Arc welding duplex stainless steels - a guide to best practice
by Kay
Bridges and Paul Woollin
Contents
Section 1 - Introduction;
Section 2 - Metallurgy; Welding
procedures
Section 3 - Weld procedure qualification testing
- Development
of duplex stainless steel
- Achieving
a balanced ferrite-austenite structure in parent duplex stainless steels
- Crystal
structures
- Ferrite
prediction
- Precipitation
of potentially detrimental phases in duplex stainless steels
- Tables
Introduction
Although duplex stainless steels have good weldability,
specific welding practices are required to ensure that their intrinsic benefits
such as good toughness and corrosion resistance are not excessively reduced by
the action of welding.
To help welding engineers utilise duplex stainless steels successfully, this
best practice guide looks at key issues such as the microstructural
transformations which occur in the weld metal and heat affected zone during
welding. It provides practical information and guidance on welding procedures,
highlighting important aspects like arc energy and interpass temperature, and
the significance of shielding gas composition, for example.
Recognising that this is a relatively involved subject, the author has
included additional, detailed information. For clarity and ease of reading, this
is kept away from the main text but, for those who want in-depth information, is
conveniently available via links throughout the guide.
Section 2 - Metallurgy; Welding
procedures
Duplex stainless steels are a
family of steels with a two-phase ferritic-austenitic microstructure.
Microstructurally intermediate between the ferritic and austenitic stainless
steels, the duplex grades combine the favourable properties of both, including
good toughness and resistance to chloride stress corrosion cracking, as well as
to other forms of corrosion such as pitting, crevice and intergranular attack.
The duplex stainless steels also possess good weldability. However, in order
to approach the favourable performance attributes of the parent metals in welded
fabrications, specific welding practices must be employed. The following
provides practical information and guidance for welding engineers on welding
these materials.
Throughout this section of the best practice guide there are links to more
detailed information, for example, Achieving
a balanced ferrite-austenite structure in parent duplex stainless steels and
Historical
developments. Items of detailed information are presented at the end of this
section of the guide under Additional
information.
A duplex alloy may
be described as one which contains a two-phase structure. More definitively, it
is reserved for alloys where both phases are present in significant quantities
and in approximately equal volume fractions (as opposed to alloys in which one
constituent appears in the form of small precipitates). In practice, the term
duplex stainless steel covers ferritic-austenitic alloys with typically between
30% and 70% ferrite. Tables
1 and 2 list some of the common wrought and cast duplex grades,
respectively.
Several sub-groups exist within the ferritic-austenitic family of stainless
steels; these may be loosely characterised by the terms 'alloy lean', 'duplex',
'high alloy' and 'superduplex'. These 'families' are grouped by similar Pitting
Resistance Equivalent Number (PREN).
PREN is a formula which has been developed to
compare the resistance to chloride pitting of stainless steels,
where;
PREN = % Cr + 3.3 x (% Mo) + 16 x
(% N) Equ. 1
and;
Alloy lean
22% Cr Duplex
25% Cr High Alloy
Duplex
25% Cr Superduplex |
23-31 PREN
30-36
PREN
32-40 PREN
>40 PREN
|
Due to the introduction of tungsten to some grades of superduplex stainless
steel, an element which also improves pitting resistance, a modified form of the
PRE relationship has also been proposed:
PREW = % Cr + 3.3 x (% Mo + 0.5 x
%W) + 16 x (% N) Equ. 2
The microstructure in a wrought base metal is disrupted at fusion welds and,
in order to obtain optimum properties, welding procedures are designed to ensure
that approximately 30 to 70% ferrite fraction is achieved in the weld.
Before discussing the practical aspects of welding duplex
stainless steels it is important to outline some of the microstructural
transformations occurring in the weld metal and HAZ (heat affected zone) during
the welding process, as these have implications on the procedures that are
adopted for welding these materials.
 |
Fig.1. Cross-section through the Fe-Cr-Ni phase diagram to represent
typical duplex stainless steel compositions |
During solidification of duplex and superduplex weld metal, a completely
ferritic structure is formed. Ferrite solidification involves epitaxial
growth from the parent material at the fusion boundary, i.e. from the high
temperature HAZ (HTHAZ), which is also fully ferritic at temperatures
approaching the melting point. Initial dendrite growth is oriented in relation
to the thermal gradient and produces a columnar ferritic structure. This
provides the starting conditions for further solid state transformations upon
cooling, and will dominate the final weld metal structure.
Further cooling initiates the formation of austenite, nucleating at the
ferrite grain boundaries. Austenite precipitation starts below the ferrite
solvus temperature, which is dependent on the weld metal composition. Austenite
forms intergranularly as Widmanstätten
side plates or intragranular plates depending on the ferrite grain size and
cooling rate.
Austenite precipitation occurs by diffusion-controlled nucleation and growth,
whereby the diffusion of interstitial elements (nitrogen and carbon) is the
controlling process. Thus, cooling rate is of major importance in determining
the extent of transformation. Slow cooling rates result in more austenite
formation. Very low heat inputs should be avoided, since these give rise to
rapid cooling and insufficient time for adequate austenite development.
Welding procedures commonly specify a minimum arc energy to avoid overly
rapid cooling and high ferrite content. The necessity to avoid low arc energy is
dependent on steel composition and the higher nitrogen superduplex alloys are
normally more tolerant of low arc energy, in view of the effect of this element
in elevating the ferrite solvus temperature (Fig.1). Nevertheless,
matching composition weld metal microstructures may contain over 50%
ferrite.
Commercial consumables are of a higher nickel level than the base steel,
typically by around 4%, to promote austenite formation in the as-deposited fused
region. The use of consumables with higher nickel content is discussed in more
detail under Consumable
types. High heat input levels must be avoided, as they may encourage slow
cooling and precipitation of intermetallic phases, especially in the higher
alloy superduplex grades. A maximum heat input level and interpass temperature
are consequently specified. More information under Precipitation
of potentially detrimental phases in duplex stainless steels.
The HAZ, next to the fusion
boundary, may be considered in terms of a -
- high temperature HAZ (HTHAZ)
- - the zone adjacent to the fusion boundary which approaches the melting
point and will become completely, or almost completely, ferritic on heating,
depending on peak temperature, and a -
- low temperature HAZ (LTHAZ)
- - a zone further from the fusion boundary, where the ferrite-austenite
phase balance remains essentially unchanged
HTHAZ
The microstructure in the HTHAZ is controlled by the welding
thermal cycle and steel composition. Problems associated with almost fully
ferritic zones, i.e. inadequate solid state transformation to austenite on
cooling, have largely been solved by increasing nitrogen levels in the steels.
As a result, modern duplex stainless steels normally show ferrite levels in the
HTHAZ in the range 50-65%, provided that appropriate welding practices are used.
For a given steel composition, the microstructure in the HTHAZ is controlled
by the imposed thermal history, so, heat input, material thickness, preheat and
interpass temperatures are all influential. More specifically, it is the peak
temperature and exposure time to the single-phase ferritic region which controls
austenite dissolution and ferrite grain size. Diffusion controlled austenite
reformation, which largely involves nitrogen redistribution, occurs at the
ferrite grain boundaries and by Widmanstätten or intragranular side plate
growth, similar to the weld metal. A large ferrite grain size has a retarding
effect on austenite formation.
In the HTHAZ, lower temperature reactions may occur in rapidly cooled welds,
such as nitride
formation if insufficient austenite reformation takes place. In such
circumstances, a nitride free zone is observed in the ferrite adjacent to the
austenite phase, reflecting the fact that the nitrogen has diffused to the
austenite from this area. In interpass regions, secondary
austenite may precipitate in areas that were originally of high ferrite
content. Such secondary austenite has low nitrogen, chromium and molybdenum
contents and consequently low pitting resistance.
LTHAZ
The LTHAZ is located further away from the fusion boundary than
the HTHAZ. Thermal history in this region must be controlled to avoid formation
of intermetallic
phases. This is a problem primarily encountered in the more highly alloyed
superduplex stainless steels, but with extremely slow cooling rates (i.e. very
high heat inputs) lower alloy grades can also be affected. Intermetallic phases,
such as sigma and chi, reduce toughness and pitting resistance. These phases
form at temperatures in the range 550-950°C, so that time in this range should
be controlled.
Welding
parameters should be chosen to ensure that overall cooling conditions are slow
enough for adequate austenite formation in the HTHAZ and weld metal and fast
enough to avoid deleterious precipitation in the LTHAZ and weld metal. This is
done by recommending heat input (arc energy) ranges, i.e. maximum and minimum
values, and maximum interpass temperatures. Various
models exist to describe the ferrite-austenite reaction at welds, based on
material composition and welding thermal cycle. Welding consumables overalloyed
in nickel are selected unless post weld solution annealing is to be used.
All common arc welding
processes may be applied to duplex stainless steels. Depending on application,
the following welding processes can be applied:
| MMA (SMAW) |
Manual Metal Arc (Shielded Metal Arc Welding) |
| TIG (GTAW) |
Tungsten Inert Gas (Gas Tungsten Arc Welding) |
| MIG (GMAW) |
Metal Inert Gas (Gas Metal Arc Welding) |
| FCAW |
Flux Cored Arc Welding |
| SAW |
Submerged Arc Welding |
| PAW |
Plasma Arc Welding |
Some details on each of these processes and their welding characteristics are
presented in Table
3. MMA (SMAW) and TIG (GTAW) are commonly regarded as the most
versatile processes for welding duplex steels, although duty cycles are lower
than for some other manual processes, e.g. FCAW, which is becoming more
common.
TIG (GTAW) welding is the most commonly used process for producing weld root
passes, specifically for pipe welding, where completion of the weld is carried
out typically with MMA (SMAW), TIG (GTAW) or SAW. Currently, MIG (GMAW) is not
widely used for welding duplex grades, however mechanised pulsed MIG procedures
and use of flux cored wires are becoming more accepted. Welding processes
characterised by very low or very high welding heat input, which may give rise
to unacceptable phase balance without additional controls or procedures, are
used only for duplex alloys in specialised applications. These processes include
resistance welding (RW), laser welding (LW), electron beam welding (EBW) and
electroslag welding (ESW). Friction welding (FW) has demonstrated acceptable
weld microstructures and properties.
Joint design
Welding without filler addition (autogenous welding) is not
typically recommended for duplex stainless steel, unless a solution anneal is
envisaged after welding. Consequently, it is essential to incorporate a root gap
in the joint design, to ensure adequate filler addition when TIG (GTA) welding.
For uniformity of joint gap, machining, rather than grinding, is recommended
during joint preparation to optimise the fit-up. If grinding has to be
performed, stainless steel compatible tooling should be used and any grinding
burr removed to avoid incomplete fusion and lack of penetration. For austenitic
stainless steels, a skilled welder can overcome some deficiencies in joint
preparation by manipulation of the torch. For duplex stainless steel, some of
these techniques may cause a longer than expected exposure in the harmful
temperature range (550-950°C) and should be avoided.
Some typical joint designs are shown in Table
4.
Material handling and pre-weld procedures
Prior to welding, it is
important that all aspects of handling are performed in a controlled manner in
order to prevent contamination, particularly by ferrous materials, of the duplex
stainless steel components. Poor handling practice can lead to rust spots and
potentially provide sites for in-service pitting corrosion. It is for this
reason that material segregation controls are adopted, whereby duplex
fabrication is separated from areas where other materials are being worked.
Suitably protected lifting equipment, covered fork lift arms, storage racking
and adequately protected rollers and pipe stands are employed for handling.
Pre-welding procedures for duplex stainless steels should follow normal
practices applied to austenitic stainless steels. Dirt, grease, oil, paint and
sources of moisture of any sort will interfere with welding operations and may
adversely affect the corrosion resistance and mechanical properties of the
weldment and so must be removed.
Preheating
As a general rule, preheat is not necessary, but warming of
field welds is recommended, where ambient temperatures are below 5°C, in order
to dry the joint area and avoid condensation.
Arc energy and interpass temperature selection
Noting the limitations of
stipulating arc energy alone, 22% Cr steels of, say, 10-20mm thickness can be
welded at 0.5 to 2kJ/mm. For thinner walled 22% Cr duplex material, heat input
should perhaps be restricted to £1.5kJ/mm, depending on the criticality of the
application. For 25% Cr superduplex grades of 10-20mm thickness, a lower maximum
arc energy should be specified to prevent intermetallic formation. The
recommended arc energy window is typically 0.5 to 1.5 kJ/mm.
Interpass temperature will depend on the duplex grade concerned, and arc
energy employed. For 22% Cr duplex alloys S31803 and S32205, for example, an
interpass temperature of 150°C has been recommended for an arc energy of
1.5kJ/mm and a thickness of 10-12mm, although this may be increased to 225°C if
lower arc energies can be guaranteed. Interpass temperatures may also be
restricted further for thin walled material to as low as 100°C. For superduplex
grades, maximum interpass temperatures as low as 75°C may be stipulated,
particularly when welding thin walled pipes (<5mm) using TIG (GTAW), for
example. 150°C is a more common limit for thicker material.
Intermetallic
precipitation is 'additive' in multipass welds and, provided that sufficient
austenite reformation takes place, lower arc energy is preferred for filler weld
runs, for optimum results. For this reason, the pass following deposition of the
root bead is often termed the 'cold' pass, as it is deposited at a heat input
lower than that of the root bead. In particular, low arc energy for the root run
and higher arc energy for the second and subsequent passes should be avoided, as
it can lead to appreciable reduction in corrosion resistance, due to the
additive effects of adverse alloying element segregation between ferrite and
austenite and the formation of intermetallic
phases and secondary austenite.
Shielding and backing gases for gas shielded processes
Shielding gases
A good gas shield during welding is essential to protect
the weld pool from atmospheric oxidation and contamination. Most typically, this
protection is achieved using argon, but gas mixtures of Ar/He can also be
employed. Oxygen content in the gas shield on both sides of the weld should be
as low as possible, ideally <25ppm, although this is seldom achievable on the
root side in practice, and, in order to assure complete flushing of air and full
protection of the weld, gas flows should be initiated several seconds prior to
striking the arc and should be maintained for several seconds after the arc is
extinguished, ideally long enough for the weld and HAZ to cool below the
oxidation range of the stainless steel.
Nitrogen influences weld metal microstructure and corrosion properties, so,
it is important to control nitrogen during welding. Using nitrogen-free
shielding gases for gas shielded processes will typically result in the loss of
some nitrogen from the weld pool. The extent of loss is dependent on the
nitrogen content of the steel in question, the arc length (voltage) and arc
energy. To prevent this problem, binary (Ar/N2)
and tertiary (Ar/He/N2) mixtures are
commercially available for use as shielding or plasma gases (0.5-2.5%N2). Their use can prevent nitrogen loss, and even result in
increased weld metal nitrogen content. However, it should be noted that too much
nitrogen can lead to weld metal porosity. Also, high levels of N2 result in excessive wear of the TIG (GTAW) electrode.
Typically, a 1% nitrogen addition may be used on 22% Cr steel and 2% nitrogen
for superduplex grades.
Backing gases
Backing gases for internal protection of single-sided
welds (in pipes, for example) can be either pure argon, or high purity nitrogen.
Welds produced using a backing gas containing nitrogen have performed well in
pitting tests. Such gases can contain up to 100% N2, although this may cause high austenite content on the
root surface. Some manufacturers recommend using 'Formier' gas, which is a 90%
N2 + 10% H2
mixture, as backing gas. This gas provides good protection against oxidation and
improved weld appearance due to better wetting. However, some hydrogen pick-up
will arise, and cannot easily diffuse out of the weld unless solution annealing
after welding is performed. Hydrogen will embrittle weldments containing high
ferrite levels, leading potentially to hydrogen cracking, but this should not be
of concern provided the weld metal ferrite content is below say 55-60% and the
hydrogen content fairly low, say <10ppm.
Backing gas should be employed in such a way that several changes of volume
have occurred prior to commencing welding, and it is typical practice, for
multipass welds, to maintain the gas for approximately 3 or 4 passes, i.e. until
the root pass ceases to be significantly heated by subsequent welding.
It is normal industrial
practice to weld duplex alloys with 'matching' duplex stainless steel filler
material (consumable) of similar composition to the parent metal. Although
sometimes termed 'matching', the consumables are overalloyed with nickel, when
compared with equivalent parent metal specifications. Usually there is about
3-4% more nickel than in the parent metal, in order to obtain a satisfactory
phase balance in the weld metal in the as-welded condition.
Weld metal is typically less corrosion resistant than parent steel of the
same initial composition, due to nitrogen loss and element partitioning. Thus,
the welding consumables typically are slightly overalloyed with respect to
parent steel and for aggressive environments, a superduplex consumable is
sometimes employed for welding 22% Cr duplex grades. More highly alloyed duplex
consumables are not available to weld superduplex grades due to the increased
risk of intermetallic formation, although nickel alloy types with high levels of
Mo may be adopted.
There are several national and international standards or working documents
for duplex and superduplex stainless steel welding consumables, although
standardisation of these is limited: IIW (Subcomm.IIE, doc. II-E-xx-91), AWS
A5.4-xx, AWS A5.9-xx and CEN (TC121 prEN..). These draft documents primarily
refer to the all-weld metal properties. Most of the leading welding consumables
manufacturers produce and market consumables for duplex and superduplex
stainless steel grades. Table
5 provides an overview of some of the available products.
Covered electrodes are available with either a rutile or a basic covering.
The basic type electrode is good for all position welding, but rutile (sometimes
referred to as rutile/basic) electrodes also perform well in most practical
applications. Basic flux systems give better weld metal toughness due to low
oxygen levels.
Genuinely matching filler wires are available for joints which are to be
solution annealed, such that the high ferrite structure formed on welding will
be returned to a more balanced structure by the solution anneal.
Postweld heat treatment
There is normally no need to postweld heat treat
duplex stainless steel welds, provided that a nickel overalloyed consumable and
appropriate procedure have been used. Indeed, intermediate temperature 'stress
relief' heat treatments may be harmful and should be avoided, as they will tend
to precipitate intermetallic phases or alpha prime.
Any postweld heat treatment, e.g. when a matching filler wire is used, should
be a full solution anneal followed by water quenching, according to the
manufacturer's specification.
For repair welding of castings
and other structures, MMA (SMAW) is usually selected. Repair of duplex welds
should be undertaken with caution, as the repair weld operation essentially adds
another heating and cooling cycle. Excessive exposure of previous weld runs and
heat-affected zones to high temperatures should be avoided, to prevent
intermetallic formation, which may cause deterioration of mechanical properties
and corrosion resistance. A sequence of weld runs, with limited heat input (e.g.
1.5kJ/mm maximum), rather than a single pass leading to prolonged exposure to
high temperature may be preferable. Also, repair welding of thin walled areas
should be approached with caution, so that overheating and slow cooling leading
to formation of intermetallic phases does not occur.
Arc energy should be selected relative to the material composition and
thickness involved, in a similar manner to the original weld. Low heat input
runs and 'cosmetic' TIG melting should be avoided, as they may cause local high
ferrite content.
Duplex stainless
steels can be welded to other grades of duplex stainless steel, to austenitic
stainless steels, and to carbon steels and low alloy steels.
When welding duplex steels to 300 series austenitic grades, an austenitic
filler metal with low carbon and a molybdenum content intermediate between the
two steels may be used, e.g. AWS ER309 LMo. Alternatively, a duplex filler or a
nickel-based filler could be used, depending on the grades in question. For
example, a nickel filler would be favoured for welding to a high alloy
superaustenitic grade. The same fillers may also be used to join duplex
stainless steels to carbon and low alloy steels. Table
6 summarises welding consumables that can be employed to weld duplex
stainless steels to dissimilar metals.
Stray arc
strikes on the base material or on the completed welds must be avoided, as high
cooling rates and consequently high ferrite contents can be produced. Such arc
strikes may subsequently be the sites of severe local corrosion and lead to
premature failure of the whole component if not removed.
Other welding defects, weld spatter, weld discolouration, flux and slag can
also cause problems. They can act as crevices and initiate crevice corrosion in
chloride-containing environments and should be avoided or removed after welding
either by chemical cleaning, mechanical cleaning, or a combination of the
two.
Residual slag can be removed by fine grinding/abrasion. However rotary
brushing (power brushing) is not recommended due to the excessive surface
deterioration and formation of fine crevices which may occur. Light oxidation or
contamination can be successfully cleaned by pickling with common pickling
pastes or liquids. This pickling agent should be removed completely after
treatment. For maximum weld corrosion resistance, pickling should be employed,
to remove all welding oxides and restore the passive film.
Section 3 - Weld procedure qualification testing
Defining the most
appropriate weld procedure qualification (WPQ) tests has been the subject of
much debate and there is still no universally agreed testing scheme. The
guidance in ASME IX or BS EN288 should be followed, but experience has shown
that most customers require additional tests. Frequently requested tests include
ferrite content measurement and ASTM G48 corrosion testing, in addition to
Charpy tests and hardness measurement. For pipelines, selection of tests has
been formalised in BS EN4515 Part 2, but for other applications, test
requirements typically remain customer specific.
Ferrite
measurement is frequently required for weld procedure qualification. There are
two types of method for measuring the ferrite content of weld metals and parent
materials: (i) point counting techniques and (ii) magnetic methods.
Point counting
A metallographic section is required, prepared to a 1µm
finish and etched appropriately. TWI recommends etching electrolytically in 40%
KOH solution, although sulphuric or oxalic acid may be substituted. The KOH etch
stains the ferrite phase typically a brown/orange, or blue colour in contrast to
the white austenite so that the volume content of ferrite can be measured using
a point counting procedure. For very fine structures, an acid etch may be
preferable. Point counting should follow procedures given in ASTM E562-89. The
main advantage of the point counting technique is that it can be applied to all
microstructures, including the narrow HAZ. Point counting is, however, a
destructive technique unless surface replication is employed and is difficult to
perform in-situ. It is also relatively slow.
Magnetic measurement
This second group of techniques takes advantage of
the different magnetic properties of the two phases: ferrite is ferro-magnetic,
whilst austenite is not. An arbitrary 'Ferrite Number' (FN) is assigned to a
given level of magnetic attraction, defined from primary standards using a
magnetic beam balance known commercially as a MagneGage instrument. There is no
single correlation of Ferrite Number with ferrite content, as it depends on
composition, although some relationships have been proposed for duplex steel and
at low values, FN is approximately equivalent to the percentage ferrite. Use of
the MagneGage requires a flat polished specimen but hand held electromagnetic
instruments for non-destructive measurement of welds are available.
The advantages of magnetic techniques are that the equipment is portable and
they may be non-destructive. However, questions over the accuracy of the
techniques have been noted in working practice, and the International Institute
of Welding is currently investigating this. It is, therefore, important that
calibration and working practices adhere to established, reproducible methods,
e.g. as given in AWS A4.2: 1997 'Standard procedures for calibrating magnetic
instruments to measure the delta ferrite content of austenitic and duplex
ferritic-austenitic stainless steel weld metal'.
Impact testing is typically
required for weld procedure qualification. Other
mechanical properties, such as tensile properties, hardness and fracture
toughness are discussed at the end of this Section of the best practice guide.
Duplex stainless steels, despite their high strength, exhibit good ductility
and toughness. Duplex steels are not as tough as austenitic stainless steels,
which do not show a ductile-to-brittle transition and maintain good toughness
down to cryogenic temperatures. However, compared with carbon steel or ferritic
stainless steels, the ductile-to-brittle transition curve of duplex steels is
more gradual, with some duplex grades able to retain good toughness down to
-100°C.
It is normally the case that the absorbed energy is lower and ductile/brittle
transition temperature (DBTT) higher for the weld area, compared with wrought
base material, particularly in the as-welded condition. For this reason, it is
often the achievable weld metal or HAZ toughness that determines
fitness-for-purpose, especially in low temperature applications.
Flux shielded welding processes invariably have lower weld metal toughness
compared with gas-shielded methods. This is mainly a function of oxygen and
hence inclusion content. Weld metal phase balance also affects impact toughness.
High ferrite levels have a negative affect on weld toughness. The general trend
is that a slower cooling rate promotes austenite formation and increases
toughness. Multipass deposits can also show greater toughness as a result of
further austenite formation in the reheated regions. On the other hand,
excessive heat input or reheat, particularly in the superduplex grades, may
cause precipitation of intermetallic phases, which have a dramatic adverse
effect on toughness. A full solution anneal PWHT can give an improvement in
toughness, as the austenite phase level is increased and any precipitates will
be re-dissolved.
The ASTM G48A testing
procedure (10% ferric chloride test) may be used to determine critical pitting
temperature (CPT), or assess whether a weld meets a specified minimum CPT. Such
tests are frequently required for weld procedure qualification. Other corrosion
tests, e.g. intergranular tests, are rarely specified for the
ferritic-austenitic grades due to their low carbon content and duplex
microstructure.
Relatively high levels of Cr, Mo and N give the 22% Cr, high alloy and
superduplex grades good resistance to chloride pitting and crevice corrosion, as
formalised in the PREN formula. However, there
is potential for corrosion attack in the region of weldments, because of the
substantial metallurgical changes that occur during welding. Local composition
and microstructure in weld metal and HAZ are important for resistance to
corrosion in service. Surface condition is equally important for pitting
corrosion and, therefore, postweld cleaning may be very effective in improving
performance.
Tensile properties
Provided that an appropriate duplex filler is
employed, there should be no difficulty achieving tensile strength values
specified for the parent steel in a cross-weld specimen over the temperature
range normally used for duplex stainless steels.
Further, there is little change in the tensile properties of the weld metal
over a wide range of ferrite levels. Elongation of the weld metal is normally
lower than that of the base metal, although values of elongation of around 25%
can be achieved.
Hardness
The hardness of duplex weldments is greater than that of the
base material, unless it is cold worked, due to strain induced by the heating
and cooling cycle. Also, there is an effect of alloy addition, as superduplex
grades are invariably harder than lower alloy grades.
The increase in hardness is manifest in both weld metal and HAZ, particularly
in the root region. This strain-induced hardening is caused by compression of
the region during cooling and is a function of the number of weld passes.
(Multipass welds in thicker material produce greater hardness values). Hardness
of weldment and base material alike are an important consideration if the
material is intended for service in sour conditions, where NACE hardness
requirements are invoked.
In NACE MR0175-2000 it is stated that the hardness of weldments should meet
the same hardness limits as the base material quoted in the Rockwell C scale
(HRC). The size of the indentor for Rockwell hardness measurement is rather
large, and so for accurate measurements of discrete regions such as small weld
passes, or narrow HAZs, a smaller Vickers indent is preferred. Table 1
shows approximate conversions of Rockwell hardness to Vickers hardness for
several of the HRC limits quoted in NACE MR0175, depending on the duplex grade,
its condition and environmental limits. This is based on a correlation of actual
hardness data for duplex steel, collated by TWI.
Table 1. Approximate conversions of Rockwell hardness to Vickers hardness
for several of the HRC limits quoted in NACE MR0175
Rockwell hardness,
HRC |
Vickers hardness,
HV10 |
| 17 |
213 |
| 20 |
246 |
| 24 |
290 |
| 28 |
334 |
| 32 |
378 |
| 34 |
400 |
| 36 |
422 |
Fracture toughness
Toughness requirements for duplex alloys are laid
down in the pressure vessel codes: BS5500, ASME VIII and DIN AD Merkblatt HP5/2,
and the impact requirements outlined in each of these codes vary slightly.
Fracture toughness, measured via CTOD, has been shown to correlate fairly
well with Charpy toughness data. Good weld area toughness may be readily
achieved down to sub-zero temperatures around -40°C. Gas shielded processes give
better weld metal toughness than flux processes due to the lower oxygen and
inclusion contents.
Common austenitic
stainless steels show only limited resistance to stress corrosion cracking in
chloride-containing media, when compared with ferritic stainless steels, which
are virtually immune in the absence of nickel. However, ferritic stainless
steels are relatively difficult to fabricate, showing a tendency towards coarse
grain formation during welding, particularly in the heat affected zone (HAZ). An
additional amount of brittle martensite can also be encountered in the weld
zone. Both phenomena cause a drop in toughness and increased cold cracking
sensitivity in a welded joint of ferritic stainless steel.
Austenitic-ferritic duplex stainless steels were developed to overcome some
of the disadvantages of the purely austenitic and ferritic grades and have been
in production for more than 60 years. The discovery of a duplex microstructure
was first described by Bain and Griffith, in 1927, when they published data on
ferritic/austenitic structures. It was not until the 1930s, however, that duplex
stainless steels became commercially available. The first wrought duplex grades
were produced in Sweden in 1930, whilst duplex castings were produced in Finland
in the same year. Prompted by observations of resistance to intergranular
corrosion in various corrosive media, combined with superior strength in
comparison with the austenitic grades, patents were issued in France in 1935 and
1937 for the forerunner of what would eventually be known as Uranus 50, a duplex
stainless steel containing 20.5-22.5% Cr, 5.5-8.5% Ni and 2-3% Mo.
A nickel shortage brought about by the Korean War (1950-1951), led to
research into duplex alloys with relatively low nickel content. It became
apparent that a balance of ferrite and austenite had better resistance to
chloride stress corrosion cracking (SCC) than a fully austenitic microstructure.
Ever since, this has been one of the main exploited advantages of duplex over
austenitic stainless steels.
Towards the end of the 1950s, the cast alloy CD4-MCu, originally with about
70% ferrite, was marketed, but had relatively poor toughness and ductility.
Although improvements to properties were achieved by lowering the chromium
content and introducing an optimised heat treatment to the process route,
understanding of the physical metallurgy of these early grades still had not
progressed sufficiently to offer a material that was easy to manufacture and
fabricate with. At this time, duplex stainless steels had a reputation for crack
sensitivity, and this limitation confined use of the first-generation duplex
stainless steels, usually in the unwelded condition, to a few specific
applications until the 1960s.
In the late 1960s and early 1970s, there were two significant factors which
advanced the development and use of duplex alloys. Firstly, another nickel
shortage pushed up the price of austenitic alloys, combined with increased
activity in the offshore oil and gas sector, which demanded a stainless steel
able to handle aggressive saline environments. Secondly, steel production
process techniques improved dramatically with the introduction of the vacuum and
argon oxygen decarburisation (VOD and AOD) practices. These techniques led to
steels with simultaneously low carbon, sulphur and oxygen contents whilst
allowing for greater control of alloy composition, especially nitrogen. The
addition of nitrogen improved corrosion resistance and also enhanced the high
temperature stability of the duplex structure by stabilising the austenite.
Modern developments
The early 1970s saw the development of the 22%
chromium duplex grade in Germany and Sweden, i.e. UNS S31803. This grade, with
excellent chloride corrosion and stress corrosion resistance, good fabricability
and high strength, became the workhorse of the second-generation duplex grades
and was used extensively for gas gathering linepipe and process applications on
offshore platforms. However, permitted compositional ranges within the standard
specifications were too wide, and highly ferritic HAZs could be formed with an
associated reduction in corrosion resistance and toughness.
Such wide compositional ranges persist in national and international
specifications (e.g. UNS S31803), even though some producers and end-users
commonly specify closer limits to control weldability. In particular, it proved
important to keep the nitrogen content towards the upper end of the range, e.g.
the limits specified for UNS S32205, which is a restricted version of S31803.
This latest development occurred as a result of the enhanced understanding and
knowledge gained in the production and welding metallurgy of duplex stainless
steel.
During the 1980s, the superduplex grades were developed. These are more
highly alloyed ferritic-austenitic grades, which can withstand more aggressive
environments and have higher strength. The superduplex steels contain about 25%
Cr, 7% Ni, 3-4% Mo, 0.2-0.3% N, 0-2% Cu and 0-2% W.
Another duplex steel developed around this time was a lean alloy grade
exemplified by UNS S32304. The lower alloy content, in particular nickel and
molybdenum, than the most common 22% Cr grade S31803 means that it is cheaper to
produce, but has a lower pitting resistance. In this respect, it competes with
the high production austenitic grades, such as 304L and 316L, on the grounds of
its higher strength, similar pitting resistance in chloride media and better
resistance to chloride SCC.
Introduction
To obtain a stable duplex microstructure that responds well
to processing and fabrication, care must be taken to obtain the correct level of
alloying elements. The major elements are chromium, nickel, molybdenum and
nitrogen. The combined effects of these elements, in terms of their tendency to
promote ferrite or austenite formation, is summarised in the chromium equivalent
(Creq) and nickel equivalent (Nieq) equations 1 and 2, respectively.
Chromium
Chromium is a ferrite former, meaning that the addition of
chromium stabilises the body-centred
cubic (
) structure of iron. Chromium equivalent is an indication of ferrite
forming tendency.
Creq = % Cr + % Mo + 0.7 x % Nb
Equ.1
The corrosion resistance of stainless steel increases with increasing
chromium content and a minimum of about 10.5% chromium is necessary to form a
stable chromium-rich passive film that is sufficient to protect a steel against
mild atmospheric corrosion. However, higher chromium contents increase the
likelihood of forming intermetallic phases, with a commensurate reduction of
toughness.
Molybdenum
Molybdenum, like chromium, is a ferrite former (see Equ.1).
Molybdenum also enhances chloride corrosion resistance. The beneficial influence
of molybdenum on pitting and crevice corrosion resistance of an alloy in
chloride solutions has been recognised for many years. This is reflected in the
PRE relationships in which it is given a coefficient of 3.3 times that of
chromium. However, Mo also increases the tendency of a stainless steel to form
detrimental intermetallic phases during hot working.
Nickel
Counter to the ferrite stabilising effect of chromium and
molybdenum, there is another group of elements, which include nickel, that
stabilise austenite. This means that adding any of these elements to a stainless
steel promotes a change in the crystal structure from ferritic (body centred
cubic) to face-centred
cubic (
), or 'austenitic'. The following equation gives an equivalence between
nickel and other austenite-stabilisers.
Nieq = % Ni + 35 x % C + 20 x % N
+ 0.25 x % Cu Equ.2
Therefore, for a duplex stainless steel, the ferrite stabilising elements
need to be balanced with the austenite stabilisers in order to maintain an
appropriate phase balance of austenite to ferrite. For this reason, the level of
nickel addition to a given duplex alloy will depend primarily on the chromium
content.
Nitrogen
Nitrogen addition to a duplex stainless steel has multiple
effects. It increases pitting and crevice corrosion resistance, and also
increases strength and austenite content. Nitrogen is, in fact, the most
effective solid solution strengthening element.
Nitrogen is a strong austenite former (see Equ.2) and partitions
preferentially to austenite in the duplex structure due to increased solubility
in this phase. Nitrogen is typically added to duplex stainless steels almost to
its solubility limit, with the amount of nickel adjusted to achieve the desired
phase balance.
Another important property of nitrogen is its ability to stabilise duplex
alloys against the precipitation of intermetallic phases such as sigma and
chi.
Phase diagram
Numerous structural changes can occur in the duplex
stainless steels during isothermal or anisothermal heat treatments. It is worth
noting that most of these transformations are concerned with the ferrite, as
interstitial element diffusion rates are approximately 100 times faster than in
austenite. This is principally a consequence of the less compact lattice of the
body-centred cubic ferrite crystal structure.
|
Fig 2.1. Cross-section through the Fe-Cr-Ni phase diagram to represent
typical duplex stainless steel compositions |
The best way to explain the phase balance of iron-chromium-nickel alloys
(i.e. the three principal elements present in duplex stainless steels) is by
using a pseudobinary cross-section through a ternary Fe-Cr-Ni phase diagram. The
section is at 68% Fe, since the iron content of most duplex stainless steel is
about 68% (Fig 2.1). The composition of duplex steels falls in the
( + ) phase field. The phase diagram illustrates that duplex alloys
solidify as ferrite and upon further cooling, some of the ferrite transforms to
austenite.
The equilibrium transformations can occur at temperatures as low as 1000°C,
depending on alloy composition, although there is little further change in the
ferrite-austenite phase balance below this temperature. The phase diagram also
illustrates the effect of nitrogen content on transformation behaviour, showing
that increased nitrogen content raises the temperature at which austenite begins
to form from the ferrite.
It can be seen in the phase diagram that the /( + ) and /( + ) phase boundaries are not vertical, and consequently the proportion of
and varies with temperature, and at high temperature the structure becomes
more ferritic. An appropriate duplex structure is obtained in wrought alloys by
hot working and annealing in the temperature range of 1000°C to 1150°C. During
such processes, nucleation of austenite occurs solely at grain boundaries,
typically giving a 'cellular' microstructure.
Solution annealing
Following any hot or warm working of wrought alloys
it is necessary to perform a full solution anneal, followed by a rapid quench,
to fully restore the mechanical properties and corrosion resistance of the
alloy. A guide to minimum solution annealing temperatures is presented in
Table 2.1. The workpiece should be taken to a temperature slightly
above the minimum solution annealing temperature and held long enough to
dissolve any intermetallic precipitates which may be present. 22% Cr grades are
typically solution annealed at around 1060°C and superduplex grades at around
1120°C. As a conservative guideline, the holding time should be comparable to
the total time that the piece was held in the 650-980°C (1200-1800°F)
temperature range subsequent to the previous full anneal. Quenching quickly into
water from the solution annealing temperature should then be performed.
Table 2.1. Minimum solution annealing temperatures for duplex stainless
steels (source: producer Data Sheets and ASTM A 480).
| Grade |
Minimum Annealing
Temperature |
| (°C) |
(°F) |
| Alloy lean |
980 |
1800 |
| 22% Cr |
1040 |
1900 |
| High alloy 25% Cr |
1040 |
1900 |
| Superduplex (depending on grade) |
1025 to 1100 |
1875 to 2010 |
Stress relief
Stress relief heat treatments, to reduce residual stresses
within a workpiece, are generally not advisable as there is no satisfactory
temperature below the solution annealing temperature at which stress relief may
be performed without the danger of formation of intermetallic phases which will
lower corrosion resistance and reduce toughness.
Ferrite ()
Body-centred cubic
(bcc): A crystal lattice with a cubic unit cell; one atom at each
corner of the cube and one atom at the centre of the cube.
 |
Fig 3.1. Ferrite (body-centred cubic, bcc) |
Epitaxial: Having the same
crystal axes. The material growing epitaxially must have a lattice spacing and
structure close to that of the substrate.
Austenite ()
Face-centred cubic: A
crystal lattice with a cubic unit cell; One atom at each corner of the cube and
one atom at the centre of each face of the cube.
 |
Fig 3.2. Austenite (face-centred cubic, fcc) |
Widmanstätten structure: likened
to a basket-weave structure (or a mesh-like distribution), the precipitating
phase forms by solid state transformation, which occurs along preferred crystal
planes. Usually produced by rapid cooling and when the transforming phase has a
large grain size.
Various models exist
to describe the ferrite-austenite reaction at welds, based on material
composition and welding thermal cycle. However, generally speaking, widespread
acceptance of these prediction models by welding engineers has not been gained
and, instead, reliance has been placed on simple relationships between
composition and microstructure, such as the Schaeffler diagram, developed in the
1940s, and the WRC-1992 diagram. These should not be regarded as exact, because
they do not recognise cooling conditions.
 |
Fig 4.1. Schaeffler diagram |
The Schaeffler diagram (Fig 4.1) is an empirical description of the
microstructures of a wide range of weld metals, predominantly stainless steels,
that result from welding different compositions (i.e. different base metal and
filler combinations). This type of diagram has been used for many years to
predict weld metal microstructures in stainless steels, and to some extent to
optimise base metal and filler compositions. However, the Schaeffler diagram
does not predict duplex microstructures well, as it makes no allowance for
nitrogen content, and other more appropriate diagrams and relationships have
been developed, culminating in the WRC-1992 diagram.
 |
Fig 4.2. WRC-1992 diagram |
The WRC-1992 diagram (Fig 4.2) is considered to be the most accurate
constitution diagram for the prediction of phase balance in duplex stainless
steels from chemical composition. However, like any other predictive diagram,
there are limitations. As an example, the weld metal ferrite number (FN) of
superduplex alloys may be overestimated with nitrogen contents below about
0.19%, but underestimated at above some 0.25% nitrogen. Also, the diagram may be
inaccurate for some levels of manganese, silicon and molybdenum and, since there
is no martensite line, it may not be appropriate for welding some dissimilar
metals, where martensite is a consideration. The Delong diagram (Fig
4.3) takes this into account, as does Schaeffler, whilst including nitrogen
but again is less accurate for modern stainless grades.
 |
Fig 4.3. Delong diagram |
Background
A number of secondary phases may form in duplex stainless
steels and weld metals subjected to temperatures in the range 300-1100°C. These
phases tend to reduce toughness and corrosion resistance of the duplex alloy.
The tendency for precipitation is strongly influenced by the content of alloying
elements and is, therefore, most pronounced in the superduplex stainless steels
and weld metals.
Partitioning in duplex and superduplex alloys means that the ferrite is
enriched in chromium and molybdenum. Both of these elements are known to promote
the formation of intermetallic phases. Furthermore, element solubility in
ferrite falls with a decrease in temperature, increasing the probability of
precipitation during heat treatment. Detrimental brittle intermetallic phases
such as sigma (
) and chi (
), as well as alpha prime (
) and various carbides and nitrides, can form in a matter of minutes at
certain temperatures. For this reason, heat treatment temperatures must be
chosen carefully (particularly for thicker wrought alloy sections) and fast
thermal cycling is generally employed, to prevent precipitation of undesirable
phases.
Where there is the possibility that undesirable phases may have formed, then
a solution annealing and quenching heat treatment is generally employed to
re-dissolve the unwanted phases. In weldments, where a range of cooling rates
may be encountered, nitride precipitation takes place primarily in predominantly
ferritic areas of the high temperature HAZ or weld metal, and may be associated
with adjacent areas of chromium depletion. Intermetallic precipitation, on the
other hand, occurs in the HAZ in regions where peak temperatures have been too
low to cause marked alteration to the austenite/ferrite balance, around
650-1000°C. In the case of weldments, a postweld solution treatment may be
employed to re-dissolve any unwanted intermetallic phases.
 |
Fig 5.1 TTT diagram for a range of ferritic-austenitic
alloys |
Time-temperature transformation (TTT) diagrams, sometimes referred to as
isothermal precipitation diagrams, produced by isothermal heat treatment
followed by quenching, are often employed to depict the susceptibility of
different grades to precipitation, e.g. Fig 5.1.
In general, precipitation processes during welding tend to be somewhat more
rapid than indicated by conventional isothermal studies, as a result of the
conjoint action of temperature and expansion and contraction stress.
Characteristics and morphology of precipitates
Austenite precipitating
below 900°C is generally termed 'secondary austenite' (or 2). Secondary austenite can form
relatively quickly and by different mechanisms depending on the temperature. The
formation can be understood on the basis that equilibrium transformations are
rarely completed in fairly short-term thermal cycles, such as are encountered
during welding. The steel, or weld metal, is first rapidly cooled from high
temperatures where the equilibrium ferrite fraction is high and later reheated
by further welding or heat treatment.
Secondary austenite has lower contents of N, Cr and Mo compared with the
primary high temperature austenite. The morphology of the secondary austenite
varies depending on location and transformation mechanism, from the
Widmanstätten-type, found predominantly in weld metal, to the globular-type seen
both in weld metal and heat-affected parent material. In either case, 2 tends to be detrimental to the
pitting corrosion resistance (a consequence of lower N, Cr and Mo contents).
Precipitation of secondary austenite has also been reported to facilitate
nucleation of Cr-rich and Ni-poor phases such as .
Sigma () and Chi () phases
Sigma-phase is a hard embrittling precipitate, which is
essentially an Fe-Cr-Mo intermetallic compound. It forms between 650° and 1000°C
and is often associated with a reduction in both impact properties and corrosion
resistance. At a temperature of around 900°C, ferrite decomposition to sigma may
take as little as 2 minutes in superduplex alloys. The phase is enriched in elements such as Cr, Mo, Si and W and depleted in
Ni and Mn. The precipitate tends to form at 
/ grain boundaries.
 |
Fig 5.2. Sigma phase |
Chi phase is a Mo-rich intermetallic phase which forms between 700° and
900°C, although in much smaller quantities than the phase. Like , -phase often forms at / boundaries and grows into the ferrite. Chi also has detrimental
effects on corrosion and toughness properties.
Nitrogen is added to duplex alloys to stabilise austenite, and to
improve strength and pitting resistance. The solubility of nitrogen is
considerably higher in austenite than in ferrite and has been shown to partition
to the former phase. Above the solution annealing temperature (approximately
1050°C for S31803) the volume fraction of ferrite increases, until, just below
the solidus, a completely ferritic microstructure can be present (though in the
higher alloy grades some austenite may remain).
 |
Fig 5.3. Nitride precipitation in weld metal ferrite
grains |
If the composition of a steel is inadequate, e.g. giving high ferrite content
in a final weld metal or HAZ structure then, even at low nitrogen levels,
numerous nitrides form in the ferrite, precipitating intragranularly in weld
metal as needle-like Cr2N (in the temperature
range 700 - 900°C). The precipitates are thought to adversely affect pitting
resistance as well as impact toughness. Welding, however, seems to favour the
formation of another nitride in the HAZ: CrN. This is reported to form from
approximately 1100°C but in the small amounts typically encountered seems to
have little or no effect on properties.
Carbides (M23C6, M7C3)
Carbides play a limited role in modern duplex
stainless steels due to their very low carbon contents. However, precipitation
of M23C6 and
M7C3 has been
reported in some duplex stainless steels. M7C3 forms between 950° and
1050°C at / boundaries. However, as its formation takes 10 minutes, its occurrence
during welding is unlikely and during annealing it can be avoided by normal
quenching techniques. M23C6, however, precipitates rapidly between 650° and 950°C,
requiring less than one minute to form at 800°C. Precipitation occurs
predominantly at / boundaries, but can also occur at / and / boundaries and to a lesser extent inside the ferrite and austenite
grains. It has been proposed that carbides promote formation of other
detrimental phases such as -phase, by providing nucleation sites.
| UNS |
Trade name(s) |
Ele ment, wt% |
Typ ical PREN |
| C |
S |
P |
Si |
Mn |
Ni |
Cr |
Mo |
Cu |
W |
N |
| Alloy lean |
|
| S31500 |
3RE60, 1
A903,9
VLX5693 |
0.030 |
0.030 |
0.030 |
1.40-2.00 |
1.20-2.00 |
4.25-5.25 |
18.0-19.0 |
2.50-3.00 |
-- |
-- |
-- |
23 |
| S32304 |
SAF2304,1,11
UR35N,2
VLX5343 |
0.030 |
0.040 |
0.040 |
1.0 |
2.50 |
3.0-5.5 |
21.5-24.5 |
-- |
0.05-0.60 |
-- |
0.05-0.20 |
25 |
| S32404 |
UR502 |
0.04 |
0.010 |
0.030 |
1.0 |
2.0 |
5.5-8.5 |
20.5-22.5 |
2.0-3.0 |
1.0-2.0 |
-- |
0.20 |
31 |
| Standard 22% Cr |
|
| S31803 |
2205,1
UR45N,2
Falc223,8
AF22,10
VS22,3
VLX562,3
DP8,4
318LN, A903,9
1.4462 / PRES35,12
NKCr22,13
SM22Cr,4
Remanit 446214 |
0.030 |
0.020 |
0.030 |
1.00 |
2.00 |
4.50-6.50 |
21.0-23.0 |
2.50-3.50 |
-- |
-- |
0.08-0.20 |
34 |
| S32205 |
UR45N +,2
22051 |
0.030 |
0.020 |
0.030 |
1.00 |
2.00 |
4.50-6.50 |
22.0-23.0 |
3.00-3.50 |
-- |
-- |
0.14-0.20 |
35 |
| High alloy |
|
| S31200 |
UHB 44LN,
UR 47N,2
VLX5473 |
0.030 |
0.030 |
0.045 |
1.00 |
2.00 |
5.50-6.50 |
24.0-26.0 |
1.20-2.00 |
-- |
-- |
0.14-0.20 |
38 |
| S31260 |
DP34 |
0.03 |
0.030 |
0.030 |
0.75 |
1.00 |
5.50-7.50 |
24.0-26.0 |
2.50-3.50 |
0.20-0.80 |
0.10-0.50 |
0.10-0.30 |
38 |
| S32550 |
Ferr alium 255,5
UR52N2 |
0.04 |
0.030 |
0.040 |
1.00 |
1.5 |
4.50-6.50 |
24.0-27.0 |
2.9-3.9 |
1.50-2.50 |
-- |
0.10-0.25 |
38 |
| S32900 |
AISI 329,
UHB 44L,
10RE51,1
453S |
0.08 |
0.030 |
0.040 |
0.75 |
1.00 |
2.50-5.00 |
23.0-28.0 |
1.00-2.00 |
-- |
-- |
-- |
33 |
| S32950 |
7Mo Plus15 |
0.03 |
0.010 |
0.035 |
0.60 |
2.00 |
3.50-5.20 |
26.0-29.0 |
1.00-2.50 |
-- |
-- |
0.15-0.35 |
36 |
| Super duplex |
|
| S32520 |
UR52N +,2
SD405 |
0.030 |
0.020 |
0.035 |
0.8 |
1.5 |
5.5-8.0 |
24.0-26.0 |
3.0-5.0 |
0.50-3.00 |
-- |
0.20-0.35 |
41 |
| S32750 |
SAF2507,1,11
UR47N
+2 |
0.030 |
0.020 |
0.035 |
0.8 |
1.20 |
6.0-8.0 |
24.0-26.0 |
3.0-5.0 |
0.5 |
-- |
0.24-0.32 |
41 |
| S32760 |
Zeron 100,6
FALC1008 |
0.03 |
0.01 |
0.03 |
1.0 |
1.0 |
6.0-8.0 |
24.0-26.0 |
3.0-4.0 |
0.5-1.0 |
0.5-1.0 |
0.2-0.3 |
>40 |
| S39226 |
|
0.030 |
0.030 |
0.030 |
0.75 |
1.00 |
5.50-7.50 |
24.0-26.0 |
2.50-3.50 |
0.20-0.80 |
0.10-0.50 |
0.10-0.30 |
>40 |
| S39274 |
DP3W4 |
0.030 |
0.020 |
0.030 |
0.80 |
1.0 |
6.0-8.0 |
24.0-26.0 |
2.50-3.50 |
0.20-0.80 |
1.50-2.50 |
0.24-0.32 |
42* |
| S39277 |
AF918,7
25.7NCu7 |
0.025 |
0.002 |
0.025 |
0.80 |
-- |
6.5-8.0 |
24.0-26.0 |
3.0-4.0 |
1.2-2.0 |
0.80-1.20 |
0.23-0.33 |
42 |
After 'Metal and Alloys in the Unified Numbering System', SAE/ASTM, September
1996.
Values are maxima unless range given.
*PREW
Manufacturers (in no particular order)
| 1 Avesta Sheffield Ltd |
8 Krupp Stahl |
| 2 Creusot-Loire
Industries |
9 Böhler Edelstahl |
| 3 Valourec |
10 Mannesmann |
| 4 Sumitomo Metal
Industries |
11 AB Sandvik Steel |
| 5 Haynes International |
12 Fabrique de Fer |
| 6 Weir Materials Ltd |
13 Nippon Kokan |
| 7 DMV Stainless/Feroni |
14 TEW |
| |
15
Carpenter |
Table 2. Cast duplex grades listed in the Unified Numbering System, with
some tradenames and typical PREN values.
| UNS |
Trade names |
Ele ment, wt% |
Typ ical PREN |
| C |
S |
P |
Si |
Mn |
Ni |
Cr |
Mo |
Cu |
N |
Other |
| Stand ard 22% Cr |
|
| J92205 |
2205,1 U
R45N,2
FALC223,7
AF22,9
VS22,3 VLX562,3 |
0.03 |
0.020 |
0.04 |
1.00 |
1.50 |
4.5-6.5 |
21.0-23.5 |
2.5-3.5 |
1.0 |
0.10-0.30 |
-- |
32-33 |
| J93183 |
DP8,4
318LN,
A903,8
KCR-D183 |
0.03 |
0.03 |
0.040 |
2.0 |
2.0 |
4.0-6.0 |
20.0-23.0 |
2.0-4.0 |
1.0 |
0.08-0.25 |
0.5-1.5 Co |
|
| High alloy |
|
| J93345 |
Escoloy |
0.08 |
0.025 |
0.04 |
-- |
1.00 |
8.0-11.0 |
20.0-27.0 |
3.0-4.5 |
-- |
0.10-0.30 |
-- |
38 |
| J93370 |
CD4-MCu,6
UR55(M)2 |
0.04 |
0.04 |
0.04 |
1.00 |
1.00 |
4.75-6.00 |
24.5-26.5 |
1.75-2.25 |
2.75-3.25 |
-- |
-- |
37 |
| J93371 |
3A |
0.06 |
0.040 |
0.040 |
1.00 |
1.00 |
4.00-6.00 |
24.0-27.0 |
1.75-2.50 |
2.75-3.25 |
0.15-0.25 |
-- |
35 |
| J93372 |
CD4-MCuN6 |
0.04 |
0.04 |
0.04 |
1.00 |
1.00 |
4.7-6.0 |
24.5-26.5 |
1.7-2.3 |
2.7-3.3 |
0.10-0.25 |
-- |
35 |
| J93550 |
KCR-D283 |
0.03 |
0.03 |
0.040 |
2.0 |
2.0 |
-- |
23.0-26.0 |
5.0-8.0 |
1.0 |
0.08-0.25 |
0.5-1.5 Co |
49 |
| Super duplex |
|
| J93380 |
Zeron 100,5
FALC1007 |
0.03 |
0.025 |
0.030 |
1.0 |
1.0 |
6.5-8.5 |
24.0-26.0 |
3.0-4.0 |
0.5-1.0 |
0.2-0.3 |
0.5-1.0W |
>40 |
| J93404 |
Alloy 958,
4469 |
0.03 |
-- |
-- |
1.00 |
1.50 |
6.0-8.0 |
24.0-26.0 |
4.0-5.0 |
-- |
0.10-0.30 |
-- |
44 |
After 'Metal and Alloys in the Unified Numbering System', SAE/ASTM, September
1996.
Values are maxima unless range given.
Manufacturers (in no particular order)
| 1 Avesta Sheffield Ltd, AB
Sandvik Steel |
6 Alloy Casting
Institute |
| 2 Creusot-Loire Industries
(Usinor Group) |
7 Krupp Stahl |
| 3 Valinox |
8 Böhler Edelstahl |
| 4 Sumitomo Metal
Industries |
9 Mannesmann |
| 5 Weir Materials Ltd |
|
| Process |
Characteristics |
MMA
(SMAW) |
Readily available, all positions, slag on weld surface to
be removed, low deposition rate. |
TIG
(GTAW) |
Requires good skill, most suitable for pipe welding, high
effect of dilution in root runs, low deposition rate, can be
mechanised/automated, i.e. orbital welding systems. |
MIG
(GMAW) |
Requires good skill, more set-up work, metal transfer
depends on wire quality (spattering), commonly only for filling of joints,
high deposition rate, can be mechanised/automated. |
| FCAW |
Limited availability of consumables, only for filling of
joints, limited positional welding, high deposition rate, slag
protection. |
| SAW |
Only mechanised, requires set-up arrangements, only
downhand (flat) welding, high dilution affects weld properties, highest
deposition rate, slag removal in joint may be difficult. |
| PAW |
Requires complex equipment, only mechanised welding, no
filler metal added: plate composition determines weld properties, high
welding speed. |
| Joint |
Welding Process |
Thickness, t (mm) |
Gap, d (mm) |
Root, k (mm) |
Bevel, (mm) |
 |
MMA (SMAW)
TIG (GTAW)
MIG
(GMAW)
SAW |
3 - 15
2.5 - 8
3 - 12
4 - 12 |
2 - 3
2 - 3
2 - 3
2 - 3 |
1 - 2
1 - 2
1 - 2
1 - 2 |
60 - 70
60 - 70
60 - 70
70 - 80 |
 |
MMA (SMAW)
TIG (GTAW)
MIG
(GMAW)
SAW |
12 - 60
>8
>12
>10 |
1 - 2
1 - 2
1 - 2
1 - 2 |
2 - 3
1 - 2
2 - 3
1 - 3 |
10 - 15
10 - 15
10 - 15
10 - 15 |
 |
MMA (SMAW)
TIG (GTAW)
MIG (GMAW) |
>10
>10
>10 |
1.5 - 3
1.5 - 3
0 |
1 - 3
1 - 3
3 - 5 |
55 - 65
60 - 70
90 |
 |
MMA (SMAW)
TIG (GTAW)
MIG (GMAW) |
>25
>25
>25 |
1 - 3
1 - 3
0 |
1 - 3
1 - 3
3 - 5 |
10 - 15
10 - 15
10 -
15 |
| Welding consumable manufacturer |
Process |
Filler material classification
(acc. CEN classification principle) |
| |
|
X 22 9 3 L |
X 25 9 3 Cu L |
X 25 9 4 L |
| Avesta |
MMA (SMAW) |
2205-PW |
|
2507/P100 |
| |
TIG (GTAW) |
2205 |
|
2507 |
| |
GMAW (MIG) |
2205 |
|
2507 |
| |
SAW* |
2205 |
|
2507 |
| Bohler |
MMA (SMAW) |
Fox CN 22/9N |
Fox Duplex Cu3.0 |
Fox CN 26/10 N** |
| |
TIG (GTAW) |
CN 22/9-1G |
|
|
| |
GMAW (MIG) |
CN 22/9-1G |
|
|
| |
SAW* |
CN 22/9-UP |
|
|
| ESAB |
MMA (SMAW) |
OK 67.50
OK 67.53 |
|
OK 68.53
OK 68.55 |
| |
TIG (GTAW) |
OK Tigrod 16.86 |
|
OK Tigrod 16.88 |
| |
GMAW (MIG) |
OK Autorod 16.86 |
|
OK Autorod 16.88 |
| |
FCAW |
OK Tubrod 14.37 |
|
|
| |
SAW* |
OK Autorod 16.86 |
|
OK Autorod 16.88 |
| Filarc |
MMA (SMAW) |
RS 22.9.3 LCN |
|
RS 25.10.4 LCN |
| |
TIG (GTAW) |
PZ 65.17 |
|
|
| |
GMAW (MIG) |
PZ 60.17 |
|
|
| |
SAW* |
|
|
|
| Messer Lincoln |
MMA (SMAW) |
Grinox 62 |
Grinox 63 |
|
| |
|
Grinox 33 |
Grinox 37 |
|
| |
TIG (GTAW) |
Grinox T-62 |
|
|
| |
GMAW (MIG) |
Grinox S-62 |
|
|
| |
SAW* |
Grinox UP-23 8 3 NL |
|
|
| Metrode |
MMA (SMAW) |
Supermet 2205 |
Supermet 2507Cu |
Supermet 2507 |
| |
|
2205KS |
|
Zeron 100XKS*** |
| |
TIG (GTAW) |
|
|
Zeron 100X*** |
| |
GMAW (MIG) |
|
|
Zeron 100X*** |
| |
SAW* |
|
|
|
| Sandvik |
MMA (SMAW) |
22.9.3.LR |
|
25.10.4.LR |
| |
|
|
|
25.10.4.LB |
| |
TIG (GTAW) |
22.8.3.L |
|
25.10.4.L |
| |
GMAW (MIG) |
22.8.3.L |
|
25.10.4.L |
| |
SAW* |
22.8.3.L |
|
25.10.4.L |
| Lincoln Smitweld |
MMA (SMAW) |
Arosta 4462 |
Jungo 4462 |
Jungo Zeron 100X |
| |
TIG (GTAW) |
LNT 4462 |
|
LNT Zeron 100X*** |
| |
GMAW (MIG) |
LNM 4462 |
|
LNM Zeron 100X*** |
| |
FCAW |
Cor-A-Rosta 4462 |
|
|
| |
SAW* |
LNS 4462 |
|
LNS Zeron 100X*** |
| Soudometal |
MMA (SMAW) |
Soudinox S 4462 |
Soudinox S 52 |
Soudinox S100 |
| |
|
|
Soudinox S 47 |
|
| |
TIG (GTAW) |
Soudotig 22 9 3L |
|
|
| |
GMAW (MIG) |
Soudor G 22 9 3L |
|
|
| |
SAW* |
Soudor 22 9 3L |
|
|
| TEW |
MMA (SMAW) |
Thermanit 22/09W |
|
|
| |
|
Thermanit 22/09 |
|
|
| |
TIG (GTAW) |
22/09/SG |
|
|
| |
GMAW (MIG) |
22/09/SG |
|
|
| |
SAW* |
22/09/UP |
|
|
* filler wire in combination with appropriate
flux
** E 25 10 3 L
*** in addition also parent material
matching composition
| |
Alloy lean |
22% Cr |
High alloy 25% Cr |
Superduplex |
| Alloy lean |
E2209 |
E2209 |
E2209 |
E2209 |
| 22% Cr |
E2209 |
E2209 |
superduplex |
superduplex |
| High alloy 25% Cr |
E2209 |
superduplex |
superduplex |
superduplex |
| Superduplex |
E2209 |
superduplex |
superduplex |
superduplex |
| 304 |
E309L/E309LMo
E2209 |
E309L/E309LMo
E2209 |
E309L/E309LMo
E2209 |
E309L/E309LMo
E2209 |
| 316 |
E309L/E309LMo
E2209 |
E309L/E309LMo
E2209 |
E309L/E309LMo
E2209 |
E309L/E309LMo
E2209 |
Carbon Steel
Low Alloy steel |
E309L/E309LMo |
E309L/E309LMo |
E309L/E309LMo |
E309L/E309LMo |
| Superaustenitic |
Nickel alloy |
Nickel alloy |
Nickel alloy |
Nickel alloy |
REPRODUCED - COURTESY TWI-UK
