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Memorie
Acciaio
On the strength of microalloyed steels
- An interpretive review
C. I. Garcia, M. Hua, K. Cho, A. J. DeArdo
In the mid-1950s, hot rolled carbon steels exhibited high carbon contents, coarse ferrite-pearlite
microstructures, and yield strengths near 300 MPa. Their ductility, toughness and weldability were poor.
Today, a half-century later, hot rolled steels can exhibit microstructures consisting of mixtures of ferrite,
bainite and martensite in various proportions. These structures are very fine and can show yield strengths over
900 MPa, with acceptable levels of ductility, toughness and weldability.
This advancement was made possible by the combination of improved steelmaking, microalloying
technology and better rolling and cooling practices. The purpose of this paper is to chronicle some of the
remarkable progress in steel alloy and process design that has resulted in this impressive.
KEYWORDS:
Accelerated cooling, direct quenching, EBSD-IQ, HSLA steel, strengthening,
thermomechanical processing, transformation
INTRODUCTION: MICROALLOYING AND STRENGTH
The year 1980 represents a benchmark in the strength of MA steels. From the early days of the 1960s to approximately 1980, the
steels being microalloyed were low hardenability steels with ferrite-pearlite (F-P) microstructures and yield strengths up to about
420 MPa (60 Ksi). These were the steels that were used to develop
the principles and interrelationships of microalloying, controlled
rolling and air cooling. They were characterized by relatively higher carbon contents and moderate Mn levels, and exhibited ferrite-pearlite (F-P) microstructures after air cooling (1).
Around 1980, both the linepipe and the automotive industries
desired strengths in excess of the 420 MPA that could be readily supplied with fine grained F-P steels. Clearly higher
strength microstructures were required. The obvious choices
were the lower temperature transformation products: matrices
comprised of non-polygonal ferrite, acicular ferrite, the bainites
and martensite, either as monoliths or as mixtures. To achieve
these microstructures, the combination of higher hardenability
and high cooling rates was required. Furthermore, much additional research was needed to reach the required goals consistently and with uniform results.
From the processing side, the solution to this dilemma was using
water cooling after hot rolling. This was accomplished in the
mid-1980s for plate processing by interrupted accelerated cooling (IAC) and interrupted direct quenching (IDQ) in plate mills.
Runout table water spray cooling to the coiling temperature in
hot strip mills had been in practice since the 1960s, but not as
C. Issac Garcia, Mingjian Hua,
Kengun Cho, A. J. DeArdo
BAMPRI (The Basic Metals Processing Research Institute)
Department of Mechanical Engineering and Materials Science - University of Pittsburgh, Pittsburgh,
Pennsylvania 15261, USA
Anthony DeArdo
Finland Distinguished Professor,
Department of Mechanical Engineering
University of Oulu, P.O. Box 4200 (Linnanmaa),
FIN-90014, Finland
Paper presented at the 3rd International Conference Thermomechanical
Processing of Steels, Padova, 10-12 september 2008, organized by AIM
la metallurgia italiana - n. 11-12/09
FIG. 1
Evolution of plate steel for large diameter linepipe:
microstructure and mechanical properties. (2)
Evoluzione dell’acciaio per lamiere destinato a
condutture di grande diametro: proprietà
microstrutturali e meccaniche. (2)
a microstructural control tool for increasing strength. This was
because of the higher carbon contents of the steels of that era.
The benefits of faster rates of cooling and lower coiling temperatures were exploited for achieving higher strengths later with
steels of lower carbon contents.
Figure 1 shows schematically how the microstructure and properties of plate steels changed over time with advances in alloy
design and processing (2).
It is obvious from Figure 1, that the accelerated cooling after rolling was largely responsible for the very high strengths attainable, practically independent of composition. With suitable cooling
practices, yield strengths in excess of 690 MPa (X100) can be
achieved in low carbon steels containing less than 2 Wt% Mn and
with C. E. and Pcm values near 0.5 and 0.2, respectively (3, 4).
One central question is what is the role of the MAE in obtaining
these strength levels? Let’s begin with the early steels (pre1980), where air cooling of plate and high coiling temperatures
of strip were used. As noted above, these were the F-P steels with
strengths up to about 420 MPa (X60) for gauges up to 18mm (0.7
inches). The most obvious contributor to strength was grain refinement, as was clearly shown by quantitative optical micro35
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Lattice Parameter, a0, nm
[100]ppt // [100] γ
[010]ppt // [010] γ
[001]ppt // [001] γ
[100]ppt // [100] α
[011]ppt // [010] α
[0-11]ppt // [001] α
NbC
NbN
VC
VN
4.4702
25.22
25.22
25.22
55.95
10.26
10.26
4.39
22.98
22.98
22.98
53.15
8.28
8.28
4.16
16.53
16.53
16.53
45.13
2.61
2.61
4.29
20.17
20.17
20.17
49.67
5.81
5.81
scopy. There is no doubt that the MAE was responsible for this
contribution through its effect on austenite conditioning. Other
contributions included solid solution strengthening by the Mn,
Si, and others, including the MAE, when retained in solution.
Equations have been published quantifying these effects, as well
(5). The other contribution to strength claimed by researchers
studying these early steels was precipitation hardening (6).
The precipitates expected to strengthen ferrite, NbCN, VCN, TiC,
TiN, all exhibit a NaCl crystal structure, and, as such, do not fit
well in the ferrite lattice. The lattice mismatch for Nb and V precipitates in both austenite and ferrite are shown in Table 1 (7).
This explains why the MA particles are always located on crystalline defects in either the austenite or ferrite (7). The misfit
strains of several percent mean that the particles cannot be coherent. The combination of incoherency with the ferrite and the
NaCl structure means that the particles must cause strengthening by the Orowan-Ashby mechanism, Eq. 1 (8-10).
(1)
This mechanism of strengthening for the Orowan process is by: (a)
the energy required for dislocations to bow between particles, and
(b) the energy required by the cross slipping of screw segments
FIG. 2
The dependence of precipitation strengthening on
precipitate size (X) and fraction according to the
Ashby-Orowan Model, compared with experimental
observations for given microalloying additions. (10)
Dipendenza dell’ indurimento per precipitazione dalla
dimensione (X) e dalla frazione in volume dei precipitati
secondo il modello Ashby-Orowan, confrontata con
osservazioni sperimentali a seguito di aggiunte mirate di
microalliganti. (10)
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TAB. 1
Lattice mismatch for MX
precipitates in austenite and
ferrite, % Austenite: fcc, a0 =
0.35698 nm; Ferrite: bcc, a0
= 0.28664 nm.
Differenze dei parametri
reticolari per precipitati MX
nell’ austenite e nella ferrite,
% Austenite: fcc, a0 = 0.35698 nm;
Ferrite: bcc, a0 = 0.28664 nm.
or climb of edge dislocation segments in bypassing the particles
(9). The predicted increase in YS caused by this mechanism depends on the amount or volume fraction of precipitate and the
size of the particles. This is shown for particles in ferrite in Figure 2 by Gladman for several precipitation systems (10).
The data of Figure 2 must be used with caution, however. First,
the volume fraction or amount used in the calculation is that actually consumed in forming the particles. The amount lost to the
austenite in rolling and the amount remaining in solid solution
do not contribute to the strength shown in Figure 2 and must be
subtracted from the total. Second, the precipitates participating
are those actually present in the steel, and, therefore, need time
to form in the ferrite. An early example of the slow kinetics for
this precipitation was shown by Honeycombe and Sakuma (11,
12), Figure 3, and later confirmed by Thillou, et al. (13)
Third, claims of precipitation hardening and the application of
Figure 2 should be independently verified by thin foil TEM.
Fourth, the mere presence of fine particles in ferrite does not
guarantee precipitation hardening. The distribution must conform to the Orowan-Ashby model to justify claims of a certain
level of strengthening. Finally, it must be recognized that isothermal laboratory experiments do not necessarily predict the
behavior of continuously cooled commercial steels, even when
the compositions are similar.
Plots of the Orowan-Ashby equation as viewed from what must
be observed in thin foil TEM are shown in Figure 4 (14,15).
These plots show what precipitate distributions must be present,
viz. measured, to claim 10, 50 and 100 MPa increments in YS
caused by precipitation hardening. Superimposed are reasona-
FIG. 3
Schematic TTT curves for Fe-0.036Nb-0.09C and Fe0.036Nb-0.09C-1.07Mn alloys. Interphase
precipitation (IP) occurs in certain shaded areas.
Curve TTT schematiche per leghe Fe-0.036Nb-0.09C
eFe-0.036Nb-0.09C-1.07Mn. La precipitazione interfase
(IP) si verifica entro le aree tratteggiate.
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that found by solution hardening and bake hardening or strain
aging (16), and is nowhere near what can be found by dislocation
or substructure strengthening of ferrite (16-18).
In discussions of precipitation hardening in MA steels, the cooling path from the finish rolling temperature is critical to the
formation of strengthening particles (7, 11, 12). While air cooling at 1°C/sec from 750-600°C in 150 sec. might be slow
enough to form strengthening precipitates, Figure 5 (19), water
spray cooling through the temperature range 750-600°C at 1050°C/sec is probably much too fast to form effective particle distributions in rolled and cooled steels.
In summary, the major role of the MAE in strengthening the pre1980 steels was mainly by grain refinement. To this was probably added some solid solution and dislocation strengthening.
The contribution by precipitation hardening is not zero, but has
been probably over estimated. This was pointed out in several
early studies (20, 21).
The strength of these steels can be understood by Equation (2),
the expanded Hall-Petch equation.
(2)
FIG. 4 Particle dispersion characteristics for precipitate
strengthening according to Orowan-Ashby theory.
(14,15)
Caratteristiche della dispersione delle particelle
all’origine dell’ indurimento per precipitazione secondo
la teoria Orowan-Ashby. (14, 15)
FIG. 5
Interphase precipitation of NBCN in ferrite in steel
containing .09%C - .07%Nb. Specimen reheated at
1250°C, rolled at 1000°C, and air cooled to RT. (19)
Precipitazione interfase di NbCN nella ferrite di un in
acciaio contenente .09%C - .07%Nb. Campioni
riscaldati a 1250° C e raffreddati in aria fino a
temperatura ambiente. (19)
ble levels of both particle size and volume fractions. The predictions of Figures 4 and 5 are fully consistent with the data of
Figure 2. The Gladman diagram shown in Figure 2 represents
the maximum strengthening increments that can be expected
when conditions for precipitation are ideal, i.e. full precipitation
of available components. Again, the abscissa in Figure 2 is the
volume fraction actually formed, not what is predicted from the
bulk composition. In commercially processed hot rolled steels, it
is extremely rare that increments caused by precipitation hardening exceed 50-80 MPa (16). This magnitude is comparable to
where YSobs is the observed yield strength; YSP-N, YSSS, YStexture,
YSdisl, YSpptn are the stress increments caused by lattice friction
(Peierls-Nabarro), solid solution, texture, dislocations, and precipitation; and Kyd–1/2 represents the contribution by the ferrite
grain size.
For the F-P steels of the 1970s, the dominant contribution to
strength was Kyd–1/2, with much smaller contributions from YSSS,
YSdisl, and YSpptn.
MODERN STEELS
As was noted earlier, in the 1980s there was a large emphasis on
increasing the strength from the 420 MPa (API-X60) level to over
490 (API X-70). On the process side, this challenge was met by
lowering the transformation temperature of the austenite during
the cooling after hot rolling. On plate mills, this was accomplished by interrupted accelerated cooling (IAC) and later by interrupted direct quenching (IDQ). On strip mills it was achieved
by increasing the cooling rate and lowering the coiling temperature. In plate rolling, controlled rolling followed by air cooling
has been termed TMP, while controlled rolling followed by IAC
or IDQ has been called TMCP in some quarters (22).
It is well known that accelerated cooling can increase the
strength of F-P steels by reducing the ferrite grain size, as shown
in numerous studies (20,23). What is less clear is that rapid cooling leading to refined polygonal ferrite also leads to higher ferrite grain center hardness, as shown by Morikawa and
Hasegawa, Figure 6 (24).
The 0.15C-0.66Mn steel used in Figure 6 showed that the ferrite
grain center hardness, viz., a volume not thought to be strongly
influenced by grain boundaries or grain refinement, increased
substantially with cooling rate from about 100Hv at 1°C/sec to
near 140Hv at 100°C/sec. This increase was attributed to higher solute C and excess dislocations present in the rapidly cooled ferrite. No bainite was observed until cooling rates exceeded
25°C/sec. in this experiment. This extra strengthening was attributed to the combination of higher solute carbon levels trapped in rapidly cooled ferrite and to higher dislocation densities.
MULTI-PHASE MATRIX MICROSTRUCTURES
The equally important change with cooling rate involves the matrix microstructure. It is obvious from Figure 1 that the difference between the 420 and >490 MPa grades is the nature of
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Memorie
the matrix (2). As noted earlier, the 420 MPa grade shows polygonal ferrite formed at high transformation temperatures, over
perhaps 600°C. The matrix in higher strength steels shows mixtures of ferrite and bainite and/or martensite, in different proportions. In general, the higher the proportion of bainite and
martensite, the higher the strength of the steel.
As the strength level increases, the steels change character from
monolithic ferrite to complex mixtures starting with ferrite-bainite, ferrite-martensite, monolithic bainite, and finally monolithic martensite. With mixed microstructures, the steels appear
to follow the Rule of Mixtures, as shown by Davies in Figure 7 for
ferrite-martensite mixtures found in DP automotive steels (25).
The influence of MAE on the transformation characteristics of controlled rolled and cooled steels can be profound, especially at higher rates of cooling. An example of this effect is shown in Figures
8-10 (7,26) for transformation start temperatures, resulting microstructures and final mechanical properties, respectively (7, 26).
FIG. 6
Effect of cooling rate on strengthening factors of
steel 1. (24)
Effetto della velocità di raffreddamento sui fattori di
indurimento dell’acciaio 1. (24)
FIG. 7
The 0.2% flow stress and the tensile strength
as a function of percent martensite for Fe-Mn-C
alloys. (25)
Limite di snervamento allo 0,2% e carico di rottura in
funzione della percentuale di martensite per le leghe FeMn-C. (25)
FIG. 9
Effects of Nb, V and Ti on volume fraction of bainite
and ferrite grain size in accelerated cooled steels
(7, 26).
Effetti di Nb, V e Ti sulla frazione in volume della diverse
dimensioni dei grani di bainite e ferrite in acciai
sottoposti a raffreddamento accelerato (7, 26)
FIG. 8
Corrected Ar3 temperatures
of microalloyed steels with
standard austenite grain size
of 100µm (7, 26).
Temperature Ar3 corrette per
acciai microlegati con
dimensione standard del grano
austenitico (100mm) (7, 26).
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FIG. 10 Effects of addition of Nb. V and Ti on tensile strength
and Charpy V 50% FATT of (1) air cooled, (2)
accelerated cooled and (3) direct quenched steels
after controlled rolling (7, 26).
Effetti dell’aggiunta di Nb, V e Ti sulla resistenza a trazione
e sulla 50% FATT Charpy V di acciai (1) raffreddati ad aria,
(2) sottoposti a raffreddamento accelerato e (3) temprati
direttamente dopo laminazione controllata (7, 26)
The positive synergy between the MAE and accelerated cooling
is significant. Figure 10 reveals that the addition of 0.04 wt%
Nb to the base steel adds about 10% to the strength after air cooling; accelerated cooling with Nb adds about 40% and DQ with
Nb adds about 76%. Other examples show how the addition of
Nb, Figure 11 or V, Figure 12, to a 0.07C - 1.55Mn - 0.018Ti reference steel has little effect on the final microstructure after air cooling but a large effect after
accelerated cooling, especially on strength (27).
MONOLITHIC MICROSTRUCTURES
It is well-known that low carbon and ultra-low
carbon bainitic and martensitic ferrite can exhibit remarkable properties. Yield strengths in excess of 850 MPa (X120) in 12-18mm plate and
strip have been achieved in MA steels processed
using TMCP (3, 4). Two obvious questions are: (i)
what can cause the strength to essentially double
from the early 350-420MPa grades to the newer
700-850MPa grades, and (ii) what is the role, if
any, of the MAE?
It is well known that the strength of bainite and
martensite is controlled mainly by the carbon
content and the Bs or Ms temperature (28, 29),
Figure 13. The data of Figure 13 were generated
with ULCB plate steels with rich chemistries intended for heavy gauge applications (29).
Attempts have been made to relate the properties of bainite to its microstructure, but with limited success (30-32). With falling temperature
and increasing strength, the sequence of upper
bainite, granular bainite, then lower bainite is
often observed. Although these microconstituents have different appearances in the OM,
their real microstructure must be revealed by
thin foil TEM. Since this is a very tedious and expensive proposition, little of this work has been
done. What is clear from the available literature
is that with falling Bs temperature, both the soluble carbon content and the dislocation density
increase. This is why the strength increases with
falling transformation temperatures. The main
reason that the soluble carbon content increases
with falling temperature can be related to the sloping upper ferrite solvus and T0 lines on the FeFe3C phase diagram and the Hultgren
extrapolation that exist in the absence of cementite (33, 34). They predict that in the absence
of equilibrium, viz. presence of Fe3C, higher cooling rates will lead to higher carbon contents
with falling temperature. This, together with the
large solution hardening capability of carbon and
the concomitant increase in the dislocation density resulting from the combination of the volume change and the shear-type nature of the
transformation will combine to result in higher
FIG. 11 Effect of Nb content on properties of plate (27). CR denotes
controlled rolling and air cooling; ACC denotes controlled rolling
and accelerated cooling.
Effetti del contenuto di Nb sulle proprietà dei laminati (27). CR indica
laminazione controllata e raffreddamento ad aria; ACC indica
laminazione controllata e raffreddamento accelerato.
FIG. 12 Effect of V content on properties of plate (27)
Effetto del contenuto di V sulle proprietà dei laminati (27).
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Memorie
strengths with falling temperature.
One way to characterize the microstructure, one that avoids the
confusion and complexity of labeling and characterizing the observed ultra-fine detail, is through the use of EBSD-IQ, a technique recently applied to microstructure assessment in advanced
steels (35). In the EBSD–IQ technique, once the specimen surface and beam stability effects have been eliminated, the quality
of the final diffraction peaks coming from the Kikuchi lines is
measured, processed and quantified (35). Highly elastically distorted lattices yield low IQ peaks since the diffraction profile is
smeared, not unlike line broadening due to elastic strain in x-ray
diffraction (36). Near-perfect lattices yield high IQ peaks, because the peaks in the diffraction pattern are narrow and sharp.
Since both solute elements, viz. C and N, and dislocations as well
as precipitates might contribute to lattice strain, the EBSD-IQ
approach to understanding the strength of bainite and martensite appears to be promising.
Recent work using the EBSD-IQ technique has shown that multiphase microstructures can be characterized and quantified
using this approach. In this technique, the EBSD-IQ data are first
processed and then plotted using the Multi-Peak Software (35).
The resulting plots show a spectrum of multiple peaks where
the peak height is proportional to the volume fraction of that microconstituent and the location on the abscissa is related to the
inverse of the lattice distortion. As noted above, this distortion
is assumed to come from the combination of lattice strain caused by the dislocation density, solute content and particles.
Typical examples are shown for studies involving HSLA strip, Figures 14 and 15 (35,37, 38), DP steel, Figure 16 (35,39), TRIP-assisted steels (40), heat treated seamless pipe, Figure 17 (38) and
bainitic plate steels, Figure 18 (41). It is clear that this technique
can discern details of the final microstructure, including the
components of multi-phase mixtures.
The first example of applying the EBSD-IQ technique to multiphase microstructures is to an HSLA hot band structure with an
optical microstructure as shown in Figure 14. Analyzing this
complex microstructure using the EBSD-IQ approach resulted if
the multi-peak profile shown in Figure 15. Notice the several
forms of ferrite present in the microstructure.
The next example is a DP steel where the amounts of ferrite and
martensite were measured in three ways; by point counting, by
image analysis and the third by EBSD-IQ. The IQ results are
shown in Figure 16. The phase balance values determined with
the three approaches fell within a few percent (35,39).
Another example is heat treated seamless pipe that shows a mixture of autotempered and untempered martensite after tempering.
The EBSD-IQ analysis of this steel is shown in Figure 17 (38).
Finally, the EBSD-IQ approach has been applied to bainitic steel
plates. The IQ analysis of this study revealed the multi-peak result shown in Figure 18 (41).
When considering the strength of the bainite or martensite, either as a monolithic matrix or as part of mixed microstructures,
certain aspects must be considered. First, there are relatively
few high angle boundaries present in the microstructure. Therefore Hall-Petch strengthening will not be important. Second,
very little precipitation hardening can be expected in these rapidly cooled steels, at least in the as-cooled condition (42).
Hence, the strength of the bainite and/or martensite will be go-
FIG. 14 Optical micrograph of a HSLA steel hot band.
Etched with 2% Nital. (35,37)
Micrografia ottica di un nastro a caldo in acciaio HSLA.
Attacco Nital 2%. (35, 37)
FIG. 13 Comparison of measured and calculated strength
values for a given Bs temperature. (29)
Confronto dei valori di resistenza meccanica misurati e
calcolati per una determinata temperatura Bs. (29)
40
FIG. 15 The IQ analysis of the HSLA hot band
microstructure shown in Figure 14 using the MultiPeak model. (35,38)
Analisi IQ della microstruttura del nastro a caldo in
acciaio HSLA mostrato in Fig. 14 mediante modello
Multi-Peak (35, 38)
Acciaio
verned by the contributors to lattice distortion, viz. solutes, dislocations and their interaction. Early examples of these strong
dislocation effects were shown by Smith and Honeycombe (17),
Mangonon and Heitmann (18), and Repas (43).
The influence of dislocation density or subgrain size on strength
can be very much higher than those from solid solution and precipitation, where increments of 30-80 MPa are typical. Manganon and Heitmann have shown that substructure hardening of
ferrite can easily exceed 200MPa, Figure 19 (18). This is a contribution that can also be expected in bainite and cannot be overlooked.
In summary, when the strength of bainite is considered through
the lens of the expanded Hall-Petch equation, the contributions
from YSdisl and YSss are probably most important and the contributions from high angle grain boundaries and precipitation
are secondary.
have nearly doubled in strength while still maintaining adequate, if not superior levels of other important properties such
as toughness, weldability, ductility, formability, etc. These improvements have been facilitated by evolutions of steelmaking,
rolling and cooling practices. The details of the improvements
have been chronicled and are largely understood. Perhaps the
main message learned over the past 50 years, or so, is that microstructural improvement and optimization are often not a simple extrapolation of the old to the new. Sometimes we need a
new box.
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2.
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CONCLUSIONS
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FIG. 16 IQ analysis of DP steel microstructure using multipeak model. (35)
Analisi IQ di una microstruttura di acciaio DP mediante
modello Multi-Peak (35)
FIG. 17 The Image Quality (IQ) distribution analysis of the
specimen cooled by 10°C/sec, A508 Gr4N steel.
Analisi di distribuzione con Image Quality (IQ) del
provino di acciaio A508 Gr4N raffreddato a 10°C/s.
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41
Memorie
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Abstract
Rassegna sulla resistenza
degli acciai microlegati
Parole chiave: acciaio, lavorazioni plastiche, proprietà
A metà degli anni 1950, gli acciai al carbonio laminati a caldo
avevano un alto contenuto di carbonio, microstrutture a
grossi grani di ferrite-pearlite, e un limite di snervamento
intorno a 300 MPa. La loro duttilità, tenacità e saldabilità
erano scarse. Oggi, mezzo secolo più tardi, gli acciai laminati a caldo possono presentarsi con microstrutture costituite da combinazioni di ferrite, bainite e martensite in varie
proporzioni. Queste strutture sono molto fini e sono in grado
di esibire una resistenza allo snervamento superiore a 900
MPa, con livelli accettabili di duttilità, tenacità e saldabilità.
Questo sviluppo è stato reso possibile dalla combinazione dei
miglioramenti nella produzione dell’ acciaio, nella tecnica di
microalligazione e nelle procedure di laminazione e raffreddamento. Lo scopo di questo documento è quello di riportare
una cronaca di alcuni dei maggiori progressi nell’alligazione
dell'acciaio e nella progettazione dei processi, che hanno portato a questo straordinario risultato.
Memorie
Acciaio
Microstructure and mechanical properties
of thermomechanically processed TRIP steel
I.B. Timokhina, P.D. Hodgson, H. Beladi, E.V. Pereloma
The strengthening mechanism responsible for the unique combination of ultimate tensile strength and
elongation in a multiphase Fe-0.2C-1.5Mn-1.2Si-0.3Mo-0.6Al-0.02Nb (wt%) steel was studied. The
microstructures with different volume fractions of polygonal ferrite, bainite and retained austenite were
simulated by controlled thermomechanical processing. The interrupted tensile test was used to study the bainitic
ferrite, retained austenite and polygonal ferrite behaviour as a function of plastic strain. X-ray analysis was
used to characterise the volume fraction and carbon content of retained austenite. Transmission electron
microscopy was utilised to analyse the effect of bainitic ferrite morphology on the strain induced transformation
of retained austenite and retained austenite twinning as a function of strain in the bulk material. The study has
shown that the austenite twinning mechanism is more preferable than the transformation induced plasticity
(TRIP) mechanism during the early stages of deformation for a microstructure containing 15% polygonal ferrite,
while the transformation induced plasticity effect is the main mechanism when there is 50% of polygonal ferrite
in the microstructure. The bainitic ferrite morphology affects the deformation mode of retained austenite during
straining. The polygonal ferrite behaviour during straining depends on dislocation substructure formed due to
the deformation and the additional mobile dislocations caused by the TRIP effect. Operation of TRIP or
twinning mechanisms depends not only on the chemical and mechanical stability of retained austenite, but also
on the interaction of the phases during straining.
KEYWORDS: Transformation induced plasticity steel, thermomechanical processing, retained austenite,
TRIP/TWIP effects, transmission electron microscopy, atom probe tomography
INTRODUCTION
The demand for high strength and high formability steels has
recently increased. These steels have found application in the
manufacture of automotive wheels, certain brackets and, potentially, of high strength drawn bars. Multiphase steels, containing
austenite and bainite, represent a new class of steel with improved strength-ductility balance. The Transformation Induced
Plasticity (TRIP) effect has been widely cited to be solely responsible for this balance [1, 2]. However, mechanical twinning
can also occur in a steel alloyed with manganese, silicon and
aluminium [3, 4]. This could lead to an increase in plasticity
through the Twinning Induced Plasticity (TWIP) effect. The main
aim of previous investigations has been to obtain the maximum
amount of stable retained austenite [5, 6]. However, the current
state of knowledge regarding the multiphase steels has revealed certain contradictions to this concept. Firstly, the distribution
of carbon within the retained austenite crystals is inhomogeneous and depends on the position of these crystals in the multiphase structure. This leads to the formation of retained austenite
crystals with different carbon content [7]. It has been suggested
that only the retained austenite with an optimum carbon content can provide the TRIP/TWIP effect and improve the elongation [8]. Furthermore, an increase in the volume fraction of the
retained austenite leads to a decrease in the average carbon of
this phase, thereby reducing its chemical stability. Hence, the
I.B. Timokhina, P.D. Hodgson, H. Beladi
Deakin University, Australia
E.V. Pereloma
The University of Wollongong, Australia
optimum volume fraction of the retained austenite is needed to
provide the TRIP/TWIP effect [8].
The size of the retained austenite also affects the stability. Coarse retained austenite blocks have lower stability than films,
for example, and tend to transform to martensite at low strain.
Hence, retained austenite only with optimum size can provide
the TRIP effect [6]. It has been suggested [9, 10] that there is
another mechanism responsible for the unique strength-ductility balance in multiphase steels in addition to the TRIP/TWIP effects. Recent publications have revealed the importance of the
effect of all phases formed in the microstructure and their interaction during straining [9, 10].
A multiphase microstructure has usually been generated by a
two stage intercritical annealing due to the sensitivity of the microstructure to the thermomechanical processing approach. In
the current approach, however, thermomechanical processing
was used to avoid the extra step required by the intercritical annealing and develop the desirable microstructure directly after
hot rolling.
The aim of the current research is to study the effect of the volume fraction of the phases on the structure-property relationship and the complex interrelationship between the phases
during the formation of the final microstructure.
EXPERIMENTAL PROCEDURE
Steel with composition of Fe-0.2C-1.5Mn-1.2Si-0.3Mo-0.6Al0.02Nb (wt%) was studied. A laboratory rolling mill was used to
simulate rolling. The thermomechanical processing schedule
was constructed based on analysis of the continuous cooling
transformation data [11], to form 15% and 50% of polygonal ferrite and non-carbide bainitic ferrite to stabilize the retained austenite at room temperature (Fig. 1).
43
Memorie
FIG. 1
Thermomechanical processing schedule.
Schema del processo termomeccanico.
The samples with initial thickness of ~ 35mm were austenitized at 1250°C for 120s in a 15kW muffle furnace and then rolled at 1100°C, where the first deformation (ε1=0.25) took place,
followed by the second deformation (ε2=0.47) at 875°C (Fig. 1).
After that, the samples air cooled at ~1Ks-1 to the accelerated
cooling start temperatures (TA) of 780°C and 760°C to form
15% and 50% of polygonal ferrite respectively. Two spray guns
were used to cool the samples at ~20Ks-1 to 520°C to avoid pearlite formation and after that the samples were placed in a
fluidbed furnace and covered with aluminum oxide sand to
hold the samples at 470°C for 1200s to form non-carbide bainite. After holding the samples were quenched in an iced brine
solution (Fig. 1). The final thickness of the slab after processing was 7 mm.
The microstructure of the samples was characterized using optical metallography, transmission electron microscopy (TEM)
and atom probe tomography (APT). Thin foils for TEM were prepared by twin-jet electropolishing using 5% of perchloric acid in
methanol at -25°C and an operating voltage of 50V. Bright and
dark-field images and selected area electron diffraction patterns
were obtained using a PHILIPS CM 20 microscope operated at
200kV. The stability of retained austenite and the transformation behavior of the phases as a function of the plastic strain
were studied on the samples after interrupted tensile testing
using TEM.
APT analysis was performed to study the microstructural features formed after TMP, such as carbon distribution within the
phases, formation of particles, etc. The standard two-stage electropolishing procedure was used to prepare the atom probe specimens [12]. The local electrode atom probe was operated at a
pulse repetition rate of 200 kHz, a 20% pulse fraction with a sample temperature of 80K. Iso-concentration surfaces were used
for easier visualization of the phases and carbides.
X-ray diffraction (XRD) analysis was performed using a PHILIPS
PW 1130 (40kV and 25mA) diffractometer equipped with a monochromator and CuKα radiation to calculate the volume fraction of retained austenite after TMP and for the samples after
different strains. The integrated intensities of the (200)α, (211)α,
(200)γ and (220)γ peaks were used in the direct comparison method [13].
Room-temperature mechanical properties were determined
using an Instron 4500 servohydraulic tensile-testing machine
with a 100kN load cell. Subsize samples with a 25mm gage
length were used to minimize the amount of material.
44
RESULTS AND DISCUSSION
Structure-Property Relationship after TMP.
The microstructures after laboratory rolling consisted of 15±3%
ferrite (hereafter called “Steel 1”) and 50±4% ferrite (hereafter
called “Steel 2”), with 16.5±3% and 12±3% retained austenite correspondingly and remaining non-carbide bainite and martensite. The average ferrite grain size was 2.4±0.5µm for the Steel
1 and 6.0±0.5µm for the Steel 2 (Figs. 2 a, b). The average carbon content of retained austenite measured by X-ray was 1.8wt%
for Steel 1 and 1.6wt% for Steel 2.
TEM of Steel 1 revealed the formation of two bainitic morphologies: (i) granular and (ii) acicular. Granular bainite is characterized by the presence of coarse bainitic ferrite plates with
isolated crystals of retained austenite in between (Figs. 3 a, b).
Some of the retained austenite crystals showed twinning and the
retained austenite/twinned austenite constituent islands were
also present in the microstructure (Figs. 3 a, b). The acicular bainite/ferrite structure appeared to be a bainitic structure with retained austenite layers between bainitic ferrite laths (Fig. 3 c).
The thickness of the bainitic laths varied from 0.1 to 0.5µm. The
retained austenite laths had a wide range of thickness, from very
thin retained austenite films to thick retained austenite laths,
which shown in some cases twinning (Fig. 3 d). The retained austenite crystals at the polygonal ferrite/bainite interface were
not observed. It is interesting to note that the clusters of bainitic ferrite laths were oriented in different directions and in some
a
b
FIG. 2
Optical micrographs of Steel 1 (a) and Steel 2 (b).
Micrografie ottiche di Steel 1 (a) e di Steel 2 (b).
Acciaio
a
b
c
d
e
f
g
FIG. 3
TEM micrographs of Steel 1 after TMP: bright (a) and (b) dark field image of
granular bainite with twinned austenite (zone axis is [110]γ), (c) acicular ferrite,
(d) twinned austenite (zone axis is [110]γ), (e) bright and (f) dark field images of
bainitic ferrite laths oriented perpendicular to each other, arrows show the Fe3C
carbides, (g) lenticular bainitic ferrite, arrows indicate carbides. RA is retained
austenite, BF is bainitic ferrite, and TA twinned austenite.
Micrografie TEM di Steel 1 dopo processo termo meccanico: (a) immagine in campo
chiaro e (b) scuro della bainite granulare con austenite geminata (asse di zona
[110] γ), (c) ferrite aciculare, (d) austenite geminata (asse di zona [110] γ), (e)
immagine in campo chiaro e (f) scuro di lamelle di ferrite bainitica orientati
perpendicolarmente tra loro, la freccia indica i carburi di Fe3C, (g) ferrite bainitica
lenticolare, la freccia indica i carburi. RA indica l’austenite residua, BF la ferrite
bainitica e TA l’austenite geminata.
cases perpendicular to each other (Figs. 3
e, f). Rounded Fe3C carbides were observed within these laths (Figs. 3 f). Bainitic
ferrite laths with a lenticular shape and
an average thickness of 0.5µm and with
fine, plate-like Fe3C carbides were also observed in the microstructure (Fig. 3 g).
Martensite crystals were not found during
TEM observation.
TEM of Steel 2 also showed the formation
of two types of bainite with bainitic ferrite,
one in the form of parallel thin laths with
an average thickness of 0.6µm and the
other in the form of plates (Figs. 4 a, b).
Most of the retained austenite was present
as small islands, although coarse blocks
of retained austenite were also found in
the vicinity of the martensite (Fig. 4 b). A
number of the retained austenite crystals
a
b
Fig. 4
TEM micrographs of Steel 2 after TMP: (a) acicular ferrite and (b) granular
bainite. BF is bainitic ferrite, RA is retained austenite and M is martensite.
Micrografie TEM di Steel 2 dopo processo termo meccanico: (a) ferrite
aciculare e (b) bainite granulare. BF indica la ferrite bainitica, RA l’austenite
residua, e M la martensite.
45
Memorie
TAB. 1
Phase compositions
calculated using APT, (at%).
Composizioni delle fasi
calcolate mediante APT,
(% atomico).
(at%)
PF
Steel 1
BF
C
Mn
Si
0.04±0.02
1±0.2
2.7±0.5
0.4±0.2
1.3±0.5
3.0±0.2
RA
PF
Steel 2
BF
2.4±0.7 0.02±0.001 0.25±0.03
1.56±0.07 0.75±0.02
1.85±0.1
3.4±0.1
2.77±0.05
2.09±0.1
RA
2.71±0.07
1.03±0.04
4.01±0.08
FIG. 5
Representative atom maps
of C (a, d), Mo (b) and Nb (c)
showing Nb-Mo-C carbides
in retained austenite (a, b, c)
and different phases in Steel
1(d). PF is polygonal ferrite,
RA is retained austenite, BF
is bainitic ferrite.
Mappe della distribuzione degli
elementi, rappresentative di C
(a, d), Mo (b) e Nb (c) che
mostrano i carburi Nb-Mo-C
nell’ austenite residua (a, b, c)
e nelle diverse fasi entro Steel
1(d). PF rappresenta la ferrite
poligonale, RA l’austenite
residua e BF la ferrite bainitica.
showed partial decomposition to martensite. Coarse blocks of
martensite were found between the bainitic ferrite laths and at
the polygonal ferrite/bainite interface. Twinned austenite was
not observed by TEM in this steel.
The APT study showed formation of sphere-like Nb carbides and
Nb-Mo carbides in the retained austenite and Nb carbides in the
bainitic ferrite for both steels. The average size of the particles
was 10±1nm (Fig. 5). The composition of phases calculated using
APT is shown in Table 2. The carbon concentration in polygonal
ferrite and in bainitic ferrite of Steel 1 was higher than in Steel
2 due to the difference in cooling schedules, which affected the
temperature intervals for phase transformations. It leads to
lower carbon content of retained austenite in Steel 1 compared
to Steel 2 (Table 1). The detailed explanation of the solute distribution within the phases in Steel 2 was reported elsewhere
[11]. Fe3C and Fe4C carbides were also observed in the microstructures of both steels using APT (Fig. 5d)
The microstructures formed after TMP control the combination
of strength and ductility in the TRIP steel. While the presence
of ferrite and retained austenite leads to high elongation, martensite and bainite are responsible for strength. The Steel 1 had
a higher ultimate tensile strength (UTS) 1300±20MPa and yield
strength (YS) 600±30MPa than Steel 2 with UTS of
1000±40MPa and YS of 400±40MPa, while the total 25±3% and
uniform 17±3% elongations of Steel 1 were lower than Steel 2,
with a total elongation of 29±2 and uniform elongation of 23±1%
(Fig. 6). The lower elongation in the Steel 1 could be due to the
lower volume fraction of polygonal ferrite. On the other hand,
the higher volume fraction of retained austenite in Steel 1
should lead to higher elongation. In order to understand structure-property relationship in these steels, the behaviour of the
microstructures during straining was studied using interrupted tensile tests.
46
FIG. 6
Representative true stress-strain curves of Steel 1
and Steel 2.
Curve rappresentative del rapporto reale caricodeformazione di Steel 1 e Steel 2.
Microstructural Behavior under Applied Strain.
X-ray analysis of Steel 1 after ~0.04 strain showed a decrease in
the retained austenite volume fraction from 16.5±3% to 11±2%,
which remained unchanged up to ~0.08 strain. Further decrease
in the retained austenite volume fraction to 5% was observed at
~0.17 of strain. TEM revealed extensive twinning of the retained austenite crystals after a strain of 0.04 (Fig. 7a) with further development of this structure at a strain of 0.08. An
increase in strain to 0.17 led to the formation of coarse marten-
Acciaio
a
b
c
d
FIG. 7
TEM micrographs of austenite twinning at 0.04 (a) (zone axis is [114]γ) and 0.17
(b) of strain, and formation of cell dislocation structure in ferrite at 0.04 (c) and
0.17 (d) of strain in Steel 1.
Micrografie TEM della geminazione dell’ austenite: (a) dopo 0.04 di deformazione
(asse di zona [114]γ) e (b) dopo 0.17 di deformazione; formazione della struttura di
dislocazioni a cella nella ferrite: (c) dopo 0.04 di deformazione e(d) dopo 0.17 di
deformazione in Steel 1.
site crystals, although twinned austenite crystals were still observed (Fig. 7b). This suggests that austenite twinning is the preferred deformation mechanism at low strains in Steel 1. It
appeared that the formation of the higher volume fraction of bainite is responsible for this behavior, i.e. during straining the bainitic ferrite laths could accommodate the stress and prevent
transformation of retained austenite to martensite and, thus. promote the formation of austenite twinning at the early stages of
straining.
Most of the retained austenite transformed to martensite at an
intermediate strain level (0.17) due to an increase in the dislocation density of bainitic ferrite and interaction between the rigid
bainitic ferrite laths and retained austenite. This leads to reduction of total elongation in Steel 1. Polygonal ferrite showed the
partial formation of dislocation cells after a strain of 0.04 and parallel deformation bands after a strain of 0.17 (Figs. 7 c, d).
The transformation of the retained austenite to martensite during straining of Steel 2 occurred gradually – ~8% of retained
austenite and ~12% of martensite at the ~0.08 strain; ~5% of retained austenite and ~15% of martensite at ~0.27 strain. The
preferred deformation mechanism for retained austenite at all
strains was TRIP effect (Figs. 8 a and b). However, ~4-5% of retained austenite was trapped between the bainitic ferrite laths
and remained in the microstructure of the fractured tensile sample (Fig. 8 c).
This behavior could be explained by the combined effect of the
inhomogeneous carbon distribution within the retained austenite and the effect of the size of the retained austenite on its stability. The carbon distribution within the retained austenite is
not homogeneous and some coarser islands of retained austenite were less enriched than smaller ones. These coarse blocks
of austenite tend to transform to martensite at a lower strain.
The microstructure of Steel 2 contained a high volume of relatila metallurgia italiana - n. 11-12/09
vely coarse austenite crystals, which did not contribute significantly to the TRIP effect. On the other hand, as a result of the
strain-induced transformation of high numbers of austenite
blocks stress transfers to the soft ferrite matrix leading to dislocation strengthening of the neighbouring regions, which, in
principal, can improve strength-ductility balance. The fine islands of austenite that are trapped between the plates of bainitic ferrite in a sheaf are much more stable because of the higher
carbon concentration and also because of the physical constraint
to transformation due to the close proximity of plates in all directions [14]. The contribution of their strain-induced transformation to the improved ductility is higher than the contribution
of coarse crystals. On the other hand, a number of supersaturated retained austenite crystals remained in the microstructure
after fracture and did not contribute to an increase in elongation.
CONCLUSIONS
The analysis of microstructure-property relationships in thermomechanically processed multiphase steels with different
amounts of phases has been conducted. The results have shown
that the strengthening mechanism in these complex multiphase
microstructures is determined not only by the amount of retained austenite but also by the volume fraction of other phases in
the microstructure and their interaction during deformation.
ACKNOWLEDGEMENTS
The authors would like to acknowledge Professor S.P. Ringer
from Australian Key Centre for Microscopy and Microanalysis
for providing access to the local electrode atom probe. One of
the authors (PDH) acknowledges the support of the ARC Federation Fellowship scheme.
47
Memorie
a
FIG. 8
c
b
TEM micrographs of partial transformation of retained austenite to martensite at 0.02 strain, zone axis is [310]γ (a),
partial transformation of retained austenite to martensite at 0.08 strain, zone axis is [110]γ (b) and retained austenite
island after fracture, zone axis is [116]γ (c).
Micrografie TEM della trasformazione parziale in martensite dell’austenite residua (a) dopo 0.02 di deformazione (asse di
zona [310]γ) , (b) trasformazione parziale in martensite dell’austenite residua dopo 0.08 di deformazione (asse di zona
[110]γ) e (c) isola di austenite residua dopo rottura ( asse di zona [116]γ).
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Abstract
Microstruttura e proprietà meccaniche degli acciai trip sottoposti
a processi termomeccanici
Parole chiave:
acciaio, processi termomeccanici, caratterizzazione materiali
Nel presente lavoro è stato studiato il meccanismo di innalzamento delle caratteristiche meccaniche frutto della combinazione di
allungamento e carico di rottura di un acciaio multifase Fe-0.2C-1.5Mn-1.2Si-0.3Mo-0.6Al-0.02Nb (% in peso). Le microstrutture
con diverse frazioni in volume di ferrite poligonale, bainite e austenite residua sono state realizzate mediante trattamento termomeccanico controllato. La prova a trazione interrotta è stata utilizzata per studiare il comportamento della ferrite bainitica, dell’austenite residua e della ferrite poligonale in funzione della deformazione plastica. Per caratterizzare la frazione in volume e il
contenuto di carbonio dell’ austenite residua è stata utilizzata l’analisi mediante raggi X.
La microscopia elettronica a trasmissione è stata utilizzata per analizzare l'effetto della morfologia della ferrite bainitica sulla trasformazione dell’ austenite residua, indotta da deformazione, e sulla geminazione dell’austenite residua sempre in funzione della
deformazione nel materiale. Lo studio ha dimostrato che il meccanismo di geminazione dell’austenite è preferibile rispetto al
meccanismo della plasticità indotta da trasformazione (Transformation Induced Plasticity - TRIP) durante le prime fasi di deformazione di una microstruttura contenente il 15% di ferrite poligonale, mentre l’effetto dal TRIP è il meccanismo principale quando
è presente il 50 % di ferrite poligonale nella microstruttura.
La morfologia della ferrite bainitica influisce sulla modalità di deformazione dell’ austenite residua durante la deformazione.
Il comportamento della ferrite poligonale durante la deformazione dipende dalla sotto-struttura delle dislocazioni dovuta alle deformazioni e dalle ulteriori dislocazioni mobili causate dall'effetto TRIP. Il verificarsi dei meccanismi di TRIP o di geminazione
dipende non solo dalla stabilità chimica e meccanica dell’ austenite residua, ma anche dall'interazione delle fasi durante la deformazione.
48
Memorie
Acciaio inossidabile
Ferritic Nb-alloyed Cr-Steel
in simulated strip casting process
S. Lachmann, C. Klinkenberg, A. Weiss, P. R. Scheller
Nb alloyed ferritic Cr-steel is usually produced by continuous casting with following hot and cold
rolling procedure. In the laboratory scale the possible new route via strip casting was studied.
The scope of the investigation in simulated process route was the development of microstructure and
precipitations. In the experiments process parameters similar to those of the real strip caster were
chosen, then those of hot rolling and cold rolling of such cast strips. The quickly solidified layer was
produced by immersion of a steel substrate under vacuum into melt. The microstructure showed
small niobium precipitates in the grain matrix and at the grain boundaries. Their size and
distribution was evaluated for different niobium contents and cooling rates in the as-solidified
structure. The diffusion controlled change of the precipitate morphology was also analysed after
preheating and rolling. Reprecipitation and precipitate growth, as well as dissolution of
precipitations at the grain boundaries were observed. The effect of various cooling rates and niobium
content on the shape and formation of niobium containing precipitates and on the grain boundary is
discussed. Thermodynamic calculations using FactSage were carried out in order to predict the
precipitation of Nb-rich phases in ferritic stainless steels. The effect of the chemical composition and
temperature on the thermodynamic stability of these precipitates was evaluated.
KEYWORDS:
niobium, niobium carbide, precipitates, ferritic steel, strip casting, rapid solidification
INTRODUCTION
Niobium is a fundamental alloying element in the steelmaking industry. Many of today’s construction steels contain niobium as a microalloying element to reach a desired strength level, high ductility
and creeping strength by precipitation hardening. In stainless steels
niobium is added to prevent a chromium carbide precipitation and
therefore to improve the corrosion resistance and in other applications to improve the high-temperature strength (1-3).
For steels with higher amount of precipitates the strip casting technology offers an interesting possibility for material processing.
Characteristic for this technology is the high cooling and solidification rate of the strip as well as the integrated casting and rolling
process. According to this, laboratory experiments were made to
get new information about the fast solidification of niobium-alloyed, ferritic stainless steel. It was tested which parameters are critical for this procedure. Thus the laboratory experiment
parameters were chosen to be similar to the real strip casting procedure.
The main issue of this investigation was to describe the precipitation behaviour of niobium-containing phases during the casting
and rolling process. Size and shape of them are very important
factors influencing the mechanical properties of the final product.
From the results of this study it can be stated, how niobium precipitations in the ferritic matrix can be affected by process parameters.
Stefan Lachmann, Andreas Weiss, Piotr R. Scheller
Institute of Iron and Steel Technology, Freiberg University of Mining
and Technology, Germany - E-mail: [email protected]
Christian Klinkenberg
SMS-Siemag, Düsseldorf , Germany
THERMODYNAMIC CALCULATIONS
In the first stage thermodynamic calculations of niobium precipitation in ferritic stainless steel were performed. For this purpose the FactSage program was used to identify and model the
stable phases during cooling at temperatures between 1550°C
and 1000°C. FactSage is based on the calculation of Gibbs energies of all possible phases. By minimising this Gibbs energy the
most stable phase composition of the system is calculated. It is
important to note that kinetic aspects are not considered – so the
real microstructure can differ from these calculations.
In order to approximate real solidification behaviour, segregation
of the alloying elements was included in the calculations. Niobium for example segregates strongly to the melt which can lead
to precipitation directly from the liquid. Its partition coefficient
kNb which is defined as cNb, solid / cNb, liquid is around 0.3 (4).
A simplified chemical composition with 16% Cr, 0.3 to 0.9% Nb,
300ppm C, 250ppm N and Fe as balance was used for the calculation. The model calculates the stable fraction of solid steel and
precipitates at a given temperature. The fraction of liquid steel is
then used as the start composition for the next calculation at a
lower temperature. In this way segregation is simulated. Additionally the formed fraction of solid steel is “cooled down” at the
same temperature steps and the amount of precipitates which
form in solid steel is calculated. The temperature step was set to
2 K. Anyway the calculations are simplified as for the model complete mixing in liquid and solid phase during cooling is assumed.
Additionally there is no possibility to change the partition coefficients ki with respect to the cooling rate when using FactSage.
The results for 0.3 and 0.9% Nb are shown in Figure 1 and Figure
2.
At each niobium content Nb(C,N)x starts to precipitate directly
from the melt as it was enriched with niobium up to 2.5% in the
49
Memorie
FIG. 1
Phase formation of solid fraction and niobium
precipitates during solidification and cooling of
ferritic stainless steel with 0.3% Nb.
Formazione della frazione solida e dei precipitati di
niobio durante solidificazione e raffreddamento
dell’acciaio inossidabile ferritico con 0.3% Nb.
FIG. 2
Phase formation of solid fraction and niobium
precipitates during solidification and cooling of
ferritic stainless steel with 0.9% Nb.
Formazione della frazione solida e dei precipitati di
niobio durante solidificazione e raffreddamento
dell’acciaio inossidabile ferritico con 0.9% Nb.
case of 0.3% Nb and up to 4.5% in the case of 0.9% Nb when first
precipitates occur. At this moment between 4.0 wt% (with 0.3%
Nb) and 8.1 wt% (with 0.9% Nb) of liquid phase still remain. With
this precipitation the amount of alloying elements Nb, C and N in
the melt is immediately reduced which enhances solidification.
Through this fact the solidification interval is reduced with increased niobium content. Precipitates forming in the liquid phase,
as for example TiN, usually have a big size (> 10 µm) and hardly
affect the grain size as they are no barrier for moving grain boundaries. On the other hand if the precipitates in the melt during
fast cooling are sufficiently small, they get entrapped into the solidification front and act as a barrier even at very high temperatures at which the grain growth is fastest. In contrast Nb(C,N)x
precipitates forming in the solid phase appear below 1182°C (with
0.3% Nb) and below 1268°C (with 0.9% Nb) when a considerable
amount of grain growth is already finished. These precipitates are
usually much smaller (< 1 µm) and appear in a higher amount
compared to phases precipitated directly from the melt. So their
contribution to grain growth limitation is also considerable. In
50
both cases increased niobium content leads to Nb(C,N)x precipitation at higher temperatures which is beneficial for decreasing
grain growth during cooling. The higher the starting temperature
for precipitation the more likely is their formation even at high
cooling rates. Especially for strip casting where no extensive reheating is made before hot-rolling a good resistance to grain
growth immediately after casting is critical. A special thermal treatment for dissolution and reprecipitation of Nb(C,N)x is not applied either.
The total amount of precipitates at room temperature only slightly
increases with niobium content as it is mainly limited by the carbon and nitrogen content. Obviously the amount of niobium in all
cases is higher than necessary for complete stoichiometric reaction with carbon and nitrogen.
Laves-Phase ([Fe,Cr]2Nb) was predicted for the all steel compositions to form at the end of solidification, but the amount of 0.025%
was very low.
In Nb(C,N)x precipitates the ratio of C/N slightly changes with falling temperature. The precipitates formed in interdendritic liquid
show a ratio of about 2 which indicates a niobium carbide, while
the precipitates formed in solid phase show a ratio of about 0.9 similar to Nb(C,N).
In this investigation the precipitation behaviour in the experimentally simulated strip casting process, hot and cold rolling was
studied after each processing step. The rolling was carried out on
laboratory scale rolling machines, while the solidification process
of steel to a strip had to be modelled with dipping of a substrate
(cold rolled sheet in the size of 30x100 mm) into the melt. The solidification experiments were performed in a vacuum induction
furnace. Three variants of AISI 430 Cb were produced by addition
of niobium to the base material of AISI 430 from industrial production (heat 0). The chemical compositions are given in Table 1.
Electrochemical etching for selective attack of precipitations was
chosen for sample preparation. The microstructure was examined
qualitatively and quantitatively using optical microscope and
image analysis software.
Additional samples were prepared by casting the melt into a triangle shaped copper mould which provided different cooling rates
for each sample. This was controlled by introduced thermocouples. The experimental procedure was similar to the one described
in Ref. (5). The temperature range between 1300 and 1400°C is
most important for the first formation of niobium precipitates in
these steels after finished solidification. In Cooling rates between
15 and 5800 Ks-1 in this temperature range were reached in samples investigated within this work. The cooling rates were either
measured (copper mould experiments) or calculated (dipping experiments).
Dipping experiments
On the surface of the substrate sheets an “inverse” solidification
compared to the casting processtakes place. The heat is transported from the outside (melt) to the inside (sheet) and partly cumulates inside the substrate sheet. In the real process the heat is
extracted from the system by the mould. Therefore in our experiments the temperature gradient decreases with dipping time. In
order to model different strip thickness the superheat of the melt
was varied, which caused steel layers with different thickness to
freeze. With increasing superheat of the melt the frozen steel layer
gets thinner.
Calculated cooling rates between 3400 and 5800 Ks-1 were achieved in the last stage of solidification at ca. 1400°C. After this initial fast cooling the steel layer warms again in the steel melt until
the samples are pulled out of the melt. The dipping time was held
Heat
%C
%Si
%Mn
%Nb
%Cr
%Ni
%N
%Al
%Ti
%P
%S
0
03
06
09
0.047
0.031
0.029
0.032
0.24
0.20
0.21
0.18
0.37
0.23
0.23
0.19
0.006
0.313
0.645
0.964
16.2
16.2
16.0
15.8
0.19
0.12
0.11
0.13
0.033
0.024
0.024
0.025
0.003
0.005
0.006
0.008
0.003
0.007
0.009
0.011
0.021
0.017
0.018
0.019
0.004
0.009
0.008
0.009
TAB. 1
Chemical composition of steel melts (0 = base material from industrial production).
Composizione chimica dell’acciaio (0 = materiale di produzione industriale).
constant of 1 s. The temperature at the contact between frozen
steel layer and substrate varied between 1280 and 1440°C. Following a fast air cooling was performed.
An example of a substrate sheet with frozen steel after dipping
experiment is shown in Figure 3. The liquidus temperature of the
steel with 0.6% Nb was measured to be 1495°C. The thickness of
the frozen steel layer depends on melt superheat and varied between 1.5 and 2.1 mm.
FIG. 3
Steel strip after dipping experiment.
Nastro di acciaio dopo l’esperimento di immersione.
FIG. 4
Steel matrix with needle-like and round
precipitates and grain boundary phase
(0.6% Nb, 3400 Ks–1).
Matrice dell’acciaio con precipitati aghiformi
e tondeggianti e fase ai bordi dei grani
(0.6% Nb, 3400 Ks–1).
Hot and cold rolling procedure
Samples were hot rolled in one pass with the deformation degree
between 22 and 29% after preheating in 10 min at 900°C. The
cold rolling was performed in three passes with total deformation
degree of 70% with respect to the as cast state. The final thickness of the layer was 0.4 to 0.8 mm.
RESULTS AND DISCUSSION
Phase identification
An example of a typical microstructure is shown in Figure 4. Most
precipitates were small with needle-like or spherical shape and
were distributed over the whole matrix. In the following evaluation they were subdivided into three groups:
• needle-like precipitates, length up to 6 µm, thickness below 1 µm;
• oblong precipitates, length up to 2.5 µm, thickness below 1 µm;
• small, spherical precipitates, diameter below 1 µm.
In all cases spherical precipitates were most frequent. Their distribution within the samples was not uniform. Inside the grains
void corridors without any precipitates occurred.
Phase identification was made using scanning (SEM) and transmission electron microscopy (TEM) for several samples. Using
SEM only very big needle-like and spherical precipitates within
the grains could be identified as NbC with nitrogen traces taking
matrix effect into account. To characterize precipitates with diameters < 1 µm carbon extraction replicas were analyzed by TEM.
An example for the precipitates found is shown in Figure 5. The
precipitates were mainly characterized as Nb(C,N) with a needle-
Fig. 5
Transmission electron
micrographs and their
analysis results by EDS
(0.9% Nb, 250 Ks-1).
Micrografie al microscopio
elettronico a trasmissione e
risultati delle analisi mediante
EDS (0.9% Nb, 250 Ks-1).
51
Memorie
like or oblong shape.
No big precipitates (> 10 µm) were found in the samples, which
means that if precipitation of Nb(C,N)x occurred even from the liquid phase, their growth was limited. This is probably attributed
to the fast solidification and the limited time for niobium diffusion.
Laves-phase Fe2Nb, as predicted by FactSage calculations and in
other publications (3, 6), was not observed. It is believed that this
phase is suppressed as a result of the high cooling rate. It normally forms after annealing at higher temperatures or because of
segregation, which is damped at higher cooling rates. Additionally
the predicted maximum amount of 0.025 wt% was probably too
low to find this phase in the samples.
Grain size
The grain size increases with increasing superheat and decreasing solidification rate, as shown in another study on the solidification of ferritic stainless steel (7). Surprisingly a pronounced
effect of the niobium content on the mean grain size was observed, that had a stronger influence than the solidification and cooling rate. It can be clearly seen from samples shown in figure 6
and Figure 7 which both were cooled with about 250 Ks-1, that increased niobium content results in finer grain. The high solidification rate resulted in a columnar structure with precipitates at
the grain boundaries and inside the grain. The mean grain size decreases with increasing niobium content regardless of the cooling
rate as shown in Figure 8. As will be discussed later on, more niobium precipitates are present with 0.3% Nb, but at 0.9% Nb their
mean size is much higher. Although fine dispersed precipitates
are usually more effective for the prevention of grain growth than
bigger ones, the main influence during the solidification process
and following cooling is obviously the starting temperature of precipitation. As shown in the thermodynamic calculations the precipitation of Nb(C,N)x starts at higher temperatures with
increasing niobium content and even in the melt. Thus at 0.9%
Nb a higher amount of Nb(C,N)x precipitates was present at the
end of solidification than at 0.3% Nb, which seems to be very effective to limit the grain growth at high temperatures. Additionally the precipitation of secondary Nb(C,N)x in solid steel starts
at higher temperatures with increasing niobium content, which
acts in the same way on limiting the grain growth.
FIG. 6
Dimensione del grano di provini come-solidificati con
0.3% Nb raffreddati a 250 Ks-1.
FIG. 7
Characterization of microstructure and precipitates
A quantitative analysis of the Nb(C,N)x amount and distribution
using optical microscopy with image analysis. A lot of these precipitates found by TEM investigation were smaller than 0.5 µm.
The amount of these very fine precipitations might be underestimated by optical microscopy. Additionally precipitates that were
identified by optical microscopy as small and round might be also
plate-or needle-like in 3D image as mainly small niobium carbonitride needles were found during TEM analysis.
Effect of cooling rate
In the as-solidified state of the samples the size and distribution of
precipitations is dependent on the cooling regime. At very low cooling rates (< 100 Ks-1) the majority of precipitates is needle-like or
oblong with a size of up to 10 µm. At cooling rates of about 3000
Ks-1 the needles were below 5 µm in length, while most of them
were around 2 µm. As cooling rate further increased the precipitates became smaller (< 2 µm in diameter). On the other hand a
great number of precipitates can be observed in the microstructure at highest cooling rates of 5800 Ks-1 – most of them below 1
µm in diameter. The shape of these niobium carbonitrides is also
influenced by the cooling rate. In the samples with the lowest cooling rate they are real plates that look like needles in cross-section.
52
Grain size of as-solidified samples with 0.3% Nb
cooled at 250 Ks-1.
Grain size of as-solidified samples with 0.9% Nb
cooled at 250 Ks-1.
0.9% Nb raffreddati a 250 Ks-1.
FIG. 8
Grain size of as-solidified samples at different
niobium contents.
diversi contenuti di niobio.
FIG. 9
Distribution of niobium
carbonitride precipitates at
different cooling rates
(0.6% Nb, as-solidified
state) (optical microscopy).
different cooling rates (0.6%
Nb, as-solidified state) (optical
microscopy).
FIG. 10
different niobium contents
(as-solidified state) (optical
microscopy).
Distribuzione dei precipitati di
carbonitruri di niobio per
diversi contenuti di niobio
(stato come-solidificato)
(microscopia ottica).
With increased cooling rate the precipitates become shorter and
more spherical. A lot of spherical niobium carbonitrides precipitates were analysed at the highest cooling rate in these solidification experiments. The results are summarised in Figure 9.
It can be seen, that the total number of niobium carbonitrides increases with increasing cooling rate. Especially the number of
fine, spherical precipitates shows a steep increase, which can improve the mechanical properties through strengthening and impeding grain growth. On the opposite the number of coarse
needle-like precipitates decreases.
The reason for an increasing number of precipitates and their decreasing mean size at higher cooling rates is their forming mechanism. Whereas at high cooling rate the driving force for matrix
precipitation is high due to strong non-equilibrium conditions the
diffusion and growth of precipitations predominate under low cooling rates. A similar trend was observed for the grain boundary
precipitates.
Effect of niobium content
Niobium content shows surprising effect as the number of precila metallurgia italiana - n. 11-12/09
pitates does not increase with higher niobium content, but decrease (see Figure 10). On the other hand at 0.9% Nb the inclusions were longer (up to 8 µm) and like needles, while at 0.3% Nb
they had a round shape and were much smaller (up to 4 µm). One
possible explanation is that according to the thermodynamic calculations at lower niobium contents the main part of precipitation occurs at lower temperatures. As the diffusivity is reduced
and the undercooling is increased, more sites for precipitation become favourable. On the other hand at high niobium contents,
most niobium precipitates at higher temperatures where the diffusivity is high. For this reason the first precipitates are enlarged
and grow in the equilibrium plate-like form. Despite the decreasing number of precipitates with increasing niobium content in
the steel, the total amount of Nb(C,N)x only slightly increases.
Effect of rolling
Samples of the solidification experiments were reheated and hotrolled. After this treatment no recrystallization was observed but
a deformation texture in rolling direction. At the grain boundaries dissolution of niobium precipitates had started, as their mean
53
Memorie
FIG. 11
Evolution of precipitate size
distribution after
solidification and hot and
cold rolling (optical
microscopy).
Evoluzione della distribuzione
della dimensione dei
precipitati dopo
solidificazione e laminazione
a caldo e a freddo
(microscopia ottica).
Figure 12: Precipitation
behaviour for different
niobium contents and
cooling rates.
Comportamento nella
precipitazione per diversi
contenuti di niobio e diverse
velocità di raffreddamento.
size decreased and the needle-like precipitates changed towards
a spherical shape. On the other hand the matrix was obviously
still supersaturated in niobium as many new small precipitates
(< 1 µm) were formed. When they reprecipitate at lower temperatures they do so on new-formed dislocations or around other precipitates in a more equilibrium form. As a result more and finer
precipitates exist in the steel matrix.
Especially at highest cooling rates the precipitates were spread
more completely in the steel matrix and the voids were smaller
after hot-rolling. Obviously the lower the cooling rate after solidification is the more completely the supersaturation of niobium in
the matrix already decreases. The potential for precipitation is
then lower when hot-rolled. On the other hand at highest cooling
rates the supersaturation rests until preheating before rolling and
creates new, fine precipitates in the steel matrix. In result only
few differences are found for samples of low cooling rate before
and after hot rolling.
After cold rolling (without any annealing) no significant difference
in the microstructure was found between hot-rolled and cold-rolled samples. As no additional heat treatment except the self-heating during rolling was applied to them, grain growth, precipitate
coarsening or reprecipitation could hardly occur. The size distri-
54
bution after hot and cold rolling compared to the as-solidified state
is shown in Figure 11.
It can be stated from this study that a short heat treatment at these
low temperatures (900°C) before hot rolling – in most cases in-line
during the casting and rolling process – has no negative effect on
the precipitation distribution and is even beneficial.
CONCLUSIONS
The evolution of the shape and distribution of niobium precipitates has been studied for ferritic stainless steels. The simulation of
the strip casting process was performed in dipping experiments
with the aim of fast solidification and high cooling rates (up to
5800 Ks-1). The change of microstructure during hot- and cold-rolling was examined. The precipitates in the matrix were identified
using SEM and TEM analyses as Nb(C,N)x, which occurred in needle-like or round shape with varying size. Fe2Nb-Laves-phase was
not detected in the microstructure.
Using image analysis niobium carbonitride occurrence depending
on niobium content and cooling rate was quantified and compared to samples with lower cooling rates (< 500 Ks-1). A short summary is shown in Figure 12. At the performed cooling rates of the
samples differences in precipitate size distribution was observed.
Generally increased cooling rates caused the formation of more
and smaller precipitates. This effect was more pronounced with
higher niobium contents, where at low cooling rates a lot of needle-like niobium precipitates with a higher size up to 10 µm form.
At higher cooling rates their relative amount is decreased. With
low content of niobium (in this case 0.3%) these long precipitates
hardly occurred and the precipitates were much finer (< 1 µm).
Surprisingly the number of precipitates in the as-solidified state
did not increase with niobium content but decrease. If the niobium content is very high (in this case 0.9%), the niobium precipitation starts at higher temperatures where the diffusivity is high
and less precipitates of higher mean size form. With lower niobium contents the precipitation starts at lower temperatures and
results in a finer distribution. This behaviour was predicted by
FactSage calculations.
With niobium contents of 0.3 to 0.9% and the steel composition
studied here a precipitation of Nb(C,N)x is also possible in the interdendritic liquid. With these precipitates and those which form
in solid steel during cooling it is possible to impede grain growth
in ferritic steels at high temperatures. This effect becomes more
pronounced with increasing niobium content.
For optimized mechanical properties a fine precipitation distribution and small ferrite grain are desirable. While the niobium
content should be as low as possible to form a sufficient amount
of fine precipitates, increasing the niobium content can limit grain
growth at high temperatures more effectively and lead to a higher
strength level through solution hardening. In any case increasing
the cooling rate improves the distribution of fine niobium carbonitrides.
ACKNOWLEDGEMENTS
The authors would like to thank the Niobium Company for their
support to maintain this work.
REFERENCES
1. N. Fujita, K. Ohmura and A. Yamamoto: Mat. Sci. Eng. A 351 (2003),
272
2. N. Fujita, H.K.D.H. Bhadeshia and M. Kikuchi: Metall. Mater. Trans. A
33A (2002), 3339
3. N. Fujita, K. Ohmura, M. Kikuchi, T. Suzuki, S. Funaki and I. Hiroshige:
Scr. Mater. 35 (1996), 705
4. J.A. Brooks, M. Li, M.I. Baskes and N.C.Y. Yang: Sci. Tech. Weld. Join. 2
(1997), 160
5. N.H. Pryds and X. Huang: Metall. Mater. Trans. A 31A (2000), 3155
6. V. Kuzucu, M. Aksoy, M.H. Korkut and M.M. Yildirim: Mat. Sci. Eng. A
230 (1997), 75
7. B. Callmer and O. Grinder: Scand. J. Metall. 9 (1980), 151
Abstract
Acciaio ferritico al Cr
legato con Nb nel processo
di Strip Casting simulato
Parole chiave:
acciaio inossidabile, simulazione, strip casting,
solidificazione
L’acciaio ferritico al Cr legato con Nb viene solitamente prodotto mediante colata continua con susseguenti procedure
di laminazione a caldo e a freddo. Nel presente lavoro è
stata studiata in laboratorio una possibile nuova modalità
mediante strip casting. Lo scopo della simulazione del processo in laboratorio era quello di esaminare lo sviluppo
della microstruttura e delle precipitazioni. Negli esperimenti sono stati scelti parametri di processo simili a quelli
realmente utilizzati nello strip casting, e in seguito parametri simili a quelli della laminazione a caldo e a freddo
dei nastri. Lo strato solidificato rapidamente è stato ottenuto mediante immersione nella fusione di un substrato di
acciaio sotto vuoto. La microstruttura ottenuta ha mostrato
piccoli precipitati di niobio entro il grano e al suo bordo. La
dimensione e la distribuzione dei precipitati è stata valutata in termini di diverso contenuto di niobio e di velocità
di raffreddamento, nella struttura con l’ acciaio allo stato
di come-solidificato. Il cambiamento della morfologia del
precipitato, controllato dalla diffusione, è stato analizzato
anche dopo preriscaldamento e laminazione. Sono stati osservati riprecipitazione e ingrossamento dei precipitati,
così come la loro dissoluzione al bordo del grano. Inoltre
sono stati messi in evidenza gli effetti delle diverse velocità di raffreddamento e del contenuto di niobio sulla formazione e la morfologia dei precipitati contenenti niobio e
sulla loro collocazione al bordo del grano.
Sono stati effettuati calcoli termodinamici, utilizzando FactSage, al fine di predire le caratteristiche della precipitazione delle fasi ricche di Nb negli acciai inossidabili
ferritici. E’ stata infine valutato l'effetto della composizione
chimica e della temperatura sulla stabilità termodinamica
di questi precipitati.
55
Memorie
Fonderia
Study of the effect of process parameters
on the production of a non-simmetric
low pressure die casting part
A. Pola, R. Roberti
Low pressure die-casting is a "near net shape" foundry process that offers a good compromise between
economical aspects, production rate and casting quality. Because of the constrained position of the gating
system, the application of traditional LPDC process is generally limited to axis-symmetric or symmetric
geometries. The aim of this work was to investigate the low pressure die-casting process in order to define the
effect of various system settings on the production of a sound non-conventional cast component. The research
was supported by the modelling of mould filling and casting solidification, in order to evaluate both the
influence of process parameters and the reliability of the modelling software in the prediction of flow pattern
and thermal history of casting as well as defects formation.
The results were compared with those obtained on an experimental die, completely instrumented, to better
understand the process, validate the calculation procedure and make more confident the use of this tool for
complex parts. Metallographic analyses were also carried out to compare the quality of simulated and real
castings, with particular reference to shrinkage and gas porosity.
KEYWORDS:
LPDC, simulation, filling, pressure influence, AlSi7Mg
INTRODUCTION
Low pressure die-casting (LPDC) is a "near net shape" foundry process in which the molten alloy is poured into a holding pressurized
furnace located below the die table. A feeding tube, called riser or
stalk tube, runs from the furnace to the bottom of the die, as shown
in Fig. 1.
The surface of metal bath in the crucible is pressed by a dry air at
relatively low-pressure (typically in the range of 0.1-1bar) in order
to overcome the difference of metallic pressure between the die
and the surface of the liquid alloy. The molten metal is, therefore,
forced through the stalk tube, feeding the die cavity with low turbulence, also because of the use of a metal filter.
The pressure ramp is increased in order to pressurize the casting
during the solidification, strongly reducing shrinkage porosity.
Once the casting is completely solidified and sufficiently cooled,
the external pressure is released, the molten metal in the riser tube
flows back down into the crucible
by gravity action and the casting is
ejected to allow the next cycle.
Compared with other permanent
mould processes LPDC provides
low levels of scrap (no risers), high
mechanical, metallurgical and technological properties (thanks to
low porosity level), dimensional accuracy, feasibility in using sand
core, etc… together with limited
Annalisa Pola, Roberto Roberti
Universitá degli Studi di Brescia,
Dipartimento di Ingegneria
Meccania e Industriale
[email protected],
[email protected]
equipment costs.
The gating system is usually positioned in the middle of the casting, correspondent to the centre of the crucible, in order to guarantee uniform pressure and, therefore, flow distribution. This
constrained arrangement imposes the use of traditional LPDC process just for axis-symmetric or symmetric geometries.
Furthermore, the low pressure die-casting subject seems to be not
widely discussed in literature.
For these reasons the present study was aimed to analyse the LPDC
of a non symmetric casting, in order to improve process performance and productivity, investigating the effect of various system
settings (pressure curve, alloy composition within the standard tolerance, melt temperature, etc...) on the production of a sound nonconventional cast component.
To verify whether LPDC may be suitable for manufacturing complex geometries, a good knowledge of metal fluid-dynamic beha-
FIG. 1 Low pressure die-casting process.
Processo di colata in bassa pressione.
57
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FIG. 2
LPDC machine and the
experimental die (cooling
channels).
Macchina LPDC e stampo
sperimentale (canali di
raffreddamento).
viour into the die cavity is fundamental. The simulation of mould
filling and casting solidification provides a highly useful and reliable tool to rapidly investigate the flow pattern or to easily modify
geometries and process parameters, evaluating their effect on castings quality.
The study was carried out by means of a commercial software (Procast ), considering a standard foundry aluminum-silicon alloy
(AlSi7Mg0.3) whose thermo-fluid dynamics properties needed for
simulation are well known for average chemical composition.
The results were compared with those obtained on an experimental die, available at LKR Laboratory, completely instrumented by
newly developed metal front and temperature sensors to validate
the calculation procedure. Deviations between trials and simulations were analysed, making more confident the use of this tool
for complex parts in the future.
Metallographic analyses were finally carried out to verify the accuracy of the simulation, with particular reference to shrinkage
and gas porosity.
The LPDC machine used for the samples production was a Kurtz
AK92 equipped with a crucible furnace of 125kg capacity (Fig. 2).
Before each tests series 2 samples for emission spectroscopy investigation were taken from the melt to verify the bath composition
and the H2 content.
The die was pre-heated by means of a gas-powered burner for 1-2
hours, in order to reach a sufficiently uniform temperature of ne-
arly 180°C.
The casting geometry consists in a plate characterized by 6 different thicknesses (3 mm, 10 mm, 20 mm, 25 mm, 15 mm and 5
mm) with a not symmetric cavity with respect to the stalk tube,
placed in the middle of the crucible (Fig. 2). This simple configuration allowed to investigate the use of LPDC for the production of
non conventional components and to verify the thermal behaviour
of the die, equipped with an ad hoc designed cooling system consisting in channels flowed with compressed air (pressure of approximately 5,5 bar).
The experimental trials were continuously monitored by means of
6 K-type thermocouples placed inside the mobile half die, 5 mm beneath the surface cavity (Fig. 3), and 8 metal front sensors placed
exactly on the surface of the fixed half mould (Fig. 4) in order to assess the metal flow pattern and to compare this with the simulated one.
At the casting ejection, the surface temperature of casting and die
cavity were also measured by a contact thermocouple.
The investigated process parameters were:
− melt temperature, between 700°C and 740°C;
− alloy composition, changing the elements percentage within
standard tolerance and also considering the effect of modifier
(Sr) and grain refiner (Ti-B) additions on metal fluidity and casting quality;
− pressure curve, in terms of different profiles, as reported in Fig.
5, according to the machine features.
Each trial was named as reported in Table I.
FIG. 3
FIG. 4
Mobile half mould geometry and thermocouples
position.
Geometria del semi-stampo mobile e posizione delle
termocoppie.
58
Sensor distribution inside the half fixed mould.
Distribuzione dei sensori all’interno del semi-stampo
fisso.
Fonderia
Samples
Chemical composition
Casting T°
Pressure
A
B
C
D
E
F
G
H
Si7.11%, Mg0.4%, Sr, TiB2
Si7.11%, Mg0.4%, Sr, TiB2
Si7.13%, Mg0.4%, Sr, TiB2
Si7.13%, Mg0.4%, Sr, TiB2
Si7.07%, Mg0.4%, Sr, TiB2
Si7.07%, Mg0.4%, Sr, TiB2
Si6.5%, Mg0.2%, Sr, TiB2
Si7.5%, Mg0.2%
740 °C
700 °C
720 °C
720 °C
720 °C
700 °C
720 °C
720 °C
0.4-0.8 bar
0.4-0.8 bar
0.4-0.8 bar
0.4-0.8 or 0.4-0.5 bar
0.4-0.5 bar
0.5-0.6 or 0.6-0.7 bar
combinations of 0.4-0.5, 0.6-0.8 bar
0.6-0.8 bar higher holding time
TAB. I
Investigated process parameters during casting trials.
Parametri di processo indagati durante le prove di colata.
MODEL
The simulation tool used in this study to evaluate the feasibility of
the non symmetric cavity die design is represented by the finite
element ProCast® software, developed by ESI Group.
A FE mesh was created using tetrahedral elements; Fig. 6 shows
shape and mesh of the component.
Temperature dependent thermo-dynamic properties of the casting
alloy (A356) and mould material (H13 steel) used in the analyses
were available in the software database.
Initially the temperature was set to 180°C for all nodes of the die,
according to the data measured by the thermocouples at the end
of the pre-heating period, and equal to 720°C for the liquid alloy,
as detected by the thermocouple dipped into the melt (Table I for
the H family of samples).
The boundary conditions for the calculations were set as follows:
− a heat transfer coefficient, h, equal to 10 W/m2K on all the external surfaces of the die in contact with ambient air;
− a heat exchange coefficient on the air cooling channels walls,
calculated by the well known equation:
Nu = 0.023.Re0.88.Pr0.33 if Re > 6000, Pr > 0.7
FIG. 5
Experimental pressure profiles.
Profili di pressione sperimentali.
[1]
where Re, Pr e Nu are respectively the Reynolds, Prandtl and Nusselt numbers;
− pressure ramp, called inlet boundary condition, at the bottom of
the ingate cylinder, established on the basis of the experimental pressure curves and in particular the pressure curve used for
samples H production (Fig. 5), which resulted to be the best condition during casting trials. It must be noticed that, since the
bottom of the ingate cylinder does not correspond to the liquid
surface in the furnace, the pressure ramp measured in the real
process should be shifted of a ∆P value (Fig. 1), according to the
following equation:
[2]
where ρ is the liquid metal density, ∆h is the height of the metal
into the riser tube with respect to the bath surface, and D and d are
respectively the crucible and riser tube diameter.
EXPERIMENTAL RESULTS
Real steady state thermal conditions were reached after some injections with air cooling enabled; no complete castings were produced in the warm-up period, until steady state conditions were
achieved. In Fig. 7 the warm-up phase of the C production is shown
as an example.
The samples cast at higher melting temperature (A) were completely filled (Fig. 8A), thanks to the increased fluidity with temperala metallurgia italiana - n. 11-12/09
FIG. 6
3D geometry and mesh.
Geometria e mesh 3D.
ture. Notwithstanding the almost complete filling, the A family
samples showed surface defects associated to shrinkage phenomena (piping in the thicker sections) as well as some distortions,
both due to the high melt temperature. Therefore, the A trials set
revealed that, in these working conditions, care must be taken to
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FIG. 7
Measured temperatures during the pre-heating
period.
Temperatura misurata durante la fase di preriscaldo
dello stampo.
obtain sound castings and to preserve the mould.
On the contrary, with a bath temperature of 700°C (B samples) the
reduced fluidity, the lower die temperature and the limited superheating (above Tliquidus) together with the absence of proper
vents could not allow a complete filling of the cavity; particularly,
the thinner section was always partially empty (Fig. 8B). It must
be noticed that the B samples were cast after to the A family, therefore the metal level within the crucible was strongly decreased
and, consequently, experiments should have been carried our
with a new pressure curve.
Based on the previous results, the melt temperature was fixed at
720°C, in order to avoid hot tearing and large shrinkage porosities
as well as early die damage.
The effect of a lower pressure was tested for the production of samples D and E, which resulted uncompleted. The surface of the thinner section (Fig. 9, left side) showed an extended cold wave; the
metal, in fact, seems to divide into two veins that are not perfectly welded, partly because of air entrapment and partly because
of early solidification, probably due to the slow filling associated
to the not proper pressure used.
The F set of samples was cast at a temperature of 700°C, to reduce shrinkage porosity and hot cracks. With this low melting
temperature the previously adopted pressure was absolutely not
suitable to fill the cavity; different pressure curves were then attempted in order to obtain sound castings, notwithstanding the
low melting temperature (i.e. low fluidity) imposed. The samples
F obtained with higher cast pressure showed a quite complete filling (Fig. 9, right side), according to the not optimized die design,
and a quite good surface appearance (very small cracks).
In the last trials, the effect of chemical composition variation on
die filling was investigated.
As well known, silicon is the main aluminum alloying element,
added to improve fluidity, castability and hot tearing resistance as
well as to reduce thermal shrinkage. During the previous tests the
Si content was always fixed at the standard tolerance average value.
For the sixth family of samples (G series) the Si percentage was
FIG. 8
Superficial appearance of first
series of tests,
A (left) and B (right).
Aspetto superficiale della prima
serie di getti prodotti,
A (sinistra)
e B (destra).
FIG. 9
Superficial appearance of
sample D (left) and F
(right).
Aspetto superficiale del
campione D (sinistra) e F
(destra).
60
Fonderia
FIG. 11 Calculated die temperatures.
Temperature dello stampo calcolate.
FIG. 10 Complete casting, family H.
Getto completo, famiglia H.
maintained at the lower limit acceptable in the standard tolerance.
Notwithstanding the good process control no sound neither complete castings were obtained, revealing the importance of the proper Si level content.
Based on these findings, the last series of tests (H family) were finally cast using the higher Si percentage together with the lower
level of Mg, which can reduce alloy fluidity due to oxidation; moreover, no modifier or grain refiner were added, in order to guarantee a high fluidity. With the imposed pressure curves
completely filled parts were obtained, as shown in Fig. 10.
The results of these experimental trials on a non symmetric are
here summarized:
− steady state thermal conditions are fundamental to guarantee
good quality castings;
− chemical composition must be strictly controlled in order to obtain a high enough fluidity;
− high pressures and/or long holding times allow a complete filling, also when the melt temperature is low or the composition
not strictly controlled. An optimized curve can, therefore, be
found as a function of the other casting parameters, but attention
must be paid at the die closing force in order to avoid excessive
flashes production or, even, dangerous outflow of liquid metal.
Obviously, a compromise between all these parameters must be
defined, depending on the specific working conditions.
SIMULATIONS VS SENSORS RESULTS
The thermal steady state conditions were reached after 12th cycles,
as shown in Fig. 11 on the left where the calculated temperature
profile in the thermocouple positions (5 mm behind the surface
cavity) is compared to that measured by the thermocouples.
It can be noticed that measured and simulated temperatures differ
by roughly 25°C/6%, except TC1 in correspondence of the thinnest section (85°C). This difference is due to the fact that the cycling modelling considers the mould cavity completely filled by the
alloy which releases latent heat of solidification, consequently increasing the die temperature. In the real warm up period the thinner section was not interested by the metal flow, because of the
combination of thin section and cold wall; therefore, this area of
mould was not interested by the metal during filling and solidification, explaining the gap between simulation and experimental
trials.
For the same reason the temperatures measured by the thermocouples, positioned exactly on the mobile die surface, and those simulated after the cycling period show a difference of very few
Celsius degrees (Fig. 12); the good correspondence of results demonstrate the proper description of material thermodynamic data
and correct choice of interfaces and boundary conditions parameters.
Fig. 13 shows the filling patterns obtained numerically at subsequent times. According to the simulation, the alloy enters the cavity through the gate forming a vein which mainly spreads along
the wall of the fixed mobile die filling in the central thicker sections
(25 mm and 20mm). The liquid metal reaches the thinner section
just at the end of the cavity filling; it is divided in two arms, one
FIG. 12 Measured and simulated temperature on the surfaces of the mobile die (H family).
Temperature Misurate e calcolare sulla superficie del semi-stampo mobile (H famiglia).
61
Memorie
FIG. 13
Simulated fluid flow.
Flusso del metallo simulato.
FIG. 14 Recorded fluid flow.
Flusso del metallo registrato dai sensori.
FIG. 15 Microstructure of the casting, family H area 8.
Microstruttura del getto, famiglia H area 8.
from top the and one from the bottom, that weld around the centre of the thinner section.
These results match well with the experimental data; as shown in
Fig. 14, in fact, there is a close agreement in the fluid flow pattern.
As confirmed by the obtainment of un-complete castings, which
present cavity in the centre of the thinner section area (Fig. 8 –
10) depending on the maximum applied pressure, the simulation
predicts almost correctly also the mould filling.
62
METALLOGRAPHIC INVESTIGATIONS
The quality of the produced casting was analysed by means of metallographic investigation; particularly, the shrinkage porosity as
well as the microstructure, in terms of secondary dendrite arms
spacing (SDAS), were analysed. The samples were cut according to
the scheme showed in Fig. 10.
In the thicker section (area 8) some shrinkage porosity were detected, as already predictable observing the surface depression of
Fonderia
FIG. 16 Calculated Secondary Dendrite Arms Spacing.
Spazio fra I rami secondary delle dendriti calcolato.
all the samples, as a consequence of an insufficient imposed pressure as well as very high temperature that induces a slow solidification. As an example, in Fig. 15 the microstructure of a sample of
the H family is reported.
The same defect was revealed by the simulation software, demonstrating the good correspondence between calculation and production, also during solidification.
Quantitative metallography measurements resulted in a SDAS of
11.7 m in the thinner section and of 36 µm in the thicker section
for the real cast sample (sections 15 and 8 Fig. 10); as shown in
Fig. 16 the calculated SDAS was very similar to the observed one
(nearly 14 µm and 42 µm respectively).
CONCLUSIONS
The production of a non symmetric component cast by low pressure die-casting process was investigated in order to define the influence of various system settings on the casting quality.
Eight set of samples were produced under different process conditions (melt temperature, pressure ramp and chemical composition within the standard tolerance range) and the process was
continuously monitored by means of thermocouples and pressure
sensors placed inside the mould.
A finite elements model was also developed for the simulation of
die filling and casting solidification.
A good correspondence between simulation and experimental
trials was observed, in terms of die temperature, mould filling and
casting quality.
The results of this study can be summarized as follows:
− simulation is in general a good tool for predicting low pressure
foundry processes,
− temperature and metal front sensors help to understand die filling and validate the simulation procedure, as well as to define
optimum casting conditions;
− thick sections behaviour can be well predicted by simulation
software, but additional phenomena occur at thin sections;
− accurate definition of cycling simulation are needed for thin section;
− small variations in chemical composition can create problems
in thin sections filling;
− LPDC thin parts can be more easily produced provided that no
modifier or grain refiner are added.
ACKNOWLEDGMENTS
The authors gratefully acknowledge the help and the suggestions
by Prof. Helmut Kaufmann and Dr. Werner Fragner, Leichtmetallkompetenzzentrum Ranshofen GmbH (LKR).
REFERENCES
1) JER-HAUR K., FENG-LIN H.and WENG-SING H., Development of an interactive simulation system for the determination of the pressure-time
relationship during the filling in a low pressure casting process,
Science and Technology of Advanced Materials 2, (2001), p.131.
2) HINES J., Determination of interfacial heat-transfer boundary conditions in an Aluminum low-pressure permanent mould test casting, Metallurgical and Material Transactions B, (2003), p 299.
3) MANILAL P.I., SINGH D.P.K. and CHEN Z. W., Computer modelling and
experimentation for thermal control of dies in permanent mold casting, AFS Transactions (2003).
Abstract
Studio dell'effetto dei parametri di processo sulla produzione
di un getto non simmetrico colato in bassa pressione
Parole chiave:
colata in bassa pressione, simulazione, riempimento, AlSi7Mg
Il processo di colata in bassa pressione (o LPDC) è una tecnologia di fonderia cosiddetta “near net shape”, ovvero tale da consentire
l’ottenimento di un pezzo di geometria “prossima alla forma finita” e che richiede un numero limitato di operazioni di finitura;
esso rappresenta un buon compromesso fra aspetti economici, velocità di produzione e qualità del getto. A causa della posizione
del sistema di alimentazione dello stampo, che risulta costruttivamente vincolata, l’applicazione del processo tradizionale LPDC
è solitamente limitato a geometrie assial-simmetriche o comunque simmetriche; tipicamente si producono cerchi in lega.
Lo scopo del presente lavoro è stato quello di studiare il processo di colata in bassa pressione al fine di definire l’effetto dei vari
parametri di produzione (pressione, temperatura del metallo, ecc..) sulla qualità di un getto non convenzionale. La ricerca è stata
supportata dalla simulazione del riempimento dello stampo e della solidificazione del getto, così da valutare sia l’influenza dei
parametri di processo sia la validità della modellazione numerica nella previsione della fluidodinamica di riempimento e della
storia termica di getto e stampo, nonché della formazione di difetti.
I risultati della simulazione sono stati confrontati con quelli ottenuti utilizzando uno stampo sperimentale, completamente strumentato, al fine di comprendere meglio il processo, validare la procedura di calcolo impostata e rendere l’utente più sicuro nell’uso di questo strumento anche per componenti dalla geometria più complessa.
Infine, sono state condotte anche delle indagine metallografiche per confrontare la qualità del getto simulato con quella del prodotto, facendo riferimento in particolare alla porosità da gas e da ritiro.
63
Memorie
Colata continua
Dynamic de-oxidation and inline alloying of Al
in continuous casting of billets and strips
D. Senk, A. Grosse, G. Gräf
The method of controlled stepwise de-oxidation and alloying of carbon steel melt with Al-wire has been
investigated. The melt is pre-deoxidized in the ladle, the main fraction of non-metallic inclusions is removed to
the ladle top slag by stirring. Final de-oxidation and alloying takes place just before solidification in the
continuous casting mould. In three steps from laboratory via a pilot facility to an industrial caster the
efficiency of that method was tested. No disadvantage could be found; the benefits are high amount of [Al]diss.,
high <Al> yield rate, better macro-cleanliness, and improved process quality by avoiding depositions and
clogging. By that method, the production of Al-killed carbon steel grades should be possible also with near-netshape casters which use in general small orifices in tundish and SEN.
KEYWORDS:
continuous casting, de-oxidation, clogging, cleanliness, near-net-shape casting
INTRODUCTION
In continuous casting the liquid steel is guided through a throttle which regulates the melt flow. The narrowest point can be either in the orifices of submerged entry nozzles (SEN) in the
mould or at the bottom of the tundish where a metering nozzle
or valve like slide gate or stopper rod is installed [1]. At that narrowest point the melt flow velocity and also the pressure conditions are changing severely. Here, small oxidic inclusions
suspending in the steel melt can come in touch with the refractory walls, and by high turbulence contact between those inclusions takes place. Heat flux from melt into refractory material
changes local temperature so that in connection with change of
pressure the local thermodynamic conditions in parts of the
steel melt are changing.
In addition, microscopic cracks in refractory material or poor
clearance in the contact areas of slide gate plates or at flanges
of SEN can allow pick-up air; by chemical reaction this oxygen
is able to form further oxidic inclusions in the steel melt. A
redox-reactions between alloyed [Al]diss. and slag, refractory or
slide gate powder can increase the amount of alumina particles
dispersed in steel melt [2].
Oxidic inclusions can stick to the refractory walls and form layers which are growing during the casting time. The layers are
able to block the melt flow with the result of stopping the casting process; parts of the layers can break and flow into the
mould where they are entrapped in solidifying steel shell and
lead to poor macroscopic cleanliness. In both cases the quality
of the as-cast steel will be diminished.
Many steel grades require a certain aluminium content which is
prescribed in technical standards. Aluminium diminishes the
amount of free, dissolved oxygen in steel so that reactions of oxygen with e. g. carbon is suppressed, and the formation of {CO}
bubbles which would lead to a weak strand shell or to rimming
in the mould is avoided by so called total oxygen killing. FurD. Senk
Dept. of Ferrous Metallurgy of RWTH Aachen University
A. Grosse
BSE Badische Stahl Engineering GmbH, Kehl, Germany
G. Gräf
Dept. of Metal Forming of RWTH Aachen University
FIG. 1
Poor macro-cleanliness by clustering of alumina
particles [3].
Scarsa macro-pulizia dovuta ad agglomerati di particelle
di allumina [3].
ther-on, alloyed aluminium is able to control grain grows in hot
rolling by (AlN) particles which are formed during hot rolling
[3] by recombination of {N2} in the later product is suppressed.
In some steel grades like dual phase steel aluminium is alloyed
to stabilize a certain fraction of ferritic grains in the final structure.
Aluminium is a chemical element with high affinity to oxygen,
and the reaction product Al2O3, called alumina, is a stabile oxide
with a melting point of appr. 2,050 °C. Those particles are formed by mechanisms mentioned above, and their initial diameter is about 1 µm; in terms of steel cleanliness the field of
micro-cleanliness is affected. By agglomeration supported by the
relatively high interfacial energy between alumina and steel
melt the particles can grow rapidly by agglomeration and form
macro-inclusions.
When their size becomes 30 µm and more by clustering (Figure
1, [3]), those particles are degrading the steel quality in the field
of macro-cleanliness, e. g. the ductility behaviour. Further-on, in
the casting process those particles can block the free cross-section in the metering nozzles or the SEN; that clogging terminates the pouring duration (Figure 2, [4]).
One method to overcome the clogging problem is the addition
65
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FIG. 2
Diminished cross sections of
SEN by deposition of alumina
[4].
Riduzione della sezione trasversale
del SEN a causa del deposito di
allumina [4].
of calcium by ladle metallurgy [5]; the formed calcium-aluminates are liquid at pouring temperature so that plating effects
are avoided. Concerning macro-cleanliness that method is not
the ultimate solution, but helps to prevent of clogging during casting sequences.
Particularly, small cross-sections of metering nozzles and SEN of
billet or strip casters are affected by clogging problems on alumina basis. In many cases aluminium addition has to be avoided,
and the oxygen in steel is killed by [Si] and [Mn]. The steel grades containing higher concentrations of aluminium cannot be
produced in those cases. To overcome this problem a method of
adding aluminium in several subsequent steps in the process
chain has been investigated; the required dissolved aluminium
content is alloyed by ‘DDA, Dynamic-Deoxidation-Alloying’ or
‘alloying on demand’.
THEORETICAL BACKGROUND
The formation of dangerous alumina particles starts spontaneously by almost homogeneous nucleation in the melt when the
solubility product of the reaction
2 [Al] + 3 [O] -> (Al2O3)
L:
[Al], [O]:
(Al2O3):
[ai]:
with
LAl2O3 = [aAl]2 · [aO]3
(1)
solubility product at chemical equilibrium
dissolved elements
solid alumina
thermodynamic activity in the steel melt, a<Al2O3> = 1
is reached. This temperature depending reaction is described by
log10KAl2O3=-64,000/T+20.57 [6]
and L=1/K
(2a)
at 1,600 °C.
(2b)
The following steps of formation of micro- or macro-inclusions on
the basis of stabile nuclei are [7]:
- growth by diffusion,
- Ostwald ripening,
- gradient, Brown, turbulence, or Stokes collisions.
The idea of control the particle formation is the application of
aluminium to the steel in the ladle up to an amount which diminishes the oxygen concentration to about [aO]=20 ppm (totally
killed steel), remove the particles as far as possible by ladle metallurgical treatment e. g. by soft bubbling Ar stirring, and add
the required dissolved aluminium by alloying in tundish, SEN,
and in mould. The dynamic alloying shall be based on the result
66
of an inline-measurement of oxygen activity e. g. using a Celox®system. Here, not the [O]diss. is minimized, but the [Al]diss. to prevent of alumina formation. The following example shows the way
of estimation of aluminium mass flow into the melt which is required to meet correctly the demanded final [Al]diss. value in the
steel just before solidification (Figure 3).
1) start-up oxygen content in the melt before de-oxidation,
[O]start: 100 ppm
demanded oxygen content at solidification, [O]final: 3 ppm
demanded diss. aluminium content in the as-cast strand,
[Al]final: 400 ppm
2) Aluminium is added into the melt to diminish [O]diss. to about
20 ppm, and the alumina is removed to the ladle top slag by
stirring: ∆[O]=80 ppm; the corresponding [Al]diss. concentration at 1,600 °C is about 0.002 wt% (resp. eq. (1) and (2)).
3) The demanded <Al> is fed by Al-wire into the melt during
pouring:
∆[Al]added = [Al]demand. - [Al]de-ox. = 400-20 = 380 ppm; (3)
if there is a leakage of oxygen the unwanted oxygen flux rate
of d[O]/dt must be taken into account.
The mass flow of aluminium is calculated by
(4)
(5)
(6)
The DDA method can be carried out by spot measurement of oxygen activity or by using a continuously working probe. For fully
automated [Al]-feeding that continuous method would be preferred but there are no commercial probes available. To overcome that lack a development of probes has been started by the
company Heraeus-Electronite and RWTH Aachen University
where the measuring period of a Celox®-probe based on electromotive forces was extended from regularly 15 s to 2 min [3].
EXPERIMENTAL SET-UP
The DDA method has been tested in 3 steps:
1) In the beginning the twin roll casting process was the goal to
approve castability and macro-cleanliness of carbon steel grades.
The near-net-shape casting process and laboratory experimental
simulations have been described in an earlier publication [8]: The
furnace, the runner system and the casting tundish including the
SEN of the RWTH Aachen twin roll caster at the Dept. of Metal
Forming were copied for pouring experiments at the Dept. of Ferrous Metallurgy. There, a 500 kg induction furnace prepared the
Colata continua
FIG. 3
Scheme of the Dynamic Deoxididation and Alloying (DDA)
method.
Schema del metodo di
Disossidazione Dinamica e
Alligazione (DDA).
FIG. 4
Laboratory trial at 0.5 t induction
furnace at IEHK of RWTH Aachen
University.
Processo di laboratorio con forno
ad induzione da 0.5 t presso l’IEHK
dell’Università RWTH Aachen.
carbon steel melt, and pouring experiments were carried-out
using the spot and continuous measurements of oxygen activity.
The Al-wire with a diameter of 4 mm has been continuously added
to the pouring melt with the dmAl/dt-value which was estimated
by eq. (3 - 6) (Figure 4). Further experiments with similar set-ups
and different heat sizes were carried out to confirm the results.
2) After those successful experiments the step of validation at
a real twin roll caster has been done. Industrial strip casters are
described in [9]. The Institute of Metal Forming (IBF) runs a Bessemer type twin roll caster in cooperation with ThyssenKrupp
Steel AG. The pilot plant-size caster consists of two water cooled
steel rolls, each with a nickel coated copper sleeve, resulting in
an outer diameter of 590 mm and a width of 150 mm [10]. Casting duration is about 3 min limited by the furnace capacity.
In twin roll casting of steel strip, steel melt is poured into the
roll nip of the counter-rotating rolls. A ‘melt pool’ is formed by
the rolls and two ceramic side dams. The melt solidifies on the
cold surfaces of the rolls, forming two layers of solid steel which
are combined at the narrowest position by a slight force. The
melt flows from the induction furnace with a capacity of 165 kg
to a runner system and subsequently into the tundish; from here
the melt is guided to the melt pool through a small SEN. To minimize re-oxidation, the melt is sealed from ambient air by liquid argon addition during the melting process and during the
flow from furnace to pool (Figure 5).
After melting and killing with [Si] and [Mn], a sample was taken
and the oxygen activity in the melt was measured by a Celox®la metallurgia italiana - n. 11-12/09
system. The result was used to calculate the demanded mass of
aluminium for further pre-deoxidation to [O]diss.=20 ppm. For
comparison, the chemical composition of a synchronous sample
from the melt was measured using a spark emission spectrometer. This procedure was repeated after 4 min to evaluate the
efficiency of the pre-deoxidation. After reaching the superheat
temperature of 1,700°C the furnace was tilted with a defined
rate to start -up the cast. The dynamic alloying of Al-wire with
adiameter of 4 mm in the tundish started 25 s after the first
melt-roll contact. The spooling-in was carried-out manually. The
feeding rate of the Al-wire was calculated as a function of steel
mass flow, the content of residual dissolved oxygen and the final
aimed [Al]diss. concentration. The feeding rate was about 1.12
cm/s, as seen in the following calculations:
(7)
(8)
(9)
A cold Al-wire will freeze-up a thin layer of steel before it starts
to smelt; the time of re-smelting is about 1…2 s depending on
the diameter of the wire and steel superheat [11].
67
Memorie
FIG. 6
Left: Billet mould with open
pouring melt stream and guiding
tube for Al-wire;
right: Al-wire feeding device.
Sinistra: stampo della billetta con
getto di metallo fuso e tubo guida
per il filo di Al;
destra: dispositivo di svolgimento
del filo di Al.
3) Following the twin roll casting tests the application of DDA to
an industrial billet caster was done to investigate the long term
reliability and metallurgical precision. The 3-strand CC machine
produces carbon steel billets with 130 mm square cross section.
The steel melt is prepared in an EAF and in a ladle furnace. In
those trials an oil lubrication in the moulds has been applied.
The metering nozzles in the tundish bottom had diameters of 18
mm. No further shrouding has been used. An aluminium wire
with 2.5 mm diameter was fed into the free falling pouring jet directly above the meniscus in the mould by a commercial spooling machine which is normally used in automatic welding
systems (Figure 6).
The pre-deoxidation of the melt with [C]=0.18 wt%, [Si]=0.24 wt%
and [Mn]=0.96 wt% has been done by Si-Mn-alloying in the ladle.
The steel temperature in the tundish was 1,577 °C, and the
Celox® measured [O]diss.-concentration was 36 ppm.
The calculation of the wire spooling velocity which was demanded
to reach the final [Al]diss. concentration of 400 ppm concerns to
the scheme mentioned above (eq. 7-9); the wire velocity of 63.78
m/min resp. 1.063 cm/s was adjusted at the spooling device.
RESULTS
1) Laboratory furnace: The results of the laboratory test of 500
kg-induction furnace equipped with a runner and tundish system have been reported in [8]. Laboratory tests using a copper
sampler for catching steel and inclusions were performed. 95 %
of the rapidly solidified steel samples show inclusion diameters
of less than 5 µm when the inline Al-wire application was used,
in comparison to a value of 14 µm at conventional [Al] alloying in
the furnace (Figure 7). The [Al] yield increased from appr. 35 %
to more than 92 %.
2) Twin roll casting: In the first trial, the dual-phase steel grade
DP600 was prepared. After melting, [O]diss. was approx. 87 ppm.
Considering the steel mass of 165 kg, the amount of <Al> needed for the pre-deoxidation was calculated to 21 g, assuming a
yield of 50 %. After 4 min the concentrations were [O]diss.=20.2
ppm oxygen (by Celox®) and corresponding [Al]diss.=18 ppm at
1,600 °C.
After 25 s the inline-alloying with Al-wire into the tundish was
started. Strip casting speed was about 40 m/min. After the process became stable, the Al content was distributed with a mean
value of 0.047 + 0.009 wt% in the as-cast strip; the yield of
[Al]diss. in the as-cast steel could be determined at 95.8 %. The deviation is explained by the not constant Al-wire feeding rate by
manual spooling. In a 2nd trial, the feeding rate of the wire was
changed in a controlled way during the casting time. The dissolved [Al] content was similar to the results of the first trial.
The [Al] concentration in the strip reacted with a small delay to
68
changes in the feeding rate.
After the trials, the surface of the as-cast low carbon strip was
inspected; no slag spots or cracks were found. No extraordinary
signs of wear or skulls were observed at refractory material of
tundish, side-dams, or SEN. The microstructure of the strips was
observed and compared to conventionally alloyed strip with
same casting conditions. The identification of size and distribution of the inclusions was carried-out by 2-dim. metallographic
analysis of steel samples on determined positions in the strip
length the images were interpreted by commercial image analysis software, which counts the number and size of inclusions
on a defined area. Energy Dispersive X-ray Analysis confirmed
that the majority of the inclusions are Al2O3 particles which have
been formed in the melt.
The investigated metallographic samples indicate that the sizes
of inclusions formed during the conventional alloying technique are bigger compared to the inline alloying technique [3]:
The DDA method made sure that alumina particles are finer dispersed as in conventionally alloying in the furnace. Furthermore, the inclusions formed during the conventional technique
tend to cluster due to collision mechanisms during the flow
from furnace to the casting rolls. In both, tundish and SEN, turbulent flow patterns dominate which promote turbulent collisions between particles and result in inclusion growth by
agglomeration.
By inline alloying of Al-wire into the tundish, the free residualoxygen was killed by the spooled-in wire. Since these newly formed particles have a short residence time, and the collision
events are restricted, too. Those particles flow directly through
the bottom orifice of the tundish through the SEN into the pool
between the rolls.
The maximum inclusion size obtained by the inline alloying technique was found to be 11.4 µm (Figure 8). With the conventional technique, clusters with sizes up to 76 µm have been
FIG. 7
Cumulative frequency of alumina grain sizes.
Frequenza cumulativa delle dimensioni dei grani di
allumina.
Colata continua
FIG. 8
Cumulative frequency of alumina grain sizes: comparison of ‚inline alloyed‘ (DDA) and conventionally de-oxidized and
alloyed samples. Left: steel grade ‚Dual Phase DP 600’. Right: steel grade ‚Low Carbon‘.
Frequenza cumulativa delle dimensioni dei grani di allumina: confronto fra provini con alligazione in linea (DDA) e con
disosossidazione e alligazione convenzionale. Sinistra: acciaio tipo bifasico DP600; Destra: acciaio tipo a basso carbonio.
detected. Inline alloying by DDA method in the tundish diminishes the total residence time in comparison to conventional alloying in the furnace, and so the agglomeration time of alumina
particles. The diagram indicates that in the inline alloyed strip
90 % of the inclusions are smaller than 4 µm and in contrast,
90% of the inclusions in the conventional technique are smaller
than 15 µm. The slight slope of the cumulative frequency curve
of conventional alloyed strip is again explained by the growth
rate of the inclusions between the furnace and the roll nip. A similar observation was carried out at samples from another
length position of the as-cast strip. During inline alloying, a
small amount of particles, which do not flow directly through
the orifice and circulate by turbulence, meet an estimated time
of residence of t=22 s [12]. According to the model of Zhang and
Lee [13] particles can grow in this time to a maximum size of
9.6 µm.
In all cases, no clogging occurred during the casting period of 3
min. The yield of [Al]diss. could be increased by the prevention of reoxidation due to inline alloying from appr. 52 % to more than 92 %.
3) Industrial billet caster: The DDA method was applied to one
mould for the duration of 6 min. No casting problems occurred.
The Al-wire was spooled exactly into the free falling melt jet to
prevent of asymmetric enrichment of [Al] in the solidifying shell.
For metallographic investigations the corresponding as-cast billet with a length of 11.3 m was investigated; three cross section
samples with 10 mm thickness from positions ‘head’, ‘middle’,
and ‘tail’ have been taken.
The chemical analysis was carried-out by emission spectroscopy
as well as by chemical analysis. The concentration of [Al]diss. was
0.0366 wt%+6 ppm versus length and +4…6 ppm in the cross
sections; compared to the aimed value of 0.040 wt% of [Al]total
the ratio of [Al]diss./[Al]tot. was 366/400*100 %=91.5 %. Those results show that the pre-calculated wire feeding rate and the application had been adjusted correctly.
The metallographic analysis resulted in fine disperse alumina
which were identified by EDX-analysis in a scanning electron
microscope in addition; in some cases particles containing Si
and Mn have been found. No cluster formation could be obserla metallurgia italiana - n. 11-12/09
FIG. 9
Cumulative frequency of alumina grain sizes:
Results of 3 different billet cross section samples.
Frequenza cumulativa delle dimensioni dei grani di
allumina: risultati di provini relativi a 3 billette con
diverse sezioni trasversali.
ved. The distribution of the non-metallic particles is similar to
results of the laboratory and strip caster experiments. The size
of the alumina inclusions were almost between 1 and 2 µm, 90%
of the particles are smaller than 4 µm (Figure 9). A single bigger inclusion reached the size of 15 µm. This result is representative for all 3 investigated cross section samples in the
billet.
CONCLUSIONS
The dynamic de-oxidation alloying (DDA method) allows the
final de-oxidation and the alloying of carbon steel melt with aluminium on the point, that means with high yield of added aluminium, and with small deviations of [Al]diss. from the required
concentrations. By this method, big particles influencing the
macro-cleanliness of the steel can by avoided since the duration
of agglomeration controlled growth is decreased significantly.
Excess oxygen after pre-deoxidation and entrapped oxygen from
leakage in the shrouding systems is bound by alloyed aluminium
just before solidification so that the Al2O3 particles stay small, al69
Memorie
most less than 4 µm, influencing only the degree of micro-cleanliness. The biggest particles have a diameter of about 15 µm by
short-time-agglomeration.
During the 3 min cast with small SEN at the twin roll caster neither clogging nor deposition of alumina could be noticed; this
behaviour is also expected for long term castings. With continuous casting of billets no clogging could occur because the <Al>
alloying took place just above the meniscus. The feeding rate of
Al-wire can pre-calculated at high accuracy and by automatic
feeding control the recent casting speed can be taken into account. The feeding point must be adapted in a proper way to
avoid super-saturation in the billet cross-section to prevent
break-outs by weak shell. The distribution of [Al] in strip and
billet was sufficient in all cases. Smelting of Al-wire and solution
of liquid [Al] in the steel melt flow worked without any problem.
Good results of DDA method application are also expected at
other CC processes using relatively small SEN like Thin Slab Casting [14], or Single Belt Casting [15].
[11]
[12]
[13]
[14]
[15]
misation of strip casting conditions and surface conditions for coating”. European Commission Technical steel research EUR 22817
EN, 2007
Bode, O.: “Verbesserung der Vergießbarkeit siliziumfreier Stähle
durch eine Calciumfülldrahtbehandlung“. Diplomarbeit Technische
Universität Clausthal (durchgeführt bei Aktiengesellschaft der Dillinger Hüttenwerke), 1991
Senk, D.; Mavrommatis, K.: “Conditions of Liquid Steel Treatment
for Near-Net-Shape Casting Processes”, steel research 74 (2003), No.
3, pp. 153-160
Zhang, J. and H.G. Lee, “Numerical Modelling of Nucleation and
Growth of Inclusions in Molten Steel Based on Mean Processing Parameters”. ISIJ Internat., 44 (2004) 10, pp. 1629-1638
Flemming, G.; Hofmann, F.; Rohde, W.; Rosenthal, D.: “Die CSP-Anlagentechnik und ihre Anpassung an erweiterte Produktionsprogramme”. Stahl und Eisen 113 (1993), 2, pp. 37-46
Kroos, J.; Evertz, T.; Dubke, M.; Urlau, U.; Reichelt, W.; Trakowski,
W.; Spitzer, K.-H.; Schwerdtfeger, K.; Nyström, R.: “The Direct Strip
Casting process”, Proc. METEC Congr., 1999, Düsseldorf
ACKNOWLEDGEMENT
The authors express their thank to TSW Trierer Stahlwerke, who
allowed and supported the industrial billet caster trials, and to
ThyssenKrupp Steel AG who agreed to the trials at the pilot-plant
twin roll caster at IBF, RWTH Aachen University. The publication was generated in the cooperation of CHAMP, the Centre of
Highly Advanced Metals and Processes at RWTH Aachen University.
REFERENCES
[1] Cramb, A.W., R. Rastogi, and R.L. Maddalena, Nozzle clogging.
Chapter 9 in Making, Shaping and Treating of Steel, 11th Ed., Vol. 5,
Casting Volume, AISE Steel Foundation, Pittsburgh, PA, 2003, pp.
9.1-9.17
[2] Toulouse, C., Petry, S.; „Stable Oxygen Isotopes for Tracing the Origin of Clogging in Continuous Casting Submerged Entry Nozzles“;
Proc. 2nd AIM-Federeracciai-VDEh Joint Meeting on Metallurg. Fundament., Sep. 30, 2009, Duisburg, ThyssenKrupp Steel AG see also:
Toulouse, C., Pack, A., Ender, A., Petry, S.; steel research int. 79
(2008) 2, pp. 149-155
[3] Grosse, A.: “Entwicklung eines dynamischen Desoxidations- und Legierungsverfahrens für die Herstellung aluminiumberuhigter Kohlenstoffstähle beim Bandgießen”. Dr.-Ing. Thesis, IEHK, RWTH Aachen, 2009
[4] Rastogi, R.; Cramb, A. W.: Inclusion Formation and Agglomeration
in Aluminium-Killed Steels. 84th Steelmaking Conference proceedings, Baltimore, March (2001), pp. 789-829
[5] Huemer, K.; Wolf, G.; Sormann, A.; Frank, G.: Auswirkungen einer
Kalziumbehandlung auf die Entstehung und Zusammensetzung
von nichtmetallischen Einschlüssen bei der Erzeugung von aluminiumberuhigten Stählen für Langprodukte. BHM 150. (2005) 7, pp.
237-242
[6] Oeters, F.; “Metallurgie der Stahlherstellung“. Berlin [u.a.]: Springer
[u.a.], 1989
[7] Zhang, L.; Pluschkell, W.: Considerations on Nucleation and Growth
Kinetics of Inclusions during Liquid Steel Deoxidation. 6th International Conference on clean steel proceedings, Balatonfüred, Hungary (2002), pp. 107-115
[8] Grosse, A., Senk, D., “Deoxidation practice in twin-roll-casting of aluminium-killed carbon steels”. Proc. 7th Internat. Conf. on Clean
Steel, Balatonfüred, Hungary, 2007, pp. 254-263
[9] a) Herbertson, J.: “The emergence of strip casting - challenges and
impacts of success“, pp. 60-69
b) Campbell, P., Wechsler, R.: “The first commercial plant for carbon
steel strip casting at Crawfordsville”, pp. 70-79
c) Legrand, H., Albrecht-Früh, U., Stebner, G., Flick, A., Hohenbichler, G.,: “Stainless steel direct strip casting”, pp. 80-89
in Proc. Dr. Manfred Wolf Memorial Symposium, May 10-11, 2002,
Zürich
[10] Nicolle, R., Schmitz, W., Senk, D., Kopp, R. Porcu, G., et al.: „Opti-
70
Abstract
Disossidazione dinamica
e alligazione in linea di Al
nella colata continua
di billette e nastri
Parole chiave:
acciaio, disossidazione, colata continua
Nel presente lavoro è stato studiato il metodo di disossidazione graduale controllata e di alligazione con filo di Al dell’
acciaio al carbonio fuso. Il bagno fuso è stata pre-disossidato
in siviera e la frazione principale di inclusioni non metalliche è stata rimossa, mediante agitazione, da parte delle scorie superiori di siviera. La disossidazione finale e
l’alligazione si svolge appena prima della solidificazione
nella lingottiera di colata continua. L'efficienza di tale metodo è stata testata a tre livelli: in laboratorio, in impianto pilota, nell’ impianto industriale. Non si sono rivelati
svantaggi; le ricadute favorevoli consistono in un’elevata
quantità di [Al]diss., in un alto rendimento di <Al>, in una
migliore macro-pulizia e in un miglioramento della qualità
del processo, che ha permesso di evitare depositi e occlusioni. Mediante il metodo studiato, dovrebbe essere possibile
anche la produzione di tipi di acciaio al carbonio calmato con
Al mediante “near-net-shape caster” che, solitamente, prevede l'utilizzo di piccoli orifizi nella paniera e nell’ugello di
entrata sommerso (SEN).
Memorie
Metallografia
Metallographic Specimen Preparation
for Electron Backscattered Diffraction
G. F. Vander Voort
Electron backscattered diffraction (EBSD) is performed with the scanning electron microscope (SEM) to
provide a wide range of analytical data; e.g., crystallographic orientation studies, phase identification and
grain size measurements. The quality of the diffraction pattern, which influences the confidence of the
indexing of the diffraction pattern, depends upon removal of damage in the lattice due to specimen
preparation. It has been claimed that removal of this damage can only be obtained using electrolytic polishing
or ion-beam polishing. However, the use of modern mechanical preparation methods, equipment and
consumables does yield excellent quality diffraction patterns. The experiments discussed here covered a wide
variety of metals and alloys prepared mechanically using three to five steps, based on straightforward methods
that generally require less than about twenty-five minutes.
KEYWORDS:
mechanical specimen preparation, diffraction patterns, deformation, relief control, flatness
INTRODUCTION
Electron backscattered diffraction (EBSD) is performed with the
scanning electron microscope (SEM) to provide a wide range of
analytical data; e.g., crystallographic orientation studies, phase
identification and grain size measurements. A diffraction pattern
can be obtained in less than a second, but image quality is improved by utilizing a longer scan time. Grain mapping requires development of diffraction patterns at each pixel in the field and is a
slower process. The quality of the diffraction pattern, which influences the confidence of the indexing of the diffraction pattern,
depends upon removal of damage in the lattice due to specimen
preparation. It has been claimed that removal of this damage can
only be obtained using electrolytic polishing or ion-beam polishing.
However, the use of modern mechanical preparation methods,
equipment and consumables does yield excellent quality diffraction patterns without use of dangerous electrolytes and the problems and limitations associated with electropolishing and
ion-beam polishing. Basically, if mechanical preparation results in
quality polarized light images of non-cubic crystal structure elements and alloys (e.g., Sb, Be, Hf, α-Ti, Zn, Zr), or color tint etching
of cubic, or non-cubic crystal structure elements or alloys produces
high-quality color images, then the surface is free of harmful residual preparation damage and EBSD patterns with high pattern
quality indexes will be obtained. Because of the acute angle between the specimen and the electron beam (70 – 74°), exceptional
surface flatness is also necessary for best results.
Polarized light image quality is dependent upon the elimination
of preparation damage and upon the quality of the microscope
optics [1]. Consequently, always check the polarized light response of metals that will respond to polarized light, to verify
preparation quality before performing EBSD. For cubic metals,
etch first with a general-purpose reagent to confirm the nature
of the expected microstructure. Then, repeat the final polishing
step and use a color tint etch [1,2] to verify freedom from damage. EBSD is best performed with an as-polished, non-etched
specimen due to the steep angle to the electron beam, as sur-
George F. Vander Voort
Buehler Ltd, 41 Waukegan Rd, Lake Bluff, Il 60044 USA
face roughness can degrade the diffraction pattern. A well-prepared, un-etched specimen will exhibit a good grain-contrast
image with a backscattered electron detector [3]; another good
test for freedom from surface damage.
DEVELOPMENT OF PREPARATION METHODS
Specimen preparation methods for metals and alloys have been
developed [4] that yield excellent results using straightforward
methods that generally require less than about twenty-five minutes. High-purity metals require more preparation time than
alloys. Automated preparation equipment is recommended, as
the methods will be performed accurately and reproducibly. Manual (“hand”) preparation cannot produce flatness, phase retention and damage removal as easily as automated processing
and is less reproducible.
Successful preparation requires that sectioning be performed
with equipment and consumables that minimize damage. Sectioning is a violent process and it can introduce massive damage.
Crystal structure does influence damage depth; face-centered
cubic metals exhibit greater damage than body-centered cubic
metals for the same preparation procedure because fcc metals
slip more readily than bcc metals. Use only abrasive blades designed for metallography that are recommended for the specific
metal/alloy in question. A precision saw yields even less damage
as the blades are much thinner and the applied loads are much
lower. Cutting with machines and blades/wheels that introduce
minimal damage is the most critical step in generating damagefree metallographic surfaces; this cannot be over-emphasized.
Then, commence grinding with the finest possible abrasive and
surface that will make all of the specimens in the holder co-planar and remove the sectioning damage in reasonable time. This
is the second critical rule for obtaining damage-free polished surfaces. The proposed methods utilize flat, woven cloths or pads
that minimize relief problems. To minimize damage, use less aggressive surfaces, such as silk, nylon, polyester or polyurethane.
The specimen preparation method must remove all scratches. If
scratches are present, so to is damage below the scratch. Scratch
depths produced in grinding and polishing are not uniform. A
deep scratch will have deep deformation below it. The preparation method must remove the scratches and the underlying da71
Memorie
mage in order to obtain high quality EBSD patterns.
The experiments discussed here covered a wide variety of metals
and alloys prepared mechanically using three to five steps. The
EBSD patterns shown were developed using both the EDAX-TSL
and Oxford Instruments HKL systems on a variety of scanning
electron microscopes (SEM) using tungsten, LaB6 and field emission electron sources. The plane-of-polish was oriented between
70 and 74° from horizontal, depending upon the system used.
The TSL system generates pattern quality indexes, PQI, and the
results shown here are the average and 95% confidence limits for
25 randomly selected grains using unetched specimens. The
high-purity metallic samples were analyzed using the HKL
Channel 5 EBSD system. These patterns were evaluated using
the band contrast data, with the average and standard deviation
calculated for a number of measurements. Several cast specimens had very large grains, so only a few EBSD patterns could
be obtained. The silicon specimen was a single crystal so all patterns were basically identical.
RESULTS
The first examples presented will be a wrought, cold worked,
high-purity (99.999%) aluminum and an Al – 7.12 % Si casting
Surface
CarbiMet
UltraPol silk
TriDent Polyester
TriDent Polyester
MicroCloth
MicroCloth
alloy. Al is a difficult EBSD subject as the low atomic number is
inefficient in generating backscattered electrons. High-purity
metals are always far more difficult to prepare than commercialpurity metals while alloys are the easiest to prepare. EBSD patterns will be more difficult to generate on a wrought,
non-recrystallized, cold worked specimen due to the resulting
distortion of the crystal lattice. So, combining both the high-purity and non-recrystallized conditions makes for an extreme test
of the preparation method. The table below presents the test method used, except that the specimen in this case was not subjected to a vibratory polish after use of the five-step preparation
method. The band contrast value averaged 151.1 after using the
five-step method. It is our experience, as discussed below that
using a 20-minute vibratory polish after the standard preparation cycle will improve the band contrast at least 10%. Longer
times will yield further improvements. When developing grain
maps, maximizing the band contrast, or the pattern quality
index, produces greater confidence in indexing; this is vital
when indexing several hundred points per second.
Shown below in Figure 1 is the cold worked microstructure of
the high-purity aluminum specimen.
The next example is the as-cast Al –7.12% Si alloy, prepared by
Abrasive Size
Load Lb (N)
Platen Speed/Direction
Time (min.)
240-grit SiC water cooled
5 (22)
240 rpm Contra**
1 per sheet
5 (22)
150 rpm Contra**
5
µm MetaDi Diamond*
1-µ
5 (22)
150 rpm Contra**
5
5 (22)
150 rpm Contra**
3
µm MasterMet
0.05-µ
5 (22)
150 rpm
3
-
VibroMet2
≥20
µm MetaDi Diamond*
9-µ
µm MetaDi Diamond*
3-µ
µm MasterMet
0.05-µ
* Add MetaDi Fluid lubricant (charge with paste and MetaDi Fluid, then add MetaDi Supreme suspension during the cycle)
** Contra means that the platen and the specimen holder rotate in opposite directions.
TAB. 1
Preparation Method for High-Purity Aluminum.
Metodo di preparazione per alluminio di elevata purezza.
a
FIG. 1
b
Microstructure of cold worked 99.999% Al; a) Keller’s reagent, Nomarski DIC; b) Barker’s reagent, 20 V dc, 2
minutes, polarized light plus sensitive tint.
Microstruttura di Al 99,999% lavorato a freddo; a) reagente di Keller Nomarski DIC; b) reagente di Barker, 20 V dc, 2 minuti,
luce polarizzata più colorazione.
72
Metallografia
b
a
FIG. 2
a) EBSD pattern for α-Al in as-cast Al – 7.12% Si – pattern quality index: 87 ± 4.2;
b) light micrograph of as-cast Al-7.12% hypoeutectic alloy etched with 0.5% HF in water.
a) diagramma EBSD per a-Al in Al – 7.12% Si as-cast – indice di qualità del diagramma: 87 ± 4.2;
b) micrografia di lega ipoeutettica Al-7.12% as-cast sottoposta ad attacco con 0.5% HF in acqua.
Abrasive/Size
Load
lbs. (N)
Speed
rpm/Direction
Time
(min.)
240 (P280) grit SiC
water cooled
6 (27)
U.P.
UltraPol or TriDent cloths
9-µm MetaDi diamond*
6 (27)
TriDent or TexMet pads
6 (27)
TriDent or TexMet cloths
6 (27)
240
Contra
150
Contra
150
Contra
150
Contra
0.05-µm MasterMet
Colloidal silica suspension
6 (27)
(7 lb/31 N for)
ChemoMet
-
Surface
CarbiMet
MicroCloth or ChemoMet pads
MicroCloth
0.05-µm MasterMet
5
5
4
150
Contra
3
VibroMet2
≥20
TAB. 2
Preparation Method for High-Purity Copper.
Metodo di preparazione applicato al rame ad alta purezza.
the same five-step method, but with only 4 minutes for the 3µm step, and without vibratory polishing. The as-cast microstructure consists of α-Al dendrites and a eutectic of α-Al and
Si. The α-Al dendrites were sampled for the EBSD patterns. As
can be seen in Figure 2, an excellent quality diffraction pattern
was obtained from the alpha-Al dendrites. Figures 1 and 2 demonstrate that mechanical preparation is capable of producing
high quality EBSD patterns when properly performed.
Pure copper is extremely ductile and malleable. Copper and its
alloys come in a wide range of compositions, including several
variants of nearly pure copper for electrical applications that are
very difficult to prepare damage free. Rough sectioning and grinding practices can easily damage copper and its alloys and the
depth of damage can be substantial. Scratch removal, particula metallurgia italiana - n. 11-12/09
larly for pure copper and brass alloys, can be very difficult. If
the scratches are not removed, there will be damage beneath.
Following the preparation cycle with a brief vibratory polish
using colloidal silica is very helpful for scratch and damage removal. Attack polishing additions have been used in the past to
improve scratch removal but are not necessary using the contemporary method followed by vibratory polishing.
Table 2 lists a five-step method for preparing copper and its alloys (vibratory polishing is an optional 6th step). It is always helpful, particularly with alloys that are difficult to prepare damage
free, to etch the specimen after the fifth step, and then repeat the
fifth step. This reduces damage and gives better EBSD patterns.
Figure 3 shows a combined EBSD grain orientation map plus
index of quality map for tough-pitch copper (Cu with about 400
73
Memorie
a
FIG. 3
b
EBSD grain orientation maps plus index of quality maps for tough-pitch copper; a) maps with twins; b) maps after
twins were removed.
Mappature dell’orientamento dei grani tramite EBSD e indice della qualità delle mappe per rame ETP; a) mappe con geminati;
b) mappe dopo la rimozione dei geminati.
a
FIG. 4
b
Microstructure of wrought, annealed tough-pitch copper; a) etched with equal parts ammonium hydroxide and
hydrogen peroxide (3% conc); b) Beraha’s PbS tint etch, polarized light plus sensitive tint illumination.
Microstruttura di rame ETP trafilato e ricotto; a) dopo attacco con parti uguali di idrossido di ammonio e perossido di
idrogeno al 3%; b) con attacco colorante PbS di Beraha, luce polarizzata e illuminazione opportuna per la colorazione.
ppm oxygen) which reveals the grain structure and annealing
twins. Figure 3 also shows the map after twins have been removed. Note that a few twins remained after image processing
that will be removed if the boundary angle requirement for a
twin is made slightly greater. This specimen was not etched. Figure 4 shows the specimen after etching for comparison. Measurement of grain size in twinned Cu and its alloys is nearly
impossible by light microscopy image analysis due to the inability to reveal all of the grain boundaries and twin boundaries,
except by color etching.
Figure 5 shows an EBSD pattern and the microstructure of
wrought cartridge brass, Cu – 30% Zn, that was cold reduced
50% in thickness and then annealed at 704 °C for 30 minutes
74
producing a coarse twinned α-Cu matrix. This is a relatively difficult alloy to prepare free of scratches and surface damage and
the EBSD pattern quality was superb. The method shown in
Table 2 was utilized to prepare this specimen except that the
times for the 3- and 1-µm steps were 4 and 3 minutes, respectively, followed by a 30 minute vibratory polish.
EBSD patterns can be developed for both phases in a two-phase
alloy, as long as preparation keeps both phases flat on the planeof-polish. If relief is present, such that one phase is recessed
below the surface, EBSD patterns will not be developed. As an
example, a specimen of Naval Brass, an α-β brass consisting
of Cu – 39.7% Zn – 0.8% Sn, was tested after etching which attacked the β phase. EBSD patterns could be generated from the
Metallografia
b
a
FIG. 5
EBSD pattern and microstructure of cartridge brass: a) EBSD pattern for Cu – 30% Zn – PQI: 221 ± 8.6; b)
microstructure of wrought, annealed Cu – 30% Zn etched with equal parts hydrogen peroxide (3%) and ammonium
hydroxide.
Diagramma EBSD e microstruttura di ottone per munizioni; a) diagramma EBSD di Cu – 30% Zn – PQI: 221 ± 8.6; b)
microstruttura di Cu – 30% Zn trafilato e ricotto sottoposto ad attacco con perossido di’idrogeno (3% conc) e idrossido di
ammonio in parti uguali.
a
c
b
FIG. 6
EBSD patterns and microstructure of Naval Brass; a) and b): EBSD patterns for the alpha and beta phase with PQIs of
118.5 ± 8.7 for α-Cu and 150.4 ± 20.7 β-Cu; c) microstructure after etching with 100 mL water, 3 g ammonium
persulfate, 1 mL ammonium hydroxide (α-Cu is the continuous phase).
Diagramma EBSD e microstruttura di ottone navale; a) e b) diagramma EBSD delle fasi alfa e beta con indici PQI di 118.5 ± 8.7,
per α-Cu e 150.4 ± 20.7, per β-Cu; c) microstruttura dopo attacco con 100 ml acqua, 3 g persolforato di ammonio, 1 ml
idrossido di ammonio (fase continua α-Cu).
α phase, but not from the recessed β phase. Re-polishing and
running the specimen unetched produced excellent results for
both the α and β phases as shown in Figure 6. The specimen
was prepared in the same manner as used for the cartridge brass
specimen.
EBSD maps can be made using a number of techniques. Figure
7 shows a grain orientation map, an index of quality map, the
combination of these two maps, and a grain-orientation map
where the colors have been assigned based on crystal orientation using an inverse pole figure.
Perhaps the most difficult metals and alloys to prepare for EBSD
have been zirconium and its alloys. Numerous approaches have
75
Memorie
a
b
c
d
FIG. 7
Various EBSD maps for the Naval Brass specimen.
Varie mappature EBSD per il provino di ottone navale.
been tried. Table 3 presents the method used that yielded excellent grain maps of high-purity Zr and Zr alloys. The SiC paper
was coated with paraffin wax before grinding. Final polishing
was performed using a 5 to 1 ratio of colloidal silica to hydrogen peroxide (30% conc.). In this experiment, the vibratory step
was used (30 minutes).
Figure 8 shows two maps of high-purity (99.99%), annealed Zr.
The first was constructed by adding an all Euler grain map with
a band contrast map; the second shows an inverse pole figure
map, plus grain boundaries, with the grains with missing pixels
(black spots in the first map) filled in. The band contrast averaged 92.34 for the area shown.
76
Table 4 summarizes PQI results for a number of metals and alloys evaluated, many of which are difficult to prepare. These results clearly show that mechanical specimen preparation, if
properly performed, is fully capable of producing damage-free
surfaces that yield acceptable EBSD patterns that can be indexed
reliably. The Ni-based superalloys (Carpenter’s Custom Age 625
Plus and the fine-grained 718) contained sub-microscopic strengthening phases (the latter also contains copious delta phase)
that make the EBSD analyses more difficult. The pure tantalum
specimen was a P/M specimen that was not fully dense.
A second set of experiments evaluated the band contrast of eighteen (18) high-purity (generally >99.95%) specimens prepa-
Metallografia
Abrasive/Size
Load
lbs. (N)
Speed
rpm/Direction
Time
(min.)
240 (P280) grit SiC water cooled
320 (P400) grit SiC water cooled
0.05-µm MasterMet Colloidal silica suspension
0.05- m MasterMet
5 (22)
5 (22)
6 (27)
6 (27)
6 (27)
6 (27)
-
240 Contra
240 Contra
200 Contra
200 Contra
200 Contra
200 Contra
VibroMet2
U.P.
1
10
7
5
7
≥20
Surface
CarbiMet
CarbiMet
UltraPol cloth
TriDent cloth
TriDent cloth
MicroCloth pad
MicroCloth
TAB. 3
Preparation Method for High-Purity Zr and Zr Alloys.
Metodo di preparazione per Zr e leghe Zr ad alta purezza.
FIG. 8
Two examples of grain maps for
high-purity (99.99%) Zr.
Due esempi di mappatura dei
grani per Zr ad alta purezza
(99.99%).
red using methods typical of those shown above, or similar methods, usually with five steps (four for Ti). These specimens varied from Mg (atomic number 12) to Bi (atomic number 83) and
covered the range of metallic crystal structures: body-centered
cubic (6), face-centered cubic (4), hexagonal close-packed (5),
diamond cubic (1) and rhombohedral/trigonal (2). Table 6 lists
the specimens prepared using our standard methods and analyzed. Results for six of these after vibratory polishing are shown
in Table 5.
Specimens of pure Sb, V and Zr were susceptible to SiC embedment, even though the grit size was coarse, e.g., 240- and 320grit. Hence, grinding was repeated after coating the paper with
paraffin wax. Attack polishing was used, mainly with 30% conc.
H2O2, for the last step for preparing Cr, Nb, Ti, W and Zr. MasterMet colloidal silica was used for the last step, except for preparing
Fe (MasterPrep alumina was used) and Mg (water-free MasterPolish was used). Oil-based diamond suspensions (9-, 3- and 1-µm)
were used to prepare the high-purity (99.999%) Mg. For the Bi and
Pb pure specimens, grinding used four steps: 240-, 320-, 400- and
600-grit SiC paper coated with paraffin wax with low loads, followed by three polishing steps using 5-, 1- and 0.3-µm alumina slurries and a final polish with MasterMet colloidal silica. All
polishing steps used MicroCloth synthetic suede cloth. Although
the Bi produced an excellent EBSD pattern, none was obtained
77
Memorie
Metal/Alloy
α-Al in Al-7.12% Si
Cu–39.7% Zn–0.8% Sn
Elgiloy (Co-based)
Si Core Fe B
2205 Duplex SS
Ni-200
Nitinol (Ni-Ti)
Fine Grain 718 (Ni-base)
Pure Nb
Pure Ta
W in W-27 Cu
Pure Pb
TAB. 4
PQI ± 95% CL
Metal/Alloy
α-Cu in Cu-30% Zn
Cu–39.7% Zn–0.8% Sn
Pure Fe
316 Stainless Steel
2205 Duplex SS
HyMu 80 (Ni-base)
CA625 Plus (Ni-base)
Pure Cr
Pure V
CP Ti ASTM F67 Gr2
Pure Bi
Pure Ru
87 ± 4.2
118.5 ± 8.7 for α
221.4 ± 7.4
199.9 ± 7.4
248 ± 15.4 for α
176.3 ± 17.6
58.7 ± 4.3
80.7 ± 4.4
166.2 ± 17.1
169.7 ± 13.0
296.9 ± 20.1
49.3 ± 3.0
PQI ± 95% CL
221 ± 8.6
150.4 ± 20.7 for β
249.6 ± 5.5
184.9 ± 8.5
207.9 ± 11 for γ
196.7 ± 7.2
200.5 ± 6.5
259.8 ± 13.1
125.9 ± 10.3
119.1 ± 4.1
86.2 ± 1.8
266.2 ± 21.8
Pattern Quality Index Values for Various Metals and Alloys.
Valori dell’indice di qualità del diagramma per diversi metalli e leghe.
High-Purity Elements
Mg
Al
Si
Ti
V
Cr
Fe
Ni
Cu
Zn
Zr
Nb
Ru
Sb
Ta
W
Pb
Bi
Atomic Number
12
13
14
22
23
24
26
28
29
30
40
41
44
51
73
74
82
83
Crystal Structure
hcp
fcc
diamond cubic
hcp
bcc
bcc
bcc
fcc
fcc
hcp
hcp
bcc
hcp
rhombohedral
bcc
bcc
fcc
rhomb./trigonal
Band Contrast (0-255)
161.2
151.2
205.75
134.0
102.2
88.27
105.4
85.0
122.6
170.8
77.3
145.6
66.0
180.2
122.8
91.6
No Pattern
255
TAB. 5 Band Contrast Values for 18 Pure Metals.
Valori della Banda di Contrasto per 18 metalli puri.
High-Purity Element
Mg
Si (single crystal)
Ti
Ni
Nb
Pb
Mean Band Contrast (0 to 255)
Standard Method
Standard + Vibratory Polish
161.2
205.75
134.0
85.0
145.6
No pattern
175.25 (+8.7%)
233 (+13.2%)
146.2 (+9.1%)
102.8 (+20.9%)
151.2 (+3.8%)
108.0
* A 60 minute vibratory polish was used for the lead specimen.
TAB. 6 Band Contrast Improvement Due to Vibratory Polishing (20 min.*).
Miglioramento della Banda di Contrasto dovuta a lucidatura a vibrazioni (20 min.*).
78
Metallografia
with the pure Pb specimen. A one-hour vibratory polish with MasterMet colloidal silica using a MicroCloth pad was required to
obtain a diffraction pattern for Pb.
A two-minute chemical polish is normally used after mechanical
polishing of Zr; so EBSD was conducted on a second specimen
after chemical polishing. Surprisingly, no pattern could be obtained on the chemically polished specimen. The chemical polish
improved polarized light response but introduced grain faceting
(excessive relief). It has been reported that using heavy pressure with the same chemical polish minimized relief and yielded
good EBSD grain maps. The result for pure Zr in Table 5 was obtained on the same specimen as illustrated above in Figure 7,
but after an earlier preparation attempt with a less effective preparation method than presented in Table 3. The average band
contrast for the high-purity Zr specimen using the method in
Table 3 was 92.34 and ~90% of the pixels produced indexable
diffraction patterns. For the results published in Table 5, the average band contrast was 77.3 and only about 20% of the pixels
yielded indexable diffraction patterns.
Five specimens were evaluated after our standard preparation
method and then after a subsequent 20 minute vibratory polish
to determine the degree of improvement that can be obtained. If
the method used to prepare the specimens is not as good as what
was used in our work, then the vibratory polish will produce a
greater improvement. Longer times will also yield greater improvements. Table 6 summarizes these test results. Vibratory
polishing improved the band contrast of the first five elements
tried by an average of 11.1%; patterns could not even be obtained
with lead without a vibratory polish.
Details on the preparation methods used to prepare these alloys,
and others, can be obtained from the author or at the web site:
www.buehler.com.
REFERENCES
[1] G. F. Vander Voort, Metallography: Principles and Practice, ASM International, Materials Park, OH, 1999; originally published by
McGraw-Hill Book Co., NY, 1984.
[2] G. F. Vander Voort, “Color Metallography,” Vol. 9 ASM Handbook,
Metallography and Microstructures, G. F. Vander Voort, ed., ASM
International, Materials Park, OH, 2004, pp.493-512.
[3] G. F. Vander Voort, “The SEM as a Metallographic Tool,” Applied Metallography, G. F. Vander Voort, ed., Van Nostrand Reinhold Publishing Co., Inc., NY, 1986, pp. 139-170.
[4] G. F. Vander Voort, et al., Buehler’s Guide to Materials Preparation,
Buehler Ltd, Lake Bluff, IL, 2004, 135 pgs.
Abstract
Preparazione di provini metallografici per diffrazione
con elettroni retrodiffusi
Parole chiave:
metallografia, microscopia elettronica, prove
La Diffrazione con elettroni retrodiffusi (EBSD) viene realizzata con il microscopio elettronico a scansione (SEM) per fornire
un’ampia gamma di dati analitici connessi, ad esempio, agli studi dell’ orientamento cristallografico, all’identificazione delle fasi
e delle dimensioni dei grani. Un diagramma di diffrazione può essere ottenuto in meno di un secondo, ma si può migliorare la
qualità dell'immagine utilizzando un tempo di scansione più lungo. La mappatura dei grani richiede lo sviluppo di diagrammi di
diffrazione per ogni pixel nel campo e si tratta di un processo più lento. La qualità del diagramma di diffrazione, che influenza
la affidabilità dell’indicazione contenuta nel diagramma stesso, dipende dall’eliminazione del danneggiamento del reticolo dovuta
alla preparazione dei provini. E 'stato affermato che la rimozione di tale danneggiamento può essere ottenuta solo con la lucidatura elettrolitica o la lucidatura a fascio ionico. Tuttavia, l'uso dei moderni metodi di preparazione meccanica, delle attrezzature
e dei materiali attualmente disponibili rende possibile la produzione di immagini di diffrazione con eccellente qualità senza
dover ricorrere all’uso di elettroliti pericolosi e senza i problemi e i limiti connessi con la lucidatura elettrolitica e con fascio ionico.
In pratica si otterranno diffrazioni EBSD con indici di alta qualità se, successivamente alle preparazioni meccaniche, risulta possibile ottenere immagini di qualità con la luce polarizzata - nel caso di elementi e leghe a struttura cristallina non cubica (ad esempio, Sb, Be, Hf, α-Ti, Zn, Zr) - oppure risulta possibile produrre immagini a colori di alta qualità a seguito di attacchi chimici
coloranti - nel caso di elementi o leghe a struttura cristallina sia cubica che non cubica - ; si ottiene così una verifica del fatto che
la superficie è priva degli effetti nocivi di un danneggiamento da preparazione. Per ottenere i risultati migliori è necessario inoltre avere un’ eccezionale planarità della superficie, a causa dell’angolo acuto tra il provino e il fascio di elettroni (70 - 74°).
I procedimenti di preparazione dei provini sono dunque fondamentali e sono state messe a punto sequenze di operazioni - sia
per i metalli che per le leghe - semplici e di breve durata (dell’ordine dei 25 minuti). La prima applicazione qui presentata riguarda
l’alluminio ad alta purezza (99.999%), deformato e incrudito. L’alluminio in queste condizioni presenta difficoltà particolari per
l’esecuzione della diffrazioni EBSD a causa del basso numero atomico che implica difficoltà nel generare elettroni retrodiffusi;
inoltre va considerato che i metalli ad alta purezza sono di per sé difficili da preparare, soprattutto se incruditi e non ricristallizzati, quindi con distorsioni del reticolo cristallino. La procedura messa a punto per l’alluminio è stata poi applicata alla lega
per fonderia Al – 7.12 % Si, bifasica e allo stato “come fuso”.
Tenuto conto che il taglio e la preparazione metallografica con abrasivi può danneggiare in profondità materiali metallici duttili
e malleabili, sono state poi indagati il rame e le sue leghe, in particolare rame ETP (Electrolitic Tough Pitch), un ottone Cu – 30%
Zn (deformato al 50% e ricotto), un ottone navale Cu – 39.7% Zn – 0.8% Sn, bifasico. Per quest’ultimo si è proceduto ad una doppia indagine: solo sulla fase α (dopo attacco selettivo della fase β) e successivamente su entrambe le fasi. Una seconda sequenza
di indagine ha riguardato 18 provini costituiti da elementi ad alta purezza (generalmente >99,95%), che andavano, come numero atomico, dal magnesio al bismuto e come reticolo cristallino dal cubico a corpo o a facce centrate all’esagonale compatto e
al romboedrico/trigonale.
79

Memorie - Gruppo Italiano Frattura

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