during continuous billet
casting
Y. Chen, Z. Peng, L. Wu, L. Zhao, M. Wang, Y. Bao..............47
Saldatura
Damage investigation on welded tubes
of a reforming furnace
E. Guglielmino, R. Pino, C. Servetto, A. Sili.........................53
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Acciaio
The effects of composition and thermal path
on hot ductility of forging steels
B. M. Connolly, J. Paules, A. DeArdo
This work examines the effects of composition and thermal handling path on the hot ductility of as-cast steel
forging ingots. Poor ductility of the as-cast structure can lead to cracking of the ingot prior to forging or the
formation of tears early during the forging process. The as-cast structure is particularly susceptible to cracking
due to the large grain size and high degree of microsegregation present.
Experiments were conducted to evaluate the ductility of the as-cast steel with varying levels aluminum and
nitrogen. Multiple thermal handling paths were followed in order to approximate the different thermal conditions
experienced approximately six inches below the surface of a large (~40 MT) steel ingot following solidification.
Hot tension testing after in-situ melting and solidification was used for quantitative measurements of the
material ductility. The majority of testing was carried out on a modified P20 mild tool steel. The experiments
indicate a significant loss of ductility for materials with high aluminum and nitrogen contents
(AlxN = 5.2x10-4) in the temperature range of 950 °C - 1050 °C upon solidification and direct cooling to the
test temperature. This behavior is not present in material with AlxN products below 1.3x10-4. All materials
tested exhibited a loss of ductility when the sample was cooled to 900 °C, immediately reheated to 1000°C and
tested. With increasing hold times at 900 °C prior to reheating to 1000 °C, the material with high aluminum and
nitrogen contents recovers ductility much more quickly than the low aluminum and nitrogen materials.
Funding in part by the Forging Industry Educational & Research Foundation and Ellwood Group, Inc.
Keywords: Steel - Hot Ductility - Aluminum Nitride
INTRODUCTION
Steel ingots remain a leading raw material for forging
facilities producing critical large cross-section parts in the
energy, mining, defense and aerospace industries. Since
the 1960’s, continuous casting has largely replaced ingot
casting due to increased yields and an improved as-cast
structure.
However, continuous casting is currently limited in crosssection to approximately 800 mm diameter maximum
which precludes the material from use in heavy section
products. Steel ingots can be cast in weights up to 600
MT and cross sections up to 4.2 meters. Steel ingots are
Brendan M. Connolly
Ellwood Quality Steels and University of Pittsburgh
John Paules
Ellwood Material Technologies
Dr. Anthony DeArdo
University of Pittsburgh
Paper presented at the 2nd Int. Conf. Ingot Casting
Rolling and Forging - ICRF 2014, Milan 7-9 May 2014
La Metallurgia Italiana - n. 1/2015
also frequently used for specialty steel grades that are not
conducive to continuous casting due to either the chemical
composition or low tonnage requirements.
Steel ingots of various compositions can experience poor
hot ductility that leads to cracking which becomes apparent
during the forging process. Hot ductility troughs have been
identified and analyzed by many researchers [2, 3, 4, 5] in
the temperature range of 700˚ - 1200˚C, with the exact
temperature range and severity of the trough varying by
investigator, composition and experimental conditions.
The ductility trough is often depicted by plotting hot
tension percent reduction of area at failure (%RA) against
test temperature for multiple samples. A rapid drop in
%RA values is commonly shown within the temperature
range noted above, while ductility is higher both above and
below the trough. This loss in ductility has been attributed
to many factors including precipitation of nitrides and/
or carbides, segregation of impurities to austenitic
grain boundaries, formation of ferrite at austenitic grain
boundaries and combinations of these with other factors.
At lower temperatures, well below the Ar3 temperature,
the ductility usually returns to high %RA values. At higher
temperatures ductility is recovered by increased grain
boundary mobility, dissolution and/or coarsening of
3
Memorie
Steel
C
Mn
P
S
Si
Ni
Cr
Mo
V
Cu
Al
N
P20-low
P20-mid
P20-high
3.5Ni
4130
.34
.34
.34
.34
.33
1.45
1.46
1.53
.73
.55
.013
.008
.011
.006
.007
.003
.002
.003
.002
.001
.29
.30
.34
.20
.32
.80
.81
.78
3.52
.14
1.97
1.99
2.03
.99
1.04
.22
.22
.22
.60
.22
.057
.055
.059
.061
.048
.19
.17
.19
.15
.17
.0075
.015
.040
.012
.015
.0081
.0085
.0130
.0073
.0099
AlxN
(x104)
.6075
1.275
5.200
.8760
1.485
Table 1: Compositions of steels examined (wt %)
grain boundary precipitates and an increased ability to
recrystallize. At temperatures approaching the solidus,
ductility drops sharply due to incipient melting.
EXPERIMENTAL WORK
Commercial heats of the desired compositions (see Table
1, compositions in wt %) were produced.
The P20 material is a mild tool steel and heats were
produced with varied (low/mid/high) aluminum and
nitrogen contents. The 3.5Ni and 4130 materials were
produced with aluminum and nitrogen contents as close
as possible to those of the P20-mid material.
Steel samples were provided by Ellwood Group, Inc. from
slices of forged ingots. No samples were taken within 150
mm of the metallurgical centerline of the slices in order to
avoid compositional variations due to macrosegregation,
namely for carbon and sulfur.
Hot ductility can be tested by several methods [6] including
hot torsion, compression, tension and bending. Hot tension
is the testing method in this work due to the quantitative
nature of the ductility measurement. In a manner outlined
by Revaux, Bricout and Oudin [2], in-situ solidification
may be performed within the tension testing apparatus.
In-situ cast specimens exhibit grain boundary segregation
of aluminum which may reduce the bulk concentrations
required to exceed the solubility limit for AlN formation
[11], as is the case for as-cast forging ingots.
In-situ solidification can introduce a shrinkage cavity into
the sample due to the volume contraction that occurs
during the liquid to austenite phase transformation. The
design of the crucible in the present work maintains
soundness in the deformation zone by using a big-endup taper of ~7% and by introducing a notch at the midheight of the crucible. The notched area, with the smallest
cross-section, will bear the strain. The solidification cavity
remains above the notched area because as solidification
progresses, the higher density solid crystallites tend to fall
through the liquid allowing a progression of the solidified
front from bottom to top. Revaux, Bricout and Oudin used
light pressure to ensure the solidification cavity was moved
outside of the deformation zone; the same method was
employed in the current work. Only slight modifications
were made to the sample configuration, and the design of
the fused quartz crucible is nearly identical to that used
by Revaux, Bricout and Oudin. The sample and crucible
dimensions for this work are shown in Figure 1.
4
Fig. 1 - Sample and crucible dimensions for this work
Equipment
An Instron testing frame was fitted with high-temperature
alloy grips, insulated heating chamber and inert gas
purging. The heating chamber was constructed from a
10” diameter steel cylinder with 1/8” wall thickness. The
cylinder is insulated with 1” thick refractory fiber and
has a 4”x4” high temperature glass viewing window for
observation during testing. The cylinder is fixed to the
upper tensile grip arm and is hinged along its vertical axis
for sample access. The heating chamber has ports for inert
gas purging, induction coil, optical pyrometer sight and
an access port for breaking the melt crucible away from
the sample. A rendering of the insulated heating chamber
mounted on the tensile frame is shown below in Figures
2a and 2b.
The three-turn water-cooled induction coil is connected
to a 10 kW Pillar MK-20 computer controlled power
supply. The induction coil leads are isolated from the
heating chamber shell using electrically insulating cloth.
Temperature readings are obtained from an Omega IR2C
dual color optical pyrometer at a rate of 1s-1.
Testing
The tensile sample halves are placed into the grips with
Acciaio
Fig. 2 - a) [left] Melt chamber
and necessary connections
and b) [right] melt chamber
in open position
the upper half in a fixed position. The lower tensile half is
given approximately 2 mm of vertical freedom by loosening
the lower grip in order to avoid the formation of any tensile
stresses due to thermal contraction. The quartz crucible
is placed on the shelf of the lower tensile half and a 0.5”
diameter x 0.375” piece of filler metal is placed in the top of
the crucible. This filler material is machined from the same
metal as the tensile specimen. The filler metal is necessary
to ensure complete filling of the crucible on melting.
The samples are positioned within 2 mm of one another
and then the heating chamber is closed.
The optical pyrometer is positioned by using the laser dot
sight which is aimed at the sample. The access port of the
heating chamber is plugged and argon is purged for 10
minutes prior to heating the sample.
The sample is rapidly heated above the liquidus and the
lower crosshead is raised until the liquid metal fills the melt
crucible completely. The upper argon inlet is then closed to
force the heat extraction mainly in the downward direction
and allow the solidification front to move upward.
This prevents the shrinkage cavity from forming within
the deformation zone. The power input to the induction
coil is manually decreased so as to achieve a cooling rate
of approximately 0.05˚C/sec through solidification. The
lower crosshead is incrementally raised in order to maintain
a light compression on the sample and avoid formation
of unwanted voids. The temperature value during this
solidification step can only be treated as qualitative data
due to the presence of the quartz crucible.
On reaching a temperature reading of 1300˚C the upper
argon inlet is turned on to avoid overheating of the upper
tensile grips. On reaching a temperature reading of 1250˚C,
the power is turned off completely and the crucible is broken
away from the solidified tensile sample via the access port.
The process of removing the crucible, from power off to
power on, takes less than 30 seconds. The temperature drop
during this time is severe, approximately 200˚C. However,
because of the sample remaining in the austenitic range
and the sluggish precipitation kinetics of AlN in austenite,
it is reasonable to assume that this brief low temperature
period has little or no effect on the results.
Once the heating power has been switched back on,
the pyrometer is adjusted to give the maximum (true)
temperature reading. The induction coil control is
immediately switched to automatic mode and the closedLa Metallurgia Italiana - n. 1/2015
loop control follows the desired thermal path. On reaching
the testing temperature, the power input to the induction
coil is set as constant to avoid potential fluctuations during
the deformation. The load cell and displacement sensor
are zeroed and then tension is applied to the sample at a
constant lower cross-head speed of 0.005 inches/second
until failure occurs.
Immediately after failure occurs, the induction coil power
is turned off and the sample is quenched via a high-velocity
helium jet.
Three basic thermal paths were followed after
solidification:
•Direct cool to varied tension testing temperature
•Undercool to varied temperature, reheat to 1000°C,
tension test
•Undercool to 900°C, hold for varied time, reheat to
1000°C, tension test
The P20 materials with varied aluminum and nitrogen
contents were tested under each of the thermal paths
while the 3.5Ni and 4130 materials were tested only with
the direct cooling path for comparison to the P20-mid
material.
RESULTS
Ductility results for the experiments performed are
presented below in the form of percent reduction of area
versus temperature. The reduction of area was measured
using the final fracture cross-section in comparison to the
as-solidified minimum cross-section of 9.8 mm diameter.
Direct Cool Test Temperature
The samples which were cooled directly to the testing
temperature are intended to provide information about the
ductility of the as-cast ingot during the solidification and
cooling period while in the ingot mold and shortly after
removing the ingot from the mold. Figure 3 shows the thermal
path for the direct cool to test temperature experiments.
Figure 4 shows the results as %reduction of area vs testing
temperature for the P20 materials and Figure 5 shows the
results for varied base compositions (P20, 3.5Ni, 4130).
5
Memorie
Fig. 3 - Thermal path for direct cooling to test
temperature experiments
Fig. 5 - Results of direct cooling to test
temperature experiments for varied base
compositions
Undercool to Varied Temperature
The undercooling experiments are intended to provide
information on the ductility of an as-cast ingot which is
being heated to forging temperatures or is thermally
insulated on the surface after the initial cool down. The
varied undercooling temperatures would correspond to
different “track times” that may be experienced in industry.
For these tests, the samples were melted and then cooled
at 0.02˚C/s to the desired undercooling temperature,
held for 240 seconds at the undercooling temperature
and then reheated at 0.2˚C/s to 1000˚C and tested. The
value of 300˚C for the lowest temperature samples is for
representation purposes only. The actual temperature
was not measured but the sample was allowed to fully
transform. Full transformation was ensured by monitoring
the load cell readout; heating to 1000˚C was not started
6
Fig. 4 - Results of direct cooling to test
temperature experiments for P20 with varied
aluminum and nitrogen
Fig. 6 - Thermal path for undercooling experiments
until 10 minutes at constant load was achieved. Figure 6
shows the undercooling thermal path, and the results of
the experiments are shown in Figure 7. The samples with an
undercooling temperature of 1000°C are the direct cooling
experiment results and are shown for comparison.
Undercool to 900°C with Varied Hold Time
The poor ductility of the mid- and low-AlxN samples for the
testing described in the previous section prompted further
investigation. Samples were tested at 1000˚C after various
hold times at 900˚C in order to examine the change in
ductility as a function of hold time. The thermal path is the
same as shown in Figure 6 and the results are presented in
Figure 8 with %RA as a function of hold time at 900˚C.
Acciaio
Fig. 8 - Results of undercooling experiments with
varied hold times at 900°C
Fig. 7 - Results of undercooling experiments
Fig. 9 - SEM images of a) intergranular, b) mixed
mode, c) microvoid coalescence and d) fully ductile
failures
METALLURGICAL ANALYSIS
The fracture surfaces of the tensile specimens were
completely intergranular for samples that exhibited less
than 20% RA. One such fracture is shown in Figure 9a. The
highlighted grain in Figure 9a was mechanically removed
for further analysis. Samples with 20% - 60% RA showed a
mixed fracture mode of intergranular failure and microvoid
coalescence as shown in Figure 9b.
Samples with greater than 60% RA failed mainly by
microvoid coalescence with occasional, small areas of
intergranular separation. A typical fracture face showing
microvoid coalescence is shown in Figure 9c. Samples
with near 100% RA do not contain enough fracture surface
Fig. 10 - SEM-EDS analysis showing a) low
magnification view of intergranular crack, b) higher
magnification view of intergranular crack, c) EDS
analysis of matrix material and d) EDS analysis of
segregated second phase
for analysis due to the necking of the sample down to a
single point or line as shown in Figure 9d.
Grain boundary separation is visible on the fracture
face shown in Figure 9a. At higher magnifications, a thin
film containing additional microcracks along this grain
boundary was visible.
EDS analysis showed a second phase of heavily segregated
composition (increased phosphorus and sufur) near the
separated grain boundary. Figures 10a through 10d show
progressively higher magnification and EDS analysis of this
area. The segregated phosphorus and sulfur are apparent.
As noted earlier, the highlighted grain in figure 9a was
mechanically removed in order to examine the facets not
exposed to any atmosphere during testing. Even with the
7
Memorie
argon atmosphere during testing, the high temperature and
long testing time allows for some very mild oxidation of the
sample fracture surface. Figure 11 shows the grain surface
below the fracture cross-section after it was removed.
The dendritic structure and a thin grain boundary film are
readily visible. The transition between the grain surface
and the grain boundary film is highlighted in the lower right
of the figure.
Many non-metallic inclusions were found on the surface
of this grain, including MnS, Mn/CuS and V(C,N). Typically
the MnS inclusions were quite large, on the order of 5 - 10
µm. The Mn/CuS particles were smaller (1 - 5 µm) and
less frequent. The vanadium carbonitride particles were
dispersed all over the grain surface and were much smaller
than the sulfide precipitates, typically all much less than
1 μm as can be seen in Figure 12. Brown et. al.[4] have
shown that V(C,N) precipitates form preferentially over AlN
at intermediate temperatures and their presence should
be expected even with the greater solubility limit of V(C,N).
This is likely due to the high misfit strain of the hexagonal
close-packed aluminum nitride particles compared to the
cubic V(C,N) precipitates.
DISCUSSION
From the fracture surfaces shown, it is obvious that poor
ductility is resultant from some form of grain boundary
weakening or embrittlement. Within the P20 material
samples, the only significant variation in composition
is that of aluminum and nitrogen contents. This implies
that the different behaviors of the P20 materials can only
be due to formation of aluminum nitride particles or the
effects of aluminum and nitrogen in solution. The variation
in aluminum and nitrogen content is significant with
respect to the solubility product of AlN (variation of over
an order of magnitude), while only subtle differences, if
any, would be expected from solid solution effects.
The direct cooling to test temperature experiments
showed results that were in reasonable agreement with
similar works. The P20-high AlxN product material shows
a significant drop in ductility at temperatures between
950°C and 1050°C. This temperature range is consistent
with the onset of aluminum nitride precipitation according
to several published solubility products. There is essentially
no ductility loss in the intermediate AlxN product material
until below 900°C and no ductility loss in the low AlxN
product material until below 800°C.
Perhaps the most interesting results of this work are the
improved ductility of high AlxN material in the undercooling
experiments as compared to the low- and mid-AlxN
materials. All samples showed that improved ductility occurs
at undercooling to temperatures of 700°C and below, but
the high-AlxN material showed excellent ductility even at
800°C. With extended holding times at the undercooling
temperature of 900°C, the rate of ductility recovery
increases with increasing AlxN product. The cause for this
improved ductility in the high AlxN material is thought to
8
Fig. 11 - SEM image showing dendritic fracture
surface with thin film
be due to aluminum nitride particle coarsening. For a given
volume fraction of grain boundary precipitate, the pinning
force is inversely related to the radius of the precipitates.
The extremely hyperstoichiometric composition of the
high-AlxN material can be expected to undergo particle
coarsening at a much faster rate than in the mid- or lowAlxN P20 material, which would explain the faster increase
in ductility. The mid- and low-AlxN P20 materials did not
exhibit poor ductility on direct cooling to 900°C or 1000°C,
but on undercooling to 900°C and reheating to the test
temperature of 1000°C both materials showed essentially
zero ductility. It has been noted [1,7,8,9] that precipitation
of aluminum nitride is very rapid upon reheating samples
that have been previously cooled to a lower temperature.
The reason for this is not perfectly clear, however it is likely
that the reduced incubation time for AlN precipitation at
lower temperatures allows for a significant amount of
precipitation to occur during the undercooling step while
the decreased temperature prevents rapid growth. Upon
undercooling, a significant amount of nuclei can form and
growth occurs during the reheating period, whereas in
the direct cooling of these samples the volume fraction of
precipitates at the time of testing remains low.
The full recovery of ductility that occurs on undercooling to
low temperatures is a result of a significant volume fraction
of ferrite precipitation. On re-austenitizing, the harmful
precipitates and segregated compositions that were
present on the prior austenite grain boundaries become
dispersed throughout the newly formed austenite grains.
The differences with varying base composition were
very subtle. This work has shown that increasing the
nickel content causes a small decrease in ductility at
temperatures above 850°C while there is a slight increase
in ductility at temperatures below 850°C. Erasmus[10]
Acciaio
in this work. At higher temperatures the ductility is slightly
reduced with increasing nickel and at lower temperatures
the ductility is slightly improved.
It must be noted that the work described in this paper
is ongoing. A deeper understanding of themechanisms
of ductility loss in the experiments presented here
will be gained via TEM analysis of thin foil and carbon
extraction replicas from the samples along with additional
metallographic examination.
REFERENCES
Fig. 12 - Showing EDS analysis of V(C,N)
precipitate
noted that additions of nickel decrease the solubility of
nitrogen in austenite which will raise the precipitation
temperature of aluminum nitride. However, in examining
the Fe-N-Ni system[11] it is seen that the effect is minimal.
For example at 1200°C, an increase from 0.0% Ni to 3.3%
Ni only decreases the soluble nitrogen from 235 ppm
down to 210 ppm.
SUMMARY
There is very little published information available
regarding the hot ductility of large-grained, as-cast,
slow-cooled steels. A thorough understanding of the
effects of thermal path and composition on ductility can
provide ingot producers and forgers with the fundamental
information necessary for well-designed thermal handling
and processing procedures. The following important points
are noted from this work:
Ingot handling within the temperature range of 950°C
- 1050°C should be avoided for materials with high
aluminum and nitrogen contents.
Poor ductility may be experienced in cast steels even with
relatively low aluminum and nitrogen contents in cases
where temperature oscillations occur in the temperature
range of 900°C - 1000°C.
Hyperstoichiometric AlxN compositions can show rapid
recovery of ductility during holding at high temperature,
however in EAF steels the necessary aluminum content
will likely cause precipitation problems at higher
temperatures.
As-cast, slowly cooled steels exhibit severe austenite
grain boundary segregation of phosphorus and sulfur
as well as grain boundary precipitation of sulfides and
carbonitrides.
Aluminum nitride is also expected on the austenitic grain
boundaries but its presence was not confirmed in this
work.
The effect of increased nickel content was found to be mild
1) F.G. Wilson and T. Gladman, “Aluminum nitride in
steel”, International Materials Reviews, 1988, Vol. 33,
No. 5, 221 – 283.
2) T. Revaux, J.P. Bricout and J. Oudin, “A New Tensile
Testing Procedure for Predicting Transverse Cracking
Susceptibility of Continuous Casting Slabs”, Journal
of Materials Engineering and Performance, April
1996, Vol. 5(2), 260 – 268.
3) W.T. Nachtrab, W.T. and Y.T. Chou, “The Effect of Sn,
Al, and N on the Hot Ductility of a Carbon-Manganese
Steel between 700˚ and 1200˚C”, Metallurgical
Transactions A, May 1988, Vol. 19A, 1305 – 1309.
4) E.L. Brown, L.J. Cuddy and A.J. DeArdo,”Aluminum
Nitride Precipitation in Microalloyed Steels”, The
Thermomechanical Processing of Microalloyed
Austenite, A.J. DeArdo, G.A. Ratz and P.J. Wray, Editors,
(Warrendale, PA, TMS-AIME: 1982), 319 – 341.
5) M. Vedani, D.Dellasega and A. Mannuccii,
“Characterization of Grain-boundary Precipitates
after Hot-ductility Test of Microalloyed Steels”, ISIJ
International, 2009, Vol. 49, No. 3, 446 – 452.
6) A. Nicholson, “Hot Workability Testing of Steels”, Iron
& Steel, June - July 1964, 290 – 294, 363 – 380.
7) U.H. Lee, T.E. Park, K.S. Son, M.S Kang, Y.M. Won,
C.H. Yim, S.K Lee, I. Kim, and D. Kim, “Assessment
of Hot Ductility with Various Thermal Histories as
an Alternative Method of in situ Solidification”, ISIJ
International, 2010, Vol. 50, No. 4, 540 – 545.
8) G.A. Wilber, R. Batra, W.F. Savage and W.J Childs,
“The Effects of Thermal History and Composition on
the Hot Ductility of Low Carbon Steels”, Metallurgical
Transactions A, September 1975, Vol. 6A, 1727 –
1735.
9) C. Spradbery and B. Mintz, “Influence of undercooling
thermal cycle on hot ductility of C-Mn-Al-Ti and C-MnAl-Nb-Ti steels”, Iron and Steelmaking, 2005, Vol. 32,
No. 4, 319 – 324.
10) L.A. Erasmus, “Effect of aluminum additions on
forgeability, austenite grain coarsening temperature,
and impact properties of steel”, Journal of the
Iron and Steel Institute, January 1964, 32 – 41.
[correspondence from Dec 1964 included].
11) P. Perrot, “Iron-Nitrogen-Nickel”, MSIT, Springer
Material Database, 2008.
9
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Acciaio
Evaluating freckle tendency for electro-slag
remelting 30CrMnSiNi2A ingots
by experiments and simulation
Y. Zhang, W. Q. Chen, L. Chen , Q. Z. Yan, C. W. Li
The mechanisms and tendency of freckle formation in industrial scale ESR ingots were studied by
thermodynamic calculation, thermo-physical property calculation, metallographic observation, and composition
analysis. The macrostructure and compositions of the freckle regions in a low alloy ultrahigh strength ESR steel
ingot were investigated to clarify the freckle formation mechanism. Combining the results with composition
analysis and thermodynamic calculation, it can be concluded that the compositions of freckles correspond
to that of liquid with a liquid fraction of 0.24 to 0.41, and the freckles were caused by the upward solute-rich
liquid flow which initiate at the lower part of mushy zone. The relative Rayleigh number (Ra), a freckle criterion
considering the effect of a tilted solidification front, was adopted to evaluate the freckle formation tendency
in the industrial scale ESR ingots. The calculated results of Ra numbers of different locations are in good
agreement with the actual distribution of the freckles in the ingot.
Keywords: Freckles, Electro-slag remelting - Thermodynamic - Rayleigh number
Introduction
Thermosolutal convection due to thermal and compositional
variation can cause the formation of freckles (channellike macrosegregation defects). Freckles are commonly
observed in cast alloy or steel ingots, such systems
generally contain two or more alloying elements with
different densities and phases [29]. Freckles appear as
long liner trails of misoriented dendrites with a composition
shift consistent with alloy segregation. Generally, freckles
in full size ingot have a size from several millimeters to
one centimeter in diameter, which at least 2 or 3 orders of
magnitude larger than that of microsegregation. Freckles
are highly undesirable in critical applications because of
their deleterious effect on mechanical performance, such
as tensile, fracture toughness, and fatigue crack growth
are significantly degraded when the longitudinal direction
of the freckle channels is perpendicular to the loading
axis. The hardness of the freckle region is much higher
than that of matrix since some alloying elements enriched
in freckle regions, then micro-cracks may formed from the
interfaces between the freckles and matrix. The freckles
cannot be removed by any heat treatment once it has
formed in a fully solidified casting [1-4].
Yang Zhang, Weiqing Chen
State Key Laboratory of Advanced Metallurgy, University
of Science and Technology Beijing, Beijing 100083, China
Lie Chen, Qingzhong Yan, Chengwei Li
Xining Special Steel Co., LTD., Xining 810005, China
It is generally agreed that the formation of the freckles is
due to a complex interaction of solute segregation, thermal
variation and dendrite morphology, all of which contribute
to the onset of thermosolutal convection in the mushy
zone [1]. The probability of the frcekle formation during the
electro-slag remelting (ESR) process is mainly dependent
on the following factors: the alloy composition, the
solidification parameters (melting rate of electrodes, radial
heat extraction, and cooling intensity, et), the defects of
electrode (for example, the perturbations of the melting rate
can be resulted from using of the electrodes with internal
cracks, which promote the thermosolutal convection in the
mushy zone), and the casting size [15-18].
Many efforts [1-20] have been made on the mechanism and
criterion of freckle formation since the macrosegregation
theory was proposed by Flemings and his co-workers [5-6].
However, few investigations have been focused on the
freckle defects of the electro-slag remelting (ESR) ingots
[4,15-20]. Previous researches ignored the effects of the
variation of liquid viscosity on the interdendritic liquid
flow, the back diffusion of interstitial elements in solid on
the liquid composition and the solidification temperature
range. For example, the solidification temperature range
calculated by Scheil-Gulliver equation (where no backdiffusion is considered) is generally 100-300°C larger than
that of the actual solidification for steels. Thus the above
two factors will be considered in this investigation.
This study aimed at enhancing the fundamental
understanding of the freckle formation mechanism and
tendency evaluation during the ESR solidification process
11
Memorie
by experimental investigations and simulated calculation.
The thermodynamic calculation software (Thermo-Calc)
and Java-based materials properties software (JMatPro)
were used to predict the variation of liquid composition,
density, and viscosity during the solidification. Finally,
a Rayleigh number criterion that represents the ratio of
the driving force for the interdendritic liquid flow and the
resistance for the flow, it is appropriate to describe the
instability convective flow phenomenon in the mushy zone
according to the theory of thermosolutal convection [20].
This criterion includes the effect of liquid composition and
viscosity variation, the slope of the solidification front,
permeability of the mushy zone, and the cooling conditions,
was adopted to evaluate the freckle tendency of industrial
scale ESR ingots.
Experimental procedures
A 720-mm-diameter ingot of 30CrMnSiNi2A was
produced on an industrial-scale ESR furnace, the chemical
composition of the steel is shown in table 1. The electrodes
with a 400-mm diameter used for the ESR process was
produced by vacuum-induction melting (VIM). A 30-mmthick slice for macro etching was sectioned along the
longitudinal centerline of the ESR ingot. The slice was
etched using 50 pct HCl-50 pct H2O to reveal freckles.
Then 25 cube samples with a dimension of 30mm × 30mm
× 30mm for dendrite etching was prepared along the radial
direction and the axial direction of the slice, the sampling
locations of the radial direction are edge, 2/3-radius, 1/2radius, 1/3-radius, and centre, and the sampling locations
of the axial direction are top, 3/4-height, middle, 1/4height, and bottom, the method is illustrated in Fig. 1. The
25 samples were etched by 5 pct trinitrophenol to reveal
the dendritic structure, and the secondary dendrite arm
spacing (SDAS) was measured using the image analysis
software ImageTool (Version 3.0, Department of Dental
Diagnostic Science at The University of Texas Health
Science Center, San Antonio, Texas). Meanwhile, two
samples for the freckles analysis were cut from the top
and the bottom of a freckle channel (as Fig. 2 shows). An
EVO 18 scanning electron microscope (SEM, Carl Zeiss
Microscopy GmbH, Jena, Germany) equipped with energy
dispersive X-ray spectrometer (EDS) was used to acquire
analytical composition data in the second electron image
(SEI) mode. And the same samples (as-polished) were also
used for measuring the microhardness by a LEICA VMHT
30M type Vickers microhardness tester with the applied
load 50 g and the loading time 15 s.
Thermodynamic and thermo-physical
approaches
Thermodynamic calculation
Based on the Scheil-Gulliver model [25,26], the traditionalScheil module has been implemented inside Thermo-Calc
12
Fig. 1 - Schematic diagram of the sampling positions
C Si Mn P
S
Cr Ni Cu
Ti
Al
0.30 1.05 1.09 0.015 0.003 1.03 1.62 0.06 0.018 0.063
Table 1 - Chemical compositions of 30CrMnSiNi2A
steel, wt%
software. The module assumes that solute elements
diffuse rapidly within the liquid phase and that diffusion
in the solid phases is negligible. Along each step in the
cooling process, local equilibrium is established at the
solidification interface where the compositions of liquid
and solid are given by the system’s phase diagram [17].
On the basis of the traditional-Scheil module, one or more
fast diffusing elements (usually interstitial elements such
as C, N, O, S, and B) can be defined in the modified-Scheil
module, thus their back diffusion in solid phases can be
taken into account during a simulation. The back diffusion
of interstitial elements in solid phases is complete, but
that of substitutional elements is negligible. Interstitial
elements are allowed to diffuse freely among the liquid and
all portions of the solid phases formed, and substitutional
elements are completed limited in the remaining liquid
and each portion of the solid phases formed [18]. And for
steels with primary ferrite, the transition from δ-ferrite
to austenite can be considered in this module, such a
transformation will change the chemical potentials of
interstitial elements and thus influence their back diffusion
in solid phases.
Then, the modified-Scheil module of Thermo-Calc software
was used to predict the composition and the density of the
liquid during the solidification, and the alloy composition
was assumed to be Fe-0.3C-1.05Si-1.09Mn-1.03Cr-1.62Ni0.015P here. Meanwhile, C and P were defined as the fast
diffusing elements, and the transition from δ-ferrite to
austenite during peritectic reaction was also considered.
Thermo-physical property calculation
Java-based materials properties software (JMatPro)
provides extensive information on how the properties
of an alloy or each individual phase may change within
its specification range, such as the liquid phase in the
mushy zone, which is usually beyond the capability of
measurement [19]. For each individual phase in multicomponent systems, properties can be calculated by
simple pair-wise mixture models.
(1)
Acciaio
Fig. 2 -Macrostructure of the freckle in the
30CrMnSiNi2A steel ingot: (a) freckle 1; (b) freckle 2
Fig. 3 - The relation of fluid flow and dendrite growth
at a tilted solidification front
Where, P is the property of the phase, P10 is the property
of the phase in the pure element, ùnij is a binary interaction
parameter dependent on the value of n, xi and are the
mole fractions of elements i and j in the phase. Both P10 and
ùnij are temperature dependent and it is possible to include
ternary or higher order effects where appropriate [20].
The General Steel solidification properties module of
JMatPro software was used to calculate the variation of
liquid viscosity during the solidification in this research. The
starting temperature of the continuous cooling process was
set at 1600 °C and a cooling rate of 0.1°Cs-1 was chosen.
In addition, C and P were defined as interstitial elements,
and the alloy composition used in this calculation is shown
in Table 1. First, the amount and composition of liquid
in this steel is calculated based on thermodynamics and
phase transformation kinetics, and the back diffusion of
interstitial elements is also considered in this calculation,
then the property of liquid can be calculated by Eq. (1). For
the case of steels, interstitial elements diffuse rapidly in
the solid state and it is possible to consider that complete
back diffusion of interstitial elements will occur. Such a
model has been implemented in JMatPro by considering
that interstitial elements will diffuse sufficiently rapidly so
that their composition in the growing austenite or ferrite
phases will be equal to that of the solid at the growing
solid/liquid interface [28].
has higher solute concentration, it can reduce the local
undercooling and, thereby retard the growth of dendrites.
As a consequence, the dendrite arms in the segregation
channel can remelt, potentially sustaining the formation of
the segregation channels. There is also an additional force
on the arms in the segregation channel, upward flow. This
flow will produce surface shear on the arms, which will
concentrate as a bending moment at the neck of dendrite
arms. By this way, both primary and secondary arms in the
channel were etched/melted.
The freckles were only found in the 1/2-radius region
from the upper part (from about 2/5-height to the top)
of the ingot, but not in the centre region. Two samples
marked by the white rectangles (Fig. 2) were cut from the
top and the bottom of a freckle channel in order to reveal
the dendritic structure and measure the composition of
the freckle, the two regions were referred to as “freckle 1”
and “freckle 2”. The right part of Fig. 2 gives the dendritic
structure of freckle 1 and freckle 2. It shows the growth
direction of the primary dendrites is basically same with
the thermal gradient direction, is from left-bottom to topright. The freckles are basically running from bottom to top
but slightly leaning from right to left.
When thermosolutal convection starts, partitioning of
the lighter solute elements cause the interdendritic fluid
to advect out of the interdendritic region. Owing to the
conservation of continuity, the channel draws segregated
liquid from the neighbouring interdendritic regions. Then
the solute-rich liquid flows upward through the mushy
zone into the liquid layer, the upward flow of solute-rich
liquid is accompanied by a downward flow of solute-lean
liquid to ensure the conservation of volume [29]. Thus the
concentration of solute elements in the channel gradually
decreases along the flow direction of solute-rich liquid,
consequently, the velocity inside the ascending segregated
liquid and the ability of solute-rich liquid to etch/remelt
dendrite arms also decreases. It can be obviously seen
from Fig. 2b that there is a misoriented dendritic structure
in the freckle 2, which is very different from the surrounding
Results and discussion
Macrostructure of freckle
During solidification, solute elements are rejected from
the dendrites and their concentration gradually increases
in the interdendritic region. Si, Mn, Cr is lighter than
Fe, and therefore natural convection is initiated and the
interdendritic liquid flows upwards. The interdendritic
segregated liquid is solute-rich, both flowing upwards
due to buoyancy forces and altering the local free energy
driving force for solidification. Since the segregated liquid
13
Memorie
Element
Freckle
Matrix
SD
Average error
Si
1.46
1.10
0.05
±0.2
Mn
1.59
1.20
0.11
±0.2
Cr
1.44
1.15
0.12
±0.1
Ni
1.89
1.61
-0.01
±0.2
HV (Mpa)
365.5
275
...
...
Tab. 2 - Compositions of freckle 2 and the surrounding area, wt-%
Fig. 4 - The elements distribution map of the freckle region (upper-left part) and the matrix: (a) Si, (b) Mn, (c) Cr,
and (d) Ni
matrix. However, the dendrite morphology in the freckle 1
is basically same as that in the surrounding matrix, only
part of dendrites in the freckle 1 (Fig. 2a) were etched/
melted. The obvious difference of the dendrite morphology
between the freckle 1 and freckle 2 confirmed that the
upward solute-rich liquid flow is the reason of freckles
formation in the 30CrMnSiNi2A ESR ingots, the liquid
flow in the bottom of a freckle channel is stronger than
that in the top, thus the original dendritic structure in the
freckle 2 was etched/melted completely, The flow during
14
remelting may be sufficient to break off the dendrite arms
in the segregation channel. These are heavier and hence
may fall, as the fragments fall the temperature around
them decreases, and some may reach equilibrium before
completely remelting, which acting as pre-existing nuclei to
form the classic misoriented dendritic structure observed
experimentally.
Fig. 3 shows the relation between fluid flow and dendrite
growth in the mushy zone. The solidification front which
has an angle θ to the horizontal plane is represented
Acciaio
top
3/4-height
middle
1/4-height
bottom
edge
46
64
50
65
65
2/3-radius
38
38
35
45
60
1/2-radius
40
35
35
40
50
1/3-radius
26
20
30
28
23
Centre
20
10
25
16
18
Tab. 3 - Solidification front angle at different locations (deg)
Fig. 5 - Liquid composition as function of Liquid
Fraction during the solidification
by the isotherm (isodensity), which is perpendicular to
the dendrite growth direction of the primary dendrites.
R is the growth velocity, g is the gravity. The flow of the
interdendritic solute-rich liquid is represented by V. For
this steel, the flow of the solute liquid was observed to
closely follow the vertical direction at a small angle of
deviation.
The formation of the freckles is a result of the thermosolutal
convection in the mushy zone. The main factors affecting
the thermosolutal convection are the thermal gradient,
the permeability in mushy zone, and the liquid density
difference. The dendrites growth direction is determined
by the thermal gradient. On the other hand, the flow of the
interdendritic fluid was limited by the dendritic structure,
considering the freckle formation as a case of solute-rich
liquid flow in a porous media, then the permeability is a key
parameter that affects the flow velocity within the mushy
zone. In addition, solute elements are ejected from solid
into liquid during the solidification, the liquid composition
in the vicinity of the dendrite bottom is different from that
of dendrite tip. Such liquid density difference plays a key
role in driving the interdendritic thermosolutal convection
and the freckle formation.
Inside the mushy zone, the buoyancy forces may become
sufficiently strong to overcome the resistance offered by
the dendritic structure as an increase in the liquid density
difference and a decrease in the liquid viscosity, and
upward flow of solute-rich liquid will occur. Because the
permeability for liquid flow perpendicular to the primary
dendrites is 2.5 to 4 times greater than that parallel to the
primary dendrites [25,26], thus the interdendritic soluterich liquid tends to flow along the path which offers lower
resistance. It is therefore suggested that, under similar
solidification conditions, a tilted solidification front should
favor of upward flow of the interdendritic liquid, while a
horizontal solidification front is freckle-free. As Fig. 3
shows, the interdendritic solute-rich liquid preferentially
flows along a low resistance path that has an angle to the
vertical direction, which is the result of the competition
between the driving force of flow offered by the liquid
density difference and the resistance from the dendritic
structure in the mushy zone. The interdendritic solute-rich
liquid would retain its composition and remelt/erode the
already formed dendritic structure when it flows upward to
higher temperature positions, because the solute diffusivity
in the liquid is much lower than the thermal diffusivity.
Microscopic characteristics of freckle
Freckle composition
A polished but unetched freckle sample was analyzed by
SEM/EDS. The compositions of the freckle region and its
surrounding matrix, with standard deviation (Standard
deviation is the difference between the compositions of the
matrix of freckle region and chemical compositions of this
steel (as Table 1 shows)), is shown in Table 2. The results
are the average compositions from the measured results
at least five different locations. Considering the average
compositions of the matrix should equal to the chemical
compositions of this steel, the measured compositions of
the freckle region can be modified by standard deviation
(SD). Fig. 4 shows the elements distribution map of the
freckle region and its surrounding matrix measured
by EDS. It can be seen from Table 2 and Fig. 4 that the
freckle region is enriched with Si, Mn, and Cr compared
with the surrounding matrix. Among these elements, the
segregation of Mn is the largest, and that of Ni is the
least.
The Vickers microhardness value of the freckle region and
the surrounding matrix are shown in the right part of Table
2. It can be seen that the microhardness in the freckle
region is 33% higher than the surrounding matrix, microcracks may form from the interfaces between the freckles
and the matrix when steel subjected to mechanical and
thermal loadings.
15
Memorie
Secondary dendrite arm spacing (SDAS) and solidification
front angle
Fig. 8a shows the measured secondary arm spacing λ2 (μm)
at different locations of the ingot. It can be seen that the
maximum SDAS occurs at the location of 1/2-radius, and
the SDAS increases as the height of the ingot increases.
The joint cooling intensity of the bottom water tank and
mold, its maximum value occurs in the bottom, thus the
dendritic structure in the bottom is the finest, and that in
the top is the coarsest. Table 3 shows the solidification front
angle at different locations of the ingot, the results were
obtained by measuring the angle between growth direction
of primary dendrites and vertical direction (Fig. 3). The
maximum error in the determination of solidification front
angles is ±1 deg. It can be concluded that the tendency
of freckle formation increases as the increase of SDAS
and that of solidification front angle, thus solidification
conditions in the 1/2-radius region from the upper part of
the ESR ingot is favor of the freckle formation.
Thermodynamic and thermo-physical
approaches to calculate the Ra
Liquid composition during the solidification
Fig. 5 shows the relation between the liquid composition
and liquid fraction, which was calculated by the modifiedScheil module of Thermo-Calc software. It can be seen
that the liquid composition has an obvious change during
the solidification. The concentrations of Mn and Si
increases dramatically as the liquid fraction decreases,
but the concentration of Ni have a minor increase. The
density of Mn, Si, and Cr is less than Fe, which being
enriched in interdendritic liquid result in the decrease in
the liquid density. Comparing the concentration profiles of
interdendritic liquid (Fig. 5) with the modified compositions
of the freckle region (Table 2), it can be observed
that the segregation level of Cr was underestimated
in the thermodynamic calculation, but the calculated
results of Mn, Si, and Ni are reasonable. Therefore, the
correspondent liquid fraction of freckle compositions was
determined from the segregation profiles of Mn, Si, and
Ni, this method is illustrated in Fig. 5. The compositions of
freckles in the 30CrMnSiNi2A ingot correspond to a range
of liquid fraction 0.24-0.41.
Density and viscosity of the liquid during the solidification
The variation of the liquid density during the solidification
can be estimated with a relative high accuracy in modifiedScheil module by considering the effects of the volume
shrinkage, the composition change, the phase transition,
and the back diffusion of interstitial elements, etc. The
difference between the density at the liquidus temperature
(liquid fraction fL = 100) and the density at a temperature
T (TS≤T≤TL, Where TL and TS are the liquidus and solidus
temperature, respectively), represented by Δρ. Fig. 6a
shows the variation of liquid density during the solidification,
it can be seen that the liquid density difference increases
16
dramatically as the liquid fraction decreases.
The effects of composition and temperature on viscosities
of liquid alloy can be estimated by the following equations
[24]:
(2)
Where rL is the liquid density of alloy (kg·m-3), is the
absolute liquidus temperature (K), M is molar mass
(kg·mol-1), and R is gas constant (8.3144J·mol-1·K-1).
Liquid viscosity affects the flow velocity of interdendritic
liquid, it is very important to obtain reliable data of the
liquid viscosity for the investigation of the liquid flow
behavior in the mushy zone. While it is impossible to
measure the liquid viscosity directly, Fig. 6b shows the
variation of the liquid viscosity calculated by the General
Steel solidification properties module of JMatPro software,
it can be seen that the liquid viscosity decreases as the
liquid fraction decreases, which means that the flow
velocity of the interdendritic liquid increases.
Calculated results of relative Rayleigh number
The permeability is a key parameter that affects the flow
velocity within the mushy zone, and is the parameter that
is associated with the region of interdendritic channels
[26]. The permeability in the mushy zone represents
the resistance to the interdendritic liquid flow, which is
offered by dendritic structure and mainly affected by the
liquid fraction and the dendritic morphology. Poirier [25]
has developed the following expressions to estimate the
permeability for a flow parallel and normal to the primary
dendrites, KP and KN, respectively.
(3)
(4)
Where, KP and KN are the permeability components in
the parallel and perpendicular directions to the primary
dendrites, respectively, fL is the liquid fraction, λ1 is the
primary arm spacing, λ2 is the secondary arm spacing.
M. EL-Bealy and B.G. Thomas [27] proposed the following
equation for estimating the primary and secondary
dendrite arm spacing of low alloy steel, the cooling rate CR
(°Cs-1) can be calculated by the Eq. (5).
(5)
(6)
Where C0 (wt pct) is the average carbon content of this
alloy, m = ﹣0.206277638, n = ﹣0.1663998.
It is difficult to measure the primary arm spacing for the
ESR ingots directly, but CR can be calculated by Eq. (5)
Acciaio
Fig. 6 - The variation of
the liquid density and
liquid viscosity during the
solidification: (a) Liquid
density; (b) Liquid viscosity
according to the measured values of the SDAS, then the
primary arm spacing can be calculated by Eq. (6).
It is well known that the freckle is caused by the flow
of the solute-rich liquid in the mushy zone during the
solidification. The flow of the solute-rich liquid in ESR ingots
is mainly affected by the variation of the liquid composition
and the liquid viscosity, the mushy zone permeability, the
solodification front angle, and the cooling conditions. A
criterion that includes the above factors developed by Jairo
Valdés et [4]. was found to predict the freckle formation
satisfactorily. Rayleigh number (Ra), as given in Eq. (7).
(7)
Where n is the dynamic or absolute viscosity of the
interdendritic liquid; R is the dendrite growth rate. In
this study, the dendrite growth rate was replaced by the
cooling rate due to no reliable data being available, thus
a simplified equation (Eq. (8)) was used to calculate the
Ra, in order to distinguish the Rayleigh number (Ra) in the
these two equations, the Rayleigh number (Ra) in Eq. (8)
was named relative Ra.
(8)
Fig. 7 gives the calculated relative Ra number of
30CrMnSiNi2A, the negative value means that the liquid
density difference is negative. The absolute value of Ra
is proportional to the tendency for freckle formation.
Along the axial direction of the ingot, the bottom had the
maximum cooling rate, and correspondingly, its relative
Rayleigh number was the lowest. The relative Rayleigh
number in the middle region of the ingot are lower than
that in the top region due to the fact that the cooling
intensity is gradually to decreases as the height increases,
it can be confirmed by the fact that the maximum SDAS
occurs in the top region, and the freckles were only found
in the upper part (from about 2/5-height to the top) of the
ingot. Along the radial direction of the ingot, the maximum
relative Rayleigh number occurs in the 1/2-radius region.
Similarly, the freckles was only found in the 1/2-radius
region but not the centre of ingot. Compare with other
regions, the upward flow of the interdendritic fluid in the
centre region is very difficult because the solidification
front is basically horizontal.
The liquid density difference increases and the liquid
viscosity decreases, respectively, as the liquid fraction
decreases, but the permeability in the mushy zone
decreases, which means that the driving force, flow
velocity, and the resistance to the interdendritic liquid
flow increase at the same time. The interdendritic liquid
could no longer flow when the liquid fraction decrease to
within a specific range, because it is fully surrounded by
solid. The tendency to form freckle in a steel is related to
maximum Ra number through solidification. For this steel,
the maximum relative Rayleigh number occur at 0.36-0.39
liquid fraction due to the competition between the driving
force and the resistance of liquid flow, this result is basically
consistent with the liquid fraction range determined by the
compositions of the freckle.
From the above analysis, it can be found that a fine
dendritic structure and a small solidification front angle
are the key factors for preventing the freckles formation.
The finer the dendritic structure, the lower the permeability
in mushy zone and the higher the dendrite growth rate.
Correspondingly, the segregation level of the interdendritic
liquid is low, thus the liquid density difference is low. On
the other hand, a small solidification front angle means
the resistance to the upward liquid flow is lower than
that of a large solidification front angle, this is due to the
permeability anisotropy in the mushy zone. For the ESR
process, a fine dendritic structure and a small solidification
front angle can be obtained by reducing the melting rate,
increasing the cooling intensity, improving heat transfer of
the radial direction, increasing the fill ratio, and reducing
the height of the ingot, etc.
It can be observed from Fig. 7 that for a 30CrMnSiNi2A
ESR ingot, it is feasible to identify a threshold value of the
Rayleigh number that gives a clear boundary between the
freckle and non-freckle conditions. The threshold value
of the Rayleigh number was defined as Ra*, below which
there is no freckles formation. For 30CrMnSiNi2A, the
estimated threshold Rayleigh number Ra* was about -0.12
according to the present calculated results.
The tendency to form freckle in an ingot is related to
maximum Ra number through solidification. Fig. 8 shows
the comparison between the secondary dendrite arm
spacing and maximum Ra numbers at different locations
of the ingot, it can be seen that the variation tendency
of maximum Ra is basically as same as that of SDAS
(secondary dendrite arm spacing), since the SDAS of as17
Memorie
Fig. 7 - Relative Ra numbers at different location of 30CrMnSiNi2A steel ingot: (a) top of the Ingot, (b) 3/4-height
of the ingot, (c) middle of the ingot, (d) 1/4-height of the ingot, (e) bottom of the Ingot, and (f) freckle 1 and
freckle
cast ingots is inversely proportional to the cooling rate,
and the cooling rate in ESR ingots is mainly determined by
the melting rate of electrodes. High melting rate causes a
deep pool profile, extends the mushy zone, and induces
a stronger electromagnetic force field. Therefore, melting
rate is the critical operating parameter for controlling the
formation of freckles.
Conclusions
18
Thermodynamic and thermo-physical property calculation
were coupled with the metallographic observation and
composition analysis of an as-cast ESR steel ingot to
investigate the mechanisms of freckle formation and
evaluate the freckle tendency in industrial scale ESR ingots,
using this method the following conclusions were drawn.
1) Freckles enriched with Mn, Si, and Cr compared with
Acciaio
Fig. 8 - SDAS and Maximum Ra Numbers at different locations of the ingot: (a) SDAS; (b) Maximum Ra Numbers
the matrix, these were only found in the 1/2-radius region
from the upper part of the 30CrMnSiNi2A ingot. And the
microhardness in defect freckle region is much higher
than that in surrounding matrix.
2) Freckles are mainly caused by the upward solute-rich
liquid flow in mushy zone which would etch/remelt the
dendrite arms inside the segregation channels, and this
mechanism is confirmed by the dendritic micrographs
of freckles in the ingot. Meanwhile, the direction of
interdendritic liquid flow is strongly influenced by the
permeability anisotropy in mushy zone.
3) The freckles in 30CrMnSiNi2A ingot has a composition
corresponding to the interdendritic liquid with a liquid
fraction of 0.24 to 0.41. This result is compared to the liquid
fraction range where the calculated maximum Ra appears,
and the calculated results have shown good agreement
with experiment, thus the freckles in ESR ingots initiate at
the lower part of mushy zone.
4) The effect of the liquid composition variation on the
density and viscosity of liquid can be explicitly analysed
by combining the thermodynamic calculation with thermophysical property calculation.
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Processi e impianti
Riduzione delle perdite al camino e delle
emissioni nocive grazie ai nuovi bruciatori
recuperativi a canali suddivisi e rigenerativi
J. G. Wünning, A. Lazzaretto, A. Milani
La riduzione delle perdite di calore al camino, spesso, rappresenta la via più efficace e conveniente per
aumentare il rendimento dei forni industriali. I bruciatori rigenerativi offrono un rendimento termico superiore,
a fronte di una spesa maggiore in termini di cicli di accensione e di aspirazione dei fumi; i nuovi bruciatori con
recuperatore di calore integrato a canali suddivisi offrono in un sistema recuperativo lo stesso rendimento
termico dei bruciatori rigenerativi. Entrambi i modelli di bruciatore utilizzano la tecnologia della combustione
senza fiamma per la riduzione delle emissioni di NOx.
Keywords: Efficienza energetica - Bruciatori recuperativi - Bruciatori rigenerativi - Combustione senza fiamma Acciaieria - Trattamenti termici - Processi - Produzione
INTRODUZIONE
La riduzione delle perdite di calore al camino, spesso, rappresenta la via più efficace e conveniente per aumentare il
rendimento dei forni industriali. In questo articolo saranno
presentati due nuovi modelli di bruciatore che consentono
di ridurre di quasi la metà la perdita di calore rispetto ai
bruciatori recuperativi con scambiatore alettato. I bruciatori rigenerativi offrono un rendimento termico superiore,
a fronte di una spesa maggiore in termini di cicli di accensione e di aspirazione dei fumi che potrebbe non essere
giustificata nel caso di bruciatori di taglia ridotta o forni
di piccole dimensioni. I nuovi bruciatori con recuperatore
di calore integrato a canali suddivisi offrono in un sistema
recuperativo lo stesso rendimento termico dei bruciatori
rigenerativi. Entrambi i modelli di bruciatore utilizzano la
tecnologia della combustione senza fiamma (“flameless”)
per la riduzione delle emissioni di NOx.
RECUPERO DEL CALORE DEI FUMI
I bruciatori recuperativi basati su scambiatori scanalati o
alettati sono lo stato dell’arte nei forni di trattamento termico e si sono diffusi dopo le crisi energetiche degli anni
‘70 e dei primi anni ‘80. Complici i prezzi netti dell’energia
primaria, stagnanti o talvolta decrescenti, il rendimento
termico delle soluzioni costruttive adottate è stato finora
migliorato solo marginalmente o, in alcuni casi, addirittura diminuito a favore di semplificazioni costruttive (scamJoachim Georg Wünning, Ambrogio Milani
WS Wärmeprozesstechnik GmbH
Alessandro Lazzaretto
ATTAS S.R.L.
biatori lisci) oppure sottodimensionando la superficie
di scambio termico per una data potenzialità. Il mercato
chiedeva soprattutto aumenti di potenzialità specifica e
di temperatura massima mediante l’impiego di materiali
ceramici. In molti impianti in tutto il mondo, in particolare, erano ancora installati bruciatori ad aria fredda senza
recupero anche in processi ad alta temperatura. Oggi, per
via dei crescenti costi energetici e grazie ad una maggiore
sensibilità verso la tutela dell’ambiente, è generalmente
riconosciuta la necessità di aumentare gli sforzi per risparmiare energia ed abbattere le emissioni nocive.
Il preriscaldo dell’aria comburente in processi ad alta temperatura è la misura più diretta ed economica per ottenere
risparmi energetici cospicui, ma il potenziale tecnologico
disponibile non è stato ancora sfruttato a fondo nella stragrande maggioranza degli impianti di trattamento termico.
Il rendimento intrinseco di uno scambiatore di calore è
descritto dal numero di unità termiche (NTU in inglese):
questo parametro è proporzionale alla superficie specifica
di scambio e ad un coefficiente di scambio termico [1].
Quando si estrae attraverso il bruciatore il 100% dei fumi,
la capacità termica dell’aria comburente è inferiore a quella degli stessi: per questa ragione, nemmeno uno scambiatore ideale sarebbe in grado di trasferire all’aria tutta
l’energia termica contenuta nei fumi. Si può notare che la
temperatura dell’aria preriscaldata si avvicina a quella dei
prodotti della combustione estratti dal forno: il recupero
energetico potenziale mediante preriscaldamento dell’aria
si può considerare pienamente sfruttato quando la temperatura fumi allo scarico è ridotta fino a circa 200 °C.
Questo corrisponderebbe a una frazione residua di calore
nei fumi di circa il 10%, ossia ad un rendimento termico
di combustione di circa il 90% (riferito al potere calorifico
inferiore del combustibile). Quanto di questo potenziale
converrà sfruttare impiegando le tecniche costruttive di21
Memorie
Fig. 1 – Rendimenti degli scambiatori di calore di
funzione delle NTU
Fig. 2 – Bruciatore rigenerativo per tubi radianti
Fig. 1 – Heat recovery of different burner types
Fig. 2 – Regenerative burner REGEMAT®
sponibili dipende in larga misura dai prezzi dell’energia.
Nel caso di estrazione dei prodotti di combustione <100%
(come nel caso a fiamma libera, cioè a riscaldo diretto) c’è
qualche complicazione in più, ma sostanzialmente si ha a
che fare con rendimenti termici del tutto paragonabili a
patto di estrarre più dell’85% dei fumi prodotti.
La Figura 1 riporta il diagramma relativo al numero di
unità termiche (NTU) da 0 a 5, che è quella di maggior
interesse tecnologico, mostrando gli effetti dell’utilizzo di
uno scambiatore in contro-corrente per quanto riguarda
il pre-riscaldo dell’aria e la temperatura dei fumi con una
temperatura dei prodotti della combustione di 1000 °C.
Con aria fredda, NTU è uguale a 0, la temperatura di scarico dei fumi è quella di partenza e le perdite al camino
sono circa il 50%. Per un tipico bruciatore a recuperatore
scanalato o alettato il numero di unità termiche è vicino a
1: quindi per prodotti di combustione a 1000 °C, l’aria si
preriscalda a poco meno di 550 °C, pari a una perdita nei
fumi di 28% circa: il risparmio rispetto a un sistema ad aria
fredda rimane pur sempre oltre il 30%. Un miglioramento
sostanziale del preriscaldamento aria mediante recuperatori scanalati più grossi non è praticabile: sia quelli metallici in leghe resistenti al calore che quelli ceramici sono
limitati dal disegno e dagli alti costi, così che un raddoppio
della superficie di scambio, a parità di diametro, comporterebbe un raddoppio della lunghezza. Un aumento significativo dell’efficienza si può invece ottenere moltiplicando,
per così dire, le unità di scambio termico.
gia, poche misure furono introdotte nel mercato per aumentare l’efficienza energetica, già dagli anni ’80, invece,
sono stati compiuti molti progressi nella riduzione delle
emissioni di ossidi di azoto [2]. La normativa “TA Luft 86”
prescriveva un limite di 250 ppm, che costituiva un grosso handicap per molti modelli di bruciatore [3], anche se
inizialmente era consentito un bonus in funzione di temperature dell’aria molto alte. Con lo sviluppo di bruciatori ad alta velocità controllati in modo sequenziale questo
limite poteva essere rispettato. Infine, grazie all’introduzione di tecniche d’iniezione selettiva (“staging”) dell’aria
di combustione e di altre misure di riduzione degli NOx,
i valori sono stati ulteriormente ridotti [2]. L’ossidazione
senza fiamma (FLOX®, marchio registrato della WS Wärmeprozesstechnik GmbH) si basa sulla ricircolazione interna
di prodotti della combustione [1]. Bruciatori recuperativi
per tubi radianti funzionanti sul principio dell’ossidazione
senza fiamma sono stati installati in gran numero per la
prima volta nel 1994 su un forno di ricottura continuo per
acciaio al silicio [4]. Entrambe le realizzazioni costruttive
illustrate qui di seguito funzionano in ossidazione senza
fiamma, ciò che consente valori molto bassi di ossidi di
azoto, nonostante le alte temperature di preriscaldamento
dell’aria comburente.
EMISSIONI DI NOx
La problematica della formazione di ossidi di azoto è strettamente legata allo sviluppo di scambiatori di calore più
efficienti. Nel caso del gas naturale è rilevante, prima di
tutto, la formazione dell’ossido di azoto termico che, come
dice il nome stesso, dipende dalla temperatura di combustione: essa sale con l’aumento del preriscaldamento
dell’aria e ciò comporta emissioni intollerabili, se non si
adottano adeguate contromisure.
Mentre negli anni ’90, a causa dei bassi costi dell’ener22
BRUCIATORI RIGENERATIVI
Il principio dello scambiatore di calore rigenerativo è noto
da qualche tempo ed è stato impiegato nella fabbricazione dell’acciaio nel forno Martin-Siemens dalla metà del
diciannovesimo secolo. Ancora oggi la maggior parte dei
forni fusori per vetro sono equipaggiati con rigeneratori.
Negli anni ’80, la British Gas studiò in laboratorio la possibilità di sfruttare questa tecnica anche per bruciatori di
piccola potenza [5]. Un problema irrisolto era costituito
dalle emissioni di ossidi di azoto, che raggiungevano rapidamente valori di migliaia di ppm a causa della temperatura dell’aria comburente. Nuovi sviluppi di bruciatori
rigenerativi in coppia si ebbero soprattutto in Giappone,
U.S.A. e in Europa, nell’intervallo di potenza di alcuni MW,
Processi e impianti
Fig. 4 – Da sinistra a destra: recuperatore alettato, a
canali suddivisi, rigeneratore
Fig. 4 – Cross sections of finned and gap flow recuperator
and regenerator
Fig. 3 – Bruciatore rigenerativo in funzionamento
“senza fiamma”
Fig.. 3 – Regenerative burner in FLOX® mode
rispetto a un recuperatore alettato è quindi del 15-20% e
di oltre il 40% rispetto a bruciatori ad aria fredda. Con i
prezzi attuali dell’energia, il sovra costo rispetto alla soluzione con recuperatore classico si può ripagare in meno di
10.000 ore di lavoro, pari a circa un anno e mezzo in un
impianto in marcia continua.
La figura 2 mostra un bruciatore rigenerativo REGEMAT
M250 da 140 kW per tubi radianti e la figura 3 lo stesso
bruciatore in funzionamento senza fiamma.
BRUCIATORI CON RECUPERATORE A CANALI SUDDIVISI
Fig. 5 – Bruciatore con recuperatore a canali suddivisi
Fig. 5 – Gap flow recuperative burner
per applicazioni in grandi forni di riscaldo e di fusione per
alluminio.
A fine 1996 un forno di ricottura continua per acciai speciali fu ammodernato con bruciatori rigenerativi di nuova
concezione. Questi bruciatori da 200 kW non funzionano
in coppie: ogni bruciatore costituisce un’unità completa a
sé stante e funziona in combustione senza fiamma; in questo modo si sono raggiunti valori <50 ppm di NOx anche
con aria preriscaldata ad oltre 900 °C [6], [7]. Per riuscire
ad installare un bruciatore rigenerativo in un tubo radiante
senza modificare tutta la struttura al contorno, occorre ridisegnare in forma molto più compatta tutti i componenti
costruttivi: ciò è stato realizzato con rigeneratori a flussi
ottimizzati e con tempi di inversione molto brevi di 10 s [8].
Rispetto a un recuperatore alettato, la superficie di scambio termico è moltiplicata molte volte e la trasmissione del
calore viene, quindi, molto migliorata. Installato in un tubo
radiante a forma di doppia P, questo bruciatore raggiunge
rendimenti termici >80%. Con prodotti di combustione in
ingresso a oltre 1000°C, i fumi di scarico sono attorno a
300 °C, pari a una perdita al camino del 15%. Il guadagno
I bruciatori rigenerativi offrono il massimo potenziale per
il recupero di calore dai prodotti della combustione, ma
con i prezzi attuali dell’energia non sono economicamente
competitivi in tutte le condizioni, almeno per potenzialità
<100 kW. L’obiettivo dello sviluppo è un nuovo sistema recuperativo per potenze <100 kW capace di efficienze vicine
ai rigenerativi, ma senza l’onere delle valvole d’inversione
richieste da tali sistemi. Il raggiungimento di quest’obiettivo è stato reso possibile con una superficie di scambio
termico maggiore e con un migliore scambio termico.
Come si vede nello schema in Figura 4, il recuperatore
alettato è una geometria sostanzialmente bidimensionale,
mentre l’utilizzazione dello spazio costruttivo disponibile è
tridimensionale sia nel caso del recuperatore a canali suddivisi che in quello del rigeneratore. Lo scambiatore a canali
suddivisi moltiplica di molte volte la superficie di scambio,
ma presenta uguale ingombro esterno e uguali perdite di
carico sui lati aria e fumi del recuperatore alettato.
Inoltre, rispetto ai bruciatori rigenerativi, non occorre applicare un’aspirazione forzata sul lato fumi. Il bruciatore a
canali suddivisi funziona sul principio dell’ossidazione senza fiamma e presenta, perciò, emissioni minime. Le temperature dei fumi estratti con ingresso a 1000 °C sono circa
300 °C, ciò che corrisponde a una perdita nei fumi <20%
e un guadagno di 10-15% in termini di rendimento termico
rispetto al recuperatore alettato. Numerosi bruciatori di
questo tipo sono in funzione regolarmente dalla primavera
2009 in diverse zone di un forno di bonifica per viti.
La Figura 5 mostra un bruciatore da Rekumat S150 da
40 kW con recuperatore a canali suddivisi e la Figura 6
mostra il funzionamento dello stesso in ossidazione senza
fiamma.
23
Memorie
RIFERIMENTI BIBLIOGRAFICI
Fig. 6 – Bruciatore con recuperatore a canali suddivisi
in funzionamento senza fiamma
Fig. 6 - Gap flow recuperative burner in FLOX® mode
[1] Wünning J.G., Milani A. “Handbook of burner technology for industrial furnaces” – Edition Heat Processing
– Vulkan Verlag, 2009. ISBN 978-3-8027-2950-8
[2] Wünning J.A. “Ein neuer NOx-armer Rekuperatorbrenner”, GWI, Band 34 (1985), Heft 2/3
[3] Wünning J.G. “Rekuperatorbrenner für die direkte Beheizung von Industrieöfen” , GWI, Band 37 (1988), Heft
10
[4] Roth W., Telger K. “Betriebserfahrung beim Einsatz von
Brennern mit flammenloser Oxidation” GWI, Band 44
(1995), Heft 7/8
[5] Cornforth J.R. “Combustion engineering and gas utilization” – British Gas, 3rd Edition, 1992
[6] Wünning J.A., Wünning J.G.”Regenerative burner using
flameless oxidation” – International Gas Research
Conference, Cannes, 1995
[7] Milani A., Salamone G.V., Wünning J.G.” Advanced regenerative design cuts air pollution” – Advanced Steel,
1998-99
[8] Georgiew A., Wünning J.G., Bonnet U. “Regenerativbrenner für Doppel-P-Strahlheizrohre in einer Feuerverzinkungslinie” , GWI, Band 56 (2007), Heft 6
CONCLUSIONI
Il preriscaldo dell’aria di combustione rappresenta un
grande potenziale per aumentare il rendimento energetico
dei forni industriali. I nuovi bruciatori con recuperatore a
canali suddivisi garantiscono altissimi rendimenti e riduzione delle emissioni nocive e possono essere utilizzati
sia in forni nuovi che per l’ammodernamento di impianti
esistenti.
Reduction of waste gas losses in industrial furnaces
through the new gap flow and regenerative burners.
Keywords: Efficiency - Recuperative burners - Regenerative burners - Flameless oxidation.
The reduction of waste gas losses is often the most effective and economic way to increase the efficiency of industrial
furnaces. The report will present two new burner models, enabling to cut waste gas losses of finned recuperative
burners almost in half. The regenerative burner achieves highest efficiency but one has accept a certain expenditure
for cyclic switching and exhaust gas suction. This might not be justified for smaller burner sizes and furnaces. The
new gap flow recuperative burner reaches almost the same efficiencies with a recuperative system. Both burner
models use flameless oxidation technology for lowest NOx emissions.
24
Alluminio e leghe
Effect of Grain Refinement on Microstructure
and Wear Behavior of Cast Al-7Si Alloys
P.C.Meena, A. Sharma, S. Singh
The present work deals with the study of wear of Al-7Si alloys without and with the addition of grain refiners
such as, Al-5Ti-1.25C, Al-5Ti-0.8C and Al-5Ti-1B individually. Grain refiner additions have shown grain
refinement in the cast Al-7Si alloys. Microstructures showed morphological changes in a- Al phases from
dendrite to equiexed. These changes in structure showed improvement in wear resistance of Al-7Si alloys.For
this, pin-on-disc test has been performed on Al-7Si alloy samples with varying addition level of grain refiners.
The worn-out surfaces of the samples were characterized by SEM studies in order to understand the wear
mechanism of Al-7Si alloys against steel disc. Worn-out surfaces of Al-7Si alloys pins exhibited different surface
morphologies. Although it has been found that wear mechanism was same for both without and with grain
refined Al-7Si alloys. However untreated samples exhibits higher wear loss than that of grain refined samples. It
was noticed that with addition of grain refiner, the samples gave smaller debris particles.
This showed less wear loss during such conditions.
Keywords: Al-7Si alloy - Grain refinement - Solidification - Microstructure - Wear
Introduction
Most common applications of Al-Si alloys are components
like connecting rods, pistons, engine blocks, cylinder liners, air conditioner compressors, brake drums etc. These
cast components are subjected to wear loss during service.
Therefore, there is a need to improve tribological properties of
Al-Si alloys. The improvement in tribological properties of Al-Si
alloys depend on number of material-related properties like
shape, size and size distribution of the second phase particles in the matrix and microstructure. In addition to the above
factors operating conditions such as sliding speed, sliding
distance, temperature, load etc. also play role on tribological
properties of the materials. The high wear resistance is mainly
attributed to the presence of hard primary silicon particles
distributed in the matrix. Due to the presence of the hard
primary silicon phase, these alloys have serious machinability problems. In order to obtain the best machinability and
low wear rate, the size of silicon phase is to be controlled
P.C.Meena
Assistant training and Placement officer Govt.
R.C.KhaitanPolytechnic college Jaipur - India
[email protected]
Ashok Sharma
Professor
Met. &Mat. Engg. MNIT Jaipur, India
[email protected]
Surendra Singh
(Ex-) Professor
Dept. of Met. & Mat. Engg. IIT Roorkee, India
[email protected]
through melt treatment. Al-Si alloys can be strengthened by
adding small amount of Cu, Ni or Mg and the presence of
silicon also provides good casting properties. Various studies
have been reported on the dry sliding wear behaviour of the
cast Al-Si alloys. A.D.Sarkar and J.Clarke[1] investigated the dry
sliding wear behaviour of Al-Si alloys using pin-on-disc wear
and friction machine and concluded that Si composition
in aluminium alloy does not appear to be a dominant factor in the calculation of wear resistance. B.N.Pramila Bai and
S.K.Biswas [2] investigated the wear behaviour of Al-Si alloys
and concluded that wear rate of an alloy without silicon is
significantly higher than the binary modified alloys containing
silicon composition between 4% and 24%. Sarkar[3] studied
the wear of Al-Si alloys against hardened steel disc and gray
cast iron and reported that the wear rate of hyper-eutectic
alloys is more than the hypo-eutectic alloys. Shivanath et
al.[4] reported that wear resistance is good for the hyper
eutectic Al-Si alloys. Somi Reddy et al.[5] investigated the wear
and seizure behaviour of Al-Si alloy containing silicon composition upto 23wt% using pin-on-disc wear and friction testing machine under various loads. It was observed from the
results that wear and seizure resistance increases with the
addition of silicon to aluminium
A.D.Sarkar and J.Clarke[1] conducted experiments on Al-Si alloys and suggested that wear fragments are produced from
the transferred material and high Young’s modulus of the
hypereutectic cast alloys increases their propensity to wear.
C.Subramanian[6] studied the effect of sliding speed on wear
behavior of Al-Si alloys and reported that the wear rate decreases with increasing sliding speed upto a critical speed.
A.S.Anasyida et al.[7] has investigated the effect of element
additions on the dry sliding wear of Al-Si alloys .P.K.Rohitgi and
B.C.Pai[8] investigated the effect of microstructure and me25
Memorie
chanical properties on the seizure resistance of aluminium
alloys and concluded that seizure resistance of aluminium
can be improved by alloying with silicon and nickel. This is
due to the precipitation of hard particles in the matrix.
Ashok Sharma et al.[9] studied the worn out test pin surface topography, sub surface damage and debris by SEM.
The worn-out test pin surfaces for all the alloys showed
a multitude of distinct topographical features in the SEM
images. It was found that a number of wear processes,
such as delamination, adhesion and abrasion take part in
removal of metal as debris, and no single wear process
was responsible for metal removal from sliding surfaces.
D.K.Dwivedi et al. [10] reported that the addition of alloying element and inoculants to the Al-Si alloy reduces the
wear rate and increases the transition load. Taking into
account all the earlier findings a systematic study of wear
rate without and with addition of grain refiners individually
in Al-7Si alloys are carried out and reported in the present
paper.
EXPERIMENT PROCEDURE
Synthesis of Cast Alloy
Experimental alloy was prepared by careful melting and
dilution of master alloy of Al-20% Si in combination with Al
of 99.99% purity in the electric resistance furnace at 720oC
provided with ± 5oC accuracy by a digital temperature
controller. Necessary allowances for melting losses were
also taken into account in computation of charges. After
proper mixing, the molten alloy was cast in metallic moulds.
The nominal compositions of base alloys were Al-7%Si and
Al-7Si-0.45Mg. All the compositions have been expressed
in wt. %.The compositions of Al-7Si alloys studied in the
present investigation are shown in table 1. The LM25 alloy
was selected in addition to Al-7Si alloy to understand the
grain refining efficiency of Al-Ti-C and Al-Ti-B master alloys
in the presence of Mg.
Al-7Si alloys were grain refined with the addition of 0.2wt%,
0.6 wt%,1wt% and 1.4wt% of Al–5Ti-0.8C, Al–5Ti-1.25 C and
Al–5Ti-1B master alloys individually in the melt.
Metallography
Samples for microstructural studies were cut from ingot
castings. Specimens were polished by standard metallographic procedure using a series of emery papers from
1/0 to 4/0 grade and finally polished on sylvet cloth using
fine alumina powder. Polished samples were etched with
Keller’s reagent. Worn out pin surfaces and wear debris
generated during wear test were subjected to SEM study.
Tab. 1 - Chemical composition
of the Al-7Si alloys used in the
present study
26
Wear test
Wear test procedure
Experiments were carried out using at the constant load of
1 Kgf and constant sliding velocity of 1m/s, for a sliding
distance of 1800m for each test which was sufficient for
highlighting reliable difference among the specimens (Fig.
1). The load was applied by weights. Wear of the test specimens was determined by the weight loss. Tests were conducted at room temperature30±3ºC and humidity 55.The
wear volume was studied under the constant sliding distance, load, sliding velocity (sliding distance 1800m, load
velocity 1M/s; load 1N/m and 170 rpm) and varying alloy
composition. Dry sliding wear tests were conducted on 6
mm diameter and 25 mm long cylindrical wear pins. Wear
tests were carried out under given sliding conditions for all
sample as shown in Fig.1. The duration of each experiment
was 30 minutes. The experiment was performed for all
type of alloy (Al-7Si alloys) castings. The flat surfaces of
both the test pins and the steel disc were ground to a surface finish about 4 mm. Before each experiment the disc
and the test sample were thoroughly cleaned with alcohol
and acetone and subsequently dried with warm air to have
identical sliding conditions. After each experiment the
steel disc was reground to restore the original surface conditions and a new test pin was used for each experiment.
Wear was determined by weighing the test samples before
and after the test with the help of single pan balance.
An electronic single pan weighing balance (mettler) accurate up to 0.1 mg was used for weighing the samples
before and after the test. A new test pin was used for each
experiment.
Fig. 1 - Schematic representation of the pin-on-disc
wear monitor
Metal/alloy
Composition (wt %)
Si
Fe Sr Cu Mg Zn Mn Ti Others
Al
Al-7Si
6.9 0.16 Balance
LM25(Al-7Si-.45Mg) 7.0 0.20 - 0.20 0.40 0.10 0.10 0.19
Balance
Alluminio e leghe
RESULT AND DISCUSSION
Microstructure
Fig.2 (a-d) show microstructures of Al-7Si alloy for both
without and with addition of grain refiners. From Fig.2 (a)
it is observed that α-phase is showing coarse columnar
grain structure (grain size 590µm). However when 0.2wt%
of different grain refiners are added individually in Al-7Si
alloy , the α-phase became finer as shown in Fig.2 (b-d).
The dendritic α-phase has converted into equiaxed dendritic structure. The maximum effect of grain refinement
is observed in case of addition of 0.2wt% of Al-5Ti-0.8C
master alloy grain refiner where grain size reduced to 170
µm. Al-5Ti-1.25C observed minimum effect on α-phase
(grain size 210 µm) while Al-5Ti-1B master alloy showed
intermediate effect on α-phase (grain size190 µm). This
was attributed to the fact that Al-5Ti-0.8C grain refiner released both TiAl3+TiC particles in the melt, which showed
both nucleant and solute effect for grain refinement. The
nucluent particles (TiAl3+TiC) released from Al-5Ti-0.8C
grain refiner have shown more powerful effect than from
TiAl3+TiB2particles released from Al-5Ti-1B grain refiner.
While Al-5Ti-1.25C grain refiner released only TiC particles. Therefore in this case solute effect was missing,
hence Al-5Ti-1.25C showed minimum grain refinement effect. Fig.3a shows microstructure of Al-7Si-0.45Mg alloy.
In this case also the a-Al phase is having coarser dendrites
(grain size 500µm) as shown in Fig.3a for Al-7Si-0.45Mg
alloy. The presence of Mg might contribute to some extent the solute effect which in turn contributes to grain
refinement effect. Fig.3 (b-d) for 0.2wt% of addition of different grain refiners in Al-7Si-0.45Mg alloys are showing
similar trend as is seen in Fig.2 (b-d) for Al-7Si alloy. The
grain size with addition of 0.2 wt %. Al-5Ti-0.8C, Al-5Ti-1
.25C and Al-5Ti-1B master alloys are 155µm, 200µm and
170µm respectively. From above results it is understood
a)
b)
c)
d)
Fig. 2 - Microstructure of Al-7Si alloy without addition of grain refiner (a) at 0 min. and with addition of 0.2wt. %
of (b) Al-5Ti-1.25C (c) Al-5Ti-0.8C (d) Al-5Ti-1B grain refiners at 2 mins. of holding time.
27
Memorie
a)
b)
c)
d)
Fig. 3 - Microstructure of Al-7Si-0.45Mg alloys without addition of grain refiner at (a) 0 min. and with addition of
0.2wt% grain refiner of (b) Al-5Ti-1.25C (c) Al-5Ti-0.8C (d) Al-5Ti-1B at 2 mins. of holding time.
that the addition of grain refiners to Al-7Si alloys resulted
in the change in the shape of the α-grains from coarse columnar to fine quiaxed [11-14] .
The hardness of cast Al-7Si alloy is 59VHN. The hardness
of cast Al-7Si -0.45Mg alloy is 61VHN.With the addition of
grain refiners improvement in the hardness is about 10%.
Tensile strength of as cast Al-7Si alloy is 160 MPa and
yield strength is 80 MPa which increased after addition of
grain refiners to about 20%. In case of Al-7Si-0.45Mg alloy,
as cast tensile strength is 166 MPa and yield strength is
90MPa which also increased after addition of grain refiners to about 20%.
Wear rate vs wt % addition level of different grain refiners
Al-7Si and Al-7Si-0.45Mg alloys to without and with addition of grain refiners are subjected to wear studied. The
results are as given hereinafter;
28
(i)
For Al-7Si alloy
Figure 4 shows wear rate vs wt. % addition level of different grain refiners in Al-7Si alloy at holding time of 2
min. Lowest wear rate is observed in case of Al-7Si grain
refined with Al-5Ti-0.8C master alloy in comparison to Al5Ti-1.25C and Al-5Ti-1B master alloys at all addition level.
It was also observed that as the wt. % addition of grain refiners increased there is decrease in wear rate because of
decrease in grain size. The wear rate is lowest at 1.4wt.%
addition level, which correlates to the lowest grain size obtained under the present grain refining condition. The trend
of the graph remains same in all the three cases, wear rate
decreases with increasing amount of grain refiner.
(ii)
For Al-7Si-0.45Mg alloy
Figure 5 shows wear rate vs wt. % addition level of different grain refiners in Al-7Si-0.45Mg alloy at holding time of
2 min. Lowest wear rate is observed in the case where AlLa Metallurgia Italiana - n. 1/2015
Alluminio e leghe
Fig. 4 - The wear rate vs varying wt% addition of
different grain refiners in Al-7Si melt at 2 min. of
holding time
Fig. 5 - The wear rate vs varying wt% addition level for
different grain refiners in Al-7Si-0.45Mg melt at 2 min.
of holding time
7Si-0.45Mg is grain refined with Al-5Ti-0.8C master alloy
in comparison to Al-5Ti-1.25C and Al-5Ti-1B master alloys
at all addition level. It was observed that as the wt. % addition of grain refiner increased there is decrease in wear
rate because of decrease in grain size. The wear rate is
lowest at 1.4 wt% addition level, which correlates to the
lowest grain size obtained under the present grain refining
condition. The trend of the graph remains same in all the
three cases, wear rate decreases with increasing amount
of grain refiner.
has shown lesser number of worn out features in comparison to Fig.6 (a,b&e).
Figure 7(a-d) shows worn out pin surface morphologies of
Al-7Si-0.45Mg alloy without any grain refiner (Fig.7a) and
with addition of different grain refiner [Fig.7 (b-d)] . From
Fig. 7 (a) it is seen that surface layer is in a state of plastic
flow with scoring marks and furrows.
The presence of oxide particles, delaminate flake, and craters are also observed. The worn out pin surface is showing deep grooves and smooth strips which are caused by
hard metallic particles. Some oxide particles are also seen
on the pin surface. The white layered structure on the pin
surface indicates that temperature at friction surface was
very high. Fig.7 (b) shows worn out surface with the addition of 1.4 wt. % of Al-5Ti-1.25C grain refiner. Fig.7(c)
shows worn out pin surface with addition of 1.4 wt. % of
Al-5Ti-0.8C grain refiner. Fig.7 (d) shows worn out pin surface grain refined with 1.4 wt. % of Al-5Ti-1B master alloy.
Some craters have been seen in Fig. 7(a & b) which might
be due to detachment of surface layer from the pin surface. Ploughing features are less prominent in Fig. 7 (b) in
comparison to Figs. 7 (a). This clearly shows that with the
addition of 1.4 wt. % addition of Al-5Ti-1.25C grain refined
wear rate has reduced. Fig.7 (c) shows rough grooves in
Al-7Si-0.45Mg alloy wear pin grain refined with 1.4wt.% of
Al-5Ti-0.8C master alloy. Fig.7 (d) shows more numbers of
grooves, debris particles in comparison to Fig.7(c).
SEM Study
In order to study the different modes of wear taking place
during pin-on-disc test (at 1m/s speed and 1N load), worn
surfaces of Al-7Si and Al-7Si-0.45Mg alloy pins were studied by using scanning electron microscope. Figure 6 (a-e)
shows worn pin surface morphologies of Al-7Si alloy without any grain refiner (Fig.6a) and with addition of different
grain refiners [Fig.6(b-e)] . Fig. 6 (a) shows worn out surface of Al-7Si alloy without adding any grain refiner. The
surface layer is seen to be in a state of plastic flow and
showing scoring marks, craters and furrows. Figure 6 (b-e)
shows various features of worn pin surfaces of Al-7Si alloy
with the addition of 1.4 wt. % addition of different grain
refiners. These exhibit oxide film, lesser number of craters
and loose particles in comparison to Fig.6 (a). White layer
is more prominent in Fig.6 (a) which is indicating that due
to frictional heat, temperature at the interface was very
high.
From Fig. 6 (b) it is observed that scoring marks are less
prominent in comparison to SEM microstructure shown in
Fig.6 (a). The magnified view of Fig.6 (c) has been shown
in Fig.6 (d), where due to edge cracking particles are seen
in a process of detachment from the edges. Fig. 6 (e) has
shown worn out pin surface of Al-7Si alloy grain refined
with 1.4wt. % of Al-5Ti-1B grain refiner. The SEM image of
pin surface of Fig.6 (e) is comparable with Fig.6(c). From
SEM images [Fig.6 (c-e)] it is observed that scoring marks,
craters and detached laminates are not so prominent as in
the case of Fig.6 (a). This shows that the wear resistance
has improved with addition of grain refiners. Fig.6(c&d)
Debris
It is observed that during sliding a black powder gradually
generated from wear pin surface and is seen to be piled-up
on both the sides of the wear tracks of steel counter face.
Since sliding generates frictional heat, oxidation of fresh
mating surfaces might have taken place in the working
atmosphere. The black powder could be a mixture of oxides of alumina and other constituents with metallic debris
along with some fine oxide debris. Representative samples
of Al-7Si and Al-7Si-0.45Mg alloy without and with addition of grain refiner are taken for observation under SEM.
For comparison purpose only samples after grain refined
with 1.4wt. % of Al-5Ti-0.8C grain refiner for Al-7Si and Al29
Memorie
a)
b)
c)
d)
Fig. 6 - SEM photo micrographs of Al-7Si alloy wear
pins (a) without grain refiner and with 1.4 wt. % of
grain refiners (b) Al-5Ti-1.25C and (c) Al-5Ti-0.8C (d)
Al-5Ti-0.8C at high magnification (e) Al-5Ti-1B
7Si-0.45Mg alloy are taken and compared with respective
alloys when no grain refiner was added.
Figures 8(a) shows debris generated from as cast worn
out pin surfaces of Al-7Si alloy (Fig.8a) and Al-7Si-0.45Mg
alloy (Fig.8b). While Fig.8(c & d) show debris generated
from worn out pin surface of Al-7Si alloy (Fig.8c) and Al-7Si-0.45Mg alloy (Fig.8d) grained refined with 1.4wt.% of Al5Ti-0.8C grain refiner. Debris generated from pin surfaces
without adding any grain refiner is larger in comparison to
respective grain refined samples. Grain refinement has reduced the wear rate which is quite clear while comparing
the debris morphologies.
e)
DISCUSSION
type of grain refiners used. Grain refinement improved the
load bearing capacity of Al-7Si alloys. The yield strength of
the sample can be estimated according to the Hall-Petch
relationship [15],
σ s = σ0 + K.d-1/2 The results show that the wear rate decreases with the
decrease in the grain size of Al-7Si alloys at a load of 1
N, sliding distance of 1800 m and sliding velocity of 1 m
s −1 under dry sliding condition. The results suggest that
wear properties also dependent on the addition level and
where σ s is the yield strength of the sample, σ0 is the yield
strength of a single crystal, K is a constant value related
with crystal structure, and d is the size of a grain. According to this equation it is very clear that larger the diameter
of the grain, lower will be its yield strength. Therefore the
30
Alluminio e leghe
a)
b)
c)
d)
Fig. 7 - SEM photo micrographs of Al-7Si-0.45Mg alloy wear pins (a) without grain refiner and with 1.4 wt. % of
grain refiners (b) Al-5Ti-1.25C (c) Al-5Ti-0.8Cand (d) Al-5Ti-1B
maximum yield strength is observed when the grain size is
the smallest. This is so because the grain boundaries act as
effective obstacle for slip dislocation on initial deformation
stage, while the stress on heavy deformation stage is predominantly affected by the interaction and across of dislocation rather than the presence of grain boundary. The
behavior of slip dislocation with decreasing grain size was
interrupted by numbers of grain boundaries. Therefore, the
yield strength and elongation were improved with decreasing grain size. Decrease in grain size increases the grain
boundary area and results in improvement of strength. Improvement in strength and hardness resist wear loss during sliding. Therefore, grain refiner additions have shown
Improvement in wear resistance of the sliding samples.
Although it has been found that wear mechanism of Al-7Si
alloys is same under without grain refiner and with grain
refiner addition. However, without grain refinement Al-7Si
alloys exhibited greater wear loss than that of grain refined
Al-7Si alloys. The wear rate decreased with increase in the
weight percentage of grain refiners. This is attributed to
the fact that the nucleating particles were able to withstand thermal softening effects due to reduced grain size
and the formation of oxidative protective transfer layer.
The wear debris generated during sliding of pin against
steel disc is mainly flakes of oxide and particles. Wear
debris formed during sliding developed a tribolayer by a
process of transfer from one surface to another, which resulted in the consolidation of these plastically deformed
and oxidized particles into a hard, protective layer that
reduced the overall wear rate. It was reported [16] that
this transition in the wear regime was more significant at
elevated temperatures, where the wear scar developed a
very smooth tribolayer. The material could be transferred
back and forth several times during sliding and eventually
produce wear debris particles. It was suggested [17-18]
that the formation of these wear particles could be a direct
result of their work-hardenability. It was further suggested
that a critical transfer layer thickness existed for a given
sliding situation. At this critical value, wear debris particles
were thought to form by delamination at or near the interface between the transferred material and the base material. During dry sliding of Al-Si alloys against steel systems,
it was shown by Antoniou et al. [19] that a finely dispersed
amphorous iron oxide formed on the wearing surfaces.
This was due to the oxidation of a significant proportion
of the steel counter face. This phase helped to stabilize
31
Memorie
a)
b)
c)
d)
Fig. 8 - SEM photo micrographs of debris (a) Al-7Si alloy without grain refiner (b) Al-7Si alloy grain refined with
1.4 wt. % of Al-5Ti-0.8C master alloy (c) Al-7Si-0.45Mg alloy without grain refiner (d) Al-7Si-0.45Mg alloy grain
refined with 1.4 wt. % of Al-5Ti-0.8C master alloy
the tribolayer on the alloy surface by pinning dislocations.
The composition of these layers consisted of an intimate
(mechanical) mixture of materials derived from both sliding materials.
The worn surfaces of the specimens were characterized
in order to understand the wear behavior of Al-7Si alloys
against steel disc. Although it has been found that wear
mechanism of Al-7Si alloys is same for both untreated
and grain refined alloys. However, earlier exhibited higher
wear loss than that of grain-refined Al-7Si alloys. Similar
studies were made by Prasad Rao et.al [18] on the influence of grain refinement on the wear mechanism of Al-7Si
alloys. The grain refinement of Al by inoculating with Al-Ti
or Al-Ti-B grain refiners leads to decrease in the size and
aspect ratio of the grains.
Heat is also generated at the interface while sliding take
place due to friction. When two asperities are in contact
at a given instant and at a location temperature rises ΔT
at the interface. This is a function of a few independent
variables, like load, sliding velocity, real area of contact
type and addition level of grain refiners and the average
thermal conductivity of the two contacting bodies. In the
32
adhesive theory of wear, as suggested by Archard [20] ,
Burwell and Strang [21] and Bowden and Tabor [22] it is
assumed that in the absence of lubricant the asperities on
the opposing surfaces adhere strongly and form asperity
junctions. Subsequent separation of the surfaces occurs
either at the interface or inside the bulk of the weaker asperity, depending on the relative strengths of the junction
and the materials. When Al-7Si alloys were slide against
steel counterface, separation occurs in the asperities and
fragments of this material are transferred to the harder
surface. Subsequently, this transferred material may become detached and form loose wear fragments. According
to Rabinowicz [23] , lump removal is due to material separation over a weak section dissimilar to the original surface
of joint. The separated material is to be transferred on to
the other surface without formation of free fragments. The
adhesion wear is considered as a fatigue process. Abrasive wear shows presence of clean furrows cut by particles
causes grooving. During fatigue, in presence of surface or
subsurface cracks accompanied by pits and spalls causes
sharp and angular edges around pits.
In the delamination theory of wear, proposed by Suh [24]
Alluminio e leghe
for unlubricated sliding situations, wear particle formation
is explained in terms of deformation and fracture of material near the sliding surface. It is suggested that surface
and sub-surface deform plastically as a result of surface
traction imposed on the sliding surfaces. Consequently,
accumulation of incremental plastic strain leads to microcrack and micro-void formation at a critical depth below
the contact surface, where some fracture criteria can be
satisfied. Propagation of these subsurface cracks parallel
to the surface as a result of stress conditions in the material leads to loose flakes, long and thin sheet-like wear particles are generated. The process of crack formation and
propagation has been analyzed by Jahanmir and Suh [25] .
Using mechanics of contacts, they found that, under sliding situations, the criteria for generation and propagation
of cracks are satisfied at a critical depth below the contact
surface. This critical depth is dependent on the material
properties, the contact load and the friction coefficient.
From above discussion, it is clear that no single wear
mechanism is responsible for the production of material loss from the sliding surface. SEM studies of worn
pin surface for all the alloys show a multitude of distinct
topographical features. There can in some cases be a direct effect where material is detached by delamination.
Delamination wear producing craters is a prominent feature of mild wear regimes and depends upon sub-surface
failure. By repeated sliding, fatigue induced cracks initiate
at a finite subsurface depth and propagates to the surface
causing delamination. Various shapes and sizes of wear
debris were found as a result of dry sliding. The debris
includes the flakes, chips, and oxide powder. The debris
essentially consists of laminates produced by fracturing of
compacted material. It shows that during sliding of wear
pin, more than one process took part in generation of
debris. The generation of laminate debris may be a more
delayed event, as it would take several sliding interactions
for cracks to initiate and propagate. Morphological observations of worn pin surface indicate that delamination is
a main mechanism responsible for removal of material in
mild oxidative wear conditions. Failure by a delamination
process is clearly indicated by the shape of the debris particles.
It is observed that Al-7Si-0.45Mg alloy grain refined with
1.4wt. % of Al-5Ti-0.8C master alloy could resist adverse
conditions of wear better in comparison to Al-7Si alloy
under similar conditions. In case of Al-7Si alloy, the presence of Si strengthens the aluminium matrix. In case of
Al-7Si-0.45Mg alloy, in addition to the presence of Si, the
Mg has also additive role in improving the strength of the
matrix. Strengthening of matrix reduces wear rate. It has
been established that the addition of grain refiners in Al7Si alloys result in the change in the shape of the grains
from coarse columnar to fine equiaxed. Due to fine grains
wear resistance of the material improved and the lower
wear rates were obtained in such cases. Numerous potent
heterogeneous nucleation particles are added through the
master alloys which act as nucleants to Al-7Si alloy. Grain
refinement of Al-7Si alloys leads to decrease in the size
of the grains [26] . This in turn increases the grain boundary area and results in improvement in strength. The grain
boundary strengthening has shown improvement in the
wear resistance of Al-7Si alloys. Hence, it is understood
that the wear resistance of grain refined Al-7Si alloys increased with the decrease in grain size. Thus finer grain
size and equiaxed grain shape both are important parameters for better wear resistance in Al-7Si alloys.
CONCLUSIONS
From the present study following conclusion are drawn;
1. The wear resistance of grain refined, Al-7Si and Al-7Si-0.45Mg alloys increased with decrease in grain size
with respect to without grain refined Al-7Si and Al-7Si0.45Mg alloys. This was attributed to the fact that due
to finer grains, grain boundary strengthing took place.
2. SEM studies of worn pin surfaces for all the alloys
showed a multitude of distinct topographical features
such as debris, scoring marks and edge cracking. The
worn pin surface is simultaneously subjected to more
than one mode of metal removal in the form of debris.
3. Debris generated from pin surfaces without adding any
grain refiner was larger in comparison to respective
grain refined samples.
4. Finer grain size and equiaxed grain shape both are important parameters for better wear resistance in Al-7Si
alloys.
References
[1] A.D. Sarkar and J. Clark, Friction and wear of Al-Si
alloys, wear 61, 1980, pp. 157-167.
[2] B.N PramilaBai, and S.K.Biswas, Characterization
of dry sliding wear of Al-Si alloys, Wear, 120, 1987,
pp.180-184.
[3] A. D.Sarkar J Clarck, Wear, 75, 1982, pp71-85.
[4] R. Shivanath, et.al, Wear of aluminium-silicon alloys,
Br. Foundry man, 70, (1977), pp.349-356.
[5] Somi Reddy et. al, Mechanism of seizure of aluminiumsilicon alloys dry sliding against steel. Wear 181-183,
1995, pp658-667.
[6] C. Subramanian, Effect of sliding speed on the
unlubricated wear behaviour of Al-12.3%Si Alloy,
Wear, Vol. 151, 1991, pp.97-110.
[7] A. S. Anasyida, A. R. Daud and M. J. Ghazali, Dry
sliding wear behaviour of Al-4Si-4Mg alloys by addition
of cerium, IJMME, Vol. 4, 2009, pp.127-130.
[8] P. K. Rohitgi and B. C. Pai, Effect of Microstructure and
mechanical properties on seizure resistance of cast
aluminum alloys”, Wear 28 , 1974, pp.353-367.
[9] Ashok Sharma & T.V. Rajan, “Scanning electron
microscopic studies of worn out leaded aluminiumsilicon alloys surfaces”, Wear 174, (1994), pp. 217228.
[10]D. K. Dwivedi, Wear behavior of cast hypereutectic
33
Memorie
aluminium silicon alloys, Materials & Design, Vol. 27,
2006, pp.610-616.
[11]P.C.meena, Surendra Singh, and Ashok Sharma,
Role of heterogeneous nucleating particles on grain
refinement of cast aluminium alloys, Transactions of
57Th IFC, 2009,pp257-264.
[12]Divya kohli, Ajay Pareek,and Ashok Sharma, Some
critical aspects of solidification and grain refinement of
aluminium alloys, Indian Foundry Journal,Vol.56,No.7,
July2010, pp31-39.
[13]P.C.meena, Surendra Singh and Ashok Sharma,
Some basic concepts of solidification in aluminium
alloys and the role of grain refiners, Foundry, 2011,
pp73-83.
[14]P.C.meena, Surendra Singh and Ashok Sharma,
Factors influencing grain refining behavior of master
alloys and their study on fading and poisoning
phenomena in Al and its alloys, Indian Foundry Journal
Vol.58,No.-2 February 2012,pp23-31
[15]C.E. Carlton, P.J. Ferreira, What is behind the inverse
Hall–Petch effect in nanocrystalline materials?, Acta
Materialia 55, 2007 3749–3756.
[16]Q. Zhuxian, Y. Yaxin, Z. Mingjie, S.K. Grotheim and H.
Kvande, Preparation of Al-Ti-B master alloys by thermal
reduction and electrolysis. Of B2O3 and TiO2 in cryolite-
34
alumina melts, Aluminium, 64, 1988, pp.1254-1257.
[17]J.B. Andrews et.al, Influence of Si content on the wear
characteristics of hypo eutectic Al-Si alloys, In K.
ludema (Ed), Wear of metals,ASME,1985, pp.180-184.
[18]A.K.Prasada Rao, K.Das, B.S.Murty, et.al, Wear 264,
2008, pp638-647.
[19]R. Antoniou and D.W. Borland, Mild wear of Al-Si binary
alloys during unlubricated sliding, N.P. Suh, Mater. Sci.
Engng, 93, 1987, pp.57-72.
[20]J.F. Archard and W.Hirst, The wear of metals under
unlubricated conditions, Proc. R. Soc. (London), 236A,
1956, pp.397.
[21]J.T.Burwell and StrangC.D.Metallic, Wear, Pro.R.Soc.
(London), 212A, May 1953, pp. 470-477.
[22]F.P. Bowden and D. Tabor, The friction and lubrication
of solids”part2, oxford University press, 1954-1964.
[23]E. Rabinowicz, Friction and wear of material, New
York, Wiley, 1965.
[24]N.P. Suh, The delamination theory of wear, Wear, 25,
1973, pp111-124.
[25]S.Jahanmir and N.P. Suh, Mechanism of surface void
nucleation indelamination wear, wear, 1977, 44, pp.1738.
[26]M.R.Rahimipour, M.Momeni,T.Naseri, IJE TransactionVol,26,No.7,2013, pp701-706.
Acciaio
Vantaggi del riscaldamento ad induzione,
nuove possibilità di efficienza e flessibilità per
i laminatoi per prodotti lunghi
A. Lainati
Il presente articolo illustra alcune soluzioni di riscaldo ad induzione per laminatoi. Le tecnologie di riscaldo ad
induzione hanno diverse opportunità di applicazione nei laminatoi per prodotti lunghi. Le evoluzioni attuali dei
sistemi di controllo di potenza e dell’automazione di processo, integrate con sempre più accurati modelli termometallurgici, offrono vantaggi operativi certi in termini di qualità del prodotto laminato, di efficienza, flessibilità
e compattezza dell’impianto.
Keywords: Acciaio - Fatica per contatto a rotolamento - Laminazione - Tratt. termici Proc. termomeccanici - Modellazione - Controllo processi - Impianti e attrezzature - Simulazione numerica Energia - Processi - Tecnologie
INTRODUZIONE
Alcune soluzioni impiantistiche che prevedono l’uso
di sistemi di riscaldo ad induzione sono già di comune
applicazione, ma hanno comunque un potenziale di
ulteriore miglioramento, altre sono tuttora oggetto di
sviluppo prototipale e potranno essere presto introdotte
su scala industriale. Tra le più significative aree di
applicazione delle tecnologie di riscaldo ad induzione si
evidenziano l’ottimizzazione del ciclo di riscaldo delle
billette, il controllo accurato del profilo termico della barra
nel treno di laminazione e il riscaldo selettivo nelle aree
di trattamento e di finitura del prodotto. Questo articolo
illustra alcune delle soluzioni progettate e applicate da
Siemens in collaborazione con ABP GmbH.
DISCUSSIONE
Il mercato per prodotti lunghi è caratterizzato da una straordinaria varietà tecnica di dimensioni e forme, con una
produzione su scala mondiale raddoppiata negli ultimi 10
anni e soggetta a continue sfide tecnologiche riguardo ai
requisiti di qualità e ai costi di trasformazione. Nel futuro ci si aspetta un’ulteriore leggera crescita dei volumi di
Alberto Lainati
Head of Technology and Innovation
Siemens Vai Metals Technologies Srl
Marnate, Italia
produzione, con un sempre più marcato spostamento di
attenzione dai volumi di produzione ad obiettivi di qualità
ed efficienza operativa.
La catena di produzione dei prodotti lunghi è sempre più
esposta alle anomale fluttuazioni del costo del materiale
e dell’energia. Di conseguenza, per mantenere sufficienti margini operativi, i produttori devono necessariamente
investire in soluzioni volte ad ottimizzare i costi di trasformazione del prodotto.
Le sfide tecnologiche ed economiche nella produzione
dei prodotti lunghi riguardano l’ottimizzazione energetica,
l’impatto ambientale, il perfezionamento della qualità del
prodotto as-rolled, la resa metallica di trasformazione e
in generale la logistica di processo. Queste sfide portano
i produttori a ricercare soluzioni per linee di produzione
compatte, equipaggiate con gruppi di macchine di processo ultra-efficienti che operano come veri e propri “centri
di lavoro” gestiti da logiche di controllo intelligenti che
consentono ad esempio settaggi in auto-adattamento dinamico.
È evidente che le varie tipologie di impianto per prodotti
lunghi hanno contenuti tecnologici a vari livelli e garantiscono diversi margini operativi in funzione di molti fattori
locali e globali. Ci sono settori basati su tecnologie molto
consolidate ma che tuttavia possiedono ulteriori margini
di miglioramento nella performance e nell’efficienza operativa, settori con ampie opportunità di espansione tecnologica e altri invece con livelli di produzione saturi e ridotti
margini di intervento. In questo scenario di continua sfida
tecnico-economica, un significativo contributo può venire
dalle competenze e dalle applicazioni di aree industriali
contigue a quelle della laminazione. Il modello di innova35
Memorie
Fig. 1 – Diagramma lifecycle delle tecnologie
di impianto per prodotti
lunghi e campi di
applicazione della
tecnologia di riscaldo ad
induzione
Fig. 1 – Life-cycle curve
of Long Product Plant
Technologies and field
of application of inline Induction Heating
Technology
Fig. 2 – Tipico layout
d’impianto con possibili
localizzazioni delle
stazioni di riscaldo ad
induzione
Fig. 2 - Reference rolling
mill layout with possible
locations of in-line induction
heating stations
zione “connect&develop” apre nuove strade alla progressione tecnologica e all’incremento della performance operativa anche per le tipologie di laminatoi più tradizionali.
A questo proposito, la tecnologia di riscaldo ad induzione
(In-line Induction Heating Technology - IIHT) è un potente
strumento per l’ottimizzazione di alcuni aspetti tecnologici
ed economici in molte tipologie di impianti per prodotti
lunghi. La figura 1 dà evidenza di questo, indicando i settori di attuale applicazione del riscaldo ad induzione: la posizione e il numero di tali settori sul diagramma life-cycle
tecnologico evidenzia l’importanza e le grandi opportunità
esistenti.
Il diagramma life-cycle sopra riportato è un prospetto generale dove possibili varianti legate soprattutto all’area geografica e a fattori di mercato possono modificare qualche
posizione. È tuttavia evidente che la tecnologia IIHT può
essere impiegata in numerose applicazioni per diversi tipi
di impianto quale potente strumento per l’ottimizzazione di
36
aspetti chiave tecnologici ed economici.
I fattori che promuovono l’efficienza in un impianto devono essere identificati e quantificati lungo tutta la linea,
dall’ingresso del treno di laminazione al prodotto finito; la
gestione termica ed energetica del prodotto laminato e la
sua resa metallica sono tipici fattori di “inefficienza” su cui
cercare ottimizzazioni.
La tecnologia IIHT, oggi disponibile ed adattabile a quasi
tutti i tipi di forma, dimensione e grado di materiale, può
essere utilizzata in diversi punti di un laminatoio, sia di
nuova installazione che negli interventi di modernizzazione sulle linee esistenti:
- nell’area a caldo all’ingresso del treno per un riscaldo totale, integrativo o di equalizzazione delle billette, dei blumi
e dei pre-sagomati, incluse le applicazioni con alimentazione diretta, sia continua che semi-continua, dalla linea di
colata (es. WinLink® e Siroll EHS);
- nelle altre aree calde lungo il treno, ad esempio per la
Acciaio
Fig. 3 – Parametri
operativi di un laminatoio
e possibili vantaggi
dell’applicazione della
tecnologia ad induzione.
Fig. 3 - Efficiency rating
factors of a rolling mill and
possible gain by use of inline induction heating
Fig. 4 – Esempi di
installazione di unità di
riscaldo ad induzione
in laminatoi per acciai
speciali
Fig. 4 – Examples of
installation of in-line
Induction Heaters in special
rolling mills
regolazione dinamica della temperatura della barra, e della
sua equalizzazione tra testa e coda (es. saldatrice billette ERT, laminatoi per acciai di qualità e speciali con treno
continuo o semicontinuo) e per la compressione e l’ottimizzazione del processo di trasformazione (es. tempra di
solubilizzazione in linea, tecnologia idRHa+ per l’indurimento della rotaia);
- nell’area a freddo all’uscita del treno per la compressione
e l’ottimizzazione dei processi di finitura (es. taglio a freddo, cassoni di raffreddamento lento, rinvenimenti e distensioni tipo pack-annealing).
Alcune soluzioni con IIHT sono già applicate industrialmente, altre sono tuttora in fase concettuale o prototipale.
A titolo esemplificativo, la figura 2 mostra le tipiche aree
di processo di un laminatoio con la possibile localizzazione delle unità di riscaldo ad induzione. La figura 3 elenca,
invece, alcuni parametri operativi di riferimento di un laminatoio per prodotti lunghi in termini di potenza termica ed
elettrica, resa metallica e layout dell’impianto, con i relativi
margini di miglioramento tecnico ed economico ottenibili
grazie all’uso della tecnologia IIHT.
Si può notare che i miglioramenti ottenibili sono piuttosto
significativi e che conseguentemente possono offrire l’opportunità di un rapido ritorno dell’investimento.
Le figure 4 e 5 illustrano i layout schematici di alcuni laminatoi progettati da Siemens dove IIHT è stata utilizzata
con successo:
- impianto per barre e vergella per acciai speciali e super
leghe, integrato con una doppia stazione IIHT per il riscaldo primario delle billette e per il riscaldo di integrazione/
equalizzazione dopo il treno sbozzatore;
- impianto combinato per barre, sezioni e vergella per acciai speciali e inossidabili, equipaggiato con una stazione
IIHT per l’integrazione termica e l’equalizzazione del riscaldo dopo il treno sbozzatore;
- un impianto WinLink® con il treno di laminazione direttamente collegato alla linea di colata, con la stazione IIHT utilizzata per l’integrazione termica e di equalizzazione della
37
Memorie
Fig. 5 – Esempi
d’installazione di unità
di riscaldo ad induzione
in un laminatoio ad
alimentazione diretta e
in uno per trattamento
termico in linea
Fig. 5 - Colata continua
billetta alimentata in modalità continua o semi-continua;
- un laminatoio per rotaie dotato di tecnologia idRHa+ per
il trattamento di indurimento in linea delle teste, con una
speciale unità IIHT per l’integrazione termica selettiva della testa e del piede della rotaia.
Siemens e ABP GmbH, azienda fra i leader mondiali nelle
tecnologie di riscaldo ad induzione, hanno attivato una proficua collaborazione per proporre applicazioni IIHT su misura per ogni tipo di laminatoio, unendo le loro competenze
al fine di sfruttare le più aggiornate evoluzioni tecnologiche
dell’hardware elettrico, della modellazione termo-metallurgica e del controllo dinamico dell’erogazione di potenza.
La progettazione hardware di unità di riscaldo ad induzione e della loro logica di controllo della potenza sono fondamentali per garantire efficienza e affidabilità operativa. Tra
gli elementi unici che garantiscono performance tecnica
superiore e reale sostenibilità economica vi sono:
- convertitori di potenza modulari progettati con la tecnologia IGBT (insulated gate bipolar transistor) con alimentazioni di potenza individuali a frequenza adattabile (tra 1000
e 2000 Hz) per ogni bobina per la massima efficienza e
accuratezza di controllo. L’efficienza elettrica del sistema
di potenza raggiunge il 95% , per un’efficienza energetica
complessiva di circa 80%;
- l’elettronica di controllo a semiconduttore garantisce un
sistema on-off affidabile e ad erogazione modulata dalle
varie unità in funzione della richiesta di integrazione termica istantanea;
- i convertitori modulari base da 250 kW possono essere
assemblati scalarmente con una tecnica “plug-in” fino ad
una potenza di 2 MW in funzione degli specifici requisiti di
ogni impianto;
- il fattore di potenza cosφ è >0,95 in tutte le condizioni
operative, minimizzando così la perdita di potenza e i disturbi alla rete elettrica;
- il convertitore e gli induttori ha un proprio sistema di raf38
freddamento ad acqua a circolo chiuso;
- il design modulare riduce i tempi di installazione e permette di avere standardizzazione dei ricambi per tutte le
unità.
Alcune importanti applicazioni industriali integrate con i
benefici della tecnologia IIHT sono illustrate qui di seguito.
Impianto WinLink® ad alimentazione diretta: l’applicazione nel campo dei prodotti lunghi della tecnologia di
laminazione diretta delle billette dalla colata continua ha
origine dalla tecnologia ESP Flat Rolling e sta prendendo piede per impianti a capacità medio-bassa per acciai
comuni normalmente laminati da billette con dimensioni
100÷160 mm.
La tecnologia WinLink® è progettata per consentire l’alimentazione diretta di billette dalla colata con modalità
continua o semi-continua. La sua caratteristica distintiva
è la colata billette a doppio filo per garantire un’alimentazione costante del treno di laminazione con massima continuità operativa e produttività della linea di colata.
I fondamentali requisiti tecnici sono la compattezza
dell’area d’impianto fra colata e treno, l’uniformità del profilo termico sia per billetta continua che per billette sequenziali che alimentano il treno, anche in presenza di una
temporale variazione del flusso di colata, e infine la totale
affidabilità operativa per lunghi periodi di produzione.
La colata e il treno sono infatti concepiti come unità tecnologiche indipendenti, ognuna con la propria logica di
funzionamento e con possibilità di avere diverse velocità
operative.
L’utilizzo della tecnologia IIHT garantisce una connessione
efficiente delle operazioni delle due unità sequenziali: la
stazione di riscaldo ad induzione agisce come una sorta
di ”accumulatore termico” intelligente in grado di compensare gli eventuali anomali gradienti termici dovuti al
Acciaio
Fig. 6 - Unità di riscaldo ad induzione: design del convertitore ABP IGBT e unità modulari del sistema
Fig. 6 - Induction heater: design of ABP IGBT converter and modular arrangement of zone-control system
disallineamento dei ritmi produttivi fra colata e treno di
laminazione (figura 7).
Le unità di riscaldo ad induzione devono poter erogare in
modo adattabile dinamicamente densità energetiche variabili in funzione della richiesta, passando con rapidità da
modalità di funzionamento normali (i.e. DT da integrare
<100°C) a picchi di erogazione per gradienti di temperatura fino a 200°C. Le caratteristiche della tecnologia ABP
IIHT, integrata da un sistema di logica sofisticato per il
controllo continuo della temperatura, si combinano perfettamente con i requisiti tecnologici.
Una taglia tipica per un riscaldatore ad induzione per l’applicazione in impianto WinLink® è tra i 2 e i 4 MW (dai 20
ai 60 kWh/t a seconda del DT massimo da compensare)
per gestire incrementi di temperatura fra 30°C e 200°C.
Alcuni casi (per esempio di impianti in India) testimoniano
che la mancata installazione di una stazione di riscaldo
ad induzione tra la colata e il treno o l’uso di tecnologie
di riscaldo ad induzione non sufficientemente evolute,
mettono seriamente a rischio l’efficienza produttiva, annullando i potenziali risparmi che deriverebbero dall’uso
dell’alimentazione diretta.
Siroll EHS-Efficient Heating System: una gestione efficiente del flusso di materiale e del bilancio termico
nell’area di impianto tra la colata e il treno può portare
significativi risparmi, talvolta sufficienti a mantenere una
buona competitività economica anche nei più difficili settori produttivi, quali ad esempio quello del rebar.
Diverse soluzioni d’impianto sono state studiate e progressivamente applicate nel corso degli ultimi vent’anni; la
maggior parte di queste soluzioni sono relative al forno di
riscaldo (principalmente a gas naturale), le cui prestazioni
in termini di efficienza di combustione e di recupero del
calore sono state oggetto di importanti miglioramenti. La
carica a caldo delle billette è oggi una pratica industriale
consolidata e in grado di garantire il 40% di risparmio di
combustibile; anche la carica mista caldo-freddo è abbastanza diffusa per garantire, oltre al risparmio di combustibile, un ulteriore aumento di produttività fino al 15÷20%.
È importante evidenziare che, nonostante i progressi della
tecnologia di riscaldo a combustibili fossili, l’ossidazione
del metallo delle billette che permangono nel forno per
lungo tempo a temperature elevate è ancora uno dei magLa Metallurgia Italiana - n. 1/2015
giori fattori di inefficienza in un laminatoio, con un tasso di
perdita di peso che raggiunge e supera lo 0,5%.
La spinta per lo sviluppo di tecnologie di riscaldo a combustibile fossile è stata ed è tuttora basata sul comune
presupposto che il combustibile fossile è la più economica
fonte di energia che possa essere usata per il riscaldo billette. Il costo dell’energia elettrica e il livello di efficienza
delle tecnologie di riscaldo ad elettricità (quali il forno Joule, il forno a resistenza e il forno ad induzione) ne ha precluso fino ad oggi il loro utilizzo come opzioni alternative
di riscaldo.
Questo scenario relativo al costo dell’energia non è ovviamente valido in tutte le aree del mondo. La scarsità di gas
naturale in alcuni Paesi (quali l’India), l’uso diffuso di fonti
rinnovabili in altri (come la Germania e l’Europa sud-ovest)
e la crescente attenzione sull’impatto ambientale, mettono in discussione il paradigma per cui il combustibile fossile debba costituire lo standard di impiego nei laminatoi. Il
sito www.eex.com (EU energy trading) riporta per esempio
dei dati molto interessanti sui livelli di prezzo dell’energia
e sulla loro fluttuazione a seconda dei vari fattori. Per citarne uno, si mostra come il prezzo dell’energia elettrica
nella Germania del Nord varia intorno ad un valore medio
di circa 45÷50 €/MWh da un massimo di 90 €/MWh ad
un minimo di 1 €/MWh; si possono anche avere valori negativi del costo dell’energia elettrica, che significa che il
consumo viene sovvenzionato. Vi è anche la possibilità di
acquistare in anticipo lotti di fornitura di corrente elettrica
a 20÷30 €/MWh.
In parallelo, una straordinaria evoluzione progettuale nelle
tecnologie di erogazione e controllo della potenza elettrica
ha permesso di colmare i limiti precedentemente esistenti
che precludevano l’utilizzo efficiente di tecnologie di riscaldo ad induzione nei laminatoi.
All’interno di questo scenario, Siemens ha sviluppato un
nuovo concetto di processo, il Siroll EHS-Efficient Heating
System (figura 8). Questo sistema garantisce una gestione
termica flessibile e ottimizzata del materiale proveniente
dalla colata e diretto al treno di laminazione, grazie ad un
uso combinato di diverse fonti di energia: calore residuo
della colata, combustibile fossile ed elettricità.
Il sistema Siroll EHS si basa su una particolare soluzione
di layout che rende possibili dei percorsi di processo alternativi con l’uso combinato della laminazione diretta, del
39
Memorie
Fig. 7 – Impianto
WinLink® ed
integrazione di
temperature attraverso
il forno modulare ad
induzione
Fig. 7 - WinLink® plant and
temperature integration
by modular Induction
Heating station (“thermal
spring”concept)
riscaldo a combustibile fossile e delle tecnologie di riscaldo ad induzione.
La figura 8 mostra la più completa configurazione di un impianto EHS con i possibili percorsi di processo e le relative
temperature di riferimento.
Il sistema Siroll EHS è concepito con un percorso di processo preferenziale per l’alimentazione diretta dalla colata al treno seguita da un’ unità di riscaldo ad induzione
all’ingresso del treno; è poi previsto in parallelo un forno
di riscaldo a combustione che può essere utilizzato per
billette fredde, calde o miste, o come stazione di accumulo quando il treno funziona con laminazione diretta dalla
colata o quando la colata ha un eccesso di output rispetto
al laminatoio.
Questa configurazione combinata garantisce la sincronizzazione ottimale delle produttività di colata e laminatoio
che solitamente è sbilanciata di circa il 15% in termini di
ore di attività (es. 7200 h/y per la colata vs. 6100 h/y per
il treno).
Un ulteriore importante vantaggio della tecnologia EHS
consiste nel fatto che le billette vengono riscaldate nel
forno a combustione ad una temperatura inferiore al campo critico per l’ossidazione di superficie (820°÷850°C). Il
riscaldo finale della billetta alla temperatura di laminazione (1000°÷1100°C) è dato infatti dal forno ad induzione
con una conseguente breve esposizione alle temperature
critiche d’ossidazione. Questo permette di ottenere un aumento di resa metallica con un ritorno economico significativo.
40
L’uso delle più moderne tecnologie IIHT ha un ruolo fondamentale nel sistema impiantistico Siroll EHS, funzionando
anche in questo caso come un “accumulatore termico”
che può rilasciare quantità di energia termica integrativa
alla billetta in funzione dei diversi requisiti di produzione.
Le applicazioni EHS più comuni studiate fino ad oggi hanno una taglia di potenza tra 3 e 6 MW a seconda della
produttività e del livello di integrazione della colata e della
linea di laminazione, ma per produttività elevate e grandi
ΔT (es. billette di 130mm a 180 t/h con un’integrazione
termica di 350°C) si può avere una potenza dell’induttore
fino a 12-14 MW.
La figura 9 mostra i benefici della combinazione del riscaldo a combustibile fossile con il riscaldo ad induzione.
Un ulteriore elemento vincente dell’impianto EHS riguarda
la flessibilità di adattamento del percorso di riscaldo della
billetta e dell’impostazione delle temperature sia in funzione di condizione pianificabili (es. costo orario variabile
dell’elettricità vs. prezzo fisso del gas, fermata pianificata
del treno) che contingenti quali ad esempio una fermata
improvvisa del treno di laminazione.
L’impianto EHS deve ovviamente essere gestito da una logica di controllo evoluta basata sia su parametri tecnologici che economici, quali ad esempio il prezzo dell’elettricità
prenotabile per determinati periodi. I fattori fisici relativi al
processo di riscaldo – per esempio le formule di efficienza/consumo del forno a gas e del riscaldatore ad induzione, il modello di formazione di scaglia vs temperatura/
tempo, etc – sono tutti modellati e inseriti nella piattaforLa Metallurgia Italiana - n. 1/2015
Acciaio
Fig. 8 – Configurazione
di un impianto con Siroll
EHS: percorsi di processo
e target termici
Picture 8 - Siroll EHS plant
configuration: process
routes and temperature
targets
Fig. 9 – Siroll EHS:
benefici del sistema di
riscaldo ad induzione
Picture 9 - Siroll EHS:
benefits of Induction
Heating route
ma di automazione insieme ai parametri economici così da
avere un sistema di gestione intelligente sulla linea.
È inoltre disponibile un simulatore di impianto con relativa valutazione del ritorno di investimento per configurare
vari scenari e offrire così ai produttori una comprensione
immediata del potenziale dell’applicazione EHS sul loro
impianto.
La figura 10 mostra il modello base di EHS con i relativi
blocchi funzionali e punti di controllo; tale modello genera una simulazione 3D dinamica dei cicli di lavoro dell’impianto sulla base delle diverse combinazioni di parametri
imposti.
La configurazione EHS mostrata in figura 8 si applica tipicamente ad un impianto completo di nuova installazione.
Tuttavia, la flessibilità della soluzione permette di studiare
applicazioni speciali su impianti esistenti. Per esempio, recentemente, un importante produttore europeo, che non
possiede la carica a caldo delle billette ma che usufruisce di prezzi molto favorevoli dell’elettricità in determinati periodi, ha chiesto di studiare un’opzione impiantistica
speciale con anche l’aggiunta di un secondo forno ad induzione per riscaldare le billette fredde con successiva alimentazione in posizione intermedia nel forno a gas.
Altri studi di fattibilità in cui l’EHS ha dimostrato di essere
una soluzione efficace sono relativi a laminatoi con operatività intermittente o con grande varietà di prodotti e conseguenti frequenti fermate per cambi.
Siemens sta attualmente sviluppando la progettazione di
due laminatoi, uno in India e uno in Thailandia, dotati di
tecnologia Siroll EHS.
Tecnologia idRHa+ per l’indurimento della testa delle
rotaie: la progressiva evoluzione dei requisiti tecnici delle rotaie ha creato una forte richiesta di nuove tecnologie
produttive nel laminatoio. Le rotaie premium sono oggi
laminate con produttività fino a 200 t/h, con estrema accuratezza dimensionale, con lunghezze fino a 120 m, con
microstrutture perlitiche o bainitiche, con indurimento superficiale con durezze >350÷370 HB e con costi di trasformazione competitivi ridotti di almeno il 30%: queste sono
le sfide per i fornitori di impianti e la risposta è un nuovo
concetto di laminatoio per rotaie integrato da una moderna linea di trattamento termico.
Siemens ha sviluppato, in collaborazione con Centro Sviluppo Materiali Spa (attualmente Rina-CSM), il sistema di
trattamento termico idRHa+ - injection dual-phase Rail Har41
Memorie
Fig. 10 - Schema di
configurazione della
piattaforma di gestione
Siroll EHS
Fig. 10 - Layout of Siroll
EHS plant management
platform
Fig. 11 – Configurazioni
d’impianto con idRHa+
Fig. 11 - idRHa+ plant
configuration
dening - per l’indurimento in linea di rotaie fino a 120 m di
lunghezza. Si tratta di una tecnologia che include anche
le ultime evoluzioni nel campo del riscaldo selettivo attraverso l’induzione elettrica. Il trattamento termico idRHa+
avviene in linea su rotaie con alimentazione diretta dal treno o con una configurazione by-pass in caso di vincoli di
layout (figura 11).
Il processo in linea sfrutta il calore residuo di laminazione per garantire un’elevata produttività e significativi vantaggi economici se paragonato alla bassa produttività dei
sistemi fuori linea in cui la testa della rotaia deve essere
riscaldata nuovamente dalla temperatura ambiente. La
linea idRHa+ è composta da un gruppo di dispositivi sequenziali, mostrati in figura 12, ognuno con uno specifico
scopo funzionale.
L’apparato di trattamento termico è composto da moduli, il
cui numero e relativa posizione dipende dalla produttività
42
d’impianto e dai gradi, lunghezze e forme di rotaia da produrre. Ciascun modulo è dotato di una serie di dispositivi
di raffreddamento intercambiabili. La temperatura superficiale della rotaia in ingresso viene mantenuta tra i 750 e
i 850 °C. Il raffreddamento è regolabile fra 0.5 e 40°C/s
in funzione delle caratteristiche microstrutturali e meccaniche desiderate. La temperatura della rotaia all’uscita
dell’apparato va dai 300°C ai 600 °C a seconda del grado
trattato.
La strategia di processo (velocità di riscaldo, velocità di
raffreddamento, profilo termico) è predefinita in funzione
delle proprietà del prodotto finale. Il sistema di controllo di
processo è basato su diversi modelli termici, meccanici e
metallurgici.
Dopo che la rotaia laminata lascia il treno, il processo
idRHa+ applica azioni di riscaldo e raffreddamento selettive sequenziali con strategie dedicate per controllare la diLa Metallurgia Italiana - n. 1/2015
Acciaio
Fig. 12 – Unità funzionali
che compongono il
sistema idRHa+
Fig. 12 – idRHa+ layout
with functional units
Fig. 13 – Configurazione
di un riscaldatore ad
induzione per rotaie
Fig. 13 – Design of
induction heater for rail
stribuzione della temperatura lungo la sezione della rotaia.
L’accuratezza e l’uniformità del controllo termico in ogni
posizione sulla sezione della rotaia (testa, anima, piede)
e su tutta la sua lunghezza (tipicamente da 25 a 120 m) è
l’essenza del processo che permette di ottenere la caratteristiche metallurgiche e meccaniche desiderate, nonché
un profilo di rettilineità ottimale.
L’uso della tecnologia IIHT offre anche in questo caso un
contributo decisivo all’efficacia del processo integrando
e controllando la temperatura della rotaia prima che entri nei moduli di raffreddamento. Il riscaldo della rotaia è
portato a termine con una serie di unità di induzione ad
alta potenza con bobine superiori ed inferiori alimentate in
modo indipendente dai convertitori IBGT.
Una tipica configurazione d’impianto idRHa+ presenta due
moduli di induzione con quattro bobine indipendenti superiori ed inferiori alimentate con diverse potenze, in grado
di fornire in modo selettivo alla testa e al piede una densità
43
Memorie
Fig. 14 –
Composizione dei
riscaldatori ad
induzione per rotaie
e simulazione FEM
per un riscaldo
selettivo delle rotaie
Fig: 14 - Arrangement
of rail induction
heaters and FEM
thermal-simulation of
a rail selective heating
di potenza fino a 30÷35 kW/t; ciò significa che per una
produttività di 150 t/h, la potenza installata è di circa 5
MW (figura 13).
L’impostazione delle unità di induzione viene predefinita
dal modello di analisi FEM e poi regolata automaticamente
durante le operazioni a seconda dell’evoluzione della temperatura della rotaia.
Le posizione delle bobine è regolabile sull’asse verticale e
orizzontale per seguire la possibile non-rettilineità della rotaia e per adattarsi alle varie forme a grandezze delle rotaie
trattate, tra cui anche le rotaie asimmetriche (figura 14).
Laminatoio combinato per barre e sezioni per acciai
speciali (rif. figura 4b): questo tipo di impianto rappresenta una delle sfide maggiori per i fornitori d’impianti a causa
della grande varietà di prodotti ( oltre 700 combinazioni di
forme e dimensioni, tra cui tondi, esagonali, quadri, piatti,
angolari, ferri a U e a T, travi) e di requisiti specifici per i
vari gradi di acciaio inossidabile.
La dimensione della billetta va da 140 a 300 mm ed è laminata da uno sbozzo reversibile sliding per sagomare le
varie prese che alimentano il treno continuo. L’accuratezza
del controllo della temperatura è fondamentale per la qualità produttiva di determinate tipologie di acciai e richiede
l’installazione lungo il treno di un sistema IIHT attivo/selettivo per integrare termicamente il materiale laminato.
L’unità modulare da 6 MW con 8 set di bobine indipendenti
(figura 15) è un apparato molto efficace per garantire l’incremento di temperatura necessario, sia quello assoluto
per recuperare il calore perso nelle operazioni a monte, sia
quello relativo sulla lunghezza della barra.
Un’altra caratteristica importante risulta essere la flessibilità di adattamento della densità energetica e dell’altezza
della linea a seconda della dimensione del prodotto per
mantenere sempre un riscaldo uniforme e una linea di laminazione costante.
CONCLUSIONI
L’impiego di unità di riscaldo ad induzione di ultima generazione può offrire notevoli vantaggi operativi in molte
tipologie di impianto di laminazione per prodotti lunghi.
Sono state sviluppate soluzioni per quasi ogni tipo di for44
Fig. 15 – Stazione ad induzione da 6 MW con 8 moduli
Fig. 15 – 6 MW induction station with eight modules
ma, dimensione e materiale di prodotto laminato, adattabili ad una vasta gamma di impianti, sia nuovi che esistenti
e oggetto di modernizzazione.
I vantaggi derivano da un trasferimento selettivo e molto
efficiente dell’energia termica al prodotto, dall’erogazione
adattabile e dalla modularità dell’installazione. Allo stesso
modo, i benefici ottenibili per qualità del prodotto e per
produttività dell’impianto sono consistenti, insieme all’opportunità di sfruttare in modo flessibile ed economicamente conveniente varie fonti di energia.
Si può concludere che la compressione del processo di
produzione è indubbiamente uno dei fattori tecnologici
più rilevanti per l’industria dell’acciaio, con impegnativi
obiettivi di efficienza energetica, impatto ambientale e riduzione dei costi di trasformazione. Il mercato e le norme
nazionali e internazionali spingono, infatti, per un’ulteriore
miglioramento dei parametri tecnici ed economici di efficienza dell’impianto. In questo scenario, la tecnologia IIHT
rappresenta un contributo fondamentale per rispondere a
tali sfide.
Acciaio
Induction heating power, new paths to efficiency and
flexibility of rolling mills for long products
Keywords: Steel - Rolling Contact Fatigue – Rolling - Thermal Treatment - Thermo-Metallurgical Process - Modeling Process Control - Numerical Simulation - Energy - Processes - Technologies
Markets are pressuring for both cost reduction and quality improvement. Long product manufacturers are therefore increasingly enticed to shift from mass production objectives to efficient and quality production targets. In particular, the anomalous fluctuations of material and energy cost push producers to pursue the utter optimization of transformation costs.
In this regard, the In-line Induction Heating Technology (IIHT) demonstrates to be a powerful “tool” to optimize some key
technological and economical aspects in many types of rolling mill for long products.
Advanced thermo-metallurgical models are integrated into state-of-the-art power control systems and process automation, and offer operational advantages of quality, efficiency and flexibility.
The technical cooperation between Siemens and ABP GmbH merges their specific know-hows to exploit the potential of IIHT
application in the long rolling field. The latest technological developments in the electronic hardware design, in the thermometallurgical modeling and in the dynamic power supply control are available for the most effective customized applications.
Among the unique features which grant an efficient and consistent operation, a superior technical performance and a
fast return of investment of IIHT installation are:
• modular power converters with Insulated Gate Bipolar Design (IGBT) technology and individual power supply at
variable frequency 1000-2000 Hz), for maximum efficiency and control accuracy;
• semiconductor control electronics, for reliable dynamic on-off switching of each zone according to the instant
demand of power;
• modular 250 kW-converters with “plug-in” scalar assembly up to 2 MW, for easy and flexible adjustment to the
specific requirements;
• high power factor cosϕ in all operating conditions, for minimized power loss;
• modular structure, for reduced installation times and standardization of spare parts.
Whether in green or brown-field installations, several areas in the mill may benefit from the adoption of IIHT. The following are but a few of typical advanced applications of IIHT to the hot rolling of long products.
Bar and wire rod rolling mills for special steels and super-alloys may be equipped with a double IIHT station for the primary heating of billets and for the integrating/equalizing heating after the rougher train.
Likewise, an effective application may be in combination bar-section-wire rod mills for stainless and special steels. Often, these high-end product mills represent a challenge due to the large variety of rolled products, the usual limited size
of campaign with the corresponding necessity of fast changes to the mill, and the specific processing requirements of
the different steel grades. IIHT’s flexibility makes it possible to adapt the energy density and the line height in accordance to product, so to keep a uniform heating and a constant pass-line.
The high efficient heat transfer and the dynamical adaptability of the modular power supply system allow to keep the
thermal balance precisely controlled in a very compact arrangement.
With WinLink® direct rolling, a two-strand billet caster is directly linked to the rolling mill. WinLink® requires a compact
arrangement, a stable temperature profile along the endless billet, even during temporary flow variation, and a high consistency of operations during extended times. IIHT is applied for thermal integration of the billet coming from the caster.
It represents an effective solution, with its high-energy density capability as well its dynamically adjustable soft-heating
mode, integrated by a sophisticated temperature monitoring logic.
Efficient Heating System (EHS) grants an optimized and flexible thermal management of the billet between the caster
and the rolling mill by a combined use of multiple energy sources: residual heat from melting, fossil fuel and electricity. An appropriate layout arrangement makes it possible to select alternative processing routes of the billet, with the
combined use of direct rolling, fossil furnace and IIHT. This permits to promptly adapt the billet route according to the
fluctuations of energy cost (eg. electricity during the day or week). An important benefit is represented by the possibility
to limit the fossil heating in time and temperature (say 870°C), and then have the induction heater to quickly bring the
billet to the required rolling temperature (say 1050°C). This allows to limit the billet surface oxidation and reduce the
amount of scale, thus increasing the overall metallic yield.
EHS is assisted by a plant configurator tool which allows simulations to be run for different scenarios and layouts, and
the corresponding ROI to be assessed. The injection and dual-phase Rail Hardening (idRHa+), developed by Siemens
and RINA-CSM, is the cutting-edge technology for in-line thermo-processing of rolled rails. The spreading construction
of heavy-load and high-speed railways call for an improved rail resistance to wear and rolling contact fatigue. Premium
rails need to be rolled with a precisely controlled microstructure and a high dimensional accuracy, in order to meet the
required mechanical characteristics of hardness, UTS and elongation.
Unlike the off-line systems, idRHa+ is capable to operate in-line immediately after rolling, allowing for sensible heat to
be utilized and high hourly productivities to be achieved. A dedicated IIHT unit integrates and selectively controls the
temperature of the rail arriving from the finishing mill. It is followed by several cooling modules, equipped with a set of
interchangeable devices (spraying nozzles with mist-atomizers or air-jet blades). The whole process is precisely monitored. The control system is assisted by a suite of embedded thermal, mechanical and metallurgical models, so that the
desired microstructure and hardness values across the rail cross section may be preset, obtained and verified.
45
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Colata continua
High-precision numerical simulation for effect
of casting speed on solidification of 40Cr
during continuous billet casting
Y. Chen, Z. Peng, L. Wu, L. Zhao, M. Wang, Y. Bao
In order to study the effects of casting speed on the solidification process and optimize the process to reduce
defects, a high-precision simulation model of solidification and heat transfer of square billet based on nailing test and
temperature measurement is presented. The experiments have proven that values calculated by the model fit well the
measured values and the relative error is maintained less than 2%. Some new rules about the relationships among
casting speed, solidification end point and shell thickness are proposed in the paper for the first time.
Keywords: 40Cr - Continuous casting - Solidification - Simulation model
Model Description and Formulation
Heat Transfer Equation
To simplify the calculations, following assumptions are
made based on the actual situation [2]: (1) Billet heat transfer process is simplified two-dimensional unsteady heat
transfer. This is quite commonly adopted considering the
length of the billet is very large compared to its cross section size. (2) Latent heat released during solidification is
dealt with by the way of equivalent heat, since carbon steels usually have large crystallization temperature range.
(3) Effect of convective motion in the liquid hole on heat
transfer is dealt equivalently by increasing thermal conductivity. (4) Steel thermal physical parameters are only
related to temperature, is not related to spatial location.
(5) Effect of mold oscillation on the solidification process
is ignored, for it has no significant thermal effect. Based
on the assumptions, differential equation of billet solidification is derived:
(1)
where r is the density, kg . m-3; cp is the specific heat,
J . (kg.°C) -1; T is the temperature, °C; t is the time, s; l is
thermal conductivity, W. (m.c°C) -1.
Initial condition:
T(x, y, 0) = Tc
Boundary Conditions [3,4]:
(1) Billet center:
Chen Ya-nan, Peng Zun, Zhao Li-hua, Bao Yan-ping
State Key Laboratory of Advanced Metallurgy-University
of Science and Technology Beijing, Beijing 100083, China
WANG Min
National Engineering Research Center of Flat Rolling
Equipment, University of Science and Technology Beijing,
Beijing 100083,China
WU Lin
Central Iron & Steel Research Institute,
Beijing 100081,China
Corresponding author e-mail: [email protected]
(2)
(3)
(2) Billet surface:
(4)
Crystallize:
(5)
Air zone:
(6)
Secondary cooling zone:
(7)
47
Memorie
where qs is surface heat flux, W.m-2; A, B is constant; e
is radiation coefficient, 0,8; s is Boltzmann constant,
5.67x10-8 W.(m2.K4)-1, Ts, Tw, Ta are the slab surface temperature, cooling water temperature and ambient temperature respectively, °C; h is heat transfer coefficient in secondary cooling zone W.(m2.°C)-1.
The heat transfer coefficient in secondary cooling zone is
one of the most important parameters to describe the heat
transfer effect in secondary cooling zone. The relationship
of heat transfer coefficient and the water flow density is
shown in Table 1. The form of these equations is basically
the same as in the references[5~8]. However, parameters
were modified to match the results of both the nail-shooting experiments and the measured surface temperature
listed in part 2. Considering that the properties of steel
and the calculation method were fixed, the boundary condition was the most influential factor. Therefore, modification of these equations should be acceptable according to
measured shell thickness and surface temperature.
Section
150 mm x 150 mm
Foot roller segment
hW = 0.32 x 0.341
First segment in
secondary cooling zone
hW = 0.3 x 0.321
Second segment in
hW = 0.28 x 0.312
Third segment in
hW = 0.26 x 0.292
Tab. 1 - The relationship of Heat transfer coefficient
and the flow-density
The real cooling behavior in secondary cooling zone is that
the billet surface temperature decrease under the spray
zone, while billet surface temperature increase outside
the region, considering the number of nozzles and spray
cooling range [9]. As a result, the slab surface outside the
spray zone is dealt by radiation heat transfer as the formula (6), while water flow density of slab surface in the
spray zone is calculated by formula (8), combining nozzle
parameters shown in Table 2.
(8)
(9)
Where W is water flow density, L.(m2.s)-1; Qw is water volume, m3.h-1; Ni is numbers of nozzles; dw is the diameter of
spray, m; dbs is distance between the nozzle and the surface of the slab, m; W is nozzle angle.
Model Parameters
Chemical composition of 40Cr steel is shown in Table 3.
The temperature of solid-liquid phase line is calculated by
empirical formula. The liquidus temperature is 1495 °C
and the solidus temperature is 1433 °C. The solid phase
48
Distance
Rows
Flat /
Numbers
nozzle from
Segment
of
Arc
per row
angle surface/
nozzles
surface
mm
Foot
2
2
Arc
60
115
roller
2
2
Flat
60
115
First
6
1
Arc
60
104
6
1
Flat
60
104
Second
7
1
Arc
60
106
7
1
Flat
60
106
Third
6
1
Arc
60
106
6
1
Flat
60
106
Diameter
of
spray/
mm
132.8
132.8
120.1
120.1
122.4
122.4
122.4
122.4
Tab. 2 - Nozzles parameters
density is taken as 7434kg·m-3; the liquid phase density
is taken as 6847 kg·m-3; the two-phase region density is
taken as 7134 kg·m-3. Latent heat of solidification is dealt
by the equivalent heat capacity, and 40Cr latent heat is
taken as 272 kJ · kg-1, compared with specific heat of solid
phase 647 J·(kg·°C)-1 and specific heat of liquid phase 828
J·(kg·°C)-1. The thermal conductivity of solid phase is 30 W
· (m ·°C) -1, while the thermal conductivity of liquid phase
is larger 1 to 4 times than solid phase, and two-phase region is the average value between solid phase and liquid
phase. Basic parameters of caster: casting machine radius
8m, casting machine metallurgical length 26m, crystallizer
effective length 0.8m, foot roller length 0.363m, length of
first segment in secondary cooling zone 0.989m, length of
second segment in secondary cooling zone 1.695m, length of third segment in secondary cooling zone 2.465m.
Element C
Si Mn
P
S
Cu Ni Cr
Content 0.405 0.27 0.65 0.0225 0.0225 0.18 0.27 0.95
Tab. 3 - Chemical composition (in mass pct) of 40Cr
Model Solution and Verification
Combined with model and parameters above, temperature field and solidified shell thickness can be obtained by
the calculation program, which is developed by VB.net.
The model is verified by nailing test and temperature measurement [10]. Nail shooting is one of the best ways to
measure the thickness of the steel solidified shell, is easy
to operate, and is reliable [11]. Sulphur carried in the steel
nail is used as a tracer agent in nail shooting to measure
the solidified shell thickness of the billet. During casting
the steel nail was shot into the billet perpendicular to the
surface by a nail shooting gun (as shown in Fig. 1). The
nails had sufficient momentum to penetrate the solid shell
and the tips would melt in the liquid steel. In this way, the
sulphur element was quickly distributed into the molten
steel. According to the morphology of the steel nail, the
thickness of the solidified shell can be measured. The steel
nails used were made of steel 60Si2MnA and the chemical
composition is shown in Table 4. The calculated melting
point of the steel nails is 1460˜1490°C.
Colata continua
Element
C
Si
Mn
Cr Other
Content 0.56˜0.64 1.60˜2.00 0.60˜0.90 ≤0.35 ——
Tab. 4 - Chemical composition (in mass pct) of steel
nails
Fig. 1 - Nail shooting gun
Fig. 3 - Temperature field and solid-liquid phase
partition of longitudinal section with different casting
speed. (a) 1.9m/min; (b)2.1m/min; (c) 2.3m/min
Fig. 2 - Photo of etched pin-shooting samples with
different casting speed.
(a) 2.10m/min; (b) 2.40m/min; (c) 2.47m/min
Pouring
steel grade Temperature
/°C
40Cr
Casting
speed
/m﹒min-1
1521
1521
1520
1518
2.01
2.10
2.4
2.47
Flow rate Flow rate Flow Flow rate Tab. 5 - Validation
tests conditions
Flow rate of
of foot
of 1st
rate of
of 3rd
3
nd
mold /(m /h) zone /
zone / 2 zone zone /
(m3/h)
(m3/h) /(m3/h) (m3/h)
122.2
5.19
5.56
4.81
3.00
122.3
5.36
5.85
5.05
3.16
122.3
6.22
6.65
5.76
3.55
122.5
6.4
6.84
5.93
3.66
In this study, the rivet nail-shooting position was the middle point of right surface 10.34m away from meniscus.
Meanwhile, surface temperature at the same point was
also measured by infrared thermometer for over 100s and
recorded every 0.5s. The validation tests conditions are
shown in Table 5. After the nail-shooting sample was lathed, etched in hot hydrochloric acid for 30min and photographed, shell thicknesses were measured by pixels in pictures, as shown in Fig. 2. Model predictions are listed and
compared with experiment results in Table 6. The results
show that values calculated by the model fit well the meaLa Metallurgia Italiana - n. 1/2015
Fig. 4 - Effect of casting speed on temperature of billet
sured values and the relative error is maintained less than
2%. Thus, the model was proved to be highly accurate.
Model results and Discussion
Effects of casting speed on the temperature field and solidification process
Temperature field and solid-liquid phase partition of longitudinal section are shown in Fig. 3 when the casting speed was taken 1.9, 2.1, 2.3 m/min, separately [12]. Fig. 3
shows visibly that with the casting speed increases, the
isotherms moves to the casting direction, liquid phase re49
Memorie
Fig. 6 - Shell thickness at the outlet of mould with
different casting speed
Fig. 8 - Effect of superheat on temperature of billet
Casting
speed/
m∙min-1
Measured
Calculated
results/mm results/mm
Error,%
2.10
56.4
55.7
1.24
2.4
51.0
50.2
1.57
2.47
44.8
43.9
2.01
Tab. 6 - Comparison between measured and calculated
shell thickness
Casting
speed/
m∙min-1
Measured
results/°C
Calculated
results/°C
Error,%
2.01
1090
1100
0.92%
2.10
1100
1115
1.36%
Tab. 7 - Comparison between measured and calculated
surface temperature
50
Fig. 7 - Solidification end point with different casting
speed
Fig. 9 - Effect of superheat on thickness of shell
gion and two-phase region expanded, liquid core length
increases significantly. Effects of casting speed on temperature of billet and thickness of shell are shown in Fig.
4 and Fig. 5. As shown in the figures, when casting speed
increase by 0.1m/min, middle temperature of the surface
at the outlet of third segment in secondary cooling zone
increase 13.75°C, the center temperature in the crystallization zone and the secondary cooling zone almost has
no change, but significantly increases in the air region,
solidification end point lengthens 0.92m or so. It can be
seen from Fig. 6 and Fig. 7 that the solidification end point
meets approximately linear relationship with casting speed. The shell thickness at the outlet of mould doesn’t meet
approximately linear relationship with casting speed, and
the shell thickness will reduce sharply with the increase of
casting speed from 2.1m/min to 2.2 m/min.
Effects of superheat on the temperature field and solidification process
Effects of superheat on temperature of billet and thickness of shell are shown in Fig. 8 and Fig. 9. As shown in
the figures, superheat almost has no effect on slab surface temperature, but has a certain impact on center temperature at the early and late period of the solidification
Colata continua
Original Method
Improved Method
Measured Predicted Measured Predicted
value/°C value/°C value/°C value/°C
Corner
temperature
Middle surface
temperature
960
940
979
953
1090
1100
1098
1103
Tab. 8 - Comparison between original method and
improved method
sting speed, and the shell thickness will reduce sharply
with the increase of casting speed from 2.1m/min to
2.2 m/min.
(3)Superheat has some effects on the solidification process to some extent. When superheat increase 10 ℃,
shell thickness at the outlet of mold decrease 0.44 mm
by average, solidification end point extended about
0.20m.
ACKNOWLEDGMENT
process. Superheat has little effect on the growth of shell
thickness, but has effect on shell thickness at the outlet
of the mold and solidification end point to some extent.
When superheat increase 10 °C, shell thickness at the
outlet of mold decrease 0.44 mm by average, solidification
end point extended about 0.20m.
Application in the Plant Operation
The model has been applied to plant operation at steelmaking plant of Hangzhou Iron and Steel Group to predict the
temperature variation and solidification process of slab.
In one case, the temperature needs to be improved on a
small scale in order to decrease some defects, such as
corner cracks and so on. To realize temperature improved
on a small scale, the model was used to investigate the
effects of operation parameters on slab temperature. The
simulated results show that corner temperature can be
improved by 13°C (Table. 8) if the flow rates in secondary cooling zone decrease by 10% (Improved Method). The
practical production result shows that corner temperature
improved about 19°C when the flow rate was by 10% at
cast speed of 2.01m/min。Experimental results were similar with the predicted results, so the model can be applied
to plant operation. When the improved trial method finally
turned to regular process in winter, the flow rate in the
third segment was further reduce by 20%. Combined this
optimized cooling strategy and more strict Nitrogen control measures in RF process, severe batch corner cracking
problems seldom happened again in winter.
CONCLUSIONS
(1)The experiments of temperature and shell thickness
measurements have proven that values calculated by
the model fit well the measured values and the relative
error is maintained less than 2%.
(2)Casting speed is the most important factor in determining the temperature field and solidification process of
40Cr during Continuous Billet Casting. With the increase of casting speed by 0.1 m/min, the solidification
end point extended about 0.92 m. The solidification
end point meets approximately linear relationship with
casting speed. The shell thickness at the outlet of mold
doesn’t meet approximately linear relationship with caLa Metallurgia Italiana - n. 1/2015
This research was supported financially by the National
Natural Science Foundation of China (No.51274029),
Doctoral Fund of Ministry of Education of China
(No.20130006110023) and State Key Laboratory of Advanced Metallurgy Foundation (No.41602014).
REFERENCES
[1] Richard A, Kai L, Atul K, et al. A Transient Simulation and
Dynamic Spray Cooling Control Model for Continuous Steel
Casting. J Metall Mater Trans. B, 2003, 34B:297
[2] Ya M, Brian G T, Heat-Transfer and solidification model of
continuous slab casting: CON1D. J Metall Mater Trans B,
2003, 34B:685
[3] Park H S, Nam H, Yoon J K. Numerical analysis of fluid flow
and heat transfer in the parallel type mold of a thin slab
caster. J ISIJ Int,2001, 41(9): 974
[4] Yang H L, Zhao L G, Zhang X Z, et al. Mathematical simulation on coupled flow, heat, and solute transport in
slab continuous casting process. J Metall Mater Trans
B,1998,29B:1345
[5] Laitinen E, Neittaanmaki P. On numerical simulation of the
continuous casting process. J J Eng Math.1988,22:335
[6] Yang H L, Zhao L G, Zhang X Z, et al. Mathematical simulation on coupled flow, heat, and solute transport in slab. J
Metall Mater Trans. B,1998,29B:1345
[7] Yoon J K.A fully-coupled analysis of fluid flow, heat transfer,
solidification and deformation behavior in continuously
cast beam blank. China-Korea Joint: Symposium on Advanced Steel Technology for Future Industry, Beijing, China,
1999:34.
[8] Kondo O, Hamada K. New dynamic spray control system
for secondary cooling zone of continuous casting machine.
1993 Steelmaking Conference Proceedings,1993, 309.
[9] López A R, López R A, Pardavé M P. Simulation of heat transfer in steel billets during continuous casting. J INT J MIN
MET MATER.2010,17(4):403
[10] Liu J J, Liu J H, Wu H J, et al. Numerical simulation of solidification process for continuous casting large size steel square
billet containing 0.45%C. J Foundry.Technol,2011,32(2):259
[11] Long M J, Chen D F, Wang Q X, et al. Determination of CC
slab solidification using nail shooting technique. Ironmak &
Steelmak, 2012,39(5):370
[12] López A R, López R A, Bello A K, et al. Simulation factors of steel continuous casting. J INT J MIN MET
MATER.2010,17(3):267
51
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Saldatura
Damage investigation on welded tubes
of a reforming furnace
E. Guglielmino, R. Pino, C. Servetto, A. Sili
In this work the creep damage of radiant tubes of a reforming furnace has been investigated. The considered furnace
contains a battery of tubes constructed by butt welding three spun cast pieces, made of ASTM 608 HP-Nb alloy.
They are designed to operate at temperatures of about 900°C, pressures of about 30 bars and times of the order of
100000 h. Tubes were inspected during the plant stops scheduled every two years, in order to identify and replace
the damaged ones with the aim to ensure conditions of safe operation in the furnace. They were selected though a
criterion based on measures of the internal diameter deformation performed in situ by Laser Optic Tube Inspection
System (LOTIS). For a verification of this method, optical and scanning electron microscopy observation, Vickers
microharndess and creep tests have been carried out on samples taken from tubes put out of service.
Keywords: Superalloys - Welding - Non-destructive testing - Mechanical testing Metallography - Electron microscopy
INTRODUCTION
During service at high temperature and for long time, metallic materials undergo to different levels of creep phenomena, which lead first to slow deformations and finally
to crack nucleation and propagation. These conditions are
particularly dangerous in the case of piping and pressure
devices. In order to verify every compliance with safety
conditions and working reliability, the operational practice
is based on a series of control activities. Investigations
on life conditions of pressure vessels are addressed both
to identify premature failures and avoid, even if their safe
working can still be fully guaranteed, to put them out service for attainment of the design life, previously evaluated
on the basis of conservative criteria.
In this work we investigate the radiant tubes of a furnace
for the hydrogen production from methane and water vapor by means of the endothermic reforming reaction that
take place thanks to a granular Ni-based catalyst.
The furnace here considered is a chamber, coated inside
by refractory material, containing 4 parallel rows of 44
vertically hung radiant tubes (Fig. 1). These tubes are constructed by butt welding three spun cast pieces, made of a
Ni-Cr-Fe alloy. Radiant tubes, heated from the outside by
E. Guglielmino, A. Sili
Dipartimento di Ingegneria Elettronica, Chimica e
Ingegneria Industriale – Università di Messina
R. Pino
Raffineria di Milazzo S.C.p.A
C. Servetto
IIS Service Srl, Genova
means of a system of gas burners, operate at temperature
of about 900°C and pressure of 33 bar, for times of the
order of 100000 hours.
Service temperature and tensile circumferential stress,
due to the internal pressure, give rise to a severe creep
conditions that led to diameter expansion [1]; inside tubes
there are carburizing conditions, however unforeseen
catastrophic failure may be accelerated by local oxidative
phenomena [2].
Two examples of tube damage are given here: the yellow
arrows in figure 2a indicates a tube section that underwent creep expansion; figure 2b shows a large longitudinal
crack across a butt welded joint.
a)
b)
Fig. 1 – Assembly of a tubes row (a) and detail of a
single tube (b). The arrows indicate the gas flows
entering and leaving the tube.
53
Memorie
a)
b)
Fig. 2 – Examples of radiant tube damage: a) creep expansion of a tube section; b) large longitudinal crack
across a butt welded joint.
Therefore good mechanical properties and corrosion resistance at high temperature are requested for radiant tube
alloys [3]. These materials have seen significant improvements during the last fifty years: in the 60s and 70s the alloy 25Cr-20Ni-Fe-0.4C, designated as HK-40, was the most
utilized, while in the following decades the HP-40 alloy
(25Cr-35Ni-Fe-0.4C-1.5Si) has been widely considered [4].
The high concentrations of Cr and Ni give great mechanical
strength and corrosion resistance at service temperature,
the presence of Si improves carburization resistance. The
mechanical reinforcement of the HP-40 alloy is obtained
by a dispersion of carbides particles with high hardness [5]
whose stability is decisive for creep behavior [6]. Moreover,
because it has been tested that microalloying elements are
able to stabilize a fine dispersion of carbides, starting from
the 90s HP-40 alloys have been developed with addition
of Nb (about 1%) [7], Ti (up to about 0.8%) [8] and Y (about
0.3% [9] to improve their creep behavior.
Tubes life is shortened by creep damage, being characterized by progressive microstructural changes [10], as carbides transformations, microcracks nucleation and formation of voids, typical of the final creep stage that precedes
fracture [11, 12].
Unfortunately, during the furnace scheduled stop, tubes
have to be decommissioned to cut metallographic samples and perform microscopic observations and mechanical test. However, because each single tube is very
expensive and its damage conditions are not necessarily representative of the other tubes, many experimental
works have been addressed to non destructive methods,
such as eddy current and ultrasonic measurements [13,
14, 15]. Moreover, being these investigations affected by
many uncertainties, studies on the relationship between
microstructural degradation and mechanical properties
are largely quoted in literature [2, 11, 16]. In this respect it
is worth noting that the residual life evaluation depends on
actual service temperatures (measured by optical pyrometers with an error of +/- 20 °C), mechanical properties
values after long time service, actual stress state and ther54
mal cycles associated to the furnace start-up and shutdown. So adequate means of investigation giving reliable
information about tubes damage are required.
Nowadays reforming furnaces rely on the Laser Optic Tube
Inspection System (LOTIS), a non-destructive control technique based on creep deformation measures performed
in situ by driving a laser probe inside tubes [17]. It is very
promising for the development of a criterion for decommissioning tubes that would be no longer safe until the
next scheduled stop [18]. So furnaces have to be regularly
inspected during each stop, in order to identify the damaged components and replace them. The radiant tubes
are visually observed with the aim of identifying macroscopic damage; creep deformations are detected through
LOTIS measures in order to select the tubes to put out of
service for safety reasons.
In previous works [18, 19] laboratory tests were carried
out to correlate mechanical properties and microstructural changes in tubes decommissioned after long time service. In this paper the experimental investigations have
been extended to the butt welds that join the tube pieces.
Samples for creep tests and microstructural observations
were cut from various sections also including welds.
MATERIALS AND METHODS
As shown in the working drawing (Fig. 3), radiant tubes are
constructed in three butt welded pieces (W1 welds); the
first one (located in the upper vertical position inside the
furnace) is welded to the upper flange (W2), the third one
is welded to the reducer (W1), in turn welded to the outlet
tube (W3); other welds concern the gas inlet nozzle (W4)
and the catalyst support grid (W5).
Radiant tubes have an internal diameter of 101.6 mm, a
nominal thickness of 10.5 mm and an overall length of
12.8 m. The design conditions are pressure of 32.7 bar
and temperature of 950°C, while the operating temperature is 900°C.
Saldatura
Fig. 3 – Working drawing: a) upper flange, b) radiant tube constructed in three pieces, c) reducer, d) outlet tube,
e) catalyst support grid, f) gas inlet nozzle.
The three pieces of the radiant tube (b), the connection
of the gas outlet (c) and the catalyst support grid (e) are
made of the ASTM 608 HP-Nb alloy, a HP alloy modified by
Nb micro additions; the upper flange (a) and the gas inlet
nozzle (f) are made of the ASTM A182 Type F22 steel; the
gas outlet tube (d) is made of the ASTM B407 UNS NO8811
Incoloy.
The composition of the ASTM 608 HP-Nb alloy is given
in table 1. The microstructure of the spun cast pieces is
characterized by radial austenitic dendrites, that are well
delineated by carbide particles precipitated in the interdendritic spaces.
Welds are performed by means of GTAW, utilizing pure Ar
as shielding gas. Fig. 4 shows the edge preparation for
W1, W2 and W3 butt welds. Welding parameters are given
in table 2.
a)
Alloy
C
Si
A 608 HP-Nb 0.45 1.5
Mn
Cr
Ni
Nb
Ti
Fe
1
25
35 1.5 add. bal.
Table 1 – Composition given by the manufacturer of
the alloy.
Root pass Subsequent Final pass
passes
Current
DC
DC
DC
Electrode polarity
negative
negative
negative
Intensity (A)
140-170
100-150
100-130
Voltage (V)
14-16
14-16
14-15
Welding speed
(cm/min)
4-7
6-9
5-8
Tungsten electrode
Type WS2 Ø 2,4 mm
Filler material
UTP A2535Nb (25Cr-35Ni-1.2Nb)
Ø 2,4 mm rod
Table 2 – Process parameters of the W1 welds.
b)
Fig. 4 – Edge preparation for butt welding of tube
pieces (a), welded joint (b).
Creep deformation of tubes is due to the hoop stress
caused by the internal pressure. The most damaged
zones are in the lower part, near the catalyst support grid,
where the gas temperature, starting from the inlet value of
500°C, reaches values around 900°C.
Creep deformation was evaluated, during the programmed
furnace stop, through LOTIS measures of the internal diameter (Di), inserting in the upper flange a probe with the
laser source and the receiving system. The probe, mechanically driven inside the examined tube, rotates up to
1800 r.p.m. and generates a helical map of internal surface, with an accuracy of 0.05 mm; so the deformation can
be measured with a precision of 0.05%, being the internal
diameter about 100 mm.
The deformation of a tube section (ε), on the basis of the
measured internal diameter (Di) and referred to the nominal diameter (Do), equal to 101.6 mm, is given by the following equation:
(1)
ε = (Di – Do) / Do
55
Memorie
Following a criterion developed in previous work [18, 19],
the tubes where deformation reached values greater than
ε = 1.5% were put out of service and utilized to take the
experimental samples.
The decommissioned tubes were visually investigated for a
first inspection of their damage conditions.
Creep tests were performed on samples cut longitudinally
from tube put out of service after about 100000 hours
and, for comparison, from samples taken from a tube in
the as cast condition. Samples with welded joint were
also machined.
For microhardness test and microscopic observations
samples were metallographically prepared with the usual
techniques and etched by a solution containing 15% of
glycerol (100% concentration), 45% of HNO3 (65% concentration) and 40% of HCl (37% concentration).
Considering that the results of LOTIS measurements
showed inhomogeneous deformations along the tubes
axis, experimental surveys were carried out on samples
cut longitudinally both from sections with high deformation (1.5-2.0%) and from undeformed sections, in order to
put in comparison their microstructure and creep strength.
Samples of welded joints were also cut.
Metallographic samples were observed optically and by
scanning electron microscopy (SEM) with energy dispersive X-ray spectroscopy (EDS).
RESULTS AND DISCUSSIONS
LOTIS measurements
The internal diameter creep expansion along each tube was
accurately recorded through LOTIS measures, performed
in situ with the aim of selecting the tubes to put out of service when ε > 1.5%. Results can be summarized as 78% of
tubes with diameter deformations less than 1.5%, while in
the remaining 22% were recorded deformations exceeding
the established limit [18].
Moreover LOTIS measurements showed that diameter deformation increases from the upper flange, where in general it is negligible, up to the lowest zone of the tube. In
the decommissioned tubes diameter deformation values
of 1.5-2.0% were recorded near the catalyst grid.
Visual inspections
The tubes put out of service after 114000 hours show
highly oxidized surfaces along all their length and zones
located in the lower part which have undergone creep
deformations detectable also by visual inspection. These
zones are mainly near the catalyst grid, where the diameter expansion is clearly observable (fig. 2a).
Creep tests
Results of creep test refer to samples in the as cast condition and samples cut longitudinally from both undeformed
and deformed (ε=1.5-2.0%) zones of a decommissioned
tube. Results of samples with welds are also considered
(cross-weld test). The test temperature (T) was in the range
56
between 920 and 980°C. In order to achieve acceptable
rupture times (between 200 and 1000 hours), the test
stress was assumed in the range 22-30 MPa, greater than
the hoop stress at the design pressure (σ = 17.5 MPa).
For a useful comparison, in fig. 5 the creep test results
are plotted on a Larson Miller diagram (σ vs. LMP), being
σ test stress and LMP = T (C + log t) / 1000, with T (K)
test temperature, t (h) rupture time and C=22.9 a constant
characteristic of the considered material [10].
Moreover both average and minimum curves characteristics of the A 608 HP-Nb alloy are given in the same diagram as reported in the manufacturer catalogue [20].
Creep tests results can be summarized as follows:
- the experimental point of the as supplied material (blue
triangle) is on the minimum curve of the manufacturer
catalogue,
- results of samples taken from undeformed zones (yellow
rhombus) are similar to the as supplied material,
- results of samples taken from deformed zones (blue circle) are well below those of the undeformed material,
- samples cut transversely to the welded joint (red circle)
behave in agreement with the previous ones.
The experimental diagram σ vs. LMP shows that specimens with welds have behavior similar to the base material with the same diameter deformation. However they
differ in creep ductility: rupture elongation in creep test
resulted around 3-5% for samples with welds, while it was
20-30% for samples of base material.
The residual life of decommissioned tubes can be obtained thanks to the curve extrapolated from experimental data: entering with the hoop stress σ = 17.5 MPa, the
corresponding LMP value, equal to about 32.6, allows to
calculate the rupture time at a given temperature. For
example, with temperatures of 900, 920 and 940°C, the
corresponding rupture times are respectively 10, 3.2 and 1
year. Then rupture time after long time service is strongly
dependent on temperature changes that are in the range
of measurement errors typical of pyrometer systems used
in reforming furnaces.
Optical microscopy and Vickers microhardness test
After long time service, the A 608 HP-Nb alloy is characterized by microstructure modifications that can be related to
the degree of diameter expansion.
The radiant tube walls undergo creep caused by circumferential stress and temperature and each zone has specific
metallurgical features.
Some considerations can be drawn with reference to the
optical micrographs shown in figure 6 at the respective
stages of a typical creep curve.
- The as cast microstructure is characterized by austenitic
dendrites, well outlined by a network of coarse carbides
precipitated in the interdendritic spaces.
- The microstructure of samples cut from undeformed
zones of a decommissioned tube is nearly similar to the
one of as cast alloy, so it can be referred to the first
creep stage.
- The microstructure of samples cut from deformed zone
Saldatura
Fig. 5 – Creep test results: Larson Miller diagram.
(ε = 1.5-2.0%) is characterized by a precipitation of fine
secondary carbides inside the austenitic grains, optically
observable at high magnification, and by nearly circular
isolated microvoids.
Isolated microvoids appear during the steady state creep.
In particular the absence of aligned microvoids, formed by
the coalescence of two or more microvoids [21], can be
considered as an indication that the final creep stage has
not yet started.
In welded sections, no evidence of particular structures
was detected at the boundary between base metal and
molten zone (fig. 7). The latter appears to be characterized by very fine carbides aligned with the heat flow direction.
Regarding microhardness survey, it was noticed a remarkable softening after 114000 hours of service compared
to values of about 400HV measured in the as cast material: microhardness values through base metal and molted
zone are around 180 HV (fig. 8).
SEM and EDS investigations
Microstructural changes in samples taken from deformed
zones of a decommissioned tube respect to the as cast
alloy are investigated by SEM observations and EDS measurements too.
In the as cast alloy, two kind of precipitates at the austenitic grain boundaries are clearly observable: a grey phase,
with greater volume fraction, and a phase in the form of
small white particles. The EDS spectra allow to recognize
the grey phase as Cr carbide, while the small white particles are rich of Nb and Si.
In samples taken from deformed zones, grey precipitates
are located both inside grains and at their boundaries (fig.
9). These particles are recognized as Cr carbide through
EDS measurements. The HP alloys underwent phase
transformations due to thermal activation during service
at temperatures between 1123 and 1325 K and primary Cr
carbides M7C3, instable at high temperature, transform into
intragranular and intergranular M23C6 precipitates [23].
Another change regards the white precipitates at grain
boundary that in deformed samples have dimension and
volumetric fraction greater than in the as cast alloy; moreLa Metallurgia Italiana - n. 1/2015
Fig. 6 – Microstructures of the ASTM 608 HP-Nb
alloy, in relation to the respective level of creep
deformation. In the case of material after 114000
hours of service it can be outlined: a) fine carbide
precipitation inside grains, b) primary carbide at grain
boundary, c) nearly circular microvoids.
Fig. 7 – Optical micrograph of the boundary between
molten zone and base metal in a welded joint (tube
decommissioned after 114000 hours).
Fig. 8 – Vickers microhardness (200g -15s) in a welded
section (tube put out of service after 114000 hours).
over EDS measurements give high concentrations of Nb,
Ni, Si. As quoted in literature [33, 34], Nb carbides are
not stable in the range of service temperatures, between
700 and 900°C, transforming into Ni-Nb silicates known
as G-phase. Unlike the Nb carbides that give good creep
properties, the G-phase coarsening is followed by reduction of creep resistance [3]: this negative effects can be
57
Memorie
REFERENCES
Fig. 9 – SEM micrograph of a sample taken form a
deformed zone (tube decommissioned after 114000
hours).
ascribed to the interface between G-phase and matrix that
becomes a preferential site for creep damage [7]. Luckily this phenomenon can be overlooked in the HP grades
microalloyed with small quantities of Ti, that reduce the
G-phase volume fraction [7].
CONCLUSIONS
The experimental investigations carried out on radiant
tubes, made of ASTM 608 HP-Nb alloy and installed in a
reforming furnace, have confirmed that LOTIS measurements of the internal diameter give good indications on
the degree of creep damage and allow to select the tubes
to put out of service.
Creep tests and metallographic observations have shown
that samples taken from undeformed section are free of
damage and have characteristics similar to those of the
as cast alloy, while samples taken from deformed section
(in the range 1.5-2%) show microstructural changes and
decay of creep strength. In any case, welds do not affect
further the strength properties of the reforming tubes: the
hardness values through them are almost similar to those
of base material; furthermore, in the same conditions of
diameter deformation, creep behavior of samples with
welds and sample of base material is comparable.
The residual life of samples taken from deformed section is
very sensitive to small temperature changes, in the range
of measurement errors, becoming dangerously close to
two years, that is the time between two scheduled stops
of the furnace. Thus is confirmed the validity of putting
out of service tubes when the internal diameter deformation is greater than 1.5%.
58
[1] J. Swaminathan, K. Guguloth, M. Gunjan, P. Roy, R. Ghosh, “Failure analysis and remaining life assessment of service exposed primary reformer
heater tubes”, Engineering Failure Analysis, 15 (2008), pp. 311-331
[2] H.M. Tawancy, “Damage analysis of catalyst tube of a reformer furnace used in hydrogen production”, Metallogr. Microstruct. Anal., 1
(2012), pp. 199-207
[3] Jaske, “Issues in life assessment of reformer tubes”, NACE Corrosion
Conference 2005, Houston, 3-7 April 2005, Anti-Corrosion Methods
and Materials, Vol. 52 Issue 5
[4] T.L. da Silveira, Iain Le May, “Reformer furnaces: materials, damage
mechanisms and assessment”, The Arabian Journal of Science and
Engineering, v. 31 n. 2C (2006), pp. 99-119
[5] T. Pardoen, J.W. Hutchinson, “An extended model for void growth and
coalescence”, Journal of the Mechanics and Physics of Solids, 48
(2000), pp. 2467-2512
[6] J. Rodriguez, S. Haro, A. Velasco, R. Colas, “A metallographic study
of aging in cast heat-resisting alloy”, Materials Characterization, 45
(2000), pp. 25-32
[7] L.H. de Almeida, A. F. Ribeiro, J. Le May, “Microstructural characterization of modified 25Cr-35Ni centrifugally cast steel furnace tubes”,
Materials Charact., 49 (2003), pp. 219-229
[8] A. Alvino, D. Lega, F. Giacobbe, V. Mazzocchi, A. Rinaldi, “Damage
characterization in two reformer heater tubes after nearly 10 years
service at different operative and maintenance conditions”, Engineering Failure Analysis, 17 (2010), pp. 1526-1541
[9] F.C. Nunes, L. H. de Almeida, J. Dille, J.L. Delplancke, J. Le May, “Microstrucutral changes caused by Yttrium on Nb-Ti-modified centrifigually
cast HP-type stainless steels”, Materials Characterization, 58 (2007),
pp. 132-142
[10]R. Voicu, J. Lacaze, E. Andrieu, D. Poquillon, J. Furtado, “Creep and
tensile behaviour of austenitic Fe–Cr–Ni stainless steels”, Mat. Science and Eng. A, 510-511 (2009), pp. 185-189.
[11] D. Alessio, G. Gonzalez, V. F. Pirrone, L. Iurman, L. Moro, “Variation of
creep properties in HP steel by influence of temperature”, Procedia
Materials Science, 1 (2012), pp. 104-109
[12] C. Maharaj, C. A.C. Imbert, J. Dear, “Failure analysis and creep remaining life of Hydrogen reformer outlet pigtail tubes”, Engineering Failure
Analysis 15 (2008), pp. 1076-1087
[13] J.R. Widrig, “New technology for inspecting critical in-plant piping service”, Process Safety Progress V. 34 (2011), No. 4, pp. 334-337
[14] J.R. Widrig, “The challenge of inspection and assessment of critical
piping systems in chemical plants”, Inspection Engineering Journal, V.
20, No. 4 (2013), pp. 13-16
[15]B. Shannon, C. Jaske, “A Comprehensive approach to reformer tube
inspection and assessment”, NDT.net, V. 9, No. 6 (2004), pp. 1-12
[16] J. Łabanowski, “Evaluation of reformer tubes degradation after long
term operation”, Journal of Achievements in Materials and Manufacturing Engineering, 43/1 (2010), pp. 244-251.
[17] R. Roberts, “Enhanced steam reformer tube inspection and remaining life assessment methodologies”, Proceedings of the International
Conference & Exhibition Nitrogenn+Syngas, Dusseldorf, 21-24 February 2011
[18]L. Bonaccorsi, E. Guglielmino, R. Pino, C. Servetto, A. Sili, “Damage
analysis in Fe-Cr-Ni centrifugally cast alloy tubes for reforming furnaces”, Eng. Failure Anal., 36 (2014), pp. 65-74
[19]L. Bonaccorsi, E. Guglielmino, R. Pino, A. Sili, G. Chiofalo, M. De
Marco, C. Servetto, “Creep damage and metallurgical characterization
of high alloyed reformer tubes after long service time”, ECCC – Creep
& Fracture Conference 2014, Rome, May 5-7, 2014
[20]Shmidt + Clemens Group, Centralloy G 4852 Micro - Material Data
Sheet, September 2009, Rev. 02
[21] Wahab Azmi Abdul, Milo V. Kral, “3D Analysis of creep voids in hydrogen reformer tubes”, Materials Science and Engineering A, 412
(2005), pp. 222-229
[22]E.A.Kenik, P.J.Maziasz, R.W. Swindeman, J. Cervenka, D. May, “Structure and phase stability in a cast modified-HP austenite after long
term ageing”, Scripta Mater., 49 (2003), pp. 117-22
[23]M. Mostafaei, M. Shamanian, H. Purmohamad, M. Amini, A. Saatchi
“Microstructural degradation of two cast heat resistant reformer
tubes after long term service exposure”, Engineering Failure Analysis,
18 (2011), 164–171
Vita associativa
2015… si volta pagina!
SOMMARIO
Non è solo la ben nota espressione metaforica, pronunciata per annunciare - in
apertura d’anno – un sacco di buone intenzioni. Questa volta sta a rappresentare proprio quello che succederà. Infatti, come molte altre realtà nel mondo dell’editoria, l’AIM è in procinto di affrontare la sfida dell’introduzione della
versione digitale, che riguarderà non solo la nostra rivista, ma anche alcuni
testi della collana tecnica. E allora una volta “voltata pagina” ...non avremo più
pagine da voltare!
VITA ASSOCIATIVA
Questa “trasmutazione” si è rivelata ormai necessaria non solo per le implicazioni economiche ma anche per le potenzialità a livello di accessibilità, di diffusione, nonché per potenziare il “fattore di impatto” dei lavori pubblicati.
Prossime manifestazioni AIM.... 61
Proprio in vista di una maggiore visibilità all’estero è nata l’esigenza di presentare - a partire dal prossimo numero - anche il nuovo sottotitolo International
Journal of the Italian Association for Metallurgy, con il duplice intento di esprimere la connotazione di internazionalità e di offrire un “biglietto da visita” più
comprensibile per chi ci legge da lontano.
D’altra parte, benché rimanga ben saldo il nostro intento di rendere la nostra
rivista vetrina e strumento di dialogo della ricerca metallurgia in Italia, essa è
diventata sempre di più risorsa in un ambito internazionale e molti sono ormai
i contributi che ci vengono sottoposti da autori stranieri, soprattutto da quando
siamo nel circuito delle riviste scientifiche recensite.
Relativamente alle manifestazioni a calendario, vi ricordiamo alcuni importanti
appuntamenti nei prossimi mesi:
• il Corso itinerante “Metallurgia Verde”, suddiviso nei due moduli Metallurgia
verde e sostenibile: come far funzionare un azienda metallurgica nel
pieno rispetto dell’ambiente (che si terrà nei giorni 10 e 24 febbraio) ed Efficienza energetica: scenari ed opportunità nell’industria metallurgica
(nei giorni 17 febbraio e 3 marzo);
• il Corso “Prove Meccaniche”, giunto alla nona edizione, che si svolgerà a
Roma nei giorni 24, 25 e 26 febbraio e organizzato in collaborazione con il
CSM;
• il nuovo appuntamento della serie di “Pillole per preposti” dal titolo “La manutenzione degli impianti”, a Brescia l’11 marzo;
• la Giornata di studio “L’impiego dell’acciaio nelle costruzioni civili”, che
si terrà a Milano il 18 marzo, i cui contenuti sono descritti nelle pagine seguenti.
L’impiego dell’acciaio nelle
costruzioni civili........................60
Efficienza energetica.................60
dai centri
Attività dei Comitati Tecnici......62
ECONOMIA E PRODUZIONE
Sfiorata quota 1,5 miliardi di
tonnellate..................................63
Italia: a novembre torna il segno
meno.........................................63
Italia: l’import cresce più
dell’export.................................64
Il preridotto corre più della ghisa.. 64
Quote sociali AIM 2015
(ANNO SOLARE)
Benemeriti (quota minima).1.750,00 €
Sostenitori (quota minima)... 750,00 €
Ordinari (solo persona)........... 70,00 €
Seniores................................... 25,00 €
Juniores.................................... 15,00 €
La quota dà diritto di ricevere la rivista
dell’Associazione La Metallurgia Italiana. Ai soci viene riservato un prezzo
speciale per la partecipazione alle manifestazioni AIM e per l’acquisto delle
pubblicazioni edite da AIM.
Per ulteriori informazioni, iscrizioni,
rinnovi:
AIM, Piazzale R. Morandi, 2
20121 Milano
Tel.: 02 76021132/76397770,
fax: 02 76020551
e-mail: [email protected]
www.aimnet.it
59
Atti e notizie
Giornata di Studio
L’IMPIEGO DELL’ACCIAIO
NELLE COSTRUZIONI
CIVILI
Milano, 18 marzo 2015
La Giornata di studio, organizzata dal
Centro di Studio “Metalli e Tecnologie Applicative” dell’Associazione
Italiana di Metallurgia, si pone come
obiettivo quello di far conoscere le
caratteristiche peculiari dell’acciaio
utilizzabili nella realizzazione di strutture civili, soffermandosi, in particolare, sugli aspetti pratici: dalle tecniche costruttive ai controlli di qualità,
passando in rassegna le principali criticità, senza trascurare l’aspetto dei
costi e dei tempi di realizzazione delle
strutture. Particolare attenzione sarà
dedicata alle metodologie di saldatura, esaminandone i principali aspetti
metallurgici, le prove di controllo non
distruttive e le procedure di processo
che, se non seguite in modo corretto, possono indurre difettosità più o
meno gravi. Tutti gli argomenti saranno trattati partendo da esempi pratici: una presentazione riguarderà la
realizzazione di una scuola post terremoto ad Haiti, un’altra la realizzazione
della torre Garibaldi a Milano. Verranno presentati anche gli aspetti pratici
dei trattamenti di protezione delle superfici, in particolare della zincatura e
della verniciatura. L’importante tema
dell’obbligo della marcatura CE per i
prodotti da costruzione in acciaio ed
alluminio sarà trattato con riferimento alle normative in vigore. Nell’organizzare questa manifestazione che è solo la prima Giornata di
studio su un argomento di grande importanza sia culturale che economica,
il principale obiettivo che il Comitato
Tecnico “Metalli e tecnologie applicative” dell’Associazione Italiana di Metallurgia si è posto è quello di favorire
l’apertura di un dialogo fra operatori
del settore e esperti di metallurgia
che possa portare a risultati significativi per un ambito il cui sviluppo risulta essenziale per uscire dall’attuale
negativa situazione congiunturale.
Il programma della Giornata di Studio
- coordinata da Andrea Rossetti - prevede i seguenti interventi:
60
•perché scegliere l’acciaio nelle
strutture civili;
•la marcatura CE delle strutture in
acciaio e alluminio: la norma UNI
EN 1090;
•difetti di saldatura: esempi pratici;
•resistenza a fatica dei giunti saldati;
•discussione dei temi trattati;
•prove non distruttive sulle strutture
saldate;
•durabilità delle strutture in acciaio:
zincatura e verniciatura;
•ricostruzione in tempi brevi: una
scuola ad Haiti;
•struttura “spire” torre Garibaldi a
Milano
Per questo evento i partecipanti potranno richiedere, in fase di registrazione, il riconoscimento di 3 Crediti
Formativi Professionali (CFP). Tale
riconoscimento è stato autorizzato
dall’Ordine Ingegneri di Milano, che
ne ha valutato anticipatamente i contenuti formativi professionali e le modalità di attuazione.
Corso
sviluppati in aziende del settore.
Il Corso è rivolto agli Energy Manager,
ai Responsabili Ambiente, ai responsabili di produzione e di manutenzione, nonché ai referenti degli uffici
tecnici e R&S aziendali.
La consolidata struttura del Corso
prevede l’integrazione tra la presentazione di alcune tematiche da un
punto di vista tecnico e normativo, la
visita agli impianti produttivi, il confronto con i tecnici che li gestiscono
attraverso l’esame di casi reali di applicazione.
Questa impostazione consente ai
partecipanti di aggiornare le proprie
conoscenze teoriche e nello stesso
tempo di osservare soluzioni applicate da alcune aziende operanti nel
settore metallurgico.
Il programma del Corso prevede i seguenti interventi:
metallurgia verde
Scenari di efficienza energetica ed
opportunità nell’industria metallurgica
L’Efficienza energetica costituisce
lo strumento più valido per ridurre i consumi, garantire la sicurezza
dell’approvvigionamento, accelerare
la diffusione di soluzioni tecnologiche
innovative, con ricadute positive in
termini economici e di competitività soprattutto in aziende energivore
quali quelle del settore metallurgico.
Per questo motivo il Centro di Studio
Ambiente e Sicurezza dell’AIM organizza un evento di due giornate, che si
svolgerà in due settimane consecutive nei giorni 17 febbraio e 3 marzo, il
cui obiettivo sarà fare il punto sullo
stato dell’arte del quadro nazionale,
normativo di sviluppo, nonché di fornire elementi pratici per la realizzazione di interventi di miglioramento
energetico a livello industriale.
Il Corso si completa, infine, con la
presentazione di progetti concreti
Per maggiori informazioni, gli interessati potranno rivolgersi alla Segreteria AIM ([email protected] / tel.
02 76021132) oppure visitare il sito
www.aimnet.it.
17 febbraio 2015 ( presso AFV BELTRAME di Vicenza)
•Efficienza energetica: inquadramento normativo ed obblighi
•Progetti di miglioramento energetico
Vita associativa
con motori elettrici ad alta efficienza, inverter, sistemi di pompaggio
•Applicazione delle tecnologie di
Controllo Avanzato di Processo
(APC) per il risparmio energetico del
forno di riscaldo
•Progetti di illuminazione ad alta
efficienza e adeguamento alla normativa – iter di analisi e studio per
garantire il risparmio energetico
•Bruciatori ad alta efficienza: analisi
comparativa e criteri di scelta
•Presentazione aziendale e visita guidata dello stabilimento
3 marzo 2015 (presso ORI Martin
di Brescia)
•Innovazione e tecnologia nel settore
siderurgico
•Recupero del calore dai fumi di un
forno elettrico ad arco: il progetto
di O.R.I. Martin Sistemi di rifasamento
•Progetti di ottimizzazione delle centrali di produzione di aria compressa
•Sistemi di monitoraggio dell’energia/Sistemi di gestione dell’energia
(ISO 50001)
•Presentazione aziendale e visita guidata dello stabilimento
Per maggiori informazioni, rivolgersi alla Segreteria AIM ([email protected] /
tel. 02 76021132) oppure visitare il sito www.aimnet.it.
LE PROSSIME MANIFESTAZIONI AIM
CORROSIONE E PROTEZIONE DEI
METALLI
Corso – Centro C
3° Modulo - Milano, 14-15 gennaio
4° Modulo - Milano, 28-29 gennaio
METALLURGIA VERDE E
SOSTENIBILE
Corso itinerante – Centro AS
10 e 24 febbraio
METALLURGIA VERDE: EFFICIENZA
ENERGETICA
Corso itinerante – Centro AS
17 febbraio – 3 marzo
PROVE MECCANICHE
9a edizione Corso – Centro CCP
Roma, 24 - 25 - 26 febbraio
Pillole per preposti: LA
MANUTENZIONE DEGLI IMPIANTI
Corso – Centro A
Brescia, 11 marzo
L’IMPIEGO DELL’ACCIAIO NELLE
COSTRUZIONI CIVILI
GdS – Centro MTA
Milano, 18 marzo
MATERIALI REFRATTARI
GdS – Centro A
Milano, 9 aprile
TRATTAMENTI TERMICI
Corso modulare – Centro TTM
Milano, 14-15, 21-22 aprile
MATERIALI PER OIL&GAS
IN SUPERLEGHE E ACCIAI
INOSSIDABILI
GdS – Centro FOR
Milano, 23-24 aprile
SOLIDIFICAZIONE E COLATA
CONTINUA
Corso itinerante – Centro A
7-8-14-15-28-29 maggio
Eur. Conf. HEAT TREATMENT &
SURFACE ENGINEERING & 22nd
IFHTSE Congress
Venezia, 20-22 maggio
http://www.aimnet.it/ht2015.htm
TRATTAMENTI TERMICI
Milano, 9-10, 16-17 giugno
XI Giornate Nazionali
CORROSIONE E PROTEZIONE
Ferrara, 15-16-17 giugno
http://www.aimnet.it/
gncorr2015.htm
STEELSIM 2015
6th Int. Conf. Modelling and
Simulation of Metallurgical
Processes in Steelmaking
Bardolino, 23-25 settembre
http://www.aimnet.it/
steelsim2015.htm
TRATTAMENTI TERMICI
Milano, 29-30 settembre
Pillole per Preposti: LA MACCHINA
FUSORIA
Corso – Centro A
Brescia, 14 ottobre
MATERIALI DI CARICA IN
ACCIAIERIA
Gds – Centro A
Milano, 18 novembre
EEC 2016 - 11th European Electric
Steelmaking Conference
Venezia, 25-27 maggio 2016
http://www.aimnet.it/eec2016.htm
Per ulteriori informazioni
rivolgersi alla Segreteria
AIM, e-mail: [email protected]
www.aimnet.it
61
Atti e notizie
ATTIVITà DEI comitati tecnici
CENTRO RIVESTIMENTI (R )
Iniziative future
(riunione del C.T. c/o U. La Sapienza, Roma – 5 marzo 2014)
- Il Presidente Vedani segnala che dal Centro Lavorazioni
Plastiche ha manifestato l’interesse per l’organizzazione
di una GdS sull’estrusione dell’alluminio; l’argomento si
discuterà in seguito.
Consuntivo di attività svolte
- Al termine del Corso “Rivestimenti spessi: placcatura
e termospruzzatura” (Roma, 4 marzo 2014) sono stati
esaminati i questionari compilati dai partecipanti e il
commento generale è stato positivo (relazioni di alto livello,
comprensibili e didatticamente chiare) anche per la visita
ai laboratori del CSM. Si prende atto dell’osservazione
riguardante le relazioni applicative presentate, molto
orientate alle problematiche dei rivestimenti per il settore
delle turbine, per cui nelle future edizioni il Corso sarà
arricchito da contributi relativi ai rivestimenti spessi in altri
settori industriali.
Manifestazioni in corso di organizzazione
- Si è deciso di rimandare al 2015 il terzo modulo del Corso
“Rivestimenti per via umida”, che verrà coordinato da:
Bestetti, Lusvarghi, Brisotto. Viene suggerito di contattare
il Presidente di Assogalvanica sia per giungere ai suoi
associati sia per sollecitare interventi in quanto la tematica
del cromo duro è molto importante e la sua descrizione
richiede di trattare non solo gli aspetti tecnici, ma anche
quelli legislativi. Il Corso può avere come sede Padova con
visita ai laboratori.
Iniziative future
- Il Presidente Bestetti segnala che si svilupperanno i
contatti con il Centro Metalli e Tecnologie Applicative per
l’organizzazione della giornata sull’alimentare.
CENTRO METALLI LEGGERI (ML)
(riunione del C.T. – 8 settembre 2014)
Consuntivo di attività svolte
- La GdS sul Semisolido al Metef ha suscitato un buon
interesse, la partecipazione è stata adeguata. Dalle
discussioni è emersa la difficoltà del processo in semisolido
ad inserirsi in un contesto industriale.
- Per la GdS “Progettare con le leghe leggere ed
ultraleggere”, prevista per il 3 dicembre 2014 a Milano, si
analizza il programma nella stesura quasi definitiva.
- Per l’evento HTDC- High Tech Die Casting, a settembre
2015 a Venezia, giunto alla 6° edizione, ci sarà la
collaborazione del Centro Pressocolata. Bonollo sarà il
coordinatore della manifestazione.
62
- Nella prossima riunione si valuteranno luogo e tempistiche
per il Corso “Igiene delle leghe di alluminio da fonderia”
(5^ edizione - data probabile ottobre 2015).
- Per il Corso “Metallografia delle Leghe Leggere” si decide
di individuare data e luogo di esecuzione sentito Rosso
(ipotesi entro aprile 2015).
- Con riferimento ad un processo importante già sviluppato
nel mondo, si valuta l’opportunità di divulgazione anche
in Italia tramite una GdS “Additive Manufacturing per
componenti metallici”.
Stato dell’arte e notizie
- È arrivata segnalazione via mail della futura organizzazione
di un Convegno sulla corrosione, articolato in quattro
giornate; si informano i presenti della richiesta di eventuali
contributi.
- Nella settimana in corso, si sta svolgendo a Rimini
il convegno AIAS coordinato dalla prof. Ceschini; le
considerazioni si faranno in sede di prossima riunione.
- Per portare a conoscenza le attività del Centro si
considera la richiesta di pubblicare, sulle due riviste AL e
P&T; un estratto del verbale del CT.
CENTRO FORGIATURA (FOR)
( riunione del C.T. – 18 settembre 2014)
- Si discute della futura GdS con argomento Oil & Gas,
che affronterà i diversi aspetti (impianti, trattamenti
termici, caratteristiche prodotto, normativa di riferimento,
materiali, teoria, prodotti e produzione) con riferimento
a due tipologie di materiali: le leghe a base di nichel e
gli acciai. Dopo un giro di tavola sui possibili argomenti
e relatori, al Presidente Rampinini vengono segnalate
le richieste di riprendere nella GdS testi riguardanti i
trattamenti degli acciai inossidabili (Corso “Forgiatori”) e
l’ottimizzazione della microstruttura in grossi forgiati ( GdS
“Trattamenti termici su forgiati di grosse dimensioni”).
Iniziative future
- Si decide di rimandare alla prossima riunione la
discussione inerente la GdS su Lean Manufacturing e
metodologie.
Economia e produzione
Sfiorata quota
1,5 miliardi di
tonnellate
Nei primi 11 mesi del 2014 le acciaierie
mondiali hanno sfornato poco meno
di 1,5 miliardi di tonnellate di prodotti
siderurgici. Questo il responso dei
numeri diffusi dalla World Steel
Association, l’associazione globale
dei produttori siderurgici, che ha
certificato un output a livello globale
di 1,498 miliardi di tonnellate, con
un incremento dell’1,8% rispetto allo
stesso periodo del 2013. Il faro della
siderurgia mondiale rimane l’Asia, con
una produzione di oltre 1,013 miliardi
di tonnellate (+2,2% rispetto al 2013),
davanti all’Unione Europea (156,7
milioni di tonnellate, +2,3%), al Nord
America (111,5 milioni di tonnellate,
+2,2%), alla CSI (96,8 milioni di
tonnellate, -2,2%), al Sud America
(41,7 milioni di tonnellate, -1,6%),
Italia: a novembre
torna il segno meno
Per la prima volta nel 2014 la produzione complessiva italiana torna
in territorio negativo. Nei primi undici mesi dell’anno, infatti, l’output
siderurgico nazionale è stato pari a
22,255 milioni di tonnellate, con un
decremento dello 0,2% (-44.000 tonnellate) rispetto allo stesso periodo
del 2014. A penalizzare il dato, oltre
alla crisi dei mercati ed al basso livello degli acquisti in filiera, anche alcuLa Metallurgia Italiana - n. 1/2015
ai Paesi europei non facenti parte
dell’Ue (33,2 milioni di tonnellate,
-0,7%), al Medio Oriente (25,7 milioni
di tonnellate, +8,1%), all’Africa (14,5
milioni di tonnellate, +0,7%) ed
all’Oceania (5,0 milioni di tonnellate,
-2,4%). Guardando ai singoli Paesi, la
Cina rimane il leader indiscusso, con
un output di 748,7 milioni di tonnellate,
ma con un tasso di crescita (+1,9%)
inferiore a quello globale. Al secondo
posto in classifica il Giappone (101,7
milioni di tonnellate, +0,4%), poi gli
Stati Uniti (81,0 milioni di tonnellate,
+1,5%), l’India (76,2 milioni di
tonnellate, +2,4%) e la Corea del Sud
(65,3 milioni di tonnellate, +8,5%). In
Europa comanda la Germania (39,7
milioni di tonnellate, +1,3%) davanti
all’Italia, alla Francia (15,0 milioni
di tonnellate, +3,7%), alla Spagna
(13,3 milioni di tonnellate, -0,3%)
ed al Regno Unito (11,2 milioni di
tonnellate, +3,6%).
ni elementi imprevisti come i guasti
che hanno colpito alcune primarie
acciaierie nazionali, costringendole
a prolungati fermi durante il mese di
novembre. Entrando maggiormente
nel dettaglio, Federacciai ha comunicato che tra gennaio ed ottobre (ultimo dato disponibile) l’output di piani
è stato pari a 10,006 milioni di tonnellate, con un incremento dello 0,4%
rispetto allo stesso periodo del 2013,
mentre quello di lunghi si è attestato
a 9,974 milioni di tonnellate (+1,2%).
A cura di Siderweb
Produzione mondiale di acciaio
(in migliaia di tonnellate, fonte: WSA)
Nov. 2014
Genn.-Nov. ‘14
Austria
671
7.252
Belgio
565
6.760
Bulgaria
50
566
Croazia
10
169
Rep. Ceca
458
4.899
Finlandia
315
3.497
Francia
1.384
15.031
Germania
3.599
39.694
Grecia
100
957
Ungheria
134
1.052
Italia
1.857
22.255
Lussemburgo
200
2.069
Olanda
592
6.362
Polonia
730
7.891
Slovacchia
415
4.281
Slovenia
50
582
Spagna
1.239
13.274
Svezia
406
4.168
Regno Unito
991
11.212
Altri Ue
490
4.760
Totale Ue
14.256
156.731
Bosnia
60
732
Macedonia
20
210
Norvegia
51
545
Serbia
59
513
Turchia
2.795
31.230
Altri Europa
2.985
33.230
Bielorussia
224
2.279
Kazakistan
339
3.441
Moldavia
29
332
Russia
5.840
64.794
Ucraina
1.822
25.261
Uzbekistan
55
687
CSI 8.
309
96.794
Canada
1.075
11.741
Cuba
30
292
El Salvador
10
106
Guatemala
35
348
Messico
1.570
17.602
Trinidad e Tobago
32
451
USA
7.209
80.957
Nord America
9.961
111.497
Argentina
461
5.042
Brasile
2.773
31.380
Cile
95
1.026
Colombia
95
1.103
Ecuador
60
623
Paraguay
5
42
Perù
100
1.023
Uruguay
10
83
Venezuela
155
1.350
Sud America
3.754
41.672
Algeria
35
380
Egitto
450
6.162
Libia
80
883
Marocco
50
481
Sudafrica
600
6.596
Africa
1.215
14.502
Iran 1.459
14.969
Qatar
242
2.777
Arabia Saudita
538
5.744
EAU
237
2.176
Medio Oriente
2.476
25.666
Cina
63.300
748.670
India6.
890
76.197
Giappone
9.175
101.667
Corea del Sud
5.911
65.324
Taiwan
1.905
20.782
Asia
87.181 1.012.640
Australia
315
4.217
Nuova Zelanda
75
801
Oceania
390
5.018
Totale mondo 130.525 1.497.750
63
Atti
Attieenotizie
notizie
Italia: l’import
cresce più
Presentata
dell’export
Economia
e produzione
Economia
e produzione
Made in Steel
Produzione mondiale di ghisa
Il preridotto
corre (in migliaia
di tonnellate, fonte: WSA)
Milano
quinta edizione
più della la
ghisa
a
Nov. 2014
Genn.-Nov. ‘14
Preridotto batte ghisa 2,5 a 1,1. della
fiera sull’acciaio
502
5.547
Entrambi in crescita, ma con
tassi
Sono questi i tassi di crescita per- Austria
diversi. Questo l’andamento nei priInternazionalizzazione
e cultura.
Saranno
mi
nove mesi del 2014
dell’import
e
queste
le
bandiere
e
i
capisaldi
dell’edidell’export siderurgico italiano. Sezione 2013 di Made in Steel, la conference
condo quanto reso noto da Federac& exhibition dedicata alla filiera siderurciai,
infatti,
il commercio
nagica in
programma
dal 3 al 5estero
aprile 2013.
zionale
acciaionovità
nel periodo
preso
La primadigrande
della quinta
ediinzione
esame
si annunciata
è incrementato
è stata
duranterispetto
la confestampa
d’apertura
tenutasi
martedì
alrenza
2013,
salendo
a 13,016
milioni
di
20 marzo nella
suggestivaedcornice
milatonnellate
per l’export
a 12,626
nese della
Fonderia Napoleonica
Eugenia,
milioni
di tonnellate
per l’import.
ovvero il trasloco da Brixia Expo-Fiera di
Entrando
maggiormente nel dettaglio,
Brescia ai padiglioni di fieramilanocity,
leuna
esportazioni
sonoquindi
saliteildel
5,0% ri-a
sfida che vede
passaggio
spetto
ai
primi
tre
trimestri
dell’anno
Milano come quasi obbligato nell’ottica di
precedente,
soprattutto
proporsi per il grazie
2013 come
«The world alle
steel
expo». Durante
la presentazione
alla
vendite
nei mercati
dell’Unione Eustampache
dell’edizione
2013, dopodell’8,1%,
il saluto e
ropea,
sono aumentate
gli auguri
di rito alla
manifestazione
da
mentre
le vendite
al di
fuori dei merparte del presidente della Camera di Comcati
comunitari sono scese dell’1,5%.
mercio di Brescia Francesco Bettoni, la paPer
che concerne
le categorie
di
rolaciò
è passata
all’ad di Made
in Steel Emaprodotti,
si segnala
nettonelincrenuele Morandi.
«Siamounpartiti
2005
per sviluppare
un evento
che non
mento
dell’export
dei prodotti
di fosse
prisolotrasformazione
una fiera in senso(+10,1%
puro e semplice
ma
ma
a 3,751
che
fosse
anche
qualcosa
di
innovativo
che
milioni di tonnellate) e dei prodotti
potesse aver al centro anche la cultura - ha
piani
(+5% a 5,300 milioni di tonneldetto Morandi -. Solo una cultura forte perlate).
lunghi aumentano
le proprie
mette I innovazione
e internazionalizzavendite
dell’1,8%
arrivando
3,166
zione, senza cultura non si va daa nessuna
milioni
di tonnellate,
perdono
invece
parte. Non
nascondo che
la scelta
di lasciare Brescia
è stata sofferta
e articolata
terreno
sia i prodotti
di seconda
traper capire quale
potesse
essere la tonnelmigliore
sformazione
(-4,5%
a 298.000
piattaforma
fieristicaeinsemilavorati,
grado di accogliere
late),
sia i lingotti
che
tutto quanto fa parte di Made in Steel. La
cedono il 2,9% fermandosi a 501.000
scelta è caduta su Milano, una città in
tonnellate.
grado di attrarre acciaio ma anche di poter
Sul
versante
dell’import,
ha
arrivare
al mondo
intero. In l’Italia
questo peravuto
un flussoscelto
all’acquisto
di 12,626
corso abbiamo
fieramilanocity
perché è una
non dispersiva,
luogo
milioni
di location
tonnellate,
con una un
crescidove
sia ancora
unaalvolta
l’incontro
ta
delci6,3%
rispetto
2013.
L’incre-di
business
ma anchedelle
dove ciimportazioni
sia la possibilità
mento
maggiore
è
di dare idee e spunti per uscire dalla crisi.
avvenuto dai Paesi non facenti parte
Dico sempre che ai propri figli si debbano
dell’Ue
5,747
milioni di tondare ali (+7,3%
e radici, aradici
in un’economia
che
nellate),
mentre
i
Paesi
europei
hanno
è ormai chiaro non può prescindere
dal
aumentato
i volumi
5,4% arrivando
manifatturiero
comedel
testimonia
anche il
commissionato
dal La
premier
arecente
6,879studio
milioni
di tonnellate.
cainglese Cameron.
Ma più
anche
le ali perché
tegoria
di prodotti
importata
daisi
possa superare
propri limiti
e proseguire
clienti
italiani èi quella
dei piani
(7,668
verso nuovi traguardi».
milioni di tonnellate, +8,8%), seguiti
A dare il benvenuto alla nuova manifestadai
lingotti
semilavorati
(2,323Enrico
mizione
ancheel’ad
di Fiera Milano
lioni
di
tonnellate,
5,9%),
dai
lunghi
Pazzali: «oggi abbiamo l’opportunità di ac(1,488
di tonnellate, e+9,4%),
coglieremilioni
questa manifestazione
aiutarla
a diventare
strumento
di politica indai
prodottiuno
di prima
trasformazione
dustriale
– ha spiegato
Pazzali+15,5%)
-. Sarà una
(1,005
milioni
di tonnellate,
e
sfida
per
noi,
ma
anche
un’opportunità:
dai prodotti di seconda trasformazioalla luce della necessità di portare le none (142.000 tonnellate, +2,2%).
64
centuale che hanno fatto registrare,
rispettivamente, i due semiprodotti
nei primi undici mesi del 2014.
Secondo i dati comunicati da World
Steel Association, infatti, l’output
mondiale di preridotto, tra gennaio e novembre, si è attestato a
55,425 milioni di tonnellate, con un
aumento di 1,364 milioni di tonnellate (pari, appunto, al +2,5%) rispetto allo stesso periodo del 2013.
La ghisa, invece, nel medesimo
periodo supera il miliardo di tonnellate, arrivando a 1,083 miliardi
di tonnellate, ma con un miglioramento dell’1,1% (+12,157 milioni di
tonnellate), rispetto al corrisponstre
aziende
ad di
essere
presenti
là dove
la
dente
lasso
tempo
dell’anno
precrescita
c’è
abbiamo
sviluppato
ad
oggi
cedente. Entrando maggiormente
circa 48 fiere all’estero. La manifestazione
nel dettaglio sul comparto ghisa, si
ha la legittima aspirazione di essere glonota
che l’Asia sale Internazionaliza quota 836,2
bale
e internazionale».
milioni
di
tonnellate
menzazione rimarcata anche (+1,2%),
dal consigliere
tre l’Europa
87,5Cuzzolin:
milioni
delegato
di Made arriva
in Steel,aLuigi
«Internazionalizzare
dire portare
di tonnellate, con vuol
una crescita
del
espositori
3,5%. esteri e attrarre visitatori. Ma
non solo questo; Made in Steel vuole diL’Italiauna
è controcorrente,
una
ventare
realtà che aiuti a con
comprendere
meglio il settore
a confrontarsi.
contrazione
di oltree 400.000
tonCon
quantopari
creato
Minellate,
al negli
6,6%anni
per da
unFiera
output
lano inoltre si potrà sviluppare un sistema
complessivo di 5,890 milioni di tona due direzioni in cui poter portare l’ecnellate.
cellenza
italiana e il suo saper fare in giro
Belgio
345
3.979
Rep. Ceca
344
3.790
Francia
918
10.007
Germania
2.225
25.473
Ungheria
95
719
Italia
524
5.890
Olanda
492
5.357
Polonia
375
4.271
Slovacchia
336
3.482
Spagna
308
3.735
Regno Unito
799
8.982
Altri Ue
598
6.312
Ue a 28
7.861
87.544
Bosnia
70
805
Serbia
56
486
Turchia
808
8.536
Altri Europa
934
9.827
Kazakistan
300
3.008
Russia
4.235
46.797
Ucraina
1.636
23.012
CSI
6.171
72.817
Canada
595
6.236
Messico
430
4.756
USA
2.300
26.788
Nord America
3.325
37.780
per il mondo».
Argentina
229
2.526
A chiudere la conferenza
d’aperBrasile
2.242 stampa
24.565
tura l’augurio e il saluto
Cile
50 di Giuseppe
536PaColombia
sini, presidente di 15Federacciai:216«La
Paraguay
5
siderurgia italiana lo sappiamo
è un61
comSud
America
27.904
parto
determinate2.541
per l’economia
nazioSudafrica
420
4.637
nale, a febbraio il tasso di crescita della
Iran
224
2.552
produzione è del 53.230
+7,3% sul 2011,
anche se
Cina
654.110
alla luce delle previsioni
riduzione
India
4.380 sulla 49.164
dei consumi credo
che noi presto
Giappone
6.945
76.926anCorea
Sud ad3.857
42.886Per
dremodel
incontro
un rallentamento.
Taiwan
1.260
13.076
il futuro ci si dovrà
aprire all’internazioAsia
70.316
nalizzazione, vedendo
l’arrivo843.351
di Made in
Australia
207
2.969
Steel aZelanda
Milano come un
Nuova
60punto di partenza
618
per collegarsi con altre
legate
Oceania
267fiere estere
3.587
ai Paesimondo
in via di 91.415
sviluppo». 1.082.810
Totale
Un contenitore indipendente di notizie,
informazioni, analisi, dati, prezzi, opinioni.
Il punto di riferimento e confronto
quotidiano per gli operatori del settore.
Prova il portale
gratuitamente per 7 giorni
collegati al sito
www.siderweb.com
Siderweb SpA
Via don Milani 5
25020 Flero, Brescia ITALY
Tel. +39 030 2540006
Fax +39 030 2540041
61
STEELSIM
2015
6th International Conference
MODELLING and SIMULATION of METALLURGICAL PROCESSES in STEELMAKING
BARDOLINO, ITALY
23-25 SEPTEMBER 2015
www.aimnet.it/steelsim2015.htm
Modelling and Simulation of metallurgical processes cover an important role in optimizing technological processes, decreasing
production costs, increasing steel quality and defining the correct design of metallurgical processes in order to improve their
sustainability even from the environmental point of view.
The fundamentals of metallurgical processes can be investigated through physical and numerical modelling following several
numerical approaches.
Traditional and new mathematical techniques applied by modern simulation facilities allow to achieve results that are useful
to understand physical interaction and to design a profitable metallurgical process.
The simulations technique can be applied to the different steps of the metallurgical production route: production and refining
of liquid metals, solidification, plastic deformation, thermo-mechanical processes, thermal treatment, verification of structural
reliability etc.
OOO CONFERENCE TOPICS
State of art and developments in modeling and simulation
in steelmaking:
• Ironmaking
• Primary metallurgy (aluminium alloys, copper alloys
titanium alloys etc.)
• Secondary steelmaking
• Refining of metal alloys
• Thermodynamic and kinetic simulation of the metallurgical
systems
• Casting and solidification
• Electrochemical processes
• Metalforming processes and thermo-mechanical treatment
• Heat treatments
• Fracture mechanics and safety criteria
• Fatigue mechanics
• Safety criteria
• Reduction of environmental impact
OOO VENUE
The Conference will be staged at the Congress Center
of Aqualux Hotel Spa Suite & Terme, in Bardolino (VR), Italy
Via Europa Unita, 24/b 37011
OOO CALL FOR PAPERS
All contributions on the theme of the Conference as
described previously are welcome. Prospective authors are
invited to submit a tentative title and an abstract of 400 words
(in English) by January 31, 2015 to the Organising Secretariat.
The abstract should provide sufficient information for a fair
assessment.
OOO EXHIBITION & SPONSORSHIP OPPORTUNITIES
SteelSim 2015 will feature an Exhibition that will enable
excellent exposure for company products, technologies,
innovative solutions or services.
Companies will also be able to become Sponsors of the
Conference.
Companies interested in taking part in the Exhibition or in
sponsoring the event may contact the Organising
Secretariat (e-mail: [email protected] / fax: +39 0276020551).
OOO DEADLINES
Deadline for submission of abstracts
Information on Acceptance
Opening of the online registration
Deadline for Full Paper Submission
January 31, 2015
March 16, 2015
March 16, 2015
June 5, 2015
Organised by
ASSOCIAZIONE ITALIANA DI METALLURGIA
ORGANIZING SECRETARIAT
AIM  ASSOCIAZIONE ITALIANA DI METALLURGIA
P.le R. Morandi, 2 · 20121 Milano · Italy · tel. +39 02 76021132 · fax. +39 02 76020551
e-mail: [email protected]
www.aimnet.it

Memorie - Analisi e Calcolo

Transcript

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