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APPLICATION OF DUPLEX STAINLESS
STEEL FOR WELDED BRIDGE
CONSTRUCTION IN AGGRESSIVE
ENVIRONMENT
G. Zilli, F. Fattorini, E. Maiorana
Paper presented at the International Conference Duplex 2007, Grado, Italy, June 2007, organised by AIM
Maintenance costs are a significant item in life cycle of steel bridges, becoming of paramount
importance in aggressive environments. The use of duplex stainless steels for bridge decks would be a
major step forward in providing durable, low maintenance structures, exploiting both their corrosion
resistance and high mechanical properties, capable of meeting in full the required structural safety
performances. A research project partially funded by the EU research programme RFCS (Research Fund for
Coal and Steel, Bridgeplex contract RFS-CR-04040) is developing technical information on the use of
duplex stainless steel in welded bridge construction via mechanical testing and numerical analyses, so as to
provide indications suitable to form the basis for an upgrade of Eurocode 3 [1] and to allow a reliable Life Cycle
Cost analysis for this kind of structures so as to address the best material choice for the future bridges.
The project is still in progress but first results are available. This paper gives an overview of the project
and summarizes results obtained, deeper detailed in other papers presented at the International
Conference Duplex 2007 ([5] and [6]). In particular the paper is concerned with:
· overview of critical details in a welded bridge deck and relevant data available in literature also on austenitic
and austeno-ferritic steels; and
· economical evaluations considering maintenance aspects and fabrication costs showing the advantages of the
application of duplex stainless steel to defined bridge typologies.
Keywords: duplex, stainless steel, bridge, construction, life cycle cost, maintenance
INTRODUCTION
Service life beyond 100 years is today the target of major infrastructure
projects in the world, such as the longer and longer
metallic suspension bridges. The capital investment involved is
very high and planned maintenance costs are of overall importance
for the return on investment. Both safety and reliability
become also of paramount importance because any temporary
closure is very expensive both in direct maintenance and
repair and in traffic interruption.
Giuliana Zilli
Centro Sviluppo Materiali s.p.a., Italy
Francesco Fattorini
Centro Sviluppo Materiali s.p.a., Italy
Emanuele Maiorana
OMBA Impianti & Engineering s.p.a., Italy
The aforementioned reasons lead to strongly consider duplex
stainless steels as construction material owing to their
expected intrinsic corrosion resistance also in very aggressive
atmosphere, assured by their chemical composition (22Cr 5Ni
3Mo 0.2N), and their high mechanical resistance due to their
austeno-ferritic microstructure.
Together with its intrinsic high cost, a major barrier to the use
of duplex stainless steel in welded bridge construction is the
lack of experimental data on both their mechanical characteristics
and technological feasibility with respect to the specific
application, properties to be assessed if compared with
the vast know-how available for traditional carbon steels.
This paper will present an overview of the whole research activity
ongoing in the frame of RFCS programme, highlighting
the aspects investigated for the promotion of the use of duplex
stainless steel in bridge construction. While specific technical
aspects related with the ability of duplex stainless steel
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Consequently many research projects and experimental activities
have been devoted in the recent past to study the
in service behaviour of this complex steelwork leading to
design and execution recommendations. But all of them
were developed and verified on traditional constructional
steel grades (i.e. S355 [8], [9], [10] and [11]).
Although the duplex basic mechanical properties are well
known, it is not enough to promote this material for huge welded
bridge construction but, because of the relevance of such a
structure, more specific investigations on structural components
typical of bridge structure are needed. Presently it is not
possible to propose stainless steels for welded bridge construction
without having a similar experimental evidence of their
applicability, although from the LCC point of view these materials
could have some advantages with respect to the
more traditional solutions, when the expected service life is
prolonged beyond two big maintenance intervention [12].
The existing bridge chosen to have a comparison between the
utilization of carbon steel and duplex stainless steel, considering
both mechanical behaviour and durability during the whole
service life with the scope of evaluating its Life Cycle Cost
(LCC), is the Verrand viaduct (Fig. 1 and Fig. 2, [13]).
The Verrand viaduct whose owner is R.A.V. spa, built in 2000 by
OMBA of Torri di Quartesolo (Vicenza, Italy), is part of the
Mont Blanc-Aosta highway, connecting Mont Blanc Tunnel
with Morgex. The finishing of this part has permit to go to the
Tunnel by an highway broad. The viaduct needed the realizations
of long length spans, to have few intermediate piers,
as for geodetics problems as to leave untouched the environmental
and panoramic view: the Dora Baltea valley.
CRITICAL DETAILS IDENTIFICATION
s
Fig. 3
Welded details in orthotropic deck bridge.
Dettagli saldati di una lastra ortotropa.
Fatigue
The bridge deck is the structural part mainly subjected to cyclic
loads (both railway and roadway actions) so as in many
cases Fatigue Limit State [1] is the relevant one in design phase.
Bridge deck can be made of different construction typologies
but orthotropic deck is the most significant one in terms of fatigue
problems: it presents a great number of welded details
and some of them are particularly complex.
An orthotropic deck consists of prefabricated deck modules
welded at factory and joined together on site also by means of
welding. The top plate joints are always welded on site, while
beam elements joints can be either bolted or welded.
In the transversal section of the Verrand bridge steel deck (Fig.
2,double-beam orthotropic deck) the transversal beams (T
shaped section) are bolted; diaphragms and braces are made of
bolted T or L profiles. Its static scheme is the continuum beam
on a few supports. In Fig. 3 are shown the welded details selected
for fatigue testing in the research project, results are presented
in the paper [6] at the International Conference Duplex
2007. Here below some of those are described also giving details
on fabrication and welding procedures adopted, all being in accordance
with bridge construction practice and needs:
- The edges of the top plate to be joined on site are usually
butt welded with a back ceramic support without backing
run, to avoid the finishing of the weld on the back side. In
that case the welding process is mixed: a first pass using
the semi-automatic MAG – FCAW and the following passes
(2nd4nth) by the automatic SAW process. Clamps are needed
to align the plates and to keep the back ceramic support. The
clamps are bolted to threaded studs welded on the
bottom edge of the top plate, close to the edges to be joined (see
Detail A.5 and A.6 of Fig. 3).
- Corresponding to the transversal top plate joint of the
deck modules it is necessary also to replace the continuity of
the longitudinal ribs of the orthotropic deck: the way is to butt
weld on site a piece of rib using a support plate (see Detail B.2
of Fig. 3).
Large effort was made in the past for assessing the fatigue design
curves of full-scale components typical of orthotropic deck,
leading also to design indications incorporated in Eurocode
3 [1] for design of steel structures. Eurocode 3 [4] proposes the
S-N curves approach for fatigue design, and it classifies a set
of structural details assigning them specific design S-N curves.
These curves were defined on the
basis of historical experimental data
collected initially for carbon steel details,
the most general were also verified
for a few stainless steel grades.
Not so for structural details typical of
orthotropic deck.
s
Fig. 4
Bridge girders with open section (left) and close section (right) stiffeners.
Travi longitudinali con anima irrigidita.
Buckling
Typical elements of steel bridges,
i.e. the main longitudinal beams
(Fig. 4), have very high web subjected
to both bending and transversal
concentrated loads.
Web buckling is a primary design
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Acciaio inossidabile
s
Fig. 5
Maintenance
timing of painting
systems for S460
of two different
performance levels.
Programma
d’ispezione e
manutenzione per
la protezione dalla
corrosione di ponti
metallici in ambiente
di categoria C5.
lowing LCCs analysis a protective system for S460 steel bridge
is considered which is among the more traditional ones due
to the easier availability of data on. Maintenance scheduling
is reported in Fig. 5.
Some effects of alternative materials highlighted during
both the fabrication of steelworks for testing and the evaluation
of test results, are economically assessed in the present
LCC analysis
As regard the shop and yard productivity, the cost of austeno
ferritic is considered 15% higher that follows by the balance
between the faster welding rate and the more expensive welding
and cutting operations (see also paper [5] of the Conference).
The total quantity of the two material is the same as for the carbon
steel bridge as for the duplex bridge in accordance with
the available mechanical test results. The increment in the
fatigue behaviour of the austeno ferritic s.s. welded details
shown by the testing activities [5] is assessed in the following
LCC evaluations by not considering repair for fatigue costs
during service life of duplex bridge. Only inspection (each
year) and cleaning (every 9 years) are considered in the LCC
evaluation of duplex alternative.
Some of the effects of alternative materials are more difficult to
quantify in monetary terms, that is the case of users costs related
with the reduction of speed or complete closure of the
bridge. For example German Steel Association evaluates for
ordinary maintenance operations 20 days of speed reduction
from 120 km/h to 60 km/h, while for exceptional maintenance
operations 40 days of speed reduction are expected. What this
means in monetary terms is also difficult to be further evaluated
but this aspect should be listed with the others and taken
into account in the final evaluations.
The LCCs of both bridge alternatives are calculated in
present-value that means all costs are discounted to the
base time (time of bridge construction). The study period is
the expected service life for the bridge that is 100 years. LCC
analyses are calculated in constant monetary value (net of general
inflation). Bridge is treated as public utility infrastructure
(non-profit building) so income tax effects are not included
in the LCC analysis. The discount rate is a very sensitive parameter
for LCCs comparisons with money savings mostly
spreaded into the future, as in the present case study.
Here two different real discount rates (net of general price inflation)
are used in the LCCs analysis:
Study period
100 years
Real discount rate 3.2% and 1.8%
Investment cost data S460 EN 1.4462
Material cost 1’100 €/t 5’500 €/t (2006 price)
3’000 €/t (2001 price)
Shop cost 320 €/t 420 €/t
Yard assembly cost 160 €/t 185 €/t
Assembly equipments 200 €/t 200 €/t
Corrosion protective coating 35 €/m 2 0
Scaffolding and protections included
Maintenance cost data
Inspection 4 €/t 4 €/t
Cleaning 50 €/t -
Top coating
(high performance system) 25 €/m 2 -
Coating renewal 35 €/m 2 0
Scaffolding and protections included
Repair for corrosion
(% of initial investment) 5.16% -
Repair for fatigue
(% of initial investment) 12.3% -
User costs related with reduction of service or closure of
the bridge during maintenance
operations are not monetary evaluated but should be taken
into account in the comparison.
End of service resale S460 EN 1.4462
30% 75%
of material cost
The results of LCC evaluations are reported and compared in
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Acciaio inossidabile
in the model due to the difficulty of their monetary evaluation
(i.e. end user costs related with the bridge closure during
maintenance operations).
We have also compared two different discount rates: supposing
the price of duplex is 3’000 €/t as in 2001, considering the less
favourable discount rate (3.2%) we obtained quite same building
cost at the end of service life while initial investment is
recovered after about 50 years of service when considering
a more favourable discount rate (1.8%) is obtained. Moreover in
the comparison user costs related with reduction bridge service
during maintenance are not monetary evaluated.
In conclusion duplex stainless steel has many attractive
characteristics for bridge construction: corrosion resistance,
high strength and also aesthetics ones. All of those where demonstrated
for the specific application during the research
project. Duplex stainless steel can be also economically
attractive when considering whole service life costs: initial
capital expense is recovered after 50 years of service, provided
that producers can keep the price into the lower level of
the last years (i.e. 3’000 €/t).
ACKNOWLEDGEMENT
The authors wish to express their deep gratitude to the
European Commission for its financial support and to the
representatives of the other partners from INDUSTEEL Le
Creusot and from RWTH Aachen for their cooperation.
REFERENCES
1] ENV 1993-1-1. Design of steel structures. General rules –
Rules for buildings.
2] ENV 1993-1-4. Design of steel structures. General rules
– Supplementary rules for stainless steels.
3] ENV 1993-1-5. Design of steel structures. General rules –
Supplementary rules for planar plated structures without
transverse loading.
4] ENV 1993-1-9. Design of steel structures. General rules – Fatigue
design
5] A. FANICA and E. MAIORANA, UNS S32205 for bridge
construction: an experience of application”, Duplex 2007 Int.
Conf. Proc. Grado, Italy (2007).
6] O. HECHLER, M. FELDMANN, R. MAQUOI and G.
ZILLI, Bridge construction made in duplex stainless steel.
Duplex 2007 Int. Conf. Proc. Grado, Italy (2007).
7] A. MIAZZON, Large span bridges: the construction of steel
plated box girders. An example: the Storebaelt East Bridge.
Costruzioni Metalliche n.6, ACAI Servizi (2004).
8] S. CARAMELLI, P.CROCE, M.FROLI and L.SANPAOLESI,
Misure ed interpretazioni dei carichi dinamici sui ponti.
ECSC Project n. 7210-SA415 (F6.7/90).
9] S.J. MADDOX, The fatigue behaviour of trapezoidal stiffener
to deck plate welds in orthotropic bridge decks. TRL Report
No. SR 96
10] K. YAMADA, A. KONDO, H. AOKI and Y. KIKUCHI,
Fatigue strength of field-welded rib joints of orthotropic steel
decks. IIW doc. XIII-1282-88, Department of Civil Engineering,
Nogoya University, Nogoya (1998).
11] J.R. CUNINGHAME, Steel bridge decks: fatigue performance
of joints between longitudinal stiffeners. Report
No. LR 1066, 1982.
12] L. BRISEGHELLA, E. MAIORANA and A. MIAZZON. Duplex
stainless steel: an alternative for structural applications.
Costruzioni Metalliche n.1, ACAI Servizi (2004).
13] A. MIAZZON, The Verrand viaduct in Courmayeur, an
orthotropic deck bridge. Design, construction, assembly and
launching. Costruzioni Metalliche n.1, ACAI Servizi (2005)
14] L. RAMPIN, A. MIAZZON and others, Fatigue design in
steel bridges. XIX CTA Conf. Genua, Italy (2003).
15] B.JOHANSSON and A.OLSSON, Current design practice
and research on stainless steel structures in Sweden.
Jour. Const. Steel Res. 54, 3-29 (2000).
16] ASTM E 917-05. Standard Practice for Measuring Life-Cycle
Costs of Buildings and Building Systems.
APPLICAZIONE DELL’ACCIAIO INOSSIDABILE
DUPLEX NELLA COSTRUZIONE DI PONTI SALDATI
IN SITUAZIONI AMBIENTALI AGGRESSIVE
Parole chiave: acc.inox, corrosione, fatica, saldatura,
selezione materiali
I costi di manutenzione sono una voce rilevante nel ciclo di vita delle infrastrutture
metalliche, specialmente quando queste sono situate in ambienti
particolarmente aggressivi, per esempio per la presenza di cloruri in elevata
concentrazione. In ambiente marino del resto vengono tipicamente costruiti
i più grandi ponti sospesi per traguardare luci sempre maggiori (Akashi
ABSTRACT
Kaikyo in Giappone, Storebaelt East in Svezia): un’aspettativa di vita di
oltre 100 anni è il parametro di progetto per tali infrastrutture. Per garantire
ciò è necessario non solo proteggere le strutture metalliche con adeguati
sistemi in fase di realizzazione (Tab. 1), ma anche programmare ispezioni
e manutenzioni in maniera da mantenere l’opera in adeguate condizioni di
sicurezza durante tutto il ciclo di vita.
L’utilizzo di acciai intrinsecamente resistenti alla corrosione è un altro
modo per garantire l’adeguatezza agli standard di progetto, in quest’ottica
l’utilizzo di acciai inossidabili austeno-ferritici (duplex), con la loro elevata
resistenza alla corrosione unita all’alta resistenza meccanica, potrebbe costituire
un notevole passo avanti verso la sicurezza e dunque l’aspettativa
di vita in esercizio.
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WELD PROPERTIES
OF SANDVIK SAF 2707 HD ®
P. Stenvall, M. Holmquist
Super duplex stainless steels have found extensive use in the oil & gas industry and in other areas in the
(petro-) chemical processing industry. The recently developed hyper duplex grade Sandvik SAF 2707HD ®
allows extension of the application range of austenitic-ferritic alloys into even more aggressive conditions.
In most applications for Sandvik SAF 2707 HD the equipment needs to be welded. Hence, weldability is of
utmost importance for a stainless steel grade of this kind. Weld documentation was made for a number of joints
to simulate various tube- and pipe applications. The welding method used was gas tungsten arc welding.
The joints were tested regarding mechanical properties, microstructure, pitting resistance and in some
cases chloride stress corrosion resistance. The filler wire used, designated Sandvik 27.9.5.L, was developed
specifically for Sandvik SAF 2707 HD.
Overlay welds were produced using submerged-arc welding and gas tungsten arc welding. The welds were
documented regarding ductility, microstructure and pitting resistance. Tube-to-tube sheet welds were also
produced to document the weld behaviour and pitting resistance.
Keywords: duplex stainless steels, gas tungsten arc welding, submerged-arc welding, pitting corrosion, stress
corrosion cracking, tensile properties, impact toughness
INTRODUCTION
Super duplex stainless steels, such as UNS S32750, have been
used for more than 15 years in various industrial segments with
great success, e.g. offshore industry, oil refineries, chemical and
petrochemical industry, and pulp and paper production [1, 2, 3,
4]. However, environmental requirements and raised productivity
demands have, in many areas, forced the end-users into recirculation
of process streams, with increased temperatures and
increased pressures leading to more aggressive process environments.
In some cases the process environment has become too
aggressive for the super duplex grades. Therefore, a new hyper
duplex stainless steel has been developed for these aggressive
conditions – Sandvik SAF 2707 HD (UNS S32707) [5, 6]. The
typical chemical composition is shown in Tab. 1. Parallel to the
development of this grade a new welding consumable has been
developed, Sandvik 27.9.5.L [7]. Typical chemical composition is
shown in Tab. 1. The composition of the filler wire is similar to
that of the base material. However, the nickel content is higher
and the molybdenum and nitrogen contents are somewhat lower
in the wire in order to optimize the weld metal properties.
Weldability is an important feature for a duplex stainless steel
intended for tubular and flat products since welding is the most
common technique – and many times the only technique – for
joining. Therefore, welding and weldability of SAF 2707 HD
has been a vital part of the development work. So far two welding
processes have been documented – TIG (GTAW) and submerged-arc
welding (SAW). Some of the results are presented in
this paper.
EXPERIMENTAL
All-weld-metal
All-weld-metals were produced with both TIG and SAW. For
mechanical testing the weld metals were produced in grooves
according to AWS A5.9 and for the corrosion testing the weld
Product
Tube/pipe
Filler
Plate*
Designation
SAF 2707 HD
27.9.5.L
S355N
C (%)
0.01
0.01
0.15
Mn (%)
1
0.8
1.5
Cr (%)
27
27
-
Ni (%)
6.5
9
-
Mo (%)
4.8
4.6
-
N (%)
0.4
0.3
-
Others (%)
Co: 1
Co: 1
-
*) Low alloy steel plate used as base for overlay welding.
Peter Stenvall
Sandvik Materials Technology, Sweden
Martin Holmquist
Sandvik Materials Technology, The Netherlands
s
Tab. 1
Nominal chemical composition of SAF 2707 HD, filler
27.9.5.L and other material included in the investigations.
Composizione chimica nominale dell’acciaio SAF 2707 HD,
del filo d’apporto 27.9.5.L e dell’altro materiale impiegato.
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Saldatura
Welding
method
TIG
SAW
Flux
n.a.
15W
Shielding
gas
Ar + 2%N 2
n.a.
Ferrite
content (%)
45
56
s
Fig. 1
Joint type tested in tube-to-tube sheet welding.
Tipo di giunzione eseguita con saldatura tubo-piastra.
2. Determination of critical pitting temperature (CPT) was made
according to ASTM G48-03 Method E modified by Sandvik. (The
same specimens were used through out the CPT determination
instead of new specimens at each temperature as stated in ASTM
G48-03 Method E.) Here the specimens were cut out from the
surface of the tube sheet containing the TIG weld but not the
tube to avoid the crevice between the tube and the tube sheet.
Two specimens were used. The temperature increment was 2.5°C
and the testing started at 40°C. The specimens were brushed and
degreased but not pickled before testing.
RESULTS AND DISCUSSION
All-weld-metal
The results in Tab. 2 show ferrite contents at reasonable levels for
s
Tab. 2
Ferrite content in all weld metal measured with
linear analysis.
Contenuto di ferrite nella saldatura misurato mediante
analisi lineare.
both all-weld-metals. The ferrite contents are somewhat lower
for the TIG weld due to the nitrogen addition in the shielding
gas leading to higher nitrogen content in the weld deposit and,
hence, lower ferrite content.
Composition and alloying vectors of all-weld-metal produced
with SAW are presented in Tab. 3. The two elements subjected to
the largest relative changes are chromium and nitrogen, which
was expected. The burn-off of chromium is normally between 0.5
and 1 percent for flux 15W. High nitrogen filler normally loose
considerable amounts of nitrogen in submerged-arc welding.
Results of tensile testing are shown in Tab. 4. The yield and tensile
strengths are very high compared to those of 25.10.4.L (filler
for SAF 2507) where typical values for Rp0.2 and Rm are around
700MPa and 860MPa respectively for TIG.
The impact toughness of all-weld-metal produced with TIG,
shown in Fig. 2, is generally good and impact toughness above
150J at -60°C is very good bearing in mind that this is a very high
Product
Chemical analysis
Alloying vector
C (%)
0.020
+0.004
Si (%)
0.5
+0.1
Mn (%)
0.6
-0.2
Cr (%)
26.7
-0.4
Ni (%)
8.8
0
Mo (%)
4.5
0
N (%)
0.25
-0.05
Co (%)
1.0
0
s
Tab. 3
Chemical analysis and alloying vectors of all-weld-metal produced with SAW using the basic flux 15W.
Analisi chimica e vettori di alligazione nel metallo deposto mediante SAW, utilizzando il flusso basico 15W.
Weld method
TIG
SAW
Rp0.2 (MPa)
805
727
Rp1.0 (MPa)
867
804
Rm (MPa)
955
905
A (%)
31
25
Z (%)
69
45
s
Tab. 4
Tensile properties of all-weld-metal of 27.9.5.L welded with Ar + 2%N 2
.
Caratteristiche tensili del metallo deposto ottenuto con materiale 27.9.5.L sotto Ar + 2%N 2
.
Welding
method
TIG
SAW
Flux
n.a.
15W
Shielding
gas
Ar + 2%N 2
n.a.
CPT (°C)
77,5
70
Location
Top
Centre
Root
Ferrite content (%)
60
54
53
s
Tab. 5
Critical pitting temperature of
all-weld-metals.
Temperatura critica di pitting
del metallo deposto.
s
Tab. 6
Ferrite contents in weld metal of girth weld in tube,
25.4 x 1.65mm.
Contenuti di ferrite nel metallo deposto con saldatura circonferenziale
in tubi 25.4 x 1.65mm.
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Test temperature (°C)
RT
Specimen no.
1
2
Rm (MPa)
970
966
Location of rupture
Weld metal
Weld metal
s
Tab. 7
Results of tensile testing transverse girth weld in tube 25.4 x 1.65mm.
Risultati delle prove di trazione trasversale in tubi ( dim 25.4 x 1.65mm ) con saldatura circonferenziale.
Specimen no.
1
2
Attack temp. (°C)
67.5
70
Location of attack
Weld metal, top and root side.
Weld metal, top and root side.
CPT (°C)
67.5
s
Tab. 8
Result of CPT determination of girth weld in tube, 25.4 x 1.65mm.
Risultato della determinazione della CPT in tubi ( dim 25.4 x 1.65mm ) con saldatura circonferenziale.
s
Fig. 5
Microstructure in centre of weld metal in pipe weld.
Pipe dim. 168 x 7,1mm. Magnification: 150x.
Microstruttura al centro del metallo deposto in una saldatura di
tubazione ( dim. 168 x 7,1mm). Ingrandimento: 150x.
Ferrite contents in weld metal measured with linear analysis are
shown in Tab. 9. The level is within the rather common interval
specified by standards and end users, 35-65% ferrite.
Results of tensile testing are shown in Tab. 10. The ruptures are
located in the parent material about 15mm from the fusion line.
Face and root bend test according to ASME IX was carried out
to 180° with approved results. One fissure measuring 1.5mm appeared
in one root bend specimen. However, according to ASME
IX this is approved.
Critical pitting temperature of the pipe weld was determined to
60°C, see Tab. 11. This value is lower than that of the tube weld
described above, but it still is higher than that of SAF 2507 welds
where the CPT is around 50°C [9, 10, 11]. With a further optimisation
of the weld procedure used, a higher CPT for this type of
multi-layer joint weld should be possible.
The results of SCC testing according to ASTM G123 with U-bend
specimens according to ASTM G30 revealed no signs of stress
corrosion cracking after testing for 1008h. These results were
expected since duplex stainless steels normally have very good
s
Fig. 6
Microstructure in HAZ and fusion line in pipe weld.
Pipe dim. 168 x 7,1mm. Magnification: 150x.
Microstruttura nella ZTA e sulla linea di fusione in una saldatura
di tubazione( dim. 168 x 7,1mm). Ingrandimento: 150x.
Location
Top
Centre
Root
Ferrite content (%)
60
46
43
s
Tab. 9
Ferrite contents in weld metal of girth weld in pipe,
168 x 7.1mm.
Contenuto di ferrite nel metallo deposto con saldatura circonferenziale
in tubazioni ( dim 25.4 x 1.65mm).
resistance to chloride induced stress corrosion cracking.
Overlay welds
The basic flux designated 15W produce a surprisingly smooth
and sound overlay weld with no signs of porosity on the surface.
Slag removal was good and no slag remnants could bee noted.
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Saldatura
3mm).
These results indicate that a basic flux is needed to obtain acceptable
ductility in the overlay produced with SAW.
Critical pitting temperatures of the overlay welds are shown in
Tab. 13. The pitting resistance of the TIG weld overlay indicate
that more than 5 runs might be required. However, it should be
borne in mind that the corrosion specimen contains both top layer
and the layer underneath. The pitting attacks were located to
one side only most likely originating from layer no 4.
The overlay welds produced with submerged-arc welding show
very high pitting resistance. Here, in contrast to the TIG overlay
weld, the top layer is rather thick and a corrosion specimen can
easily be taken from the top layer. These CPT results are very
encouraging since SAW is a more productive welding process
compared to TIG. It should also be noted that the chromium
compensated flux, 10SW, did not give better CPT than the flux
without chromium, flux 15W.
Chemical analyses of the top layers show that the dilution from
the parent material is close to nil in the TIG weld. See Tab. 14. For
the submerged-arc weld there is a small dilution. For flux 15W
16 ottobre 2008

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Saldatura
Weld method
TIG
SAW
Flux
n.a.
15W
10SW (Cr comp)
No. of layers
5
3
3
Ferrite content (%)
53
60
51
s
Tab.12
FFerrite contents of top layers in overlay welds.
Contenuti di ferrite negli strati superiori delle placcature.
Welding
method
TIG
SAW
Flux
n.a.
15W
10SW (Cr comp)
No. of
layers
s
Tab.13
Results of CPT determination of overlay welds.
Risultati delle determinazioni della CPT per le placcature.
5
3
3
Attack temp. (°C)
Specimen 1
Specimen 2
62.5
65
75
72,5
70
72,5
CPT (°C)
62,5
72,5
70
Welding
method
TIG
SAW
SAW
Flux
n.a.
15W
10SW (Cr comp)
No. of
layers
5
3
3
C (%)
0.013
0.020
0.017
Mn (%)
0.8
0.6
0.5
Cr (%)
27.0
26.4
26.2
Ni (%)
8.8
8.6
8.4
Mo (%)
4.5
4.4
4.3
N (%)
0.30
0.24
0.26
s
Tab.14
Chemical analysis of top layers welded with
filler 27.9.5.L.
Analisi chimica degli strati superficiali saldati con materiale
d’apporto 27.9.5.L.
the composition is not far from that of all-weld metal in Tab. 3. It
is also interesting to note that the chromium compensating flux
10SW is not giving any higher chromium content compared to
flux 15W. Indeed, the dilution from parent material is somewhat
larger with flux 10SW but this fact cannot explain why there was
no effect of the chromium compensation flux.
Obviously flux 15W is the best flux for this purpose, giving better
weld bead appearance, approved bend test results and pitting
resistance equal to are better than that of flux 10SW.
Tube-to-tube sheet welds
The ferrite content in the tube-to-tube sheet weld was determined
to 33%. The microstructures of tube to tube sheet weld
metals, HAZ in tube and HAZ in weld overlay are shown in Fig.
9 and 10. The microstructure in Fig. 9 and ferrite content of 33%
indicate that the nitrogen content of the shielding gas can be lowered
to get a slightly higher ferrite level.
Determination of pitting resistance in tube-to-tube sheet welds
is difficult since the crevice between the tube and the tube sheet
needs to be completely removed in order to avoid crevice corrosion
during the pitting test. Here the testing was carried out successfully
and the CPT was determined to 60°C. See Tab. 15.
CONCLUDING REMARKS
It should be noted that the welded joints were not pickled,
s
Fig. 9
Microstructure in weld metal of tube-to-tube
sheet weld (TIG). Magnification: 300x.
Microstruttura del metallo deposto nella saldatura TIG
tubo-piastra . Ingrandimento: 300.
ground or polished after welding meaning that the testing was
made at fairly severe conditions. If the welds would have been
pickled the CPT level would most likely have been even higher.
However, the conditions used in these trials are more similar to
real conditions, even though pickling of the top side of the weld
is rather common.
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RESEARCH OF THE BEST TECHNOLOGICAL
AND METALLURGICAL PARAMETERS FOR
PERFORMING THE ELECTRIC RESISTANCE
WELDING OF LOW CARBON STEELS
C. Mapelli, C. Corna
This work deals with the research of the optimal technological and metallurgical parameters in order to implement
a reliable procedure for the electric resistive welding of low carbon structural steel, in order to evaluate the
conditions which can grant the best mechanical performances. Low carbon steels must be featured by high plastic
formability properties, since the production process consists in the piping of a rolled band, followed by an Electric
Resistance Welding (ERW) of the edges. The optimal technological parameters have been identified performing
welding tests at several levels of electric power, squashing length and forward velocity of the pipe along the coil
axis. Several mechanical tests have been performed for the determination of the properties of the materials under
examination, in order to characterize the main mechanical properties, i.e. Young modulus, yield and the ultimate
stresses, yield point elongation (the strain after which the plastic behaviour takes place), anisotropy coefficients
(r m
, Δr), Vickers micro-hardness and hardening coefficient of the materials analysed, while the residual stress
induced in correspondence of the welded joining have been determined by X-ray diffraction. The microstructural
characteristics of the steels have been obtained through micrographic analyses coupled with the use of Electron
Back Scattered Diffraction techniques (EBSD). The value assumed by the hardening coefficient and by the yield
elongation point has been revealed to be a strongly significant parameter for assuring the quality of the joining in
order to avoid a very early formation of the cracks in the welding region.
Keywords: electric resistive welding, cementite precipitation, hardening coefficient, yield elongation point, residual
stresses
INTRODUCTION
This work is about the identification of the best technological parameters
of the steel properties which can grant the soundness
of pipes realized by ERW high frequency welding. This process
is based on the resistive heating of the edges of the steels which
cross a volume contained in a coil interested by a current varying
at high frequency (500-1000kHz). The time-variant magnetic
flow induced by the coils current causes a potential difference
and a related current which concentrates on the steel edges producing
an intensive and concentrated heating (Fig. 1).
Just after the heating, the strip edges are pulled against themselves
by the action of rollers. This is the system through which the
welding operation is performed exploiting the High Frequency
Carlo Mapelli, Cristian Corna
Sezione Materiali per Applicazioni Meccaniche
Dipartimento di Meccanica, Politecnico di Milano,
via La Masa 34, 20156 MILANO (ITALY)
email: carlo.mapelli@polimi.it
s
Fig. 1
Example of a simulation showing the layout of the
system and the resistive heating produced on the pipe edges
to be joined.
Esempio di una simulazione che mostra il layout del sistema e
il riscaldamento prodotto sulle estremità del tubo che devono
essere saldate.
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Saldatura
%wt
P1
P2
C
0.042
0.048
Mn
0.239
0.224
Si
0.012
0.014
S
0.010
0.012
P
0.017
0.010
Cr
0.0134
0.0274
Ni
0.0153
0.0242
Cu
0.0159
0.0319
Al
0.055
0.050
Mo
0.004
0.005
s
Tab. 1
Average chemical composition of the two examined steels.
Composizione chimica media dei due acciai esaminati.
s
Fig. 6
The hydroforming device used for testing the
welded pipes.
Strumento di idroformatura utilizzato per testare i tubi
saldati.
experimental trials have been performed on strips 2.0 and 2.5
mm thick (provided by two different suppliers indicated as
P1 and P2) in order to point out possible differences produced
by the variation of either the chemical composition within
the tolerated ranges or in the performed thermo-mechanical
processes.
The welding process has been performed in order to produce
pipes of 135mm diameter applying different combinations of
the operative parameters which can be easily controlled by the
operators:
- electric power supply: 210kW-250kW-290kW;
- forwarding velocity: 45m/min-50m/min-55m/min;
- squashing length between the edges: 0.5mm-1mm-1.5mm;
provided a starting distance of pipe edges of 0.2mm. The electric
power has been developed applying a frequency of 650kHz.
The welding region has been characterized through Vickers micro-hardness
profile. Moreover, the analysis of the morphology
of the sandglass shape of heat and deformation affected zone
(HADZ) and the inclination of the plastic flow deflection lines
of this region have been performed. Susequently, for each combination
of the operative parameters, a pipe 50mm long has
undergone a hydroforming instrumented test (Fig. 6) through
which the water has been pulled into the pipes at a rate of
8MPa/min at room temperature.
The maximum pressure reached during the test has been recorded
and assumed as the load which has led the pipe to collapse.
The higher the supported pressure the better the reliability of
the welded structure is considered. The hydroforming device
has been designed in order to avoid the induction of axial stresses
along the pipe wall.
The ERW process imposes significant plastic deformation to
the welded edges and this represents a peculiarity of such a
welding procedure. The characterization of the main properties
of the materials which undergo a plastic deformation process
after heating is a fundamental step to identify which is the
most important alloy property to be monitored and controlled
in order to realize a good and reliable design of the fabrication
process. The performed characterization is articulated in:
- chemical analyses, to establish the average composition of the
sample;
- metallographic trials to measure the grain size of the steel
sample, to detect the different phases, their distribution and
the possible presence of particular crystallographic orientation
which can affect the mechanical behaviour;
- tensile tests performed along different directions to determine
the main mechanical properties (yield stress, ultimate tensile
stress, coefficient of hardening, total elongation etc.) and microhardness
measurements to evaluate the features of heat affected
and strained zone near the welding joint;
- X-ray diffraction examination near the welded region in order
to point out the residual stresses left by the welding operation.
Chemical Analysis
The chemical analysis of steels supplied by P1 and P2 revealed
that P2 material contains a higher concentration of alloying elements,
i.e. Ni, Cr and Cu (Tab. 1).
Metallographic Analyses
This step of the analyses was performed for identifying the different
phases appearing inside the material, paying particular
attention to their sizes, shapes and distributions 7) . In this case
the samples have been etched by Picral solution (2÷4g of Picric
Acid in 100ml of Ethanol) for 7s in order to point out the
presence of the different phases and the grain boundaries. The
determination of the grain size has been performed on the realized
micrographs according to the UNI 3245 and ASTM E112-82
standards.
The cementite volume fraction featuring the microstructure of
the analysed steels has been measured through an automatic
image analyser. For each sample an area of 10mm 2 has been
examined.
The Electron Back-Scattered Diffraction (EBSD) probe mounted
on a Scanning Electron Microscope (SEM) has been applied for
the identification of the crystallographic textures 8,9,10) . For this
operation the samples, after the grinding and polishing to an
average roughness of 0.05μm - operated through the colloidal
silica (solution of 80% silica suspended within a 20%H 2
O deposited
on a rotating titanium disk) - have been inserted within
a conductive resin 11) . The microscope has been set to 20kV and
the total scanned surface to obtain the texture measure is of
100mm 2 . The samples used for this analysis are the same investigated
for the optical metallographic examination before the
application of the etching solution to avoid the alteration of the
surface characteristic which can compromise the quality and
la metallurgia italiana >> ottobre 2008 21

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s
Fig. 10
Example of the revealed deflection lines
of plastic flow revealed on a welding performed
through the correct combination of the technological
parameters.
Esempio delle linee di deflessione associate al flusso
plastico rivelate su una saldatura effettuata con il
settaggio ottimale dei parametri tecnologici.
s
Fig. 9
Morphology of the welded zone in 2.5mm thick
pipes with steel provided by P1.
Morfologia della zona saldata in un tubo di spessore
2.5mm fornito da P1.
The presence of a possible Heat Affected Deformed Zone
(HADZ) has been evaluated through the determination of Vickers
micro-hardnesses (ASTM E384) across the welded joint,
in which the measurements have been performed with a step
of 50μm between two successive measurements and applying
a load of 25g for 15s.
Determination of the residual stresses
Using an X-Ray diffractometer (X-Stress 3000) and varying
the work angle between -45° and +45°, the measurement of
the residual stresses inside the material has been performed:
the diffractometer provides the values of the two stresses σ 1
and σ 2
and the amplitude of the angle φ, representing the rotation
between the stresses measured along the fixed reference
system and the direction of the principal stresses(σ,τ). These
quantities can be opportunely elaborated to give the value of
the Von Mises equivalent stress:
where
(3)
(4)
(5)
s
Fig. 11
Example of dirty materials and oxides pulled out
from the welded joining by the squashing movement.
Esempio dello sporco e degli ossidi estratti dal giunto
saldato durante il movimento di squashing.
RESULTS AND DISCUSSION
The highest resistance level to the hydroforming pressure
has been reached for 1mm pulling length and this implies
(provided an initial edge distance of 0.2mm) that
the squashing penetration between the pulled edges is of
0.8mm (Fig. 7, Fig. 8, Fig. 9). This distance seems fundamental
to grant a correct symmetry of the sandglass shape
of HADZ and the average deflection angle of 35.1° (st.dev.
±3.1°) at the middle of thickness and of 78.3° (st.dev. ±2.9°)
near the surface in order to assure an efficient removal of
the defects produced by the presence of oxides or dirty
residuals (Fig. 10, Fig. 11). At the same time the largest
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P1 2mm
P1 2.5mm
P2 2mm
P2 2.5mm
E(GPa)
205
189
202
207
Yield stress
(MPa)
300
275
303
325
Yield point of
elongation (%)
4
4.5
3.6
3
r m
0.91
0.94
0.93
0.93
Δ r
-0.09
-0.13
-0.1
-0.18
n
0.22
0.2
0.15
0.14
s
Tab. 2
Main average mechanical characteristics revealed by
the tensile tests.
Valori medi delle principali caratteristiche meccaniche misurate
mediante prove di trazione.
s
Fig. 16
Main textures pointed out by the ODF diagram
section on correspondence of (a)ϕ 2
=0° and (b)ϕ 2
=45°at the
middle of the thickness in steel 2.5mm thick provided by
P1.
Principali tessiture emerse dalla sezione del diagramma ODF
in corrispondenza di (a)ϕ 2
=0° e (b)ϕ 2
=45° a metà profondità
in un acciaio dello spessore di 2.5mm fornito da P1.
ponents particularly suitable for a plastic deformation process,
actually a prominence of the components in γ-fibre
in all the samples under examination has been revealed;
the only difference is the greater dispersion of components
featuring the P2 samples, joined together with a lower intensity
of favourable textures characterized by the planes
{111} and {110} of the body centred cubic lattice lying parallel
to the rolling plane (Fig. 16, Fig. 17). Moreover, P2
steel shows a more intense {001} Cube component
which is usually detrimental for the formability attitude.
Thus, this situation can cause a worse formability attitude,
which seems to produce considerable variation on the hardening
coefficient.
The tensile tests carried out indicated that P2 steels are featured
by higher values of Young modulus and yield stress,
if compared to the values typical of P1 materials (Tab. 2).
On the contrary, P1 steels present yield point elongations
slightly higher than P2 ones, even if the values are very
close and correspond to few percents. The presence of significant
yield point elongation is a peculiarity of the low
carbon steels and it can represent a ductility parameter of
the material, although an excessive value of this parameter
may cause the appearance of the so called ‘Lüders bands’
on the surface and on the layer immediately under it. This
phenomenon can be detrimental for the surface quality of
the component, but in this case the performed industrial
trials have not revealed this problem.
The average normal anisotropy parameter (r m
) and the
one describing the planar anisotropy (Δr) turned out to be
practically similar in all the analysed samples and the difference
pointed out cannot be the responsible for the formation
of the micro-cracks developed in P2 steel.
On the contrary, the hardening coefficient and the yield
elongation point assume significantly higher values in
the steels provided by P1 than in the ones from P2. Thus,
this parameter seems to cover an important role in order
to avoid the start up and the development of the cracks
s
Fig. 17
Main textures pointed out by the ODF diagram
section on correspondence of (a)ϕ 2
=0° and (b)ϕ 2
=45°
at the middle of the thickness in steel 2.5mm thick
provided by P2.
Principali tessiture emerse dalla sezione del diagramma ODF
in corrispondenza di (a)ϕ 2
=0° e (b)ϕ 2
=45° a metà profondità
in un acciaio dello spessore di 2.5mm fornito da P2.
s
Fig. 18
Example of the comparison of the average
measured micro-hardness profile in the steel provided
by P1 and P2.
Esempio del confronto dei profili medi di microdurezza negli
acciai forniti da P1 e P2.
la metallurgia italiana >> ottobre 2008 25

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Memorie >>
Saldatura
[4] J. W. Elmer, T. A. Palmer, W. Zhang, B. Wood and T. DebRoy:
Acta Mater., 51 (2003), 3333.
[5] B. H. Chang and Y. Zhou: J. Mater. Process. Technol., 139
(2003), 635.
[6] A. De, L. Dorn and O. P. Gupta: Sci. Technol. Weld. Joining,
5 (2000), 49;
[7] Y. Watanabe and I. Momose: Ironmaking Steelmaking,
31 (2004), 265;
[8] Internet site: www.ebsd.com;
[9] W. B Hutchinson and M. Hatherley: An Introduction to
Texture in Metals, Monograph 5, The Institution of Metallurgists,
London, (1979), 255.
[10] U. F. Kocks, C. N. Tomè and H.-R. Wenk: Texture and
Anisotropy, Cambridge University Press, Cambridge,
(2000), 421;
[11] W. F. Hosford and R. M. Caddell: Metal Forming: Mechanics
and Metallurgy, 2nd Ed., PTR Prentice Hall, New
York, (1993), 286
[12] R. K. Ray, J. J. Jonas and R. E. Hook: Int. Mater. Rev., 39
(1994), 129;
[13] Standard UNI EN 10002, ‘Materiali metallici: prova di
trazione a temperatura ambiente’ (1992);
[14] Standard ASTM E517-00: ‘Standard test method for plastic
strain ratio for sheet metal’ (August 2000);
[15] W. T. Lankford, S. C. Snyder and J. A. Bauscher: Trans.
Am. Soc. Met., 42 (1950), 1197.
[16] M. R. Barnett: Modern LC and ULC Sheets Steels for
Cold Forming: Processing and Properties, ed. by W. Bleck,
Aachen University of Technology, Aachen, (1998), 61;
[17] M. R. Barnett and J. J. Jonas: ISIJ Int., 39 (1999), 856;
[18] H. J. Bunge: Texture Analysis in Materials Science-Mathematical
Methods, Butterworths, London, (1982), 145;
[19] U. F. Kocks, C. N. Tomè and H.-R. Wenk: Texture and
Anisotropy, Cambridge University Press, Cambridge,
(2000), 421.
LIST OF SYMBOLS
E
r m
Δr
K
n
σ V.M.
ε w
ε t
ε p
l 0
l f
r
r m
X m
X n
w 0
w f
Young modulus [GPa]
average normal anisotropy coefficient
planar anisotropy coefficient
coefficient of strengthening in the Hollomon relation [MPa]
hardening coefficient
Von Mises Equivalent Stress [MPa]
width deformation
thickness deformation
plastic component of the deformation
initial length of the specimen used for the tensile test [m]
final length of the specimen used for the tensile test [m]
normal anisotropy coefficient
average normal anisotropy coefficient
average value of the generic mechanical parameter X
value of the generic mechanical parameter X along a
direction rotated by n from the rolling direction
initial width of the specimen used for the tensile test [m]
final width of the specimen used for the tensile test [m]
RICERCA DEI PARAMETRI TECNOLOGICI E
METALLURGICI OTTIMALI PER L’ESECUZIONE DELLA
SALDATURA PER RESISTENZA ELETTRICA DEGLI
ACCIAI A BASSO CARBONIO
Parole chiave: saldatura per resistenza elettrica, precipitazione
della cementite, coefficiente di incrudimento,
deformazione allo snervamento, sforzi residui
Il presente lavoro tratta la ricerca dei parametri tecnologici e metallurgici
ottimali per implementare un processo affidabile di saldatura elettrica per
resistenza degli acciai strutturali a basso tenore di carbonio (Tabella 1) e
per stabilire le condizioni in grado di garantire le migliori prestazioni dal
punto di vista meccanico. Gli acciai in esame devono possedere elevate
capacità di deformazione plastica in quanto il processo produttivo prevede
l’avvolgimento di un nastro laminato, seguito dalla saldatura delle estremità
per resistenza elettrica (ERW – Electric Resistance Welding) (Figure
1 e 2). I parametri tecnologici ottimali sono stati evidenziati mediante
ABSTRACT
l’esecuzione di test di saldatura a diversi livelli di potenza elettrica, lunghezza
di schiacciamento e velocità di avanzamento del tubo lungo gli assi
delle bobine. Per la misura delle proprietà del materiale considerato sono
stati eseguiti diversi test meccanici allo scopo di caratterizzare le principali
proprietà meccaniche, quali il modulo di Young, i carichi di snervamento
e di rottura, l’allungamento al punto di snervamento (lo sforzo oltre
il quale comincia il comportamento plastico), i coefficienti di anisotropia
(r m
, Δr), le microdurezze Vickers e i coefficienti di incrudimento (Tabella
2); gli sforzi residui indotti in corrispondenza dei giunti saldati sono stati
determinati per mezzo della diffrazione di raggi X (Tabella 3). Le caratteristiche
microstrutturali degli acciai sono state ottenute attraverso analisi
micrografiche accoppiate all’utilizzo di tecniche di diffrazione EBSD (diffrazione
degli elettroni retrodiffusi) (Figure 16 e 17). Si è riscontrato che i
valori dei coefficienti di incrudimento e dei punti di yield elongation sono
da ritenersi un parametro particolarmente significativo per assicurare la
qualità della saldatura ed evitare la prematura formazione di cricche in
prossimità dei giunti saldati (Figure 13 e 19) a seguito delle operazioni di
compressione o espansione sulle superfici laterali dei tubi.
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Corrosione
CORROSION AND PROTECTION OF
FRICTION STIR WELDS IN AEROSPACE
ALUMINIUM ALLOYS
C. G. Padovani, A. J. Davenport, B. J. Connolly, S. W. Williams,
A. Groso, M. Stampanoni, F. Bellucci
Keywords: aluminium alloys, welding, corrosion
INTRODUCTION
Friction stir welding [1] (FSW) offers the opportunity of obtaining
high quality welds in the traditionally poorly weldable high
strength aluminium alloys of the 2XXX and 7XXX series. Due to
the excellent quality of the welded joints, aircraft manufactures
are considering the introduction of this technology in aircraft
components. Friction stir welding has been used with success in
joining primary structures in the Eclipse 500 jet [2], and will be
applied to join external fuel tanks in the NASA Space Shuttle [3].
A review of recent investigations on the properties of FSW has
been compiled by Mishra and Ma. [4].
The corrosion performance of the welds has been analysed in a
number of studies, which show that the thermal cycle produced
by welding leads to significant changes in the microstructure of
the metal, leading to enhanced corrosion susceptibility [5-24]. In
aerospace alloys of the 2XXX and 7XXX series, this causes concerns
related to the corrosion-fatigue of FSW components, as
the onset of localised corrosion in aluminium alloys is known
to be able to decrease this parameter (e.g. [25]). Recent work on
AA2024 T351 [16, 17] showed the correlation between welding
parameters and precipitation of the age-S phase, while for 7XXX
alloys changes in electrochemical behaviour have been attributed
to the precipitation of η phase.
Due to the sensitisation of the weld region, it may be desirable
to improve the corrosion performance of friction stir welds by
the use of appropriate post treatments. The use of post weld heat
treatments to increase and homogenise the corrosion resistance of
the weld had limited success [22, 26-30] and tend to be restricted
by physical limitations related to the size of the components to
be treated.
Laser surface melting is able to increase the corrosion resistance
of aluminium by dissolving the detrimental constituent particles
present in commercial alloys [31] and can be considered for the
treatment of FSW due to its ability of forming, in appropriate
conditions, corrosion resistant, precipitate free layers. This has
been obtained with Excimer lasers [32-39], in which the short duration
of the thermal cycle induced by laser irradiation leads to
limited microsegregation in the molten and resolidified layer.
The use of laser surface melting to increase the corrosion resistance
of friction stir welds has been recently investigated [5, 6, 10,
11, 40, 41]. Apart from increasing the corrosion resistance of the
parent material and of the weld region, the use of laser surface
melting to increase the corrosion resistance of welds might offer
AA2024
AA7449
Si
0.50
0.12
Fe
0.50
0.15
Cu
3.8-4.9
1.4-2.1
Mn
0.3-0.9
0.20
Mg
1.2-1.8
1.8-2.7
Cr
0.10
-
Zn
0.25
7.5-8.7
Ti + Zr
0.15
0.25
Al
bal
bal
s
Tab. 1
Nominal chemical composition of AA2024 and AA7449.
Composizione chimica nominale delle leghe AA2024 and AA7449.
C. G. Padovani, A. J. Davenport, B. J. Connolly
University of Birmingham, Metallurgy and Materials, Birmingham (UK)
S. W. Williams
Cranfield University, Welding Engineering Research Centre, Cranfield (UK)
A. Groso, M. Stampanoni
Swiss Light Source, Paul Scherrer Institut, Villigen PSI, (Switzerland)
F. Bellucci
Università degli studi di Napoli Federico II, Dipartimento di Ingegneria
dei Materiali, Napoli (Italia)
the ulterior benefit of reducing galvanic coupling effects between
different weld regions that can occur if wetting of the metal with
a relatively conductive electrolyte takes place. This paper discuss
the application of laser treatment with Excimer laser to increase
the corrosion resistance of friction stir welds in AA2024-T351 and
AA7449 T7951.
EXPERIMENTAL METHOD
AA2024-T351 and AA7449-T7951 laser surface melted friction stir
welds were supplied by BAE SYSTEMS in the form of 4.0 mm
and 12.2 mm thick plates respectively; the nominal chemical composition
of the alloys is reported in Tab. 1. Friction stir welding
was performed with a Triflute carbon steel tool piece at rotation
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Corrosione
a
s
Fig. 1
Excimer laser treated AA2024-T351 FSW; (a)
and (b) optical micrographs showing surface morphology
after the treatment; (c) and (d) SEM micrographs
(secondary electron mode) showing absence of precipitation
on the treated surface.
Saldatura FSW in lega 2024-T351 dopo trattamento
con Excimer laser; (a) e (b) micrografie che mostrano la
morfologia della superficie dopo il trattamento; (c) e (d)
micrografie SEM (secondary electron mode) che mostrano
l’assenza di precipitati sulla superficie trattata.
b
s
Fig. 2
Cross section SEM micrographs (backscattered
electron mode) showing melted constituent particles
in the LSM layer produced on (a) parent material, (b)
FSW HAZ and, (c) FSW nugget on AA2024-T351 laser
treated FSW.
Micrografie SEM della sezione trasversale (backscattered
electron mode) che mostrano la dissoluzione delle particelle
costituenti nello strato LSM su (a) parent material,
(b) FSW HAZ e, (c) FSW nugget su saldature FSW in
lega AA2024-T351.
the sample were exposed to the corrosive solution in addition to
the laser treated surface. The samples were glued to the stainless
steel holders with a continuous layer of glue in order to prevent
the simultaneous exposure to the electrolyte of aluminium and
stainless steel which would have resulted in undesired galvanic
coupling effects. On each in situ sample, analysis before and during
immersion (after 24 hours) was carried out. These samples
were analysed to investigate the mechanism of corrosion propagation
in laser treated layers.
EXPERIMENTAL RESULTS
Laser-treated layer morphology
Fig. 1a shows an optical micrograph of a AA2024 T351 laser
treated friction stir weld; the characteristic pattern produced
s
Fig. 3
EDX elemental analysis of untreated and laser
treated parent material; (a) AA2024 T351; (b) AA7
449-T7951. The laser treated material shows slight
enrichment in Cu (a) and Cu and Zn (b) relative to the
untreated material. The nominal chemical composition
of the alloys is also plotted.
Analisi EDX su parent material trattato laser e non trattato;
(a) lega AA2024-T351; (b) lega AA7449 T7951.
Il materiale trattato laser mostra arricchimento in Cu (a)
e Cu e Zn (b) della superficie rispetto al materiale non
trattato. La composizione chimica nominale delle leghe è
anche riportata.
on the metal surface after the LSM treatment is visible from
the magnified view displayed in Fig. 1b. Higher magnification
SEM micrographs of the treated surface show the absence of
the characteristic micron-sized constituent particles found in
AA2024-T351 (Figs. 1c and 1d). SEM micrographs of the cross
section of the same sample show dissolution of the bright,
micron-sized constituent particles and formation of a 3 5 μm
thick precipitate-free layer in any weld region (Fig. 2). Similar
morphology was found for the AA7449-T7951 (not shown),
although less contrast elemental between LSM layer and substrate
was visible in this case in the SEM backscattered images.
Fig. 3 shows the elemental composition of the LSM layer ob-
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Corrosione
a
b
c
d
Optical and SEM microscopy examination after immersion
in 0.1 M NaCl solution
To verify whether the LSM treatment increases the corrosion
resistance of FSWs and to understand whether the presence
of potential scratches in the treatment would lead to significant
dissolution in the scratched area, immersion for 20 days
of scratched laser treated and untreated welds in 0.1 M NaCl
solution was performed. Post immersion analysis was performed
both in the scratched area and in areas ‘away’ from
the scratch.
Fig.6 shows the appearance of the AA2024-T351 untreated
weld after 20 days immersion in 0.1 M NaCl followed by cors
Fig. 5
Cathodic reactivity of laser treated and untreated FSWs in 0.1 M NaCl. (a) and (b) AA2024-T351; (c) and
(d) AA7449-T7951. (a) and (c) Are typical cathodic polarisation curves in parent material comparing the reactivity of
the laser treatment with the reactivity of the untreated metal. (b) and (d) Are cathodic current densities at 900 mV
vs. Ag/AgCl as a function of position relative to the weld centre for laser treated (dipped in nitric acid) and untreated
FSW (polished). A = ‘advancing’ side of the weld; R = ‘retreating’ side of the weld.
Caratteristica catodica di saldature FSW dopo trattamento laser in 0.1 M NaCl. (a) e (b) Lega 2024-T351; (c) e (d) lega
7449-T7951. (a) e (c) Sono tipiche curve di polarizzazione catodica nel parent material che confrontano la reattività del
trattamento laser con quella del metallo non trattato. (b) e (d) Sono le correnti catodiche nominali valutate al potenziale di
900 mV vs. Ag/AgCl in funzione della posizione rispetto al centro della saldatura. A = parte ‘advancing’ della saldatura; R
= parte ‘retreating’ della saldatura.
weld were scattered and not uniform across the whole sample,
while that measured o the untreated weld show lower
values in the weld region, indicative of enhanced susceptibility
to anodic attack.
Cathodic polarisation curves and cathodic currents measured
on laser treated and untreated welds for both alloys are shown
in Fig. 5. The graphs show typical cathodic polarisation curves
in parent material and the values of the cathodic current at a
fixed potential of 900 mV vs. Ag/AgCl, which was used to
compare the reactivity across the weld region.
It is clear that the laser treatment can increase the corrosion
resistance of both alloys by reducing the cathodic reactivity.
The laser treated material (broken line) shows lower cathodic
reactivity in the whole weld region and more uniform reactivity
in comparison with the untreated weld (solid line), where,
for both AA2024 and AA7449, a cathodic current density peak
is observed in the weld nugget.
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s
Fig. 7
Laser treated AA2024-T351 FSW after 20 days immersion in 0.1 M NaCl and removal of corrosion products in concentrated
nitric acid; (a) weld surface micrograph; (b), (c) and (d) optical micrographs of surface ‘away’ from the scratch in nugget,
HAZ and parent material respectively; (e), (f) and (g) optical micrographs of cross section ‘away’ from the scratch showing typical
localised corrosion sites in nugget, HAZ and parent material respectively.
Saldatura FSW in lega AA2024-T351 trattata laser dopo immersione per 20 giorni in 0.1 M NaCl e rimozione dei prodotti di corrosione
in acido nitrico concentrato; (a) micrografia della superficie; (b), (c) e (d) micrografie ottiche della superficie in zone lontane dall’intaglio
in nugget, HAZ e parent material rispettivamente; (e), (f) e (g) micrografie ottiche della sezione trasversale in zone lontane dall’intaglio
che mostrano tipici attacchi corrosivi in nugget, HAZ e parent material rispettivamente.
regions in the untreated weld. However, laser treatment was
beneficial in decreasing the reactivity of the HAZ, in which superficial
attack (Fig. 10c and 10f) was found in place on fairly
deep pits (Fig. 9c and 9f).
Fig. 11 shows optical micrographs of the scratched area in
different regions of the AA7449-T791 laser treated weld after
20 days immersion. Contrarily to what observed for AA2024
T351, the extent of attack in the scratched area was found to
be much lower than on the laser treated surface. The number
of pits in parent material (Fig. 11a) and nugget (Fig. 11c), for
example, was much lower in the scratched area than on the
intact LSM surface and much lower that that found on the untreated
weld.
Open circuit potential measurements on AA2024-T351 and
AA7449-T7951 laser treated and untreated parent material
were employed to explain the behaviour of the scratched laser
treated material. Measurements performed in 0.1 M NaCl on
intact and scratched laser treated parent material and on untreated
parent material are shown in Fig. 12. For AA2024 T351
(Fig.12a), the measurements show higher OCP of the intact
laser treated material in comparison with the untreated and
scratched laser treated material. For AA7449-T7951 (Fig.12b),
in contrast, the OCP of the LSM layer was lower that that
observed on intact parent material and similar to that of the
scratched LSM material. Considerations on the OCP measurements
are presented in the discussion.
X-ray microtomography examination of ex-situ samples
In order to study corrosion propagation in damaged laser
treated layers, X-ray microtomography was used to analyse
ex situ samples cut out from a scratched AA7449-T7951 laser
treated weld after immersion in 0.1 M NaCl for 5 days. The
corrosion products were not removed before examination.
Surface observation of the weld (Fig.11) had highlighted attack
of the LSM surface in all weld regions but virtually no attack
of the underlying substrate in the scratched area. X ray microtomography
was used to gain a better characterisation of the
corrosion damage. The observation that little attack develops
in the scratched area of LSM AA7449 when exposed to NaCl is
significantly strengthen by the set of micrographs displayed in
Fig. 13, which show “slices” parallel to the LSM layer extracted
form a 3D volume reconstruction of a sample cut out from
the HAZ region of a LSM weld. Significant generalised attack,
penetrating to a depth of about 30 μm, is visible on the surface
of the sample. In contrast, no attack is visible in the scratched
area of this sample.
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s
Fig. 10
Laser treated AA7449-T7951 FSW after 20 days immersion in 0.1 M NaCl and removal of corrosion products in concentrated
nitric acid; (a) weld surface micrograph; (b), (c) and (d) optical micrographs of surface ‘away’ from the scratch in nugget,
HAZ and parent material respectively; (e), (f) and (g) optical micrographs of cross section ‘away’ from the scratch showing typical
localised corrosion sites in nugget, HAZ and parent material respectively.
Saldatura FSW in lega AA7449-T7951 trattata laser dopo immersione per 20 giorni in 0.1 M NaCl e rimozione dei prodotti di corrosione
in acido nitrico concentrato; (a) micrografia della superficie; (b), (c) e (d) micrografie ottiche della superficie in zone lontane
dall’intaglio in nugget, HAZ e parent material rispettivamente; (e), (f) e (g) micrografie ottiche della sezione trasversale in zone lontane
dall’intaglio che mostrano tipici attacchi corrosivi in nugget, HAZ e parent material rispettivamente.
on “pins” cut from the nugget, HAZ and parent material of
a laser treated weld to investigate this effect. In this case, differently
from the “ex situ” samples, the cut untreated surfaces
were exposed together with the laser treated surface.
Fig. 14 shows X ray microtomography “slices” perpendicular
to the axis of the “pin” sample acquired on LSM AA2024 T351
in situ before (Fig. 14a) and after (Fig. 14b) 24 hours exposure
of a parent material sample in 0.1 M NaCl. The distribution
of constituent particles clearly identifies the two slices as the
same section of the sample. It is evident how delamination
of the LSM layer took place during corrosion propagation in
the laser treated material. The results obtained on HAZ and
nugget samples, however, did not show any sign of delamination
after 24 hours exposure, suggesting that this phenomenon
might take place only on some areas of a laser treated surface.
Similar results were found on AA7449 T7951 (not shown).
DISCUSSION
Electrochemical measurements and immersion tests indicated
a higher corrosion susceptibility of the weld region in comparison
with the parent material for untreated FSWs in both
AA2024 T351 and AA7449-T7951. These results are in agreement
with the findings of other studies that highlighted the
decrease in corrosion resistance often obtained in heat treatable
aluminium alloys as a consequence of friction stir welding
[5-24].
Laser surface melting produced the formation of a homogeneous,
3-5 μm thick laser treated layer across weld region and
parent material. Thermal dissolution of constituent particles
and fine precipitates occurred in the LSM weld, leading to the
formation of a precipitate free layer. The dissolution of constituent
particles was enhanced in the nugget region (e.g. Fig.
2b), as in these area the constituent particles are fragmented
into smaller pieces by the action of the FSW tool [16, 17]. The
morphology of the laser treated layer observed in this study is
consistent to that observed by other studies after laser surface
melting aluminium alloys with Excimer lasers [32-39].
Electrochemical measurements indicated that laser surface
melting with an Excimer laser can improve the corrosion resistance
of AA2024-T351 friction stir welds by decreasing cathodic
reactivity and increasing the breakdown potential in weld
region and parent material. Furthermore the electrochemical
measurements showed that laser treating the weld can produce
a certain homogenisation of the reactivity, with consequent reduction
of galvanic coupling effects that could occur if wetting
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s
Fig. 13
X-ray microtomography “slices” of scratched laser treated AA7449-T7951 FSW in HAZ region after ex-situ immersion for
5 days in 0.1 M NaCl. The slices show planes parallel to the laser treatment at different depths below the surface: (a) 7 μm; (b)
15 μm; (c) 31 μm. Although significant corrosion is observed on the sample, little attack developed in the scratched area.
‘Sezione’ di microtomografia ai raggi X di un campione di saldatura FSW in lega AA7449-T7951 trattata laser nella HAZ dopo immersione
per 5 giorni in 0.1 M NaCl. Le sezioni mostrano piani paralleli al trattamento laser a diverse profondità sotto la superficie: (a) 7 μm;
(b) 15 μm; (c) 31 μm. Sebbene l’attacco corrosivo osservato sulla superficie del campione sia notevole, l’entità della corrosione nell’intaglio
è limitata.
fect or as a consequence of corrosion development over time)
may occur. In this scenario, considerations related to the exposure
of damaged (scratched) laser treated samples and to
potential galvanic coupling effects between the LSM layer and
the substrate become important.
The results shown in this paper indicate that, for AA2024 T351,
the intact laser treated layer has higher OCP than the untreated
parent material. This suggests that, if the substrate is exposed,
galvanic coupling effects between laser treated layer and substrate
tend to drive corrosion preferentially in the substrate.
The OCP of the scratched laser treated sample, however, is
similar to that of the untreated material indicating that the
galvanic couple formed between the LSM layer and the substrate
is corroding at the potential that the uncoupled substrate
alone would exhibit during free corrosion. This suggests that,
at least for the anode/cathode ratio used in this study, the low
cathodic reactivity of the LSM layer is unable to significantly
polarise the substrate and that galvanic coupling between the
substrate (anode) and the LSM layer (cathode) does not result
in accelerated corrosion rate of the substrate. For AA7449
T7951, in contrast, the incorporation of Zn into the LSM layer
ensured a relatively high anodic reactivity of the laser treated
surface. The OCP of the laser treated layer was lower than that
of the untreated substrate, ensuring that galvanic coupling of
s
Fig. 14
X-ray micro-tomography “slices” of a parent material
laser treated sample collected in situ before and after
immersion for 24 hours in 0.1 M NaCl. The slices show the
same plane perpendicular to the pin axis direction (a) before
immersion and (b) during immersion (24 hours) and highlight
delamination of the laser treated layer during exposure to
the electrolyte. The in-situ samples were extracted from a
pristine, non scratched, laser treated AA2024-T351 FSW.
’Fetta’ di microtomografia ai raggi X di un campione di parent
material trattato laser acquisita in situ prima e dopo immersione
per 24 ore in 0.1 M NaCl. La fetta mostra lo stesso piano
perpendicolare all’asse del campione (a) prima dell’immersione
e (b) durante l’immersione (24 ore) ed evidenzia delaminazione
dello strato LSM durante esposizione all’elettrolita. I campioni
per misure in situ sono stati estratti da saldature FSW trattate
laser in lega AA2024-T351 non intagliate
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la metallurgia italiana >> ottobre 2008 41

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Refrattari
CORROSION MECHANISMS
OF ZIRCONIA/CARBON BASED
REFRACTORY MATERIALS
BY SLAG IN PRESENCE OF STEEL
Filippo Cirilli, Antonello Di Donato, Umberto Martini, Patrizia Miceli,
Philippe Guillo, Jose Simoes, Yi Jie Song
Zirconia is usually utilised in Submerged Entry Nozzle (SEN) in the slag contact zone, because of its high
resistance to corrosion. However inconsistency of component performance and apparently erratic behaviours,
in terms of corrosion rate, are frequently experienced. An important cause of the unexplained variability of
component performance is the typical trial-and-error approach used to develop materials for the specific applications,
and the “Darwinian selection” for the choice of the most suitable material despite the fact that a number
of studies are available in literature. As a matter of fact, although almost all the mechanisms that have been
proposed are based on some form of cyclic mechanism where the oxide is attacked by the slag and the exposed
graphite is then attacked by the metal, contradictory conclusions can be often found about specific features. It is
not to be excluded that contradictory results could be dependant on the experimental conditions used.
In this paper laboratory experiments have been carried out, using together slag and steel, in order to clarify
their role on the global corrosion mechanism. The results showed that, besides the dissolution of carbon in steel
and oxide in slag, other phenomena contribute to the corrosion. In particular the experiments put in evidence
the critical role of steel in dissolving the products of reactions between slag components and carbon, pushing
the attack of slag to carbon. The consequence is that the corrosion phenomenon is complex, and parameters
such as activity of slag components, porosity of refractory matrix, characteristics of carbon material are involved
in the tendency of the carbon to react with slag, hence on the global corrosion rate.
KEYWORDS: zirconia, continuous casting, Submerged Entry Nozzle, SEN, corrosion
INTRODUCTION
Zirconia is usually utilised in Submerged Entry Nozzle (SEN)
in the slag contact zone because of its high resistance to corrosion.
The occurrence of SEN corrosion is often the phenomenon
determining the duration of the casting sequence. The steelmaker
need is the availability of refractory materials at high
resistance against corrosion, in order to make long sequences
avoiding unforeseen stops of the casting operations. However
inconsistency of component performance and apparently
erratic behaviours, in terms of corrosion rate, are frequently
experienced.
Filippo Cirilli, Antonello Di Donato, Umberto Martini, Patrizia Miceli
Centro Sviluppo Materiali, Rome Italy
Philippe Guillo, Jose Simoes
Vesuvius International, Feignies, France
Yi Jie Song
Vesuvius Research, Pittsburgh, United States of America
Several corrosion mechanisms of zirconia/carbon refractories
are available in the literature, taking into account the role of
the two main different refractory components, zirconia and
graphite.
All the mechanisms that have been proposed for attack of
SENs are based on some form of cyclic mechanism [1,2,3]: the
oxide component of the nozzle (zirconia) dissolves into the
slag; as a consequence graphite remains exposed. Then a change
in mould level brings this graphite into contact with the
steel where it dissolves very rapidly, leaving refractory oxides
exposed. The process then starts again leading to global refractory
corrosion.
Hauck and Potschke [4] found two weak points in this type of
cyclic mechanism:
1) fluctuations in the meniscus are less than the extent of the
wear zone on nozzles
2) graphite dissolves more readily in the steel than the oxide in
the flux; for this reason corroded nozzles would be expected to
exhibit a network of exposed alumina or faster erosion in the
steel than in slag.
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Refrattari
Slag 1
Slag 2
Slag 3
CaO [%]
13
22
32
SiO 2
[%]
56
33
25
Al 2
O 3
[%]
9
38
32
MgO [%]
2
2
4
MnO [%]
17
0
0
Na 2
O [%]
3
5
7
s
Tab. 3
Chemical composition of the slags used for the tests.
Composizione chimica delle scorie usate per i test sperimentali.
Slag 1
Slag 2
Slag 3
a CaO
0.002
0.005
0.082
a SiO2
0.694
0.208
0.128
a Al2O3
0.081
0.991
0.736
a MgO
0.012
0.013
0.035
a MnO
0.064
-
-
a Na2O
3 E-8
3 E-6
5 E-6
s
Tab. 4
Calculated activity of slags components at 1550°C referred to the standard state of pure oxide (by Thermo-Calc TM ).
Attività dei componenti della scoria calcolate a 1550°C e prendendo come stato standard l’ossido puro ( i calcoli sono stati fatti
conThermo-Calc TM ).
not been fixed with the objective to reproduce the composition
of casting powder, but to put in evidence the role of slag
properties. According to this concept, three slags have been
produced, having the following characteristics:
1. high SiO 2
and MnO activity
2. high SiO 2
3. high CaO activity
The complete slag compositions are reported in Tab. 3.
The chemical activity of the slags components, referred to the
standard state of pure oxides, has been calculated with the
thermodynamic code Thermo-CalcTM at the test temperature
of 1550°C. The calculated values are reported in Tab. 4.
Description of experimental apparatus and procedure
The experimental tests were carried out in an electrical furnace,
with graphite heating elements, under Ar atmosphere.
The refractory samples were cut as rods of 2 cm of diameter
and 5 cm length. For each test, an alumina crucible was filled
with pure iron and heated up to the temperature of 1550°C.
When the iron was completely melted, the furnace was open
for adding the slag to the crucible and for putting the sample
inside the furnace up to 10 cm above the crucible, to be preheated
before submerging. Then, after complete slag melting
and sample pre-heating (typically 5 minutes), the refractory
rod was lowered inside the crucible so to be in contact with the
liquid iron and the slag.
Fig. 1 shows a scheme of the experimental apparatus.
The duration of each test was 30 minutes. At the end of the test,
the furnace was switched off, the sample left submerged and
cooled under Ar flow.
After cooling and solidification, the crucible was cut and samples
of refractory in contact with iron and slag were taken and
submitted to Scanning Electron Microscopy (SEM) and Energy
Dispersive Spectroscopy (EDS) investigation.
RESULTS
s
Fig. 1
Scheme of the experimental apparatus used for
the experimental tests.
Composizione del refrattario usato per i test sperimentali.
The investigation has been focused on the type and extent of
the predominant interaction that occurs at the interface between
the refractory and molten phases depending on the slag
used. As already published in literature, the following phenomena
have been observed on the refractory material after all
the performed tests. They are:
- Graphite consumption: this occurs in general where the refractory
is in contact with the metallic phase. A layer is formed
in which the slag takes the place of the graphite and surrounds
the zirconia grains. In what follows, this layer is called “decarburised
layer”.
- Slag penetration: the slag can penetrate through the refractory
carbonaceous matrix.
- Structure degradation of ZrO 2
grains: this takes place in the
grains that are in contact with the slag and can be observed in
different forms, like simple cracks of the grain or complete
crushing.
The extent of each phenomenon was different, depending on
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Results of test with slag 3
Fig. 10 shows the appearance of the refractory interface in contact
with slag 3.
Zirconia grains are attacked by the slag, but the extent of the
interaction is less evident respect that observed with slags 1
and 2. The slag analysis carried out on near the zone of the
interface (see the zone 1) shows that the slag composition did
not change in a significant way. The presence of dissolved ZrO 2
in the slag (up to 4÷5% wt.) has been remarked.
It is certainly caused by the degradation phenomena that affect
the ZrO 2
grains that are in the zone of the refractory borderline.
Anyway, with slag 3 only the smaller grains are attacked
by the slag, while the larger ones are not significantly modified
after the experimental test.
Fig. 11 reports the appearance of the refractory border in the
liquid iron/slag zone. The layer of slag penetration is in the
s
Fig. 5
Penetration of slag 1 inside the refractory below the
liquid iron/slag contact level. EDS analyses performed on
penetrated slag.
Penetrazione della scoria N. 1 all’interno del refrattario al di sotto
della zona di contato con la scoria. L’analisi EDS è stata fatta sulla
scoria penetrata a diverse profondità all’interno del refrattario.
s
Fig. 7
Penetration of slag 2 inside the refractory at the
slag contact level.
Penetrazione della scoria N. 2 all’interno del refrattario.
s
Fig. 6
Refractory sample appearance after test with
slag 2 at the slag contact level.
EDS analysis performed on points 1, 2 and 3 of the slag.
Aspetto del refrattario nella zona di contatto con la scoria
N. 2 dopo il test sperimentale.
trated slag in contact with the grain. Analysis of the slag in
point 2 of Fig. 8 indicates CaO concentrations of about 30 %,
while the starting value was about 20 %.
Fig. 9 reports the appearance of the refractory below the liquid
iron/slag contact level. In this case the average thickness of
slag impregnation layer is less than 200 µm, and the extent of
structure degradation is less than that remarked with slag 1.
s
Fig. 8
Structure degradation at the border of a coarse
ZrO2 grain in contact with slag 2 at the slag contact
level. EDS analysis performed on the grain and on points
1 and 2 of the slag.
Decadimento della struttura dei bordi dei grani di zirconia
dopo interazione con la scoria N. 2. L’analisi EDS è stata
fatta sui punti (1) e (2) indicati in figura.
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As a matter of fact, the degradation of the grain structure
plays an important role in the corrosion of the material. In
fact, grain degradation is accomplished by the slag penetration
through the material and the loss of portions of ZrO 2
grains after their crushing.
Grains degradation occurred with all the three slags. Results
from SEM observations showed that in general the grain degradation
is associated with the loss of the stabilising agent
CaO [12] as confirmed by the absence of CaO in the crushed
grains and by the enrichment in CaO of the slag surrounding
them. It follows that the extent of CaO dissolution from
the grain into the slag can depend on slag characteristics, in
particular on CaO activity in the slag or, in other words, on
basicity index CaO/SiO 2
.
In our tests, the extent of grain degradation is significantly
different depending on the slag used: structure degradation
occurs in the whole ZrO 2
grain in the case of slag 1, while it
takes place mainly on the border of the grain in the case of
slag 2 and with even less extent in the case of slag 3.
From a qualitative evaluation, the extent of grains degradation
has the following order:
Extent of grain degradation: Slag 1 >> Slag 2 > Slag 3
that is in agreement with the increasing basicity index of the
three slags.
CaO and SiO 2
activities reported in Tab. 4 for the three slags
have the following orders:
CaO activity: Slag 1 < Slag 2 < Slag 3
SiO 2
activity: Slag 1 >> Slag 2 > Slag 3
Again slag 1 results to be the most aggressive also regarding
the extent of grain structure degradation due to the high silica
activity and low calcia activity.
Slag penetration
The slag can penetrate through the refractory matrix.
At this stage, slag penetration cannot be directly put in relation
to refractory corrosion, but it should be considered part
of the global corrosion mechanism since most of the grains
reached by the penetrated slag are partially or even totally
degraded.
Tab. 5 reports the maximum values of slag penetration
depth observed with the three slags. It is expected that the
extent of penetration of a slag depends on slag viscosity and
interfacial tension between slag and ZrO 2
. In this case, the
interfacial tension can be considered as a first approximation
depending on the characteristics of the slags used, that
is on the slag surface tension.
However there is no agreement between slag penetration
and calculated [16] slag viscosity and surface tension [17] values
reported in Tab. 6. This can be explained by considering
that the chemical composition of the penetrated slag can be
modified by reactions like decarburation and dissolution of
stabilising agent CaO. The reaction with the graphite matrix
typically causes a decrease of MnO and SiO 2
, the reaction
with the ZrO 2
grains typically leads to an increase of CaO
that is lost from the grains. This leads to the consideration
that slag penetration could depend on characteristics of the
modified penetrated slag rather than on the starting slag
composition used.
CONCLUSIONS
Zirconia is usually utilised in Submerged Entry Nozzle (SEN)
in the slag contact zone, because of its high resistance to corrosion.
The occurrence of SEN corrosion is often the phenomenon
determining the duration of the casting sequence.
An activity has been carried out to investigate the corrosion
mechanism of calcia stabilised zirconia based refractory in
presence of slag and steel. Slags with different activity of its
constituents have been used.
The carried out activity individuate three main phenomenon
operating at the same time:
1. Graphite consumption: the graphite of the refractory may
be lost not only by direct dissolution into the steel, but also
for the reaction with slag constituents. The reactions between
slag components as SiO 2
and MnO that can oxidise the
graphite needs the presence of the metallic phase to take
place. The higher are the activity values of the above mentioned
species the more is the level of decarburization of the
refractory. Of course, a higher decarburization level of the
refractory implies a higher global corrosion rate.
2. Zirconia grains degradation: this is associated with the
dissolution of the stabilising agent CaO. A correlation between
the “capacity” of the slag to dissolve CaO and the extent
of degradation of the zirconia grains has been found. Slags
with high SiO 2
and low CaO activities cause high levels of
zirconia grains degradation up to a complete crushing, thus
concurring to a faster global corrosion of the material.
3. Slag penetration: the slag penetrates through the refractory
matrix. The penetrated slag interacts with the zirconia
grains in the inner parts of the refractory beyond the borderline
of the decarburised layer. The grains interacted with
this penetrated slag are often partially or even totally degraded.
This means that also the phenomenon of slag penetration
can participate at the global corrosion mechanism. In
general, the extent of slag penetration can be put in relation
with slag properties like viscosity, but it must be taken into
account that the composition of the penetrated slag can vary
depending on the reactions involved in the interaction mechanism.
This work demonstrated that the same ZrO 2
/C refractory
material underwent corrosion with different extents when
Penetration (μm)
Slag 1
400
Slag 2
400
Slag 3
500
s
Tab. 5
Depth of slag penetration inside zirconia refractory.
Profondità di penetrazione delle tre scorie nel refrattario.
Viscosity (Pa·s)
Surface tension (mN/m)
Slag 1
2.2
350
Slag 2
4.6
362
Slag 3
1.5
345
s
Tab. 6
Calculated viscosity according to Ref. 16 and calculated
surface tension according to Ref. 17 for the three slags
used in the experimental tests.
Viscosità calcolate usando il modello del rif. 6 e tensioni superficiali
calcolate secondo il modello riportato nel rif. 17 per le tre
scorie usate nei test sperimentali.
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APPLICATION OF OPTICAL BASICITY
PARAMETER TO FOAMING OF SLAGS
Y. A. A. Murali Krishna, T. Sowmya, S. Raman Sankaranarayanan
Metallurgical slags play an important role in the melting and refining of metals. Efforts are being made, by
many researchers, to understand the factors influencing the properties of slags. Optical basicity is a chemical
parameter which has been applied to glasses and slags, and, is a more comprehensive representation of slag
composition than conventional basicity. Foaming is an important phenomenon in steelmaking, but limited
information is available on the effect of slag composition on foaming. Optical basicity values, for different
slags, were calculated from the chemical composition – following the approach of Duffy and co-workers. The
calculated values were then applied to follow the trends in foaming, bath smelting and ladle slags. The results
demonstrate the potential use of optical basicity in this area, but the trends could be investigated further with
respect to structure and the ionic concentrations.
KEYWORDS: metallurgical slags, foaming, chemical composition, optical basicity
INTRODUCTION
Selection and performance of slags is very critical for many
operations in melting, refining and casting of metals. Chemical
properties of metallurgical slags such as chemical composition
and basicity as well as physical properties such as fusion temperature,
viscosity, foaming index have a strong influence on
the performance of slags [1]. However, physical properties of
slags need to be measured at elevated temperatures and often
difficulties are encountered in the same. Hence, the need
to predict properties of slags based on chemical composition
and certain empirical relations. Optical basicity, a parameter
based on the ionic nature of oxides, has been used for predicting
the properties of glasses and slags. The present work is an
attempt to track the variations in foaming behaviour of slags,
as function of optical basicity. The approach has been used for
studying the behaviour of three different types of slags used in
ironmaking and steelmaking.
FOAMING
Foam is a system consisting of a concentrated dispersion of
gas bubbles in a liquid. Foam properties depend primarily
on chemical composition, interfacial characteristics, rheology,
pressure and temperature. Foaming has been observed in
metallurgical processes such as oxygen steelmaking, but has
become a critical phenomenon in the newer process modifications.
Experimental investigations, based on actual foam
Y. A. A. Murali Krishna, T. Sowmya, S. Raman Sankaranarayanan
Department of Metallurgical and Materials Engineering
National Institute of Technology
Tiruchirappalli – 620 015 India
e-mail: raman@nitt.edu,
ramantech19811985@yahoo.com
measurements and physical models have been reported in the
literature [2]. Viscosity has been cited as an important influencing
variable, but not much work has been done on the relation
between chemical composition and foaming. This becomes
significant as the experimental measurement of viscosity is a
difficult proposition.
CONCEPT OF OPTICAL BASICITY
Oxide slags used in melting and refining are considered ionic
in nature and the behaviour of the slag is strongly influenced
by the chemical composition, structure and nature of ions/
ionic charges. Parameters such as basicity do not take into
consideration the presence of many oxide species (other than
lime and silica) and also the ionicity is itself a function of the
chemical composition. The relation between the ionic structure
and optical basicity for salts, glasses and slags as well as the
significance of optical basicity in metallurgical processes has
been described in the literature [3-7]. Procedures for calculation
of optical basicity have been described, in detail, in the
literature. Calcium Oxide is taken as the anchor point with an
optical basicity value of 1 and different numerical values have
been assigned to the other oxides. Therefore, the optical basicity
value of a slag can be simply calculated from the chemical
composition (expressed in equivalent fractions of ions) and the
polarizing powers of different ions. The optical basicity (∧) of
a slag is given by:
∧ = ∧ X + ∧2X2 + …….
1 1
where ∧ i
is the optical basicity of the pure oxide i, and X i
is the
equivalent fraction of oxide i.
PROBLEM FORMULATION AND APPROACH
Physical properties of slags – such as foaming index and viscosity
have been experimentally measured by other researchers
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1] R. H. Tupkary, “Introduction to Modern Steel Making”,
Khanna Publishers, 1997.
2] Kimihisa Ito, R. J. Fruehan, “Study on the foaming of CaO –
SiO2 – FeO slags: Part I. Foaming parameters and Experimental
Results”, Met. Trans. B, 1989, vol. 20 B, pp. 509 – 514.
3] J. A. Duffy and M. D. Ingram, “Establishment of an Optical
Scale for Lewis Basicity in Inorganic Oxyacids, Molten salts
and Glasses – III”, Journal of American Chemical Society, Dec
1, 1971, pp. 6448 – 6454.
4] J. A. Duffy and M. D. Ingram, “Lewis Acid – Base interactions
in inorganic Oxyacids, Molten salts and Glasses – III “, J.
Inorg. Nucle. Chem., 1974, vol.36, pp. 43 - 47.
5] J. A. Duffy and M. D. Ingram, “Optical Basicity - IV: Influence
of electro negativity on the Lewis Basicity and solvent pros
Fig. 3
Relation between calculated optical basicity and measured
foaming index of bath smelting slags in CaO – SiO 2
– MgO – Al 2
O 3
– FeO system.
Relazione fra la “Optical Basicity” calcolata e l’indice di formazione
di schiuma misurato per schiume di bagno di fusione in
sistemi CaO – SiO 2
– MgO – Al 2
O 3
– FeO.
rical values for foaming index, in all the four cases, support
this interpretation.
The analysis was then extended to ladle slags9. In this case (3
slags), the foaming index was found to decrease steadily with
increasing optical basicity values (0.75 to 0.77), with an excellent
correlation (R 2 = 0.98) (Fig. 4). Good correlation between
surface tension and optical basicity (R 2 = 0.89)11 was observed
in this case also. The reverse trend (foaming Vs optical basicity)
is attributed to the presence of CaF2 in these slags, which
could considerably alter the silicate structure and reduce the
viscosity. This interpretation is supported by the fact that the
foaming indices are lower in this system than the previous system.
Presence of oxide particles/precipitates can have a significant
impact on the behaviour of slags. Slags containing di-calcium
silicate additions, as reported by Jiang and Fruehan [9], were
then investigated. In this system (5 points) (Fig. 5), foaming
index (1-4) was found to increase steadily with increasing values
of optical basicity (0.65 – 0.67). Correlation was very good,
with R 2 value of 0.87. In this case, the presence of oxide particles
would have increased the slag viscosity (R 2 = 1.0) [11] and
this, in turn, would have stabilized the foam – resulting in the
relatively higher values observed for foaming index.
CONCLUDING REMARKS
s
Fig. 4
Relation between calculated optical basicity and
measured foaming index of ladle slags.
Relazione fra la “Optical Basicity”calcolata e l’indice di formazione
di schiuma misurato per le scorie.
The concept of optical basicity, which is much more comprehensive
of the slag composition than basicity, has been applied
to study the trends in foaming of slags. The exercise
has been useful as the potential for the use of optical basicity
has been demonstrated. It could also be seen that the effect
of oxide composition on slag structure, involving Al 2
O 3
and CaF 2
, has a strong influence on foaming. The relation
between slag composition and structure has been reported
elsewhere [12,13]. A more rigorous analysis of slag structure
(Vs composition) can result in an improved understanding of
slag behaviour.
ACKNOWLEDGEMENT
The authors wish to acknowledge the management of National
Institute of Technology – Tiruchirappalli and the Department
of Metallurgical and Materials Engineering, for permission to
carry out the said work. SRS is grateful to the MHRD, for financial
support of research in process metallurgy. Suggestions
made by the referee, towards improving the manuscript, are
much appreciated.
REFERENCES
s
Fig. 5
Effect of addition of 2CaO - SiO2 particles on the
optical basicity and measured foaming index of slags.
Effetto dell’aggiunta di 2CaO - SiO2 sull’ “Optical Basicity”e l’indice
di formazione di schiuma misurato per le scorie.
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