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WeldingTechnology2 English
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2005
ISF – Welding and Joining Institute RWTH – Aachen University
Lecture Notes
Welding Technology 2 Welding Metallurgy
Prof. Dr. –Ing. U. Dilthey
Table of Contents Chapter Subject Page
1. Weldability of Metals 3
2. TTT - Diagrams 8
3. Residual Stresses 21
4. Heat Treatment and its
Function During Welding 31
5. Welding Plain and
Low Alloy Steels 44
6. Welding High Alloy Steels 70
7. Welding of Cast Materials 89
8. Welding of Aluminium 96
9. Welding Defects 108
10. Testing of Welded Joints 126
1.
Weldability of Metals
1. Weldability of Metals 4
DIN 8580 and DIN 8595 classify welding into production technique main group 4 "Joining“,
group 3.6 "Joining by welding“, Figure 1.1.
Weldability of a component is determined
by three outer features according to DIN
8528, Part 1. This also indicates whether a
given joining job can be done by welding,
Figure 1.2.
Figure 1.1
Figure 1.2
1. Weldability of Metals 5
Material influence on weldability, i.e. welding
suitability, can be detailed for a better un-
derstanding in three subdefinitions, Figure
1.3.
The chemical composition of a material and
also its metallurgical properties are mainly set
during its production, Figure 1.4. They have a
very strong influence on the physical
characteristics of the material.
Process steps on steel manufacturing, shown
in Figure 1.4, are the essential steps on the
way to a processible and usable material.
During manufacture, the requested chemical
composition (e.g. by alloying) and metallurgi-
cal properties (e.g. type of teeming) of the
steel are obtained.
Another modification of the mate-
rial behaviour takes place during
subsequent treatment, where the
raw material is rolled to processi-
ble semi-finished goods, e.g. like
strips, plates, bars, profiles, etc..
With the rolling process, material-
typical transformation processes,
hardening and precipitation proc-
esses are used to adjust an opti-
mised material characteristics
(see chapter 2).
Figure 1.4
Figure 1.3
1. Weldability of Metals 6
A survey from quality point of view about the influence of the most important alloy elements
to some mechanical and metallurgical properties is shown in Figure 1.5.
Figure 1.6 depicts the deci-
sive importance of the car-
bon content to suitability of
fusion welding of mild steels.
A guide number of flawless
fusion weldability is a carbon
content of C < 0,22 %. with
higher C contents, there is a
danger of hardening, and
welding becomes only pos-
sible by observing special
precautions (e.g. pre- and
post-weld heat treatment).
Figure 1.5
Figure 1.6
1. Weldability of Metals 7
In addition to material behaviour, weldability is also essentially determined through the design
of a component. The influence of the design is designated as welding safety, Figure 1.7.
The influence of the manufac-
turing process to weldability is
called welding possibility,
Figure 1.8. For example, a
pre- and post-weld heat
treatment is not always possi-
ble, or grinding the weld sur-
face before welding the
subsequent pass cannot be
carried out (narrow gap weld-
ing).
Figure 1.7
Figure 1.8
2.
TTT - Diagrams
2. TTT – Diagrams 9
An essential feature of low
alloyed ferrous materials is
the crystallographic trans-
formation of the body-
centred cubic lattice which
is stable at room tempera-
ture (a-iron, ferritic struc-
ture) to the face-centred
cubic lattice (?-iron, aus-
tenitic structure), Figure
2.1. The temperature,
where this transformation
occurs, is not constant but
depends on factors like
alloy content, crystalline structure, tensional status, heating and cooling rate, dwell times,
etc..
In order to be able to
understand the basic
processes it is necessary to
have a look at the basic
processes occuring in an
idealized binary system.
Figure 2.2 shows the state
of a binary system with
complete solubility in the
liquid and solid state.
If the melting of the L1 alloy
is cooling down, the first
crystals of the composition
c1 are formed with reaching
the temperature T1. These crystals are depicted as mixed crystal a, since they consist of a
compound of the components A (80%) and of B (20%). Further, a melting with the composi-
tion c0 is present at the temperature T1. With dropping temperature, the remaining melt is en-
Figure 2.1
Figure 2.2
2. TTT – Diagrams 10
riched with component B, following the course of line Li (liquidus line, up to point 4). In paral-
lel, always new and B richer a-mixed crystals are forming along the connection line So
(solidus line, points 1, 2, 5). The distribution of the components A and B in the solidified struc-
ture is homogeneous since concentration differences of the precipitated mixed crystals are
balanced by diffusion processes.
The other basic case of complete solubility of two components in the liquid state and of com-
plete insolubility in the solid state shows Figure 2.3 If two components are completely insolu-
ble in the solid state, no mixed crystal will be formed of A and B. The two liquidus lines Li cut
in point e which is also designated as the eutectic point. The isotherm Te is the eutectic line.
If an alloy of free composition solidifies according to Figure 2.3, the eutectic line must be cut.
This is the temperature (Te) of the eutectic transformation:
S ? A+B (T = Te = const.).
This means that the melt at a constant temperature Te dissociates in A and B. If an alloy of
the composition L2 solidifies, a purely eutectic structure results. On account of the eutectic
reaction, the temperature of the alloy remains constant up to the completed transfo rmation
(critical point) (Figure 2.2).
Eutectic structures are normally fine-grained and show a characteristic orientation between
the constituents. The alloy L1 will consist of a compound of alloy A and eutectic alloy E in the
solid state.
You can find further in-
formation on transforma-
tion behaviour in relevant
specialist literature.
The definite use of the
principles occurs in the
iron-iron carbide diagram.
Transformation behaviour
of carbon containing iron
in the equilibrium condi-
tion is described by the Figure 2.3
2. TTT – Diagrams 11
stable phase diagram iron-graphite (Fe-C). In addition to the stable system Fe-C which is
specific for an equilibrium-close cooling, there is a metastable phase diagram iron cementite
(Fe-Fe3C). During a slow cooling, carbon precipitates as graphite in accord with the stable
system Fe-C, while during accelerated cooling, what corresponds to technical conditions,
carbon precipitates as cementite in agreement with the metastable system (Fe-Fe3C). Per
definition, iron carbide is designated as a structure constituent with cementite although its
stoichiometric composition is identical (Fe3C). By definition, cementite and graphite can be
present in steel together or the cementite can decompose to iron and graphite during heat
treatment of carbon rich alloys. However, it is fundamentally valid that the formation of ce-
mentite is encouraged with increasing cooling rate and decreasing carbon content. In a dou-
ble diagram, the stable
system is shown by a
dashed, the metastable by
a solid line, Figure 2.4.
The metastable phase
diagram is limited by the
formation of cementite with
a carbon content of 6,67
mass%. The strict
stoichiometry of the
formed carbide phase can
be read off at the top X-
coordinate of the molar
carbon content. In accordance with the carbon content of Fe3C, cementite is formed at a mo-
lar content of 25%. The solid solutions in the phase fields are designated by Greek charac-
ters. According to convention, the transition points of pure iron are marked with the character
A - arrêt (stop point) and distinguished by subjacent indexes. If the transition points are de-
termined by cooling curves, the character r = refroidissement is additionally used. Heat-up
curves get the supplement c - chauffage. Important transition points of the commercially more
important metastable phase diagram are:
- 1536 °C: solidification temperature (melting point) δ-iron,
- 1392 °C: A4- point γ- iron,
Figure 2.4
2. TTT – Diagrams 12
- 911 °C: A3- point non-magnetic α- iron,
with carbon containing iron:
- 723 °C: A1- point (perlite point).
The corners of the phase fields are designated by continuous roman capital letters.
As mentioned before, the system iron-iron carbide is a more important phase diagram for
technical use and also for welding techniques. The binary system iron-graphite can be
stabilized by an addition of silicon so that a precipitation of graphite also occurs with
increased solidification velocity. Especially iron cast materials solidify due to their increased
silicon contents according to the stable system. In the following, the most important terms
and transformations should be explained more closely as a case of the metastable system.
The transformation mechanisms explained in the previous sections can be found in the bi-
nary system iron-iron carbide almost without exception. There is an eutectic transformation in
point C, a peritectic one in point I, and an eutectoidic transformation in point S. With a tem-
perature of 1147°C and a carbon concentration of 4.3 mass%, the eutectic phase called Le-
deburite precipitates from cementite with 6,67% C and saturated γ-solid solutions with 2,06%
C. Alloys with less than 4,3 mass% C coming from primary austenite and Ledeburite are
called hypoeutectic, with more than 4,3 mass% C coming from primary austenite and Lede-
burite are called hypereutectic.
If an alloy solidifies with less than 0,51 mass percent of carbon, a δ-solid solution is formed
below the solidus line A-B (δ-ferrite). In accordance with the peritectic transformation at
1493°C, melt (0,51% C) and δ-ferrite (0,10% C) decompose to a γ-solid solution (austenite).
The transformation of the γ-solid solution takes place at lower temperatures. From γ-iron with
C-contents below 0.8% (hypoeutectoidic alloys), a low-carbon α-iron (pre-eutectoidic ferrite)
and a fine-lamellar solid solution (perlite) precipitate with falling temperature, which consists
of α-solid solution and cementite. With carbon contents above 0,8% (hypereutectoidic alloys)
secondary cementite and perlite are formed out of austenite. Below 723°C, tertiary cementite
precipitates out of the α-iron because of falling carbon solubility.
2. TTT – Diagrams 13
The most important distinguished feature of the three described phases is their lattice struc-
ture. α- and δ-phases are cubic body-centered (CBC lattice) and γ-phase is cubic face-
centered (CFC lattice), Figure 2.1.
Different carbon solubility of solid solutions also results from lattice structures. The three
above mentioned phases dissolve carbon interstitially, i.e. carbon is embedded between the
iron atoms. Therefore, this types of solid solutions are also named inte rstitial solid solution.
Although the cubic face-centred lattice of austenite has a higher packing density than the cu-
bic body-centred lattice, the void is bigger to disperse the carbon atom. Hence, an about 100
times higher carbon solubility of austenite (max. 2,06% C) in comparison with the ferritic
phase (max. 0,02% C for α-iron) is the result. However, diffusion speed in γ-iron is always at
least 100 times slower than in α-iron because of the tighter packing of the γ-lattice.
Although α- and δ-iron show the same lattice structure and properties, there is also a differ-
ence between these phases. While γ-iron develops of a direct decomposition of the melt (S
→ δ), α-iron forms in the solid phase through an eutectoidic transformation of austenite (γ →
α + Fe3C). For the transformation of non- and low-alloyed steels, is the transformation of δ-
ferrite of lower importance, although this δ-phase has a special importance for weldability of
high alloyed steels.
Unalloyed steels used in industry are multi-component systems of iron and carbon with alloy-
ing elements as manganese, chromium, nickel and silicon. Principally the equilibrium dia-
gram Fe-C applies also to
such multi-component sys-
tems. Figure 2.5 shows a
schematic cut through the
three phase system
Fe-M-C.
During precipitation, mixed
carbides of the general
composition M3C develop.
In contrast to the binary
system Fe-C, is the three
Figure 2.5
2. TTT – Diagrams 14
phase system Fe-M-C characterised by a temperature interval in the three-phase field α + γ +
M3C. The beginning of the transformation of α + M3C to γ is marked by Aclb, the end by Acle.
The indices b and e mean
the beginning and the end
of transformation.
The described equilibrium
diagrams apply only to low
heating and cooling rates.
However, higher heating
and cooling rates are pre-
sent during welding, con-
sequently other structure
types develop in the heat
affected zone (HAZ) and in
the weld metal. The struc-
ture transformations during
heating and cooling are described by transformation diagrams, where a temperature change
is not carried out close to the equilibrium, but
at different heating and/or cooling rates.
A representation of the transformation
processes during isothermal austenitizing
shows Figure 2.6. This figure must be read
exclusively along the time axis! It can be
recognised that several transformations
during isothermal austenitizing occur with e.g.
800°C. Inhomogeneous austenite means
both, low carbon containing austenite is
formed in areas, where ferrite was present
before transformation, and carbon-rich
austenite is formed in areas during
transformation, where carbon was present
before transformation. During sufficiently long
annealing times, the concentration differences
are balanced by diffusion, the border to a ho-
Figure 2.6
Figure 2.7
2. TTT – Diagrams 15
mogeneous austenite is passed. A growing of the austenite grain size (to ASTM and/or in
µm) can here simultaneously be observed with longer annealing times.
The influence of heating rate on austeniti zing is shown in Figure 2.7. This diagram must only
be read along the sloping lines of the same heating rate. For better readability, a time pattern
was added to the pattern of the heating curves. To elucidate the grain coarsening during aus-
tenitizing, two microstructure photographs are shown, both with different grain size classes to
ASTM.
Figure 2.8 shows the rela-
tion between the TTA and
the Fe-C diagram. It's obvi-
ous that the Fe-C diagram
is only valid for infinite long
dwell times and that the
TTA diagram applies only
for one individual alloy.
Figure 2.9 shows the dif-
ferent time-temperature
passes during austeniti zing
and subsequent cooling
down.
The heating period is com-
posed of a continuous and
an isothermal section.
During cooling down, two
different ways of heat con-
trol can be distinguished:
1. : During continuous
temperature control a
cooling is carried out with a
constant cooling rate out of
Figure 2.8
Figure 2.9
2. TTT – Diagrams 16
the area of the homogeneous and stable austenite down to room temperature.
2. : During isothermal temperature control a quenching out of the area of the austenite is
carried out into the area of the metastable austenite (and/or into the area of martensite), fol-
lowed by an isothermal holding until all transformation processes are completed. After trans-
formation will be cooled down to room temperature.
Figure 2.10 shows the
time-temperature diagram
of a isothermal transforma-
tion of the mild steel Ck 45.
Read such diagrams only
along the time-axis! Below
the Ac1b line in this figure,
there is the area of the me-
tastable austenite, marked
with an A. The areas
marked with F, P, B, und M
represent areas where fer-
rite, perlite, Bainite and
martensite are formed. The
lines which limit the area to the left mark the beginning of the formation of the respective
structure. The lines which limit the area to the right mark the completion of the formation of
the respective structure. Because the ferrite formation is followed by the perlite formation, the
completion of the ferrite formation is not determined, but the start of the perlite formation.
Transformations to ferrite and perlite, which are diffusion controlled, take place with elevated
temperatures, as diffusion is easier. Such structures have a lower hardness and strength, but
an increased toughness.
Diffusion is impeded under lower temperature, resulting in formation of bainitic and marten-
sitic structures with hardness and strength values which are much higher than those of ferrite
and perlite. The proportion of the formed martensite does not depend on time. During
quenching to holding temperature, the corresponding share of martensite is spontanically
formed. The present rest austenite transforms to Bainite with sufficient holding time. The right
Figure 2.10
2. TTT – Diagrams 17
detail of the figure shows the present structure components after completed transformation
and the resulting hardness at room temperature.
Figure 2.11 depicts the graphic representation of the TTT diagram, which is more important
for welding techniques. This is the TTT diagram for continuous cooling of the steel Ck 15.
The diagram must be read along the drawn cooling passes. The lines, which are limiting the
individual areas, also depict the beginning and the end of the respective transformation.
Close to the cooling curves, the amount of the formed structure is indicated in per cent, at the
end of each curve, there is the hardness value of the structure at room temperature.
Figure 2.12 shows the TTT
diagram of an alloyed steel
containing approximately
the same content of carbon
as the steel Ck 15. Here
you can see that all trans-
formation processes are
strongly postponed in rela-
tion to the mild steel. A
completely martensitic
transformation is carried
out up to a cooling time of
about 1.5 seconds, com-
pared with 0.4 seconds of
Ck 15. In addition, the
completely diffusion con-
trolled transformation proc-
esses of the perlite area
are postponed to clearly
longer times.
The hypereutectoid steel C
100 behaves completely
different, Figure 2.13. With
this carbon content, a pre-
Figure 2.11
Figure 2.12
2. TTT – Diagrams 18
eutectoid ferrite formation cannot still be car-
ried out (see also Figure 2.3).
The term of the figures 2.9 to 2.11 "austeniti z-
ing temperature“ means the temperature,
where the workpiece transforms to an austen-
itic microstructure in the course of a heat
treatment. Don’t mix up this temperature with
the AC3 temperature, where above it there is
only pure austenite. In addition you can see
that only martensite is formed from the aus-
tenite, provided that the cooling rate is suffi-
ciently high, a formation of any other
microstructure is completely depressed. With
this type of transformation, the steel gains the
highest hardness and strength, but loses its
toughness, it embrittles. The slowest cooling
rate where such a transformation happens, is
called critical cooling rate. Figure 2.13
Figure 2.14 Figure 2.15
2. TTT – Diagrams 19
Figure 2.14 shows schematically how the TTT diagram is modified by the chemical compo-
sition of the steel.
The influence of an increased austenitizing temperature on transformation behaviour shows
Figure 2.15. Due to the higher hardening temperature, the grain size of the austenite is
higher (see Figure 2.6 and 2.7).
This grain growth leads to
an extension of the diffu-
sion lengths which must be
passed during the trans-
formation. As a result, the
"noses" in the TTT diagram
are shifted to longer times.
The lower part of the figure
shows the proportion of
formed martensite and
Bainite depending on cool-
ing time. You can see that
with higher austenitizing
temperature the start of
Bainite formation together
with the drop of the mart-
ensite proportion is clearly
shifted to longer times.
As Bainite formation is not
so much impeded by the
coarse austenite grain as
with the completely diffu-
sion controlled processes
of ferrite and perlite forma-
tion, the maximum Bainite
proportion is increased
from about 45 to 75%.
Figure 2.16
Figure 2.17
2. TTT – Diagrams 20
Due to the strong influence of the austenitizing temperature to the transformation behaviour
of steel, the welding technique uses special diagrams, the so called Welding-TTT-diagrams.
They are recorded following the welding temperature cycle with both, higher austenitizing
temperatures (basically between 950° and 1350°C) and shorter austenitizing times.
You find two examples in Figures 2.16 and 2.17.
Figure 2.18 proves that the
iron-carbon diagram was
developed as an equilib-
rium diagram for infinite
long cooling time and that
a TTT diagram applies al-
ways only for one alloy.
Figure 2.18
3.
Residual Stresses
3. Residual Stresses 22
The emergence of residual
stresses can be of very
different nature, see three
examples in Figure 3.1.
Figure 3.2 details the
causes of origin. In a pro-
duced workpiece, material-
, production-, and wear-
caused residual stresses
are overlaying in such a
way that a certain condition
of residual stresses is cre-
ated. Such a workpiece
shows in service more or
less residual stresses, and it will never be stress-free!
Figure 3.3 defines residual stresses of 1., 2., and 3. type. This grading is independent from
the origin of the residual stresses. It is rather based on the three-dimensional extension of the
stress conditions.
Based on this definition, Fig-
ure 3.4 shows a typical distri-
bution of residual stresses.
Residual stresses, which
build-up around dislocations
and other lattice imperfections
(s III), superimpose within a
grain causing stresses of the
2nd type and if spreading
around several grains, bring
out residual stresses of the 1st
type.
The formation of residual
stresses in a transition-free
Figure 3.1
Figure 3.2
3. Residual Stresses 23
steel cylinder is shown in Figures 3.5. and 3.6. During water quenching of the homogeneous
heated cylinder, the edge of the cylinder cools down faster than the core. Not before 100
seconds have elapsed is the temperature across the cylinder's cross section again
homogeneous. The left part of
Figure 3.5 shows the T-t-
curve of three different meas-
urement points in the cylinder.
Figure 3.6 shows the results
of quenching on the stress
condition in the cylinder. At
the beginning of cooling, the
cylinder edge starts shrinking
faster than the core (upper
figure). Through the stabilising
effect of the cylinder core,
Figure 3.3 Figure 3.4
Figure 3.5
3. Residual Stresses 24
tensile stress builds up at the edge areas while the core is exposed to pressure stress. Re-
sulting volume differences between core and edge are balanced by elastic and plastic defor-
mations. When cooling is completed, edge and core are on the same temperature level, the
plastically stretched edge now supports the unstressed core, so that pressurestresses are
present in the edge areas and tensile residual stresses in the core.
These changes are principally shown once again in Figure 3.7 with the 3-rod model. A warm-
ing of the middle rod causes at first an elastic expansion of the outer rods, the inner rod is
exposed to pressure stress (line A-B). Along the line B-C the rod is plastically deformed, be-
cause pressure stresses have exceeded the yielding point. At point C, the cooling of the rod
starts, it is exposed to tensile stress due to shrinking. Along the line D-E the rod is plastically
deformed due to the influence of the counter members beeing in tension. At the point E the
system has cooled down to its initial temperature. This point represents the remaining resi-
dual stress condition of this construction. If heating is stopped before point C is reached and
cooled down to the initial temperature, then stress increase in the centre rod will be in parallel
Figure 3.6 Figure 3.7
3. Residual Stresses 25
with the elastic areas. Starting with point B, the same residual stress condition is present as
in a case of heating up to a temperature above 600°C.
Figure 3.8 divides the development of residual stresses in welded seams in three different
mechanisms.
Shrinking stresses: these are stresses formed through uniform cooling of the seam.
Caused by expansion restriction of the colder areas at the edge of the weld and base mate-
rial , tensile stresses develop along and crosswise to the seam.
Quenching stresses: If cooling is not homogenous, the surface of the weld cools down
faster than the core areas. If the high-temperature limit of elasticity is exceeded due to build -
up stress differences, pressure stresses will be present at the weld surface after cooling. In
contrast, the core shows tensile stresses in cold condition (see also Figure 3.6).
Transition stresses: Transitions in the ferrite and perlite stage cause normally only residual
stresses, because within this temperature range the yield strength of the steel is so low that
generated stresses can be undone by plastic deformations.
This is not the case with transitions in the Bainite and martensite stage. A transition of the
austenite causes an increase in volume (transition cfc in cbc, the cfc lattice has a higher den-
sity, additional volume increase through la t-
tice deformation). In the case of a homoge-
nous transition, the weld will consequently
unfold pressure stresses. If the transition of
the edge areas happens earlier than the tran-
sition of the slower cooling core, plastic de-
formations of the core area may be present
similar to quenching (see above: quenching
stresses). In this case, the weld surface will
show tensile stresses after cooling.
Generally these mechanisms cannot be
separated accurately from each other, thus
the residual stress condition of a weld will
represent an overlap of the cases as shown
in the 3rd figure. This overlap of the different
mechanisms makes a forecast of the remain-
ing residual stress condition difficult.
Figure 3.8
3. Residual Stresses 26
Figure 3.9 shows the building-up of residual
stresses crosswise to a welded seam in anal-
ogy to the 3-rod model of Figure 3.7. This fig-
ure considers only shrinking residual stresses.
Before application of welding heat, the seam
area is stress-free (cut A-A). At the weldpool
the highest temperature of the welding cycle
can be found (cut B-B), metal is liquid. At this
point, there are no residual stresses, because
molten metal cannot transmit forces at the
weldpool. Areas close to the joint expand
through welding heat but are supported by
areas which are not so close to the seam.
Thus, areas close to the joint show compres-
sion stress, areas away from the joint tensile
stress. In cut C-C the already solidified weld
metal starts to shrink and is supported by
areas close to the seam, the weld metal
shows tensile stresses, the adjacent areas
compression stresses. In cut D-D is the tem-
perature completely balanced, a residual
stress condition is recognised as shown in
the lower right figure.
Figure 3.10 shows how much residual
stresses are influenced by constraining ef-
fects of adjacent material. The resulting
stress in the presented case is calculated
according to Hooke:
s = e ? E
Elongation e is calculated as ? l/a (? l is the
length change due to shrinking). With con-
Figure 3.9
Figure 3.10
3. Residual Stresses 27
stant joint volume will shrinking and ? l always have the same value. Thus the elongation e
depends only on the value a. The smaller the a is chosen, the higher are the resulting
stresses.
Effects of transition on cooling can be estimated from Figure 3.11. Here curves of tempera-
ture- and length-changes of ferritic and austenitic steels are drawn. It is clear that a ferritic
lattice has a higher volume than an austenitic lattice at the same temperature.
A steel which transforms from austenite to one of the ferrite types increases its volume at the
critical point. This sudden rise in vo lume can be up to 3% in the case of martensite formation.
To record the effects of this behaviour on the stress condition of the weld, sample welds are
carried out in the test device outlined in Figure 3.12. Thermo couples measure the T-t – curve
at the weld seam, a force sensor records the force which tries to bend the samples.
The lower picture shows the results of such a test.
The temperature behaviour at the fusionline as well as the force necessary to hold the sam-
ple over the time is plotted.
Figure 3.11 Figure 3.12
3. Residual Stresses 28
In the temperature range above 600°C the force sensor registers a tensile force which is
caused by the shrinking of the austenite. Between 600 and 400°C a large drop in force can
be seen, which is caused by the transition of the austenite. The repeated increase of the
force is based on further shrinking of the ferrite.
With the help of TTT diagrams
of base material and welding
consumable, the transition
temperatures and/or tempera-
ture areas for the individual
zones of the welded joint can
be determined. With these
data and with the course of
temperature it can be clearly
determined in which part of
the curve the force drop is
caused by the transition of the
welding consumable and in
which part by transition in the heat affected
zone (HAZ).
These results can be used to determine the
longitudinal residual stresses transversal to
the joint, as shown in Figure 3.13. During
welding of austenitic transition-free materials
only tensile residual stresses are caused in
the welded area according to Figure 3.8. If an
austenitic electrode is welded to a StE 70,
transitions occur in the area of the heat af-
fected zone which lead to a decrease of ten-
sile stresses. If a high-strength electrode
which has a martensitic transition, is welded
to a StE 70, then there will be pressure resid-
ual stresses in the weld metal and tensile re-
sidual stresses in the HAZ.
Figure 3.13
Figure 3.14
3. Residual Stresses 29
If parts to be welded are not fixed, the shrinking of the weld will cause an angular distortion of
the workpieces, Figure 3.14 . If the workpieces can shrink unrestricted in this way, the re-
maining residual stresses will be much lower than in case with firm clamping.
Methods to determine resid-
ual stresses can be divided
into destructive, non-
destructive, and condition-
ally destructive methods.
The borehole and ring core
method can be considered
as conditionally destructive,
Figures 3.15 and 3.16.
In both cases, present re-
sidual stresses are released
through partial material re-
moval and the resulting de-
formations are then
measured by wire strain gauges. An essential advantage of the borehole method is the very
small material removal, the diameter of the borehole is only 1 to 5 mm, the bore depth is 1- to
2-times the borehole diameter.
The disadvantage here is that only surface elongations can be measured, thus the results are
limited residual stresses in the surface area of the workpiece.
With the ring core method,
a crown milling cutter is
used to mill a ring groove
around a three-axes wire
strain gauge. The core is
released from the force
effects and stress-relieved.
At the time when the resil-
ience of the core is meas-
ured, the detection of the
residual stress distribution
Figure 3.15
Figure 3.16
3. Residual Stresses 30
across the depth is also possible.
Both methods are limited in their suitability for measuring welding residual stresses, because
steep strain gradients in the HAZ may cause wrong measurements.
The table in Figure 3.17
shows a survey of meas-
urement methods for re-
sidual stresses and what
causes residual stresses
to be picked-up when us-
ing one of the respective
methods.
Figure 3.18 shows a sur-
vey of the completely de-
structive procedures of
residual stress recognition.
Figure 3.17
Figure 3.18
4.
Heat Treatment and
its Function During Welding
4. Heat Treatment and its Function During Welding 32
When welding a workpiece, not only the weld
itself, but also the surrounding base material
(HAZ) is influenced by the supplied heat
quantity. The temperature-field, which ap-
pears around the weld when different welding
procedures are used, is shown in Figure 4.1.
Figure 4.2 shows the influence of the material
properties on the welding process. The de-
termining factors on the process presented in
this Figure, like melting temperature and -
interval, heat capacity, heat extension etc,
depend greatly on the chemical composition
of the material. Metallurgical properties are
here characterized by e.g. homogeneity,
structure and texture, physical properties like
heat extension, shear strength, ductility.
Structural changes, caused by the heat input
(process 1, 2, 7, and 8), influence directly the mechanical properties of the weld. In addition,
the chemical composition of the weld metal and adjacent base material are also influenced
by the processes 3 to 6.
Based on the binary system,
the formation of the different
structure zones is shown in
Figure 4.3. So the coarse
grain zone occurs in areas of
intensely elevated
austenitising temperature for
example. At the same time,
hardness peaks appear in
these areas because of
greatly reduced critical cooling
rate and the coarse austenite
Figure 4.1
Figure 4.2
4. Heat Treatment and its Function During Welding 33
grains. This zone of the weld is the area,
where the worst toughness values are found.
In Figure 4.4 you can see how much the for-
mation of the individual structure zones and
the zones of unfavourable mechanical prop-
erties can be influenced.
Applying an electroslag one pass weld of a
200 mm thick plate, a HAZ of approximately
30 mm width is achieved. Using a three pass
technique, the HAZ is reduced to only 8 mm.
With the use of different procedures, the
differences in the formation of heat affected
zones become even clearer as shown in
Figure 4.5.
These effects can actively be used to the ad-
vantage of the material, for example to adjust
calculated mechanical properties to one's choice or to remove negative effects of a welding.
Particularly with high-strength fine grained steels and high-alloyed materials, which are spe-
cifically optimised to achieve special quality, e.g. corrosion resistance against a certain at-
tacking medium, this post-weld heat treatment is of great importance.
Figure 4.6 shows areas in
the Fe-C diagram of differ-
ent heat treatment meth-
ods. It is clearly visible that
the carbon content (and
also the content of other
alloying elements) has a
distinct influence on the
level of annealing tempera-
tures like e.g. coarse-grain
Figure 4.3
Figure 4.4
4. Heat Treatment and its Function During Welding 34
heat treatment or normalising.
It can also be seen that the start of martensite formation (MS-line) is shifted to continuously
decreasing temperatures with increasing C-content. This is important e.g. fo r hardening
processes (to be explained later).
As this diagram does not
cover the time influence,
only constant stop-
temperatures can be read,
predictions about heating-up
and cooling-down rates are
not possible. Thus the indi-
vidual heat treatment meth-
ods will be explained by
their temperature-time-
behaviour in the following.
Figure 4.5 Figure 4.6
Figure 4.7
4. Heat Treatment and its Function During Welding 35
Figure 4.7 shows in the detail to the right a T-t course of coarse grain heat treatment of an
alloy containing 0,4 % C. A coarse grain heat treatment is applied to create a grain size as
large as possible to improve machining properties. In the case of welding, a coarse grain is
unwelcome, although unavoidable as a consequence of the welding cycle. You can learn
from Figure 4.7 that there are two methods of coarse grain heat treatment. The first way is to
austenite at a temperature close above A3 for a couple of hours followed by a slow cooling
process. The second method is very important to the welding process. Here a coarse grain is
formed at a temperature far above A3 with relatively short periods.
Figure 4.8 shows sche-
matically time-temperature
behaviour in a TTT-
diagram. (Note: the curves
explain running structure
mechanisms, they must not
be used as reading off ex-
amples. To determine t8/5,
hardness values, or micro-
structure distribution, are
TTT-diagrams always read
continuously or isother-
mally. Mixed types like
curves 3 to 6 are not a llowed for this purpose!).
The most important heat treatment methods can be divided into sections of annealing, hard-
ening and tempering, and these single processes can be used individually or combined. The
normalising process is shown in Figure 4.9. It is used to achieve a homogeneous ferrite -
perlite structure. For this purpose, the steel is heat treated approximately 30°C above Ac3
until homogeneous austenite evolves. This condition is the starting point for the following
hardening and/or quenching and tempering treatment. In the case of hypereutectoid steels,
austenisation takes place above the A1 temperature. Heating-up should be fast to keep the
austenite grain as fine as possible (see TTA-diagram, chapter 2). Then air cooling follows,
leading normally to a transformation in the ferrite condition (see Figure 4.8, line 1; formation
of ferrite and perlite, normalised micro-structure).
Figure 4.8
4. Heat Treatment and its Function During Welding 36
To harden a material, aus-
tenisation and homogeni-
sation is carried out also at
30°C above AC3. Also in
this case one must watch
that the austenite grains
remain as small as possi-
ble. To ensure a complete
transformation to marten-
site, a subsequent quench-
ing follows until the
temperature is far below
the Ms-temperature, Figure
4.10. The cooling rate dur-
ing quenching must be high enough to cool down from the austenite zone directly into the
martensite zone without any further phase transitions (curve 2 in Figure 4.8). Such quenching
processes build-up very high thermal stresses which may destroy the workpiece during hard-
ening. Thus there are variations of this process, where perlite formation is suppressed, but
due to a smaller temperature gradient thermal stresses remain on an uncritical level (curves
3 and 4 in Figure 4.8). This
can be achieved in practice
–for example- through stop-
ping a water quenching
process at a certain tem-
perature and continuing the
cooling with a milder cooling
medium (oil). With longer
holding on at elevated tem-
perature level, transforma-
tions can also be carried
through in the bainite area
(curves 5 and 6).
Figure 4.9
Figure 4.10
4. Heat Treatment and its Function During Welding 37
Figure 4.11 shows the quenching and tempering procedure. A hardening is followed by an-
other heat treatment below Ac1. During this tempering process, a break down of martensite
takes place. Ferrite and cementite are formed. As this change causes a very fine micro-
structure, this heat treat-
ment leads to very good
mechanical properties like
e.g. strength and tough-
ness.
Figure 4.12 shows the pro-
cedure of soft-annealing.
Here we aim to adjust a
soft and suitable micro-
structure for machining.
Such a structure is charac-
terised by mostly globular
formed cementite particles, while the lamellar structure of the perlite is resolved (in Figure
4.12 marked by the circles, to the left: before, to the right: after soft-annealing). For hypoeu-
tectic steels, this spheroidizing of cementite is achieved by a heat treatment close below A1.
With these steels, a part of the cementite bonded carbon dissolves during heat treating close
below A1, the remaining cementite lamellas transform with time into balls, and the bigger
ones grow at the expense of
the smaller ones (a transfor-
mation is carried out because
the surface area is strongly
reduced ? thermodynami-
cally more favourable condi-
tion). Hypereutectic steels
have in addition to the lamel-
lar structure of the perlite a
cementite network on the
grain boundaries.
Figure 4.11
Figure 4.12
4. Heat Treatment and its Function During Welding 38
Spheroidizing of cementite is achieved by making use of the transformation processes during
oscillating around A1. When exceeding A1 a transformation of ferrite to austenite takes place
with a simultaneous solution of a certain amount of carbon according to the binary system Fe
C. When the temperature drops below A1 again and is kept about 20°C below until the trans-
formation is completed, a
re-precipitation of cemen-
tite on existing nuclei takes
place. The repetition of this
process leads to a step-
wise spheroidizing of
cementite and the frequent
transformation avoids a
grain coarsening. A soft-
annealed microstructure
represents frequently the
delivery condition of a ma-
terial.
Figure 4.13 shows the principle of a stress-relieve heat treatment. This heat treatment is
used to eliminate dislocations which were caused by welding, deforming, transformation etc.
to improve the toughness of a workpiece. Stress-relieving works only if present dislocations
are able to move, i.e. plastic structure deformations must be executable in the micro-range. A
temperature increase is
the commonly used
method to make such de-
formations possible be-
cause the yield strength
limit decreases with in-
creasing temperature. A
stress-relieve heat treat-
ment should not cause any
other change to properties,
so that tempering steels
Figure 4.13
Figure 4.14
4. Heat Treatment and its Function During Welding 39
are heat treated below tempering temperature.
Figure 4.14 shows a survey of heat treatments which are important to welding as well as their
purposes.
Figure 4.15 shows princi-
pally the heat treatments in
connection with welding.
Heat treatment processes
are divided into: before,
during, and after welding.
Normally a stress-relieving
or normalizing heat treat-
ment is applied before
welding to adjust a proper
material condition which for
welding. After welding, al-
most any possible heat treatment can be carried
out. This is only limited by workpiece dimen-
sions/shapes or arising costs. The most important
section of the diagram is the kind of heat treatment
which accom-panies the welding. The most impor-
tant processes are explained in the following.
Figure 4.16 represents the influence of different
accompanying heat treatments during welding,
given within a TTT-diagram. The fastest cooling is
achieved with welding without preheating, with
addition of a small share of bainite, mainly mart-
ensite is formed (curve 1, analogous to Figure 4.8,
hardening). A simple heating before welding with-
out additional stopping time lowers the cooling rate
according to curve 2. The proportion of martensite
is reduced in the forming structure, as well as the
Figure 4.15
Figure 4.16
4. Heat Treatment and its Function During Welding 40
level of hardening. If the material is hold at a temperature above MS during welding (curve 3),
then the martensite formation will be completely suppressed (see Figure 4.8, curve 4 and 5).
To explain the temperature-time-behaviours
used in the following, Figure 4.17 shows a
superposition of all individual influences on
the materials as well as the resulting T-T-
course in the HAZ. As an example, welding
with simple preheating is selected.
The plate is preheated in a period tV. After
removal of the heat source, the cooling of the
workpiece starts. When tS is reached, welding
starts, and its temperature peak overlays the
cooling curve of the base material. When the
welding is completed, cooling period tA starts.
The full line represents the resulting tempera-
ture-time-behaviour of the HAZ.
The temperature time course during welding
with simple preheating is shown in Figure
4.18. During a welding time
tS a drop of the working
temperature TA occurs. A
further air cooling is usually
carried out, however, the
cooling rate can also be
reduced by covering with
heat insulating materials.
Another variant of welding
with preheating is welding
at constant working
temperature. This is
Figure 4.17
Figure 4.18
4. Heat Treatment and its Function During Welding 41
achieved through further
warming during welding to
avoid a drop of the working
temperature. In Figure 4.19
is this case (dashed line,
TA needs not to be above
MS) as well as the special
case of isothermal welding
illustrated. During isother-
mal welding, the workpiece
is heated up to a working
temperature above MS
(start of martensite forma-
tion) and is also held there
after welding until a transformation of the austenitised areas has been completed. The aim of
isothermal welding is to cool down in accordance with curve 3 in Figure 4.16 and in this way,
to suppress martensite formation.
Figure 4.20 shows the T-T course during
welding with post-warming (subsequent heat
treatment, see Figure 4.15). Such a treatment
can be carried out very easy, a gas welding
torch is normally used for a local preheating.
In this way, the toughness properties of some
steels can be greatly improved. The lower
sketch shows a combination of pre- and post-
heat treatment. Such a treatment is applied to
steels which have such a strong tendency to
hardening that a cracking in spite of a simple
preheating before welding cannot be avoided,
if they cool down directly from working tem-
perature. Such materials are heat treated
immediately after welding at a temperature
between 600 and 700°C, so that a formation
Figure 4.19
Figure 4.20
4. Heat Treatment and its Function During Welding 42
of martensite is avoided and welding residual stresses are eliminated simultaneously.
Aims of the modified step-
hardening welding should
not be discussed here, Fig-
ure 4.21. Such treatments
are used for transformation-
inert materials. The aim of
the figure is to show how
complicated a heat treatment
can become for a material in
combination with welding.
Figure 4.22 shows tempera-
ture distribution during multi-
pass welding. The solid line
represents the T-T course of a point in the HAZ
in the first pass. The root pass was welded
without preheating. Subsequent passes were
welded without cooling down to a certain tem-
perature. As a result, working temperature in-
creases with the number of passes. The
second pass is welded under a preheat tem-
perature which is already above martensite
start temperature. The heat which remains in
the workpiece preheats the upper layers of the
weld, the root pass is post-heat treated through
the same effect. During welding of the last
pass, the preheat temperature has reached
such a high level that the critical cooling rate
will not be surpassed. A favourable effect of
multi-pass welding is the warming of the HAZ
of each previous pass above recrystallisation
temperature with the corresponding crystallisa-
Figure 4.21
Figure 4.22
4. Heat Treatment and its Function During Welding 43
tion effects in the HAZ. The coarse grain zone with its unfavourable mechanical properties is
only present in the HAZ of the last layer. To achieve optimum mechanical values, welding is
not carried out to Figure 4.22. As a rule, the same welding conditions should be applied for all
passes and prescribed t8/5 – times must be kept, welding of the next pass will not be carried
out before the previous pass has cooled down to a certain temperature (keeping the inter-
pass temperature). In addition, the workpiece will not heat up to excessively high tempera-
tures.
Figure 4.23 shows a nomogram where working temperature and minimum and maximum
heat input for some steels can be interpreted, depending on carbon equivalent and wall thick-
ness.
If e.g. the water quenched and tempered fine grain structural steel S690QL of 40 mm wall
thickness is welded, the following data can be found:
- minimum heat input between 5.5 and 6 kJ/cm
- maximum heat input about 22 kJ/cm
- preheating to about 160°C
- after welding, residual stress relieving between 530 and 600°C.
Steels which are placed in
the hatched area called
soaking area, must be
treated with a hydrogen re-
lieve annealing. Above this
area, a stress relieve anneal-
ing must be carried out. Be-
low this area, a post-weld
heat treatment is not re-
quired.
Figure 4.23
5.
Welding Plain and
Low Alloy Steels
5. Welding Plain and Low Alloy Steels 45
© ISF 200 4b r-er05-01.cdr
Einteilung n ach der chem ischen Zusam mensetzung� unlegie rte Stähle
� nichtro stende Stähle
Einteilung n ach Hauptgüt eklassen� unlegi erte Stähle - unlegier te Qualitätsstä hle - unlegier te Edelstähle � nichtro stende Stähle
Definitio n des Begriff es Stahl
� andere legierte Stäh le
� andere legierte Stäh le - legierte Qualitätsstähl e
- legierte Edelstähle In the European Standard DIN EN
10020 (July 2000), the designations
(main symbols) for the classification of
steels are standardised. Figure 5.1
shows the definition of the term „steel“
and the classification of the steel
grades in accordance with their
chemical composition and the main
quality classes.
In accordance with the chemical compo-
sition the steel grades are classified into
unalloyed, stainless and other alloyed
steels. The mass fractions of the individ-
ual elements in unalloyed steels do not
achieve the limit values which are indi-
cated in Figure 5.2.
Stainless steels are grades of steel with
a mass fraction of chromium of at least
10,5 % and a maximum of 1,2 % of car-
bon.
Other alloyed steels are steel grades
which do not comply with the definition of
stainless steels and where one alloying
element exceeds the limit value indicated
in Figure 5.2.
Figure 5.1
Definition for theclassification of steels
© ISF 2004br-er05-01.cdr
Classification in accordance with the chemical composition:
l
l
l
unalloyed steels
stainless steels
other, alloyed steels
Classification in accordance with the main quality class:
·
·
·
unalloyed steels - unalloyed quality steels- unalloyed special steels
stainless steels
other, alloyed steels - alloyed quality steels- alloyed special steels
Definition of the term “steel”
Steel is a material with a mass fraction if iron which is higherthan of every other element, ist carbon content is, in general,lower than 2% and steel contains, moreover, also otherelements. A limited number of chromium steels might contain acarbon content which is higher than 2%, but, however, 2% is thecommon boundary between steel and cast iron [DIN EN 10020(07.00)].
Figure 5.2
Boundary between unalloyedand alloyed steels
© ISF 2004br-er05-02.cdr
Determined elementlimit value
Mass fraction in %
a) If just the highest value has been determined for
mangenese, the limit value us 1,80% and the 70%-rule
does not apply.
Al aluminium
B boron
Bi bismuth
Co cobalt
Cr chromium
Cu copper
La lanthanides
(rated individually)
Mn manganese
Mo molybdenum
Nb niobium
Ni nickel
Pb lead
Se selenium
Si silicon
Te tellurium
Ti titanium
V vanadium
W tungsten
Zr zirconium
Others (with the exception
of carbon, phosphorus,
sulphur, nitrogen)
(Each)
5. Welding Plain and Low Alloy Steels 46
As far as the main quality classes are concerned, the steels are classified in accor-
dance with their main characteristics and main application properties into unalloyed,
stainless and other alloyed steels.
As regards unalloyed steels a distinction is made between unalloyed quality steels
and unalloyed high-grade steels.
Regarding unalloyed quality steels, prevailing demands apply, for example, to the
toughness, the grain size and/or the forming properties.
Unalloyed high-grade steels are characterised by a higher degree of purity than
unalloyed quality steels, particularly with regard to non-metal inclusions. A more
precise setting of the chemical composition and special diligence during the manufac-
turing and monitoring process guarantee better properties. In most cases these
steels are intended for tempering and surface hardening.
Stainless steels have a chromium mass fraction of at least 10,5 % and maximally
1,2 % of carbon. They are further classified in accordance with the nickel content and
the main characteristics: corrosion resistance, heat resistance and creep resistance.
Other alloyed steels are classified into alloyed quality steels and alloyed high-grade
steels.
Special demands are put on the alloyed quality steels, as, for example, to toughness,
grain size and/or forming properties. Those steels are generally not intended for
tempering or surface hardening.
The alloyed high-grade steels comprise steel grades which have improved properties
through precise setting of their chemical composition and also through special manu-
facturing and control conditions.
5. Welding Plain and Low Alloy Steels 47
The European Standard DIN EN 10027-1 (September 1992) stipulates the rules for
the designation of the steels by means of code letters and identification numbers.
The code letters and identification numbers give information about the main applica-
tion field, about the mechanical or physical properties or about the composition.
The code designations of the steels are divided into two groups. The code designa-
tions of the first group refer to the application and to the mechanical or physical
properties of the steels. The code designations of the second group refer to the
chemical composition of the steels.
According to the utilization of the
steel and also to the mechanical or
physical properties, the steel grades
of the first group are designated with
different main symbols (Fig. 5.3).
Figure 5.3
Classification of steels in accordancewith their designated use
© ISF 2004br-er05-03.cdr
l
l
l
l
l
l
l
l
l
l
l
e.g. S235JR, S355J0
P =e.g. P265GH, P355M
L =e.g. L360A, L360QB
E =e.g. E295, E360
B =e.g. B500A, B500B
Y =e.g. Y1770C, Y1230H
R =e.g. R350GHT
H =
e.g. H400LA
D =e.g. DD14, DC04
T =
e.g. TH550, TS550
M =e.g. M400-50A, M660-50D
S = Steels for structural steel engineering
Steels for pressure vessel construction
Steels for pipeline construction
Engineering steels
Reinforcing steels
Prestressing steels
Steels for rails (or formed as rails)
Cold rolled flat-rolled steels with higher-strengthdrawing quality
Flat products made of soft steels for cold reforming
Black plate and tin plate and strips and also speciallychromium-plated plate and strip
Magnetic steel sheet and strip
5. Welding Plain and Low Alloy Steels 48
An example of the code designation structure with reference to the usage and the
mechanical or physical properties for “steels in structural steel engineering“ is ex-
plained in Figure 5.4.
Figure 5.4
5. Welding Plain and Low Alloy Steels 49
For designating special features of the steel or the steel product, additional symbols
are added to the code designation. A distinction is made between symbols for spe-
cial demands, symbols for the type of coating and symbols for the treatment con-
dition. These additional symbols are stipulated in the ECISS-note IC 10 and depicted
in Figures 5.5 and 5.6.
© ISF 2004br-er-05-06.cdr
1
2
))The symbols are separated from the preceding symbols by plus-signs (+)In order to avoid mix-ups with other symbols, the figure T may precede,for example +TA
Symbol ) )
+ A+ AC+ C
+ Cnnn+ CR+ HC+ LC+ Q+ S+ ST+ U
1 2 treatment condition
softenedannealed for the production of globular carbideswork-hardened (e.g., by rolling and drawing), also a distinguishingmark for cold-rolled narrow strips)cold-rolled to a minimum tensile strength of nnn MPa/mm²cold-rolledthermoformed/cold formedslightly cold-drawn or slightly rerolled (skin passed)quenched or hardenedtreatment for capacity for cold shearingsolution annealeduntreated
Symbols for the treatment condition
Figure 5.6
© ISF 2004br-er-05-05.cdr
Symbol ) )
+ A+ AR+ AS+ AZ+ CE+ Cu+ IC+ OC+ S+ SE+ T+ TE+ Z+ ZA+ ZE+ ZF+ ZN
1 2 Coating
hot dippedaluminium, cladded by rollingcoated with Al-Si alloycoated with Al-Tn alloy (>50% Al)electrolytically chromium-platedcopper-coatedinorganically coatedorganically coatedhot-galvanised
upgraded by hot dipping with a lead-tin alloyelectrolytically coated with a lead-tin alloyhot-galvisedcoated with Al-Zn alloy (>50% Zn)electrolytically galvaniseddiffusion-annealed zinc coatings (galvannealed, with diffused Fe)nickel-zinc coating (electrolytically)
electrolytically galvanised
1
2
))The symbols are separated from the preceding symbols by plus-signs (+)In order to avoid mix-ups with other symbols, the figure S may precede,for example +SA
Symbols for the coating type
Figure 5.5
5. Welding Plain and Low Alloy Steels 50
Figure 5.7 shows an example of the novel designation of a steel for structural steel
engineering which had formerly been labelled St37-2.
Figure 5.8 depicts the chemical composition and the mechanical parameters of dif-
ferent steel grades. The figure explains the influence of the chemical composition on
the mechanical properties.
Figure 5.7
© ISF 2002br-er-05-07.cdr
S = steels for structural steel engineeringP = steels for pressure vessel constructionL = steels for pipeline constructionE = engineering steelsB = reinforcing steels
The steel St37-2 (DIN 17100) is, according to the new standard (DIN EN 10027-1),designated as follows:
S235 J 2 G3
Steel for structural steel engineering
R 235 MPa/mmeH
2³
further property(RR = normalised)
test temperature 20°C
impact energy ³ 27 J
Steel designation in accordance with DIN EN 10027-1
Stahl C Si Mn P S Cr Al Cu N Mo Ni Nb VS355J0(St 52-3)S500N(StE500)P295NH(HIV)S355J2G1W(WTSt510-3)S355G3S(EH36)
Stahl
S355J2G3(St 52-3)S500N(StE500)P295NH(HIV)S355J2G1W(WTSt510-3)S355G3S(EH36)
Kerbschlagarbeit AV
[J]
Zugfestigkeit Rm
[N/mm²]BruchdehnungA
[%]
StreckgrenzeReH
[N/mm²]0°C -20°C
27
610-780 500 16 31-47
27355510-680 20-22
285
355
355
>18
22
>22
49(bei +20°C)
76(bei -10°C)
21-39
460-550
510-610
400-490
£0,18
£0,55
£0,35
£0,1-0,35
£0,50
0,1- 0,6
£0,26
£0,15
0,21
£0,20 £1,60 0,040
1- 1,7 0,035
³0,6 £0,05
0,5- 1,3 0,035
0,7- 1,5 £0,05
0,040 /
0,030 0,30
£0,05 /
0,0350,40-0,80
£0,05 /
/ /
0,020 0,20
/ /
/0,25-0,5
/ /
£0,009 /
0,020 0,1
/ /
/ £0,30
/
/ /
1 0,05
/ /
£0,65 /
/
0,22
/
0,02-0,12
// / /
Chemical composition and mechanicalparameters of different steel sorts
© ISF 2004br-er-05-08.cdr
impact energy AVelongation after fracture Ayield point ReHTensile strength RmSteel
Steel
Figure 5.8
5. Welding Plain and Low Alloy Steels 51
The steel S355J2G2 represents the basic type of structural steels which are nowa-
days commonly used. Apart from a slightly increased Si content for desoxidisation it
this an unalloyed steel.
S500N is a typical fine-grained structural steel. A very fine-grained microstructure
with improved tensile strength values is provided by the addition of carbide forming
elements like Cr and Mo as well as by grain-refining elements like Nb and V.
The boiler steel P295NH is a heat-resistant steel which is applied up to a temperature
of 400°C. This steel shows a relatively low strength but very good toughness values
which are caused by the increased Mn content of 0,6%.
S355J2G1W is a weather-resistant structural steel with mechanical properties similar
to S355J2G2. By adding Cr, Cu and Ni, formed oxide layers stick firmly to the work-
piece surface. This oxide layer prevents further corrosion of the steel.
S355G3S belongs to the group of shipbuilding steels with properties similar to those
of usual structural steels. Due to special quality requirements of the classification
companies (in this case: impact energy) these steels are summarised under a special
group.
5. Welding Plain and Low Alloy Steels 52
The steel grades are classified into four subgroups according to the chemical com-
position (Fig 5.9):
● Unalloyed steels (except free-cutting steels) with a Mn content of < 1 %
● Unalloyed steels with a medium Mn content > 1 %, unalloyed free-cutting
steels and alloyed steels (except high-speed steels) with individual alloying
element contents of less than 5 percent in weight
● Alloyed steels (except high-speed steels), if, at least for one alloying element
the content is ≥ 5 percent in weight
● High-speed steels
The unalloyed steels with Mn con-
tents of < 1% are labelled with the
code letter C and a number which
complies with the hundredfold of the
mean value which is stipulated for the
carbon content.
Unalloyed steels with a medium Mn
content > 1 % are labelled with a
number which also complies with a
hundredfold of the mean value which
is stipulated for the carbon content, the
chemical symbols for the alloying
elements, ordered according to the
decreasing contents of the alloying
elements and numbers, which in the
sequence of the designating alloying
elements give reference about their
content. The individual numbers stand
for the medium content of the respective alloying element, the content had been
multiplied by the factor as indicated in Fig. 5.9/Table 5.1 and rounded up to the next
whole number.
Codes accordingto the chemical composition
© ISF 2004br-er05-09.cdr
Unalloyed steels (Mo content < 1%)
Unalloyed steels (Mn content > 1%)
Alloyed steels (content of alloying element > 5%)
X10CrNi18-10
Legiert C=10/100=0,1% Cr=18% Ni=10%
10CrMo9-10
C=10/100=0,10% Cr=9/4=2,25% Mo=10/10=1%
C45
Carbon 0,45% Carbon
element factor
Cr, Co, Mn, Ni, Si, W
Al, Be, Cu, Mo, Nb, Pb, Ta, Ti, V, Zr
C, Ce, N, P, S
B
4
10
100
1000
High-speed steels
HS 2-9-1-8
Mo=9% Co=8%W=2% V=1%
Table 5.1
Figure 5.9
5. Welding Plain and Low Alloy Steels 53
The alloyed steels are labelled with the code letter X, a number which again com-
plies with the hundredfold of the mean value of the range stipulated for the carbon
content, the chemical symbols of the alloying elements, ordered according to de-
creasing contents of the elements and numbers which in sequence of the designating
alloying elements refer to their content.
High-speed steels are designated with the code letter HS and numbers which, in the
following sequence, indicate the contents of elements:: tungsten (W), molybdenum
(Mo), vanadium (V) and cobalt (Co).
The European Standard DIN EN 10027-2 (September 1992) specifies a numbering
system for the designation of steel grades, which is also called material number
system..
The structure of the material number is as follows:
1. XX XX (XX)
Sequential number The digits inside the brackets are intended for possible future demands.
Steel group number (see Fig. 5.10)
Material main group number (1=steel)
5. Welding Plain and Low Alloy Steels 54
Figure 5.10 specifies the material numbers for the material main group „steel“.
Figure 5.10
5. Welding Plain and Low Alloy Steels 55
The influence of the austenite grain size on the transformation behaviour has been
explained in Chapter 2. Figure 5.11 shows the dependence between grain size of the
austenite which develops during the welding cycle, the distance from the fusion line
and the energy-per-unit length from the welding method. The higher the energy-per-
until length, the
bigger the austen-
ite grains in the
HAZ and the width
of the HAZ in-
creases. Such
coarsened austen-
ite grain decreases
the critical cooling
time, thus increas-
ing the tendency of
the steel to harden.
With fine-grained structural steels it is tried to suppress the grain growth with alloying
elements. Favourable are nitride and carbide forming alloys. They develop precipita-
tions which suppress undesired grain growth. There is, however, a limitation due to
the solubility of these precipitations, starting with a certain temperature, as shown in
Figure 5.12. Steel 1 does not contain any precipitations and shows therefore a con-
tinuous grain growth related to temperature. Steel 2 contains AIN precipitations which
are stable up to a temperature of approx. 1100°C, thus preventing a growth of the
austenite grain.
Influence of the energy-per-unitlength on the austenite grain size
13
11
9
7
5
30 0,2 0,4 0,6 0,8 1,0
Au
ste
nite
gra
in s
ize
ind
ex
acc
ord
ing
to D
IN 5
06
01
Distance of the fusion linemm
Energy-per-unit length in kJ/cm
9 12 18 36
© ISF 2004br-er-05-11.cdr
Figure 5.11
5. Welding Plain and Low Alloy Steels 56
With higher temperatures, these
precipitations dissolve and cannot
suppress a grain growth any more.
Steel 3 contains mainly titanium car-
bonitrides of a much lower grain-
refining effect than that of AIN. Steel 4
is a combination of the most effective
properties of steels nos. 2 and 3.
The importance of grain refinement
for the mechanical properties of a
steel is shown in Figure 5.13. Pro-
vided the temperature keeps con-
stant, the yield strength of a steel
increases with decreasing grain size.
This influence on the yield point Rel is
specified in the Hall-Petch-law:
dKR
iel
1⋅+= σ
According to the
above-mentioned
law, the increase of
the yield point is
inversely propor-
tional to the root of
the medium grain
diameter d. σi
stands for the inter-
nal friction stress of
the material. The
grain boundary
resistance K is a
measure for the
influence of the grain size on the forming mechanisms. Apart from this increase of the
yield point, grain refinement also results in improved toughness values. As far as
Austenite grain size as a functionof the austenitization temperature
Steel % C % Mn % Al % N % Ti
1 0,21 1,16 0,004 0,010 /
2 0,17 1,35 0,047 0,017 /
3 0,18 1,43 0,004 0,024 0,067
4 0,19 1,34 0,060 0,018 0,140
900 1000 1100 1200 1300 1400°C
Austenitization temperature
18
6
4
2
10-1
8
6
4
2
10-2
6 10-3
8
mm
Mediu
m fib
re le
ngth
Gra
in s
ize in
dex
acc
ord
ing to D
IN 5
0601
-4
-2
0
2
4
6
8
10
12
Steel 1Steel 2Steel 3Steel 4
© ISF 2004br-er05-12.cdr
Figure 5.12
Connection betweenyield point and grain size
900
800
700
600
500
400
300
200
N/mm²
10 2 3 4 5 6 7 8 10mm-1/2
Yie
ld p
oin
t or
0,2
boundary
Grain size d-1/2
Temperature in °C:
-193
-185
-180
-155
+20
-40
-100
-170
© ISF 2004br-er-05-13.cdr
Figure 5.13
5. Welding Plain and Low Alloy Steels 57
structural steels are concerned, this means the improvement of the mechanical prop-
erties without any further alloying. Modern fine-grained structural steels show im-
proved mechanical properties with, at the same time, decreased content of alloying
elements. As a consequence of this chemical composition the carbon equivalent
decreases, the weldability is improved and processing of the steel is easier.
The major advan-
tages of microal-
loyed fine-grained
structural steels in
comparison with
conventional struc-
tural steels are
shown in Figure
5.14. Due to the
considerably better
mechanical proper-
ties of the fine-
grained structural
steel in comparison
with unalloyed structural steel, substantial savings of material are possible. This
leads also to reduced joint cross-sections and, in total, to lower costs when making
welded steel constructions.
Based on the
classification of
Figure 5.2, Fig-
ure 5.15 divides the
steels with regard
to their problematic
processes during
welding. When it
comes to unalloyed
steels, only ingot
Figure 5.14
Influence of the steel selection on theproducing costs of welded structures
S235JR S355J2G3 S690Q S890Q S960Q Verhältnis
(St37-2) (St52-3) (StE690) (StE890) (StE960) S235JR - S960Q
Streckgrenze N/mm2215 345 690 890 960 1 : 5
Blechdicke mm 50 31 14,4 11 10 5 : 1
Nahtquerschnitt mm2870 370 100 60 50 17 : 1
Schweißdraht ø 1.2 mm SG2 SG3 NiMoCr X 90 X 96 -
Schweißdrahtkosten Verhältnis 1 1 2,4 3,2 3,3 1 : 3,3
Stahlkosten Verhältnis 1 1,2 1,9 2,3 2,4 1 : 2,4
Schweißgutkosten Verhältnis 5,3 2,3 1,5 1,16 1 5,3 : 1
Spez. Schweißnahtkosten Verhältnis 12 5,1 1,8 1,18 1 12 : 1
Kostenverhältnis inklusiveGrundwerkstoffe
5 : 1
Randbedingungen: Schweißverfahren = MAG
Abschmelzleistung = 3 kg Schweißdraht / h, Nahtform X - 60°
Lohn- und Maschinenkosten = 60 DM / h
Spez. Schweißnahtkosten = Schweißzusatzwerkstoffe + Schweißen
Berechnungsgrundlage =szul = Re / 1.5
Stahlsorte
© ISF 2004br-er-05-14.cdr
Yield point
Plate thickness
Weld cross-section
Welding wire Ø 1.2
Welding wire costs
Steel costs
Weld metal costs
Special weld costs
Costs ratio inclusive basematerials
Ratio
Ratio
Ratio
Ratio
Boundary condition: welding process = MAG
Deposition rate = 3 kg welding wire/h, weld shape X -60°
Costs of labour and equipment = 30€/h
Special weld costs = weld filler materials + welding
Calculation base = = Re/1.5szul
Steel type Ratio
Figure 5.15
Classification of steels withrespect to problems during welding
low-alloyed high-alloyed
hardeningspecial properties areachieved, for example:
heat resistance,tempering resistant,
high-pressure hydrogen resistance,toughness at low temperatures,
surface treeatment condition, etc.
corrosionresistant steels
tool steels
Hardening,special properties
are achieved
steels
unalloyed alloyed
mild steel higher-carbon steel
HardeningUnderbead cracking
rimmed steel killed steel duplex killed steel
cutting ofsegregation
zones
cold brittleness(coarse-grained recrystallization
after critical treatment)stress corrosion crackingsafety from brittle fracture
ferritic pearlitic-martensitic austenitic
grain desintegrationstress corrosion
cracking hot cracks(sigma phaseembrittlement)
hardeningembrittlement
formationof chromium
carbide
grain increase inthe weld interfaces
Post-weld treatment forhighest corrosion resistance
© ISF 2004br-er-05-15.cdr
5. Welding Plain and Low Alloy Steels 58
casts, rimmed and semi-killed steels are causing problems. “Killing” means the re-
moval of oxygen from the steel bath.
Figure 5.16 shows cross-sections of ingot blocks with different oxygen contents.
Rimming steels with increased oxygen content show, from the outside to the inside,
three different zones after solidification: 1.: a pronounced, very pure outer envelope,
2.: a typical blowhole formation (not critical, blowholes are forged together during
rolling), 3.: in the
centre a clearly
segregated zone
where unfavourable
elements like sul-
phur and phospho-
rus are enriched.
During rolling, such
zones are stretched
along the complete
length of the rolling
profile.
Figure 5.17 shows important points to be observed during welding such steels. Due
to their enrichment with alloy elements, the segregation zones are more transforma-
tion-inert than the
outer envelope
and are inclined to
hardening. In
addition, they are
sensitive to hot-
cracking, as, in
these zones, the
elements phospho-
rus and sulphur
are enriched.
Figure 5.16
Ingot cross-sectionsafter different casting methods
Figures: mass content of oxygen in %
fully killed steel semi-killed steel rimmed steel
0,003
0,012
0,025
© ISF 2004br-er-05-16.cdr
Figure 5.17
Example of unfavourable (a) andfavourable (b) welds
a b
B CD
E
© ISF 2004br-er-05-17.cdr
5. Welding Plain and Low Alloy Steels 59
Therefore, “ touching” such segregation zones during welding must be avoided by all
means.
In the case of low-
alloy steels, the
problem of HAZ
hardening during
welding must be
observed. Fig-
ure 5.18 shows
hardness values of
various microstruc-
tures. The highest
hardness values
can be found with
martensite and
cementite. Hardness values of cementite are of minor importance for unalloyed and
low-alloy steels because its proportion in these steels remains low due to the low C-
content.
However, hardening because of martensite formation is of greatest importance as the
martensite proportion in the microstructure depends mainly on the cooling time.
Figure 5.19 shows
the essential influ-
ence of the mart-
ensite content in
the HAZ on the
crack formation of
welded joints.
Hardening through
martensite forma-
tion is not to be
expected with pure
carbon steels up to
about 0,22%,
Hardness of Several Microstructures
Microstructures Average Brinell Hardness (Approximately)
Ferrite 80
Austenite 250
Perlite (granular) 200
Perlite (lamellar) 300
Sorbite 350
Troostite 400
Cementite 600 - 650
Martensite 400 - 900
© ISF 2004Br-er-05-18.cdr
Figure 5.18
Influence of Martensite Content
strength,calculated at
max. hardness
with maximummartensite
contentHV HRC N/mm2 %
root crackingpresumable
400 41 1290 70
root crackingpossible
400 - 350 41 - 36 1290 - 1125 70 - 60
no root cracking 350 36 1125 60
sufficient operational safetywithout heat treatment
280 28 900 30
maximum hardness
If too much martensite develops in the heat affected zone during welding (below or next to the weld),a very hard zone will be formed which shows often cracks.
© ISF 2004Br-er-05-19.cdr
Figure 5.19
5. Welding Plain and Low Alloy Steels 60
because the critical cooling rate with these low C-contents is so high that it normally
won’t be reached within the welding cycle. In general, such steels can be welded
without special problems (e.g., S. 235).
In addition to car-
bon, all other alloy
elements are im-
portant when it
comes to marten-
site formation in
the welding cycle,
as they have sub-
stantial influence
on the transforma-
tion behaviour of
steels (see
Fig. 2.12 ). It is not
appropriate just
to take the carbon content as a measure for the hardening tendency of such steels.
To estimate the weldability, several authors developed formulas for calculating the
so-called carbon equivalent, which include the contribution of the other alloy ele-
ments to hardening tendency, (Fig. 5.20). As these approximation formulas are em-
pirically determined
and as for the
hardening tendency
the general condi-
tions like plate
thickness, heat
input, etc., are also
of importance, the
carbon equivalent
cannot be a com-
mon limit value for
the weldability.
For the determina-
Figure 5.20
Definition of C - Equivalent
C-Äqu.= carbon equivalent (%) PLS = pipeline steels PCM = (%)cracking parameters
IIW
Stout
Ito and Bessyo
Mannesmann
Hoesch
Thyssen
15
NiCu
5
VMoCr
6
MnCÄqu.C
++
++++=-
40
Cu
20
Ni
10
MnCr
6
MnCÄqu.C ++
+++=-
5B10
V
15
Mo
60
Ni
20
CrCuMn
30
SiCPCM ++++
++++=
40
Ni
20
CuCr
10
MoMnCCET +
++
++=
20
VMoNiCrCuMnSiCÄqu.C
+++++++=-
15
V
40
Mo
60
Ni
20
Cr
16
CuMn
25
SiCÄqu.C PLS ++++
+++=-
© ISF 2002Br-er-05-20.cdr
Mo
Figure 5.21
Quelle: DIN EN 1011-2br-er05-21.cdr
Calculation of the preheating temperatures
Tp =697 CET + 160 tanh (d/35) + 62 HD + (53 CET - 32) Q - 3280,35
-100
-80
-60
-40
-20
0
20
40
0 0,5 1 1,5 2 2,5 3 3,5 4 4,5 5
Wärmeeinbringen Q [kJ/mm]
delt
aT
p[°
C]
delta Tp = (53 CET - 32) Q - 53 CET + 32
d = 50 mmHD = 8
CET = 0,4 % CET = 0,2 % CET = 0,2 %
delta Tp = (53 CET - 32) Q - 53 CET + 32
CET = 0,4 % CET = 0,2 % CET = 0,2 %
d = 50 mmHD = 8
0
20
40
60
80
100
0 5 10 15 20 25
Wasserstoffgehalt HD des Schweißgutes [%]
de
lta
Tp
[°C
]
delta Tp = 62 HD 0,35 - 100
CET = 0,33 %d = 30 mmQ = 1 kJ/mm
delta Tp = 62 HD - 1000,35
CET = 0,33 %dQ = 1 kJ/mm
= 30 mm
0
50
100
150
200
250
0,2 0,3 0,4 0,5
Kohlenstoffäquivalent CET [%]
Tp
[°C
]
Tp = 750 CET - 150
d = 30 mmHD = 4Q = 1 kJ/mm
Tp = 750 CET - 150
d = 30 mmHD = 4Q = 1 kJ/mm
0
10
20
30
40
50
60
0 20 40 60 80 100
Blechdicke d [mm]
de
lta
Tp
[°C
]
delta Tp = 160 tanh (d/35) - 110
CET = 0,4 %HD 2Q = 1 kJ/mm
delta Tp = 160 tanh (d/35) - 110
CET = 0,4 %HD = 2Q = 1 kJ/mm
© ISF 2005
Heat input
Hydrogen content of the weld metalCarbon aquivalent
Plate thickness
Source:
5. Welding Plain and Low Alloy Steels 61
tion of the preheating temperature Tp, the formula as shown in Fig. 5.21 is used. The
effects of the chemical composition which is marked by the carbon equivalent CET,
the plate thickness d, the hydrogen content of the weld metal HD and the heat in-
put Q are considered.
The essential factor
to martensite forma-
tion in the welding
cycle is the cooling
time. As a measure
of cooling time, the
time of cooling from
800 to 500°C (t8/5) is
defined (Fig. 5.22).
The temperature
range was selected
in such a way that it
covered the most
important structural transformations and that the time can be easily transferred to the
TTT diagrams.
Figure 5.23 shows
measured time-
temperature distri-
butions in the vicin-
ity of a weld. Peak
values and dwell
times depend obvi-
ously on the loca-
tion of the
measurement and
are clearly strongly
determined by the
heat conduction
conditions.
Figure 5.22
Definition of t8/5
Te
mp
era
ture
T
Time t
Tmax
°C
800
500
t t s800 500
DT
t8/5
© ISF 2004br-er-05-22.cdr
Figure 5.23
Temperature-time curvesin the adjacence of a weld
2000
°C
1500
1000
500
00 50 100 150 200 250 s 300
Time t
Tem
pera
ture
T
A
B
C
10mm
A
B
C
© ISF 2004br-er-05-23.cdr
5. Welding Plain and Low Alloy Steels 62
With the use of thinner plates with complete heating of the cross-section during weld-
ing, the heat conductivity is only carried out in parallel to the plate surface, this is the
two-dimensional heat dissipation.
With thicker plates, e.g. during welding of a blind bead, heat dissipation can also be
carried out in direction of plate thickness, heat dissipation is three-dimensional.
These two cases
are covered by the
formulas given in
Figure 5.24, which
provide a method
of calculating the
cooling time t8/5 of
low-alloyed steels.
In the case of a
three-dimensional
heat dissipation,
t8/5 it independent
of plate thickness.
In the case of two-dimensional heat dissipation it is clear that t8/5 becomes the shorter
the thicker the plate thickness d is. Provided, the cooling times are equal, the plate
thickness can be calculated from these relations where a two-dimensional heat dissi-
pation changes to a three-dimensional heat dissipation.
Figure 5.25 shows
the influence of the
welding method on
the heat dissipa-
tion. With the same
heat input, the
energy which is
transferred to the
base material
depends on the
Figure 5.24
Calculation equation for two- andthree-dimensional heat dissipation
3 - dimensional:
2 - dimensional:
© ISF 2004br-er-05-24.cdr
÷÷ø
öççè
æ
--
-×
××
××=
00
5/8800
1
500
1
2 TTv
IUt
lph
( ) 3
00
04
5/8800
1
500
110567,0 N
TTv
IUTt ×¢×÷÷
ø
öççè
æ
--
-×
×××-= - h
úúû
ù
êêë
é÷÷ø
öççè
æ
--÷÷
ø
öççè
æ
-××÷
ø
öçè
æ ××
××××=
2
0
2
02
22
5/8800
1
500
11
4 TTdv
IU
ct
rlph
( ) 22
2
0
2
02
2
05
5/8800
1
500
11103,4043,0 N
TTdv
IUTt ×¢×
úúû
ù
êêë
é÷÷ø
öççè
æ
--÷÷
ø
öççè
æ
-××÷
ø
öçè
æ ×××-= - h
÷÷ø
öççè
æ
-+
-×
××¢×
×-×-
=-
-
0004
05
800
1
500
1
10567,0
103,4043,0
TTv
IU
T
Td
üh
K3
universal formula:
extended formulaFor low-alloyed steel:
universal formula:
extended formulaFor low-alloyed steel:
K2
formula for the transitionthickness of low-alloyed steel:
Figure 5.25
Relative thermal efficiency degreeof different welding methods
0 0,1 0,2 0,3 0,4 0,5 0,6 0,7 0,8 0,9 1
SA welding
Manual arc welding
MAG-(CO )-2 welding
MIG-(Ar)-welding
TIG-(Ar)-welding
TIG-(He)-welding
welding methods
Relative thermal efficiency degree ‘h
© ISF 2004Br-er-05-25.cdr
5. Welding Plain and Low Alloy Steels 63
welding method. This dependence is described by the relative thermal efficiency ŋ’.
The influence of
the groove ge-
ometry is covered
by seam factors
according to
Fig. 5.26. Empiri-
cally determined,
these factors were
introduced for an
easier calculation.
For other groove
geometries, tests
to measure the
cooling time are recommended.
Fig. 5.27 shows the transition of the two-dimensional to the three-dimensional heat
dissipation for two different preheating temperatures in form of a curve according to
the equation of Fig. 5.24. Above the curve, t8/5 depends only on the energy input, but
not on the plate thickness, heat dissipation is carried out three-dimensionally.
Figure 5.26
Weld factors for differentweld geometries
Type of weldweld factor
2-dimensionalheat dissipation
3-dimensionalheat dissipation
1
0,45 - 0,67
0,9
0,9
1
0,67
0,67
0,9
© ISF 2004br-er-05-26.cdr
Figure 5.27
Transition From Two to ThreeDimensional Heat Flow
Heat input E. .N [kJ/cm]h n
0 10 20 30 40 50
5
cm
3
2
1
0
Pla
te thic
kness
cooling time t [s]10 15 20 25
8/5
3040
60100
2-dimensional
3-dimensional
T =20°CA
0 10 20 30 40 50
cooling time t [s]10 20 30 40 50
8/5
2-dimensional
3-dimensional
T =200°CA
60
80
100
150
© ISF 2004Br-er-05-27.cdr
5. Welding Plain and Low Alloy Steels 64
Fig. 5.28 shows the
possible range of
heat input depend-
ing on the elec-
trode diameter. It is
clear that a rela-
tively large working
range is available
for arc welding
procedures. A
variation of the
energy-per-unit
length can be
carried out by alteration of the welding current, the welding voltage and the welding
speed.
Fig. 5.29 depicts variations of the heat
input during manual metal arc weld-
ing. The shorter the fused electrode
distance, i.e., the shorter the ex-
tracted length, the higher the energy-
per-unit length.
Figure 5.28
br-er-05-28.cdr
Heat Inputs ofVarious Welding Methods
3,25 4 5 6 0,8 1,0 1,2 1,6 2,5 3,0 4,0 5,0
20
kJ/cm
12
8
4
He
at
inp
ut
Manual metal arc welding MAGC-, MAGM-method
SA-welding
-short arc
-sprayarc
© ISF 2004
Figure 5.29
© ISF 2004br-er05-29.cdr
35
kJ/cm
25
20
15
10
5
0
Energ
y-per-
unit
length
0 50 100 150 200 250 300 350 400 450 500 mm 600
run-out length
Stick electrode(mm)
Current intensity (A)
Current intensity (A)
2,5
90
75
3,25
135
120
4,0
180
140
5,0
235
190
6,0
275
250
Æ6,0mm x 450mm
Æ5,0mm x 450mm
Æ4,0mm x 450mm
Æ3,25mm x 350mm
Æ2,5mm x 350mm
Energy-per-unit length as afunction of the run-out length
5. Welding Plain and Low Alloy Steels 65
In order to minimize calculation efforts in practice, the specified relations were
transferred into nomograms from which permissible welding parameters can be read
out, provided some additional data are available. Fig. 5.30 shows diagrams for two-
dimensional heat dissipation, where a dependence between energy-per-unit length,
cooling time and preheating temperature is given, depending on the plate thickness. .
If a fine-grained structural steel is to be welded, the steel manufacturer presets a
certain interval of cooling times, where the steel characteristics are not too negatively
affected. The user lays down the plate thickness and, through the selection of a
welding method, a specified range of heat input E. Based on the data E and t8/5 the
diagram provides the required preheating temperature for welding the respective
plate thickness.
Figure 5.30
br-er05-30.cdr
Dependence of E, t andd During SA - Welding
8/5
Heat input E5 6 7 8 9 10 15 20 30 kJ/cm 50
504030
20
10
7
504030
20
10
7
504030
20
10
7
504030
20
10
7
Coolin
g tim
e t
in s
8/5
d = 7,5 mm
d = 10 mm
d = 15 mm
d = 20 mm
transition to3-dimensional
heat flow
T 200°C150°C100°C
20°C
0
T 200°C150°C100°C
20°C
0
T 200°C150°C100°C
20°C
0
T 200°C150°C100°C
20°C
0
© ISF 2004
5. Welding Plain and Low Alloy Steels 66
With the transition to thicker plates,
the diagrams in Fig. 5.31 apply. The
upper part of the figure determines
whether a two-dimensional or a three-
dimensional heat dissipation is pre-
sent. For the three-dimensional heat
dissipation, the lower diagram applies
where the same information can be
determined, independent of plate
thickness, as with Fig. 5.30.
The relation be-
tween current and
voltage for MAG
welding is shown
in Fig. 5.32 and
the used shielding
gas is one of the
parameters. Weld-
ing voltage and
welding current, or
wire feed speed,
determine the type
of arc.
Figure 5.31
br-er05-31.cdr
Dependence ofE, T , t And d0 8/5 Ü
Heat input E
50s
40
30
20
15
109
87
Co
olin
g t
ime
t 8
/5
5 6 7 8 9 10 15 20 30 kJ/cm 50
T250
°C
200°C
150°C
100°C
20°C
0
Heat input E
50mm
40
30
20
15
109
87
Tra
nsi
tion
th
ickn
ess
dÜ
5 6 7 8 9 10 15 20 30 kJ/cm 50
aera of3-dimensional
heat flow
area of2-dimensional
heat flow
T250 °C 200 °C
150 °C 100 °C
20 °C
0
© ISF 2004
Figure 5.32
br-er-05-32.cdr
Dependence of Current And Voltage DuringMAG-Welding, Solid Wire, 1.2 mmÆ
35V
30
25
20
15
Weld
ing v
olta
ge
Welding current
Wire feed
150 200 250 A 300
3,5 4,5 5,5 7,0 8,0 9,0 10,5 m/min
C1
M21
M23
gas composition:C1 100% COM21 82% Ar + 18% COM23 92% Ar + 8% O
2
2
2
short arc
contact tube distance ~15mm contact tube distance ~19mm
mixed arc spray arc
© ISF 2004
5. Welding Plain and Low Alloy Steels 67
The diagram in Fig. 5.33 demon-
strates the dependence of plate thick-
ness, heat input E and cooling time
t8/5 for fillet welds at a preheating
temperature of T0 = 150°C. If d and
t8/5 are given, the acceptable range of
heat input can be determined with the
help of this diagram. The kinks of the
curves mark the transition between
two-dimensional and three-
dimensional heat dissipation.
Fig. 5.34 shows the same depend-
ence for butt welds with V groove
preparation..
Figure 5.33
br-er05-33.cdr
Permissible E-RangeDuring SA - And MAG - Welding
hh
' = 1' = 0,85
d = 32 mmd = 15 mm
UP
MAG
U max
U min
F = 0,67F = 0,67
3
2
t = 30 st = 6 s8/5 max
8/5 min
E = 66 kJ/cmE = 14 kJ/cm
max
min
60
kJ/cm
50
45
40
35
30
25
20
15
10
5
0
70
kJ/cm
59
53
47
41
35
29
23
18
12
6
0
He
at
inp
ut
ES
A-
we
ldin
g
He
at
inp
ut
EM
AG
- w
eld
ind
Plate thickness
0 5 10 15 20 25 30 mm 40
cracking tendency
toughness affection
fillet weldsT = 150 °C0 30s
25s
20s
15s
10s
6s
© ISF 2004
Figure 5.34
br-er05-34.cdr
Permissible E-RangeDuring SA - And MAG - Welding
hh
' = 1' = 0,85
d = 34 mmd = 15 mm
UP
MAG
U max
U min
F = 0,9F = 0,9
3
2
t = 30 st = 6 s8/5 max
8/5 min
E = 49 kJ/cmE = 10 kJ/cm
max
min
60
kJ/cm
50
45
40
35
30
25
20
15
10
5
0
70
kJ/cm
59
53
47
41
35
29
23
18
12
6
0
Heat
inp
ut
ES
A-
weld
ing
Heat
inp
ut
EM
AG
- w
eld
ing
Plate thickness
0 5 10 15 20 25 30 mm 40
cracking tendency
toughness affection
butt weldsT = 150 °C0
30s
25s
20s
15s
10s
6s
© ISF 2004
5. Welding Plain and Low Alloy Steels 68
The curve family in Fig. 5.35 shows the dependence of the heat input from the weld-
ing speed as well as the acceptable working range. The parameters of the curves 1
to 8 in the table
have been taken
from Figures 5.32
and 5.34 and apply
only for related
conditions like wire
diameter, wire
feed, welding
voltage, etc.
Figure 5.36 shows
a reading example
for such diagrams
(according to DVS-
Reference Sheet
Nr. 0916).
In this example, a
plate thickness of
15 mm and a cool-
ing time t8/5 be-
tween 10 and 20 s
are given. In this
case, the maximum
cooling time for MAG welding is 15 s. A solid wire with a diameter of 1.2 mm at 29V
and 300A is used.
The left diagram provides heat input values between 13 and 16 kJ/cm, based on the
given data. Using these values, the acceptable range of welding speeds can be
taken from the diagram on the right.
Figure 5.35
br-er-05-35.cdr
E as a Function of Welding Speed,Solid Wire, 1.2mmÆ
MAG/ M21 (82% Ar, 18% CO)
25kJ/cm
20
15
10
5
010 15 20 25 30 35 40 45 50 cm/min 60
Welding speed vS
Heat in
put E
working range
12
34
56
7
8
curve
V
A
v (m/min)Z
29
300
10.5
27
275
9.0
24
250
8.0
22
225
7.0
20
200
5.5
19
175
4.5
18
150
3.5
17
125
3.0
1 2 3 4 5 6 7 8
© ISF 2004
Figure 5.36
br-er-05-36.cdr
Determination of Welding Speedfor MAG - Welding
curve
V
A
v (m/min)Z
29
300
10.5
27
275
9.0
24
250
8.0
22
225
7.0
20
200
5.5
19
175
4.5
18
150
3.5
17
125
3.0
1 2 3 4 5 6 7 860
kJ/cm
50
45
40
35
30
25
20
15
10
5
0
70
kJ/cm
59
53
47
41
35
29
23
18
12
6
0
He
at
inp
ut
E
SA
- w
eld
ing
He
at
inp
ut
E
MA
G -
we
ldin
g
Plate thickness0 5 10 15 20 25 30 mm 40
cracking tendency
toughness affection
butt weldsT = 150 °C0
30s
25s
20s
15s
10s
6s
30s
25s
20s
15s
10s
6s
1613
25kJ/cm
20
15
10
5
010 15 20 25 30 35 40 45 50 cm/min 60
Welding speed vS
he
at
inp
ut
E working range
12
34
56
7
8
16
13
33 41
© ISF 2004
5. Welding Plain and Low Alloy Steels 69
Fig. 5.37 presents a simplification of
the determination of the microstruc-
tural composition and cooling time
subject to peak temperatures which
occur in the welding cycle. In the
lower diagram, the point of the plate
thickness at the top line is linked with
the point of heat input at the lower
line. The point of intersection of the
linking line with the middle scale
represents the cooling time t8/5 .
If the peak temperature of the welding
cycle is known, one can read from the
middle diagram in which transition
field the final microstructures are
formed. The advantage of the deter-
mination of microstructures compared
with the upper TTT diagram is that
a TTT diagram applies only for exactly one peak temperature, other peak tempera-
tures are disregarded. The disadvantage of the PTCT diagram (peak temperature
cooling time diagram) is the very expensive determination, therefore, due to the
measurement efforts a systematic application of this concept to all common steel
types is subject to failure.
Figure 5.37
© ISF 2004
Peak temperature/cooling time– diagram for the determination
of t and the structure8/5
bie5-37.cdr
1400
°C
1200
1000
800
600
800
°C
700
600
500
400
300
200
1 10 100 1000Te
mpera
ture
Peak
tem
pera
ture
B
M
M
Arc3
Arc1
B+M F+B
300 200HV30=400
F+P
F
P
s t8/5
40 30 25 20 15 10 9 8 7 6 5 mm 4plate thickness
300 100200three-dimensional
1 2 3 5 10 20 50 100 200 400 s 1000two-dimensional
0 100 °C 200preheating temperature
6 8 10 20 30 40 50 kJ/cm 70energy-per-unit length
t8/5
1000°C1400°C
Peak temperature
6.
Welding High Alloy Steels
6. Welding High Alloy Steels 71
Basically stainless steels are characterised by a chromium content of at least 12%. Figure
6.1 shows a classification
of corrosion resistant
steels. They can be sin-
gled out as heat- and
scale-resistant and
stainless steels, depend-
ing on service tempera-
ture. Stainless steels are
used at room temperature
conditions and for water-
based media, whilst heat-
and scale-resistant steels
are applied in elevated
temperatures and gaseous
media.
Depending on their microstructure, the alloys can be divided into perlitic-martensitic, ferritic,
and austenitic steels. Perlitic-martensitic steels have a high strength and a high wear resis-
tance, they are used e.g. as knife steels. Ferritic and corrosion resistant steels are mainly
used as plates for household appliances and other decorative purposes.
The most important group are austenitic steels, which can be used for very many applications
and which are corrosion resistant against most media. They have a very high low tempera-
ture impact resistance.
Based on the simple Fe-C
phase diagram (left figure),
Figure 6.2 shows the ef-
fects of two different
groups of alloying elements
on the equilibrium diagram.
Ferrite developers with
chromium as the most im-
portant element cause a
strong reduction of the aus-
Classification of Corrosion-Resistant Steels
non-stabilized
(austenite withdelta-ferrite)X12CrNi18-8
stabilized
(austenite withoutdelta-ferrite)
X8CrNiNb16-13
ferritic austenitic
stainlesssteels
scale- and heat-resistantsteels
corrosion-resistant steels
semi-ferritic ferritic-austenitic
X40Cr13 X10Cr13 X8Cr13 X20CrNiSi25-4
perliticmartensitic
© ISF 2002br-er-06-01e.cdr
Figure 6.1
Modifications to the Fe-C Diagramby Alloy Elements
ChromiumVanadiumMolybdenumAluminiumSilicon
NickelManganeseCobalt
Alloy elements in %Alloy elements in % Alloy elements in %
T
A4
A3
T
A4
A3
T
A4
A3
gg g
aa
a(d)
d d
© ISF 2002br-er-06-02e.cdr
Figure 6.2
6. Welding High Alloy Steels 72
tenite area, partly with downward equilibrium line according to Figure 6.2 (central figure).
With a certain content of the related element, there is a transformation-free, purely ferritic
steel.
An opposite effect provide austenite developers. In addition to carbon, the most typical mem-
ber of this group is nickel.
Austenite developers cause an extension of
the austenite area to Figure 6.2 (left figure)
and form a purely austenitic and transforma-
tion-free steel.
The table in Figure 6.3 summarises the ef-
fects of some selected elements on high alloy
steels.
The binary system Fe-Cr in Figure 6.4 shows
the influence of chromium on the iron lattice.
Starting with about 12% Cr, there is no more
transformation into the cubic face-centred
lattice, the steel solidifies purely as ferritic. In
the temperature range between 800 and
500°C this system contains the intermetallic
σ-phase, which decomposes in the lower
temperature range into a low-chromium α-solid solution and a chromium-rich α’-solid solution.
Both, the development of the σ-phase and of the unary α-α’-decomposition cause a strong
Effects of Some Elementsin Cr-Ni Steel
Element Steel type, no. Effect
Carbon
l
l
l
All types
l
l
l
Increases the strength, supports development
of precipitants which reduce corrosion
resistance, increasing C content reduces
critical cooling rate
Chromium
l
All types
l
Works as ferrite developer, increases
oxidation- and corrosion-resistance
Nickel
l
l
All typesWorks as austenite developer, increases
toughness at low temperature, grain-refining
Oxygen
lSpecial types l
Works as strong austenite developer
(20 to 30 times stronger than Nickel)
Niobium
l
1.4511,1.4550,
1.4580 u.a.
Binds carbon and decreases tendency to
intergranular corrosion
Manganese
l
l
All types
l
l
Increases austenite stabilization, reduces hot
crack tendency by formation of manganese
sulphide
Molybdenum
l
l
1.4401,1.4404,
1.4435 and others.
l
Improves creep- and corrosion-resistance
against reducing media, acts as ferrite
developer
Phosphorus,
selenium, or
sulphur l
1.4005, 1.4104,
1.4305
l
l
Improve machinability, lower weldability,
reduce slightly corrosion resistance
Silicon l
l
All types
l
l
Improves scale resistance, acts as ferrite
developer, all types are alloyed with small
contents for desoxidation
Titanium
l
l
1.4510, 1.4541,
1.4571 and others
l
Binds carbon, decreases tendency to
intergranular corrosion, acts as a grain refiner
and as ferrite developer
Aluminium
l
Type 17-7 PH
l
Works as strong ferrite developer, mainly
used as heat ageing additive
Copper
l
l
l
Type 17-7 PH,
1.4505, 1.4506
l
l
Improves corrosion resistance against certain
media, decreases tendency to stress
corrosion cracking, improves ageing
© ISF 2002br-er06-03e.cdr
Figure 6.3
Binary System Fe - Cr
Te
mp
era
ture
1800
0
200
400
600
800
1000
1200
1400
1600
°C
30Fe 10 20 40 50 60 70 80 90 Cr%
Chromium
S S+a
a
a'a
dd+a d+a'
g+ag
© ISF 2002br-er06-04e.cdr
Figure 6.4
6. Welding High Alloy Steels 73
embrittlement. With higher alloy steels, the diffusion speed is greatly reduced, therefore both
processes require a relatively long dwell time. In case of technical cooling, such embrittle-
ment processes are suppressed by an increased cooling speed.
Nickel is a strong austenite developer, see Figure 6.5 Nickel and iron develop in this system
under elevated temperature a complete series of face-centred cubic solid solutions. Also in
the binary system Fe-Ni
decomposition processes
in the lower temperature
range take place.
Along two cuts through the
ternary system Fe-Cr-Ni,
Figure 6.6 shows the most
important phases which
develop in high alloy steels.
A solidifying alloy with 20%
Cr and 10% Ni (left figure)
forms at first δ-ferrite. δ-
ferrite is, analogous to the
Fe-C diagram, the primary
from the melt solidifying
body-centred cubic solid
solution. However α-ferrite
is developed by transfor-
mation of the austenite, but
is of the same structure
from the crystallographic
point of view, see Figure
6.4.
Binary System Fe - Ni
30Fe 10 20 40 50 60 70 80 90 Ni
0
200
400
600
800
1000
1200
1400
1600
Te
mp
era
ture
°C
%
Nickel
Fe Ni3
Fe N
i 3
a a+g
g
dd+g
S+dS+g
© ISF 2002br-er-06-05e.cdr
Figure 6.5
Sections of the Ternary System Fe-Cr-Ni
700
800
900
1000
1100
1200
1300
1400
1500
1600
0 5 10 15 20 % Ni
% Cr30 25 20 15 10
70 % Fe
0 5 10 15 20 25700
800
900
1000
1100
1200
1300
1400
1500
1600
40 35 30 25 20 15
% Ni
% Cr
60 % Fe
Tem
pera
ture
°C °C SS
S+d+g
d+gd+g
d+g+s
dd
d+s
d+s
gg
g+sg+s
S+gS+gS+d
S+d
d+g+
s
S+d+g
© ISF 2002br-er-06-06e.cdr
Figure 6.6
6. Welding High Alloy Steels 74
During an ongoing cooling, the binary area ferrite + austenite passes through and a transfor-
mation into austenite takes place. If the cool-
ing is close to the equilibrium, a partial trans-
formation of austenite into the brittle α-phase
takes place in the temperature range below
800°C. Primary ferritic solidifying alloys show
a reduced tendency to hot cracking, because
δ-ferrite can absorb hot-crack promoting ele-
ments like S and P. However primary austen-
itic solidifying alloys show, starting at a certain
alloy content, no transformations during cool-
ing (14% Ni, 16% Cr, left figure). Primary aus-
tenitic solidifying alloys are much more
susceptible to hot cracking than primary fer-
ritic solidifying alloys, a transformation into the
σ-phase normally does not take place with
these alloys.
Figure 6.7 shows some typical compositions
of certain groups of high alloy steels.
The diagram of Strauß and Maurer in Figure 6.8 shows the influence on the microstructure
formation of steels with a C-content of 0,2%. The classification of high-alloy steels in Figure
6.1 is based on this dia-
gram. If a steel only con-
tains C, Cr and Ni, the
lowest austenite corner will
be at 18% Cr and 6% Ni.
And also other elements
than Ni and Cr work as an
austenite or ferrite devel-
oper. The influence of
these elements is de-
scribed by the so-called
chromium and nickel
Typical Alloy Content ofHigh-Alloy Steels
4.Aus
teni
tic-fe
rritic
ste
els
3.Aus
teni
tic s
teel
s
2. M
arte
nsitic
stee
ls
1. F
errit
ic s
teel
s
C
Si
Mn
Cr
Mo
Ni
Cu
Nb
Ti
Al
V
N
S
£0.1
£
.0 1
£0.1
0.11.2
max.1.0
max.1.0
max.1.0
max.1.0
max.1.0
max.1.5
max.2.0
max.2.0
1518
1218
1726
2428
up to2.0
up to1.2
up to5.0
up to2.0
£1.0
£2.5
726
47.5
up to2.2
+
+
+
+
+
+
+
+
+
+
+ indicates that the alloyelements can be added ina defined content to achievevarious characteristics
© ISF 2002br-er06-07e.cdr
Figure 6.7
Maurer - Diagram
0
4
8
12
16
20
24
28
%
0 2 4 6 8 10 12 14 16 18 20 22 24 26%
ferrite / perlite
martensite / troostite / sorbite
austenite / martensite
martensite / ferrite
austenite / martensite / ferrite
austenite / ferrite
austenite
Nic
kel
Chromium
© ISF 2002br-er-06-08e.cdr
Figure 6.8
6. Welding High Alloy Steels 75
equivalents. The Schaeffler diagram reflects additional alloy elements, Figure 6.9. It repre-
sents molten weld metal of high alloy steels and determines the developed microstructures
after cooling down from very high temperatures. The diagram was always prepared consider-
ing identical cooling conditions, the influence of different cooling speeds is here disregarded.
The areas 1 to 4 in this diagram limit the chemical compositions of steels, where specific de-
fects may occur during welding.
Depending on the composition, purely ferritic chromium steels have a tendency to embrittle-
ment by martensite and therefore to hot cracking (area 2) or to embrittlement due to strong
grain growth (area 1).
A cause for this strong grain
growth during welding is the
greatly increased diffusion
speed in the ferrite com-
pared with austenite. After
reaching a diffusion-start
temperature, Figure 6.10
shows that ferritic steels
have a considerably
stronger grain growth than
austenites. Therefore high
alloyed ferritic steels are to
be considered as of limited
weldability.
The area 3 marks a possible
embrittlement of the material
due to the development of
σ-phase. As explained in
6.6, this risk occurs with in-
creased ferrite contents,
increased chromium con-
tents, and sufficiently slow
cooling speed.
Schaeffler Diagram With Border Lines ofWeld Metal Properties to Bystram
0%Ferri
t
5%
10%
40%
80%
100%
20%
ferrite
martensite
M + FF+M
A+M+F
A + F
A +M
austenite
0 2 4 6 8 10 12 14 16 18 20 22 24 26 28 30 32 34 36 38 40
30
28
26
24
22
20
18
16
14
12
10
8
6
4
2
0
hardening crack susceptibility(preheating to 400°C!)
sigma embrittlementbetween 500-900°C
hot cracking susceptibility above 1250°C grain growth above 1150°C
Chromium-equivalent = %Cr + %Mo + 1,5x%Si + 0,5x%Nb
Nic
kel-equiv
ale
nt =
%N
i +
30x%
C +
0,5
x%
Mn
© ISF 2002br-er-06-09e.cdr
Figure 6.9
Grain Size as a Function of Temperature
0 200 400 600 800 1000 1200
1000
2000
3000
4000
5000
6000
gra
in s
ize
temperature
°C
m²
ferritic steel
austenitic steel
© ISF 2002br-er-06-10e.cdr
Figure 6.10
6. Welding High Alloy Steels 76
Finally, area 4 marks the strongly increased tendency to hot cracking in the austenite. Rea-
son is, that critical elements responsible for hot cracking like e.g. sulphur and phosphorous
have only very limited solubility in the austenite. During welding, they enrich the melt residue,
promoting hot crack formation (see also chapter 9 - Welding Defects).
There is a Z-shaped area in the centre of the diagram which does not belong to any other
endangered area. This area of chemical composition represents the minimum risk of welding
defects, therefore such a composition should be adjusted in the weld metal. Especially when
welding austenitic steels one tries to aim at a low content of δ-ferrite, because it has a much
greater solubility of S and P, thus minimising the risk of hot cracking.
The Schaeffler diagram is not only used for determining the microstructure with known
chemical composition. It is also possible to estimate the developing microstructures when
welding different materials with or without filler metal. Figures 6.11 and 6.12 show two exam-
ples for a determination of the weld metal microstructures of so-called 'black and white' joints.
Application Example ofSchaeffler - Diagram
0 4 8 12 16 20 24 28 32 360
4
8
12
16
20
24
28
Chromium-equivalent
Nic
ke
l-e
qu
iva
len
t
F
A
A+M
M
M+F
A+M+F
A+F
² ·: =1:1
²
·
10
20
40
80
100
%
F+
9
·
30%
Weld metal under 30 % dilution (= base metal amount)
² ·
·
9
S235JR (St 37)
Welding consumable
X8Cr17 (W.-Nr. 1.4510)
21% Cr, 14% Ni, 3% Mo
1
2
3
© ISF 2002br-er06-11e.cdr
Figure 6.11
·9
Application Example ofSchaeffler - Diagram
0 4 8 12 16 20 24 28 32 360
4
8
12
16
20
24
28
Chromium-equivalent
Nic
kel-equiv
ale
nt
F
A
A+M
M
M+F
A+M+F
A+F² ·: =1:1
²
10
20
40
80
100
%
F
+
20%
Weld metal under 30 % dilution (= base metal amount)
² ·
·
9
S235JR (St 37)
Welding consumable
X10CrNiTi18-9 (W.-No. 1.4541)
21% Cr, 14% Ni, 3% Mo
123
·
© ISF 2002br-er06-12e.cdr
Figure 6.12
6. Welding High Alloy Steels 77
The ferrite content can only be measured with a relatively large dispersal, therefore DeLong
proposed to base a measurement procedure on standardized specimens. Such a system
makes it possible to measure comparable values which don't have to match the real ferrite
content. Based on these measurement values, the ferrite content is no longer given in per-
centage, but steels are grouped by ferrite numbers. In addition to ferrite numbers, DeLong
proposed a reworked Schaeffler diagram where the ferrite number can be determined by the
chemical composition, Figure 6.13. Moreover, DeLong has considered the influence of nitro-
gen as a strong austenite developer (effects are comparable with influence of carbon). Later
on, nitrogen was included into the nickel-equivalent of the Schaeffler diagram.
The most important feature
of high alloy steels is their
corrosion resistance start-
ing with a Cr content of
12%. In addition to the
problems during welding
described by the Schaeffler
diagram, these steels can
be negatively affected with
view to their corrosion re-
sistance caused by the
welding process. Figure
6.14 shows schematically
the processes of electro-
lytic corrosion under a
drop of water on a piece of
iron. In such a system a
potential difference is a
precondition for the devel-
opment of a local element
consisting of an anode and
a cathode. To develop
De Long Diagram
16 17 18 19 20 21 22 23 24 25 26 27
Chromium-equivalent = %Cr + %Mo + 1,5 x %Si + 0,5 x %Nb
Nic
kel-e
qu
iva
len
t =
%N
i + 3
0 x
%C
+ 3
0 x
%N
+ 0
,5 x
%M
n
21
20
19
18
17
16
15
14
13
12
11
10
austenite
Schaeffler-austenite-martensite-line
austenite + ferrite
form
erly m
agnetically
measu
red
ferri
te c
ontents
in v
ol.-%
ferri
te n
umber
2%
4%
6%
7,6%
9,2%
10,7%
12,3%
13,8%
0%
0
2
4
68
10
12
14
16
18
© ISF 2002br-er-06-13e.cdr
Figure 6.13
Corrosion Under a Drop of Water
air
water
Fe(OH)3
iron
2Fe +O+H O 2Fe +2OH++ +++ -
2 ®
H O2
O
OH-
cathode
anode
2Fe 2Fe +4e®++ -
4e-
O +2H O+4e 4OH2 2
- -®
O2 OH
Fe+++
2Fe++
© ISF 2002br-er-06-14e.cdr
Figure 6.14
6. Welding High Alloy Steels 78
such a local element, a different orientation of grains in the steel is sufficient. If a potential
difference under a drop of water is present, the chemically less noble part reacts as an an-
ode, i.e. iron is oxidised here and is dissolved as Fe2+-ion together with an electron emission.
Caused by oxygen access through the air, a further oxidation to Fe3+ takes place. The ca-
thodic, chemically nobler area develops OH- ions, absorbing oxygen and the electrons. Fe3+-
and OH--ions compose into the water-insoluble Fe(OH)3 which deposits as rust on the sur-
face (note: the processes here described should serve as a principal explanation of electro-
chemical corrosion mechanisms, they are, at best, a fraction of all possible reactions).
If the steel is passivated by chromium, the corrosion protection is provided by the develop-
ment of a very thin chromium oxide layer which separates the material from the corrosive
medium. Mechanical surface damages of this layer are completely cured in a very short time.
The examples in Figure 6.15 are more critical, since a complete recovery of the passive layer
is not possible from various reasons.
passive layerpassive layer
passive layerpassive layer
activedissolution
pitting corrosion
tensile stress
active dissolutionof the crack base
active dissolution of the gap
crevice corrosion
activly dissolvedgrain boundary
chromium zones
grain boundarycarbides
depleted
intergranular corrosion
stress corrosion cracking
© ISF 2002br-er06-15e.cdr
Figure 6.15
gap
incorrect correct
© ISF 2002br-er06-16e.cdr
Figure 6.16
6. Welding High Alloy Steels 79
If crevice corrosion is pre-
sent, corrosion products built
up in the root of the gap and
oxygen has no access to
restore the passive layer.
Thus narrow gaps where the
corrosive medium can ac-
cumulate are to be avoided
by introducing a suitable de-
sign, Figure 6.16.
With pitting corrosion, the
chemical composition of the
attacking medium causes a
local break-up of the passive layer. Especially salts, preferably Cl—ions, show this behaviour.
This local attack causes a dissolution of the material on the damaged points, a depression
develops. Corrosion products accumulate in this depression, and the access of oxygen to the
bottom of the hole is obstructed. However, oxygen is required to develop the passive layer,
therefore this layer cannot be completely cured and pitting occurs, Figure 6.17.
Stress-corrosion cracking occurs when the material displaces under stress and the passive
layer tears, Figure 6.18. Now the unprotected area is subjected to corrosion, metal is dis-
solved and the passive
layer redevelops (figures 1-
3). The repeated displace-
ment and repassivation
causes a crack propaga-
tion. Stress corrosion
cracking takes mainly
place in chloride solutions.
The crack propagation is
transglobular, i.e. it does
not follow the grain
boundaries.
Pitting Corrosion of a Storage Container
Steel
br-er-06-17e.cdr
Figure 6.17
Model of Crack PropagationThrough Stress Corrosion Cracking
1 2 3 4 5 6
121110987
offset; passive layer; metal surface; dislocation
br-er-06-18e.cdr
Figure 6.18
6. Welding High Alloy Steels 80
Figure 6.19 shows the expansion-rate dependence of stress corrosion cracking. With very
low expansion-rates, a curing of the passive layer is fast enough to arrest the crack. With
very high expansion-rates, the failure of the specimen originates from a ductile fracture. In
the intermediate range, the material damage is due to stress corrosion cracking.
Figure 6.20 shows an example of crack propagation at transglobular stress corrosion crack-
ing. A crack propagation speed is between 0,05 to 1 mm/h for steels with 18 - 20% Cr and 8 -
20% Ni. With view to welding it is important to know that already residual welding stresses
may release stress corrosion cracking.
The most important problem in the field of welding is intergranular corrosion (IC).
It is caused by precipitation of chromium carbides on grain boundaries.
Although a high solubility of carbon in the austenite can be expected, see Fe-C diagram, the
carbon content in high alloyed Cr-Ni steels is limited to approximately 0,02% at room tem-
perature, Figure 6.21.
TransgranularStress Corrosion Cracking
© ISF 2002br-er06-20e.cdr
Figure 6.20
Influence of Elongation Speed onSensitivity to Stress Corrosion Cracking
SpRK
completecover layer tough fracture
Sensitiv
ity to s
tress c
orr
osio
n c
rackin
g
Elongation speed e
e2 e1
· ·
·
T=RT
© ISF 2002br-er06-19E.cdr
Figure 6.19
6. Welding High Alloy Steels 81
The reason is the very high affinity of chro-
mium to carbon, which causes the precipita-
tion of chromium carbides Cr23C6 on grain
boundaries, Figure 6.22. Due to these precipi-
tations, the austenite grid is depleted of
chromium content along the grain boundaries
and the Cr content drops below the parting
limit. The diffusion speed of chromium in aus-
tenite is considerably lower than that of car-
bon, therefore the chromium reduction cannot
be compensated by late diffusion. In the de-
pleted areas along the grain boundaries (line
2 in Figure 6.22) the steel has become sus-
ceptible to corrosion.
Only after the steel has been subjected to
sufficiently long heat treatment, chromium will
diffuse to the grain boundary and increase the
C concentration along the
grain boundary (line 3 in
Figure 6.22). In this way, the
complete corrosion resis-
tance can be restored (line 4
in Figure 6.22).
Figure 6.23 explains why the
IC is also described as in-
tergranular disintegration.
Due to dissolution of de-
pleted areas along the grain
boundary, complete grains
break-out of the steel.
Carbon Solubility ofAustenitic Cr - Ni Steels
0 0.05 0.1 0.15 0.2 0.25 % 0,3
Carbon content
600
700
800
900
1000
1100
°C
1200
A
He
at
tre
atm
en
t te
mp
era
ture
to Bain and Aborn
© ISF 2002br-er06-21e.cdr
Figure 6.21
Sensibility of a Cr - Steel
Chro
miu
m c
onte
nt of auste
nite
resistance limit
1 - homogenuous starting condition2 - start of carbide formation3 - start of concentration balance4 - regeneration of resistance limit
1
2
3
4
Distance from grain boundary© ISF 2002br-er-06-22e.cdr
Figure 6.22
6. Welding High Alloy Steels 82
The precipitation and re-
passivation mechanisms
described in Figure 6.22
are covered by intergranu-
lar corrosion diagrams ac-
cording to Figure 6.24.
Above a certain tempera-
ture carbon remains dis-
solved in the austenite
(see also Figure 6.21).
Below this temperature, a
carbon precipitation takes
place. As it is a diffusion
controlled process, the
precipitation occurs after a
certain incubation time
which depends on tem-
perature (line 1, precipita-
tion characteristic curve).
During stoppage at a con-
stant temperature, the
parting limit of the steel is
regained by diffusion of
chromium.
Figure 6.25 depicts characteristic precipitation curves of a ferritic and of an austenitic steel.
Due to the highly increased diffusion speed of carbon in ferrite, shifts the curve of carbon
precipitation of this steel markedly towards shorter time. Consequently the danger of inter-
granular corrosion is significantly higher with ferritic steel than with austenite.
Grain Disintegration
© ISF 2002br-er-06-23e.cdr
Figure 6.23
Area of Intergranular Disintegrationof Unstabilized Cr - Steels
¬R
ecip
roca
l o
f h
ea
t tr
ea
tme
nt
tem
pe
ratu
re 1
/T
oversaturatedaustenite
austenite -chromium carbide (M C )
no intergranular disintegration23 6
unsaturated austenite
Heat treatment time (lgt)
1 incubation time2 regeneration of resistance limit3 saturation limit for chromium carbide
1
2
3
austenite + chromium caride (M C )
to intergranular disintegration23 6 sensitive
© ISF 2002br-er-06-24e.cdr
Figure 6.24
6. Welding High Alloy Steels 83
As carbon is the element that triggers the intergranular corrosion, the intergranular corrosion
diagram is relevantly influenced by the c con-
tent, Figure 6.26.
By decreasing the carbon content of steel,
the start of carbide precipitation and/or the
start of intergranular corrosion are shifted
towards lower temperatures and longer
times. This fact initiated the development of
so-called ELC-steels (Extra-Low-Carbon)
where the C content is decreased to less
than 0,03%
During welding, the considerable influence of
carbon is also important for the selection of
the shielding gas, Figure 6.27. The higher the
CO2-content of the shielding gas, the
stronger is its carburising effect. The C-
content of the weld metal increases and the
steel becomes more susceptible to inter-
granular corrosion.
An often used method to
avoid intergranular corro-
sion is a stabilisation of the
steel by alloy elements like
niobium and titanium, Fig-
ure 6.28. The affinity of
these elements to carbon is
significantly higher than
that of chromium, therefore
carbon is compounded into
Nb- and Ti-carbides. Now
carbon cannot cause any
chromium depletion. The
Precipitation Curves of VariousAlloyed Cr Steels
Tempering time
Te
mp
erin
g t
em
pe
ratu
re
quenchtemperature
18-8-Cr-Ni steel17% Cr steel
precipitation curves for
cooling curve
© ISF 2002br-er06-25e.cdr
Figure 6.25
Figure 6.26
Influence of C-Contenton Intergranular Disintegration
101
102
103
104
105
106
Times
400
500
600
700
800
900
1000
Te
mp
era
ture
°C
0.07%C0.05%C
0.03%C
0.025%C
© ISF 2002br-er-06-26e.cdr
6. Welding High Alloy Steels 84
proportion of these alloy elements depend on the carbon content and is at least 5 times
higher with titanium and 10 times higher with niobium than that of carbon. Figure 6.28 shows
the effects of a stabilisation in the intergranular corrosion diagram. If both steels are sub-
jected to the same heat treatment (1050°C/W means heating to 1050°C and subsequent wa-
ter quenching), then the area of intergranular corrosion will shift due to stabilisation to
significantly longer times. Only with a much higher heat treatment temperature the inter-
granular corrosion accelerates again. The cause is the dissolution of titanium carbides at suf-
ficiently high temperature. This carbide dissolution causes problems when welding stabilised
steels. During welding, a narrow area of the HAZ is heated above 1300°C, carbides are dis-
solved. During the subsequent cooling and the high cooling rate, the carbon remains dis-
solved.
If a subsequent stress relief treatment around 600°C is carried out, carbide precipitations on
grain boundaries take place again. Due to the large surplus of chromium compared with nio-
bium or titanium, a partial chromium carbide precipitation takes place, causing again inter-
Influence of Shielding Gason Intergranular Disintegration
Shield ing gas Ar [% ] C O2 O2
S 1 99 / 1
M 1 90 5 5
M 2 82 18 /
C omposition
0,2 0,5 1 2,5 5 10 25 50 100 250 h 1000400
450
500
550
600
°C
700
0.058 % C0.53 % NbNb/C = 9
0.030 % C0.51 % NbNb/C = 17 0.018 % C
0.57 % NbNb/C = 32M2
M1
S1
Heat treatment time
Heat tr
eatm
ent te
mpera
ture
© ISF 2002br-er06-27e.cdr
Figure 6.27
Influence of Stabilizationon Intergranular Disintegration
800
700
650
600
550
500
450
°C
Heat tr
eatm
ent te
mpera
ture
0,3 1 3 10 30 100 300 1000 h 10000Time
1050°C/W
X5CrNi18-10 unstabilized
800
700
650
600
550
500
450
°C
Heat tr
eatm
ent te
mpera
ture
0,3 1 3 10 30 100 300 1000 h 10000Time
1300°C/W
1050°C/W
X5CrNiTi18-10 stabilized
W.-No.:4301 (0,06%)
W.-No.:4541
© ISF 2002br-er06-28e.cdr
Figure 6.28
6. Welding High Alloy Steels 85
granular susceptibility. As this susceptibility is limited to very narrow areas along the welded
joint, it was called knife-line attack because of its appearance. Figure 6.29.
In stabilised steels, the chromium carbide represents an unstable phase, and with a suffi-
ciently long heat treatment to transform to NbC, the steel becomes stable again. The stronger
the steel is over-stabilised, the lower is the tendency to knife-line corrosion.
Nowadays the importance
of Nickel-Base-Alloys in-
creases constantly. They
are ideal materials when it
comes to components
which are exposed to spe-
cial conditions: high tem-
perature, corrosive attack,
low temperature, wear re-
sistance, or combinations
hereof. Figure 6.30 shows
one of the possible group-
ing of nickel-base-alloys.
Materials listed there are selected examples, the total number of available materials is many
times higher.
Group A consists of nickel
alloys. These alloys are
characterized by moderate
mechanical strength and
high degree of toughness.
They can be hardened only
by cold working. The alloys
are quite gummy in the an-
nealed or hot-worked con-
dition, and cold-drawn
material is recommended
for best machinability and
smoothest finish.
Knife-Line Corrosion
br-er-06-29e.cdr
Figure 6.29
© ISF 2002br-er-06-30e.cdr
Alloy Chem. composition Alloy Chem. Composition
Group A Group D1
Nickel 200 Ni 99.6, C 0.08 Duranickel 301 Ni 94.0, Al 4.4, W 0.6
Nickel 212 Ni 97.0, C 0.05, Mn 2.0 Incoloy 925 Ni 42.0, Fe 32.0, Cr 21.0, Mo 3.0, W 2.1, Cu 2.2, Al 0.3
Nickel 222 Ni 99.5, Mg 0.075 Ni-Span-C 902 Y2O3 0.5, Ni 42.5, Fe 49.0, Cr 5.3, W 2.4, Al 0.5
Group B Group D2
Monel 400 Ni 66.5, Cu 31.5 Monel K-500 Ni 65.5, Cu 29.5, Al 2.7, Fe 1.0, W 0.6
Monel 450 Ni 30.0, Cu 68.0, Fe 0.7, Mn 0.7 Inconel 718 Ni 52.0, Cr 22.0, Mo 9.0, Co 12.5, Fe 1.5, Al 1.2
Ferry Ni 45.0, Cu 55.0 Inconel X-750 Ni 61.0, Cr 21.5, Mo 9.0, Nb 3.6, Fe 2.5
Group C Nimonic 90 Ni 77.5, Cr 20.0, Fe 1.0, W 0.5, Al 0.3, Y2O3 0.6
Inconel 600 Ni 76.0, Cr 15.5, Fe 8.0 Nimonic 105 Ni 76.0, Cr 19.5, Fe 112.4, Al 1.4
Nimonic 75 Ni 80.0, Cr 19.5 Incoloy 903 Ni 39.0, Fe 34.0, Cr 18.0, Mo 5.2, W 2.3, Al 0.8
Nimonic 86 Ni 64.0, Cr 25.0, Mo 10.0, Ce 0.03 Incoloy 909 Ni 58.0, Cr 19.5, Co 13.5, Mo 4.25, W 3.0, Al 1.4
Incoloy 800 Ni 32.5, Fe 46.0, Cr 21.0, C 0.05 Inco G-3 Ni 38.4, Fe 42.0, Cu 13.0, Nb 4.7, W 1.5, Al 0.03, Si 0.15
Incoloy 825 Ni 42.0, Fe 30.0, Cr 21.5, Mo 3.0, Cu 2.2, Ti 1.0 Inco C-276 Ni 38.4, Fe 42.0, Cu 13.0, Nb 4.7, W 1.5, Al 0.03, Si 0.4
Inco 330 Ni 35.5, Fe 44.0, Cr 18.5, Si 1.1 Group E
Monel R-405 Ni 66.5, Cu 31.5, Fe 1.2, Mn 1.1, S 0.04
Typical Classification of Ni-Base Alloys
Figure 6.30
6. Welding High Alloy Steels 86
Group B consists mainly of those nickel-copper alloys that can be hardened only by cold
working. The alloys in this group have higher strength and slightly lower toughness than
those in Group A. Cold-drawn or cold-drawn and stress-relieved material is recommended for
best machinability and smoothest finish.
Group C consists largely of nickel-chromium and nickel-iron-chromium alloys. These alloys
are quite similar to the austenitic stainless steels. They can be hardened only by cold working
and are machined most readily in the cold-drawn or cold-drawn and stress-relieved condition.
Group D consists primary of age-hardening alloys. It is divided into two subgroups:
D 1 – Alloys in the non-aged condition.
D 2 – Aged Group D-1 alloys plus several other alloys in all conditions.
The alloys in Group D are characterized by high strength and hardness, particularly when
aged. Material which has been solution annealed and quenched or rapidly air cooled is in the
softest condition and does machine easily. Because of softness, the non-aged condition is
necessary for trouble free drilling, tapping and all threading operations. Heavy machining of
the age-hardening alloys is best accomplished when they are in one of the following condi-
tions:
1. Solution annealed
2. Hot worked and quenched or rapidly air cooled
Group E contains only one material: MONEL R-405. It was designed for mass production of
automatically machined screws.
Due to the high number of possible alloys with different properties, only one typical material
of group D2 is discussed here: Material No. 2.4669, also known as e.g. Inconel X-750.
The aluminium and titanium containing 2.4669 is age-hardening through the combination of
these elements with nickel during heat treatment: gamma-primary-phase (γ') develops which
is the intermetallic compound Ni3(Al, Ti).
During solution heat treatment of X-750 at 1150°C, the number of flaws and dislocations in
the crystal is reduced and soluble carbides dissolve. To achieve best results, the material
6. Welding High Alloy Steels 87
should be in intensely worked condition before heat treatment to permit a fast and complete
recrystallisation. After solution heat treatment, the material should not be cold worked, since
this would generate new dislocations and affect negatively the fracture properties.
The creep rupture resistance of X-750 is due to an even distribution of the intercrystalline γ'
phase. However, fracture properties depend more on the microstructure of the grain bounda-
ries. During an 840°C stabilising heat treatment as part of the triple-heat treatment, the fine γ'
phase develops inside the grains and M23C6 precipitates onto the grain boundaries. Adjacent
to the grain boundary, there is a γ' depleted zone. During precipitation hardening (700°C/20
h) γ' phase develops in these depleted zones. γ' particles arrest the movement of disloca-
tions, this leads to improved strength and creep resistance properties.
During the M23C6 transformation, carbon is stabilised to a high degree without leaving chro-
mium depleted areas along the grain boundaries. This stabilisation improves the resistance
of this alloy against the attack of several corrosive media.
With a reduction of the precipitation temperature from 730 to 620°C – as required for some
special heat treatments – additional γ' phase is precipitated in smaller particles. This en-
hances the hardening effect and improves strength characteristics.
Further metallurgical discussions about X-750, can be taken from literature, especially with
view to the influence of heat treatment on fracture properties and corrosion behaviour.
The recommended processes for welding of X-750 are tungsten inert gas, plasma arc, elec-
tron beam, resistance, and pressure oxy arc welding.
During TIG welding of INCONEL X-750, INCONEL 718 is used as welding consumable. Joint
properties are almost 100% of base material at room temperature and about 80% at 700° -
820°C. Figure 6.31 shows typical strength properties of a welded plate at a temperature
range between -423° and 1500°F (-248 – 820°C).
Before welding, X-750 should be in normalised or solution heat treated condition. However, it
is possible to weld it in a precipitation hardened condition, but after that neither the seam nor
the heat affected zone should be precipitation hardened or used in the temperature range of
precipitation hardening, because the base material may crack. If X-750 was precipitation
hardened and then welded, and if it is likely that the workpiece is used in the temperature
range of precipitation hardening, the weld should be normalised or once again precipitation
hardened. In any case it must be noted that heat stresses are minimised during assembly or
welding.
6. Welding High Alloy Steels 88
X-750 welds should be solution heat treated before a precipitation hardening. Heating-up
speed during welding must be from the start fast and even touching the temperature range of
precipitation hardening only as briefly as possible. The best way for fast heating-up is to in-
sert the welded workpiece into a preheated furnace.
Sometimes a preheating before welding is advantageous – if the component to be welded
has a poor accessibility, or the welding is complex, and especially if the assembly proves to
be too complicated for a post heat treatment. Two effective welding preparations are:
1. 1550°F/16 h, air cooling
2. 1950°F/1 h, furnace cooling with 25°-100°F/h up to 1200°F, air
A repair welding of already fitted parts should be followed by a solution heat treatment (with a
fast heating-up through the temperature range of precipitation hardening) and a repeated
precipitation hardening.
A cleaning of intermediate layers must be
carried out to remove the oxide layers which
are formed during welding. (A complete isola-
tion of the weld metal using gas shielded
processes is hardly possible). If such films
are not removed on a regular basis, they can
become thick enough to cause material sepa-
rations together with a reduced strength.
Brushing with wire brushes only polishes the
surface, the layer surface must be sand-
blasted or ground with abrasive material. The
frequency of cleaning depends on the mass
of the developed oxides. Any sand must be
removed before the next layer is welded.
X-750 can be joined also by spot-, projection-
, seam-, and flash butt welding. The welding
equipment must be of adequate performance.
X-750 is generally resistance welded in nor-
malized or solution heat treated condition.
© ISF 2002br-er06-31e.cdr
tensile strength
0.2% yield stress
elongation in 1/2”
elongation in 2”
Elo
ngation, %
S
tress, 1000 p
si
Temperature, F°
220
200
180
160
140
120
100
80
60
20
0
10
30
10
20
0-423 0 200 400 600 800 1000 1200 1400 1600
Mechanical Properties ofa Typical Ni-Base Alloy
Figure 6.31
7.
Welding of Cast Materials
7. Welding of Cast Materials 90
Figure 7.1 pro-
vides a summary
of the different
cast iron materi-
als. In this con-
nection it is only
referred to cast
iron, cast steel
and malleable
steel, as special
cast materials,
due to their poor
weldability, are of
no importance in
welding.
Figure 7.2 shows the designation of
the cast material in accordance with
DIN EN 1560. A distinction is made
between the designation “according to
the material code” and the designa-
tion “according to the material num-
ber”. In Figure 7.2, examples of two
materials are specified.
Table of the cast Iron Materials
cast materials
perlitic
alloyed
perlitic
ferritic
cast steel
unalloyed
ferritic austenitic
decarburized not decarburized
malleable iron
decarburizedannealed
malleable cast iron
not decarburizedannealed
malleable cast iron
ferritic perlitic perliticferritic austenitic
ferritic perlitic austenitic
cast iron
lemellar graphitecast iron
nodular graphitecast iron
high alloyedlow alloyed
special cast iron(G...)
hard castiron
clear chillcasting
low Ccontent
high C-content
ledeburitic graphite
Cr-castiron
otherelements
Si-castiron
Al-castiron
plastics, gypsum and s.th.similar
non-iron-metalcast materials
metalliccast materials
non-metalliccast materials
iron-carbon-cast materials
© ISF 2002br-er-07-01e.cdr
Figure 7.1
Designation according to the material number
e.g.: EN- J L 1271
1 2 3 4,5,6
Position 1: EN - standardised materialPosition 2: J - cast materialPosition 3: L - graphite structure (lamellar graphite)Position 4: 1 - number for the main characteristicPosition 5: 27 - material identification numberPosition 6: 1 - special requirement
Designation of Materials
© ISF 2004br-er07-02e.cdr
Designation according to the material code (DIN EN 1560)
e.g.: EN-GJ L F – 150
1 2 3 4 5
Position 1: EN -Position 2: GJ -Position 3: L -Position 4: F -Position 5: 150 - (R = 150 N/mm )
-Position 6: -
m
2
standardised materialcast materialgraphite structure (lamellar graphite)microstructure (ferritic)mechanical properties
chemical composition (high alloyed)optionally
Figure 7.2
7. Welding of Cast Materials 91
Figure 7.3 depicts a survey of the mechanical properties and the chemical composi-
tions of several customary cast materials. As to its analysis and mechanical proper-
ties which are very different from other cast materials, cast steel constitutes an
exception to the rule.
In Figure 7.4 the stable and the metastable iron-carbon diagram are shown. The dif-
ferences between
the cast material
are best explained
this way. Cast iron
with lamellar and
spheroidal graph-
ite has carbon
contents of be-
tween 2,8 and
4,5%. Through the
addition of alloying
elements, above
all Si, these mate-
rials solidify fo llow-
ing the stable sys-
tem, i.e., the car-
bon is precipitated
in the form of
graphite. Malleable
cast iron shows
similar C-contents,
the solidification
from the molten
metal, however,
follows the me-
tastable system.
The C-contents of
cast steel, on the
Figure 7.3
Figure 7.4
7. Welding of Cast Materials 92
other hand, comply with those of
common structural steels, i.e., they
are, as a rule, below 0,8% C.
The structure of a normalised cast iron
which is composed of ferrite (bright)
and pearlite (dark) is shown in Figure
7.5. Since the properties are similar to
those of structural steels these materi-
als are weldable, constructional weld-
ing is also possible. It is recommended
to normalise the cast steel parts before
welding. Through this type of heat
treatment, on the one hand the trans-
formation of the cast structure is ob-
tained, the residual stresses inside the
workpiece are, on the other hand, re-
duced.
From a C-content in the steel cast of
0,15% up, it is recommended to carry
out preheating during welding, the
preheating temperature should follow
the analysis of the material, the work-
piece geometry and the welding
method. After welding the cast work-
pieces are subject to stress-relief an-
nealing.
Figure 7.6 shows the structure of cast
iron with lamellar graphite (grey cast
iron). Apart from their carbon content,
these materials are characterised by
increased contents of S and P which
Figure 7.5
Figure 7.6
7. Welding of Cast Materials 93
improves castability. Besides the poor mechanical properties (elongation after frac-
ture of approx. 1%), these chemical properties also impede welding with ordinary
means. It is not possible to carry out constructional welding with grey cast iron. Re-
pair welds of grey cast iron are, in contrast, carried out more frequently as damaged
cast parts are not easily replaceable. For those repair welds, the cast parts must be
preheated (entirely or partly) to temperatures of approx. 650°C. Heating and cooling
must be done very slowly as the cast piece may be destroyed already by the thermal
stresses. The highly liquid weld metal also constitutes a problem, and thus the molten
pool must be supported by a carbon pile. Welding may be carried out with similar
filler material (materials of the same composition as the base). If grey cast iron is to
be welded without any preheating, the filler material must, as a rule, be dissimilar (of
different composition to the base metal). During this type of welding, there are always
strong structural changes in the region of the weld which lead to high hardening and
high residual stresses. For the minimisation of these structural changes, a highly duc-
tile filler material is applied. The heat input into the base material should be as low as
possible.
Figure 7.7 depicts
the structural con-
stitution of spher-
oidal graphite cast
iron. The graphite
spheroidization is
achieved by the
addition of magne-
sium and cerium.
As, with this type
of graphite, the
notch actions are
considerably
lesser than this is
the case with grey cast iron, this type of cast iron is characterised by substantially
better mechanical parameters with a considerably higher elongation after fracture
and improved ductility. For this reason, the risk of material failure caused by weld
residual stresses or thermal stresses is considerably reduced for spheroidal graphite
Figure 7.7
7. Welding of Cast Materials 94
cast iron. Frequently, nickel-based
alloys are used as filler material. Prob-
lems occur in the HAZ where, besides
the ledeburite eutectic alloy system,
also Ni-Fe-martensite is frequently
formed. Both structures lead to ex-
treme hardening in the HAZ which
can be removed only by time-
consuming heat treatment.
Figures 7.8 and 7.9 show the structures of
Carburized Annealed Malleable Cast Iron
(7.7) and of Decarburized Annealed Malle-
able Cast Iron (7.9). The composition of the
malleable cast iron is thus that during solidifi-
cation, the total of carbon is bound in cemen-
tite and precipitated. During a subsequent
annealing process, the iron carbide disinte-
grates into graphite and iron.
Figure 7.8
Figure 7.9
7. Welding of Cast Materials 95
If annealing is carried out in neutral
atmosphere, the structure of Carbur-
ized Annealed Malleable Cast Iron
develops. Annealing in oxidising at-
mosphere leads to the decarburisa-
tion of the workpiece surfaces and
Decarburized Annealed Malleable
Cast Iron is developed, Figure 7.10.
Carburized Annealed Malleable
Cast Iron is not weldable. Decarbur-
ized Annealed Malleable Cast Iron,
in contrast, is weldable.
You can see in Figure 7.11 that, also
with malleable cast iron, hardening in
the region of the HAZ occurs. For car-
rying out constructional welds made of
malleable cast iron parts, a special
material quality has been developed.
Figure 7.11 shows that this material,
EN-GJMW-400-12, is characterised by
considerably less hardening. This ma-
terial is weldable without any problems
up to a wall thickness of 8 mm.
Figure 7.11
Figure 7.10
8.
Welding of Aluminium
8. Welding of Aluminium 97
Figure 8.1 compares basic physical properties
of steel and aluminium. Side by side with dif-
ferent mechanical behaviour, the following
differences are important for aluminium weld-
ing:
- considerably lower melting point compared
with steel
- three times higher heat conductivity
- considerably lower electrical resistance
- double expansion coefficient
- melting point of Al203 considerably higher
than that of Al; metal and iron oxide melt ap-
proximately at the same temperature.
Figure 8.2 compares some mechanical prop-
erties of steel with properties of some light
metals. The important advantages of light
metals compared with steel are especially
shown in the right part of the figure. If a comparison should be based on an identical stiff-
ness, then the aluminium supporting beam has a 1.44 times larger cross-section than the
steel beam, however only about 50% of its weight.
Figure 8.3 compares quali-
tatively the stress-strain dia-
gram of Aluminium and
steel. In contrast to steel,
aluminium has a fcc (face
centred cubic)-lattice at
room temperature. This is
why there is no distinct yield
point as being the case in a
bcc (body centred cubic)-
lattice. Aluminium is not
subject to a lattice trans-
Deflexions and Weights of Cantilever Beams Under Load
br-er-08-02.cdr
Figure 8.2
Property Al Fe
Atomic weight [g/Mol] 26.9 55.84
Specific weight [g/cm³] 2.7 7.87
Lattice fcc bcc
E-module [N/mm²] 71*10³ 210*10³
R PO,2 [N/mm²] ca. 10 ca. 100
R m [N/mm²] ca. 50 ca. 200
spec. Heat capacity [J/(g*°C)] 0.88 0.53
Melting point [°C] 660 1539
Heat conductivity [W/(cm*K)] 2.3 0.75
Spec. el. Resistance [nWm] 28-29 97
Expansion coeff. [1/°C] 24*10-6
12*10-6
FeO
Oxydes Al2O3 Fe3O4
Fe2O3
1400
Melting point of oxydes [°C] 2050 1600
(1455)
Basic Properties of Al and Fe
pO,2
m
© ISF 2002br-er08-01.cdr
Figure 8.1
8. Welding of Aluminium 98
formation during cooling, thus there is no structure transformation and consequently no
danger of hardening in the heat affected zone as with steel.
Figure 8.4 illustrates the effect of the considerably higher heat conductivity on the welding
process compared with steel. With aluminium, the temperature gradient around the welding
point is considerably smaller than with steel. Although the peak temperature during Al weld-
ing is about 900°C below steel, the isothermal curves around the welding point have a clearly
larger extension. This is due to the considerably higher heat conductivity of aluminium com-
pared with steel.
This special characteristic of Al requires a input heat volume during welding equivalent to
steel.
Figure 8.5 lists the most important alloy elements and their combinations for industrial use.
Due to their behaviour during heat treatment can Al-alloys be divided into the groups harde-
nable and non-hardenable (naturally hard) alloys.
Comparison of Stress-ElongationDiagrams of Al and Steel
Elongation
Al-alloy
Steel
Str
ess
© ISF 2002br-er08-03.cdr
Figure 8.3
Isothermal Curves of Steel and Al
4
2
-2
-4
4
2
-2
-4
low carbon steel
aluminium
400
600
200°C
800
10001200
1500
-6
-8
cm
8
6
-18 -16 -14 -12 -10 -8 -6 -4 -2 0 2 cm 6
500
600
400300
200
100°Ccm
© ISF 2002br-er08-04.cdr
Figure 8.4
8. Welding of Aluminium 99
Figure 8.6 shows typical applications of some
Al alloys together with preferably used weld-
ing consumables.
Aluminium alloys are often welded with con-
sumable of the same type, however, quite
often over-alloyed consumables are used to
compensate burn-off losses (especially with
Mg and Zn because of their low boiling point)
and to improve the mechanical properties of
the seam.
The classification of Al alloys into two groups
is based on the characteristic that the group
of the non-hardenable alloys cannot increase
the strength through heat treatment, in con-
trast to hardenable alloys which have such a
potential.
The important hardening mechanism for this
second group is explained by the figures 8.7 und 8.8. Example: If an alloy containing about
4.2% Cu, which is stable at room temperature, is heat treated at 500°C, then, after a suffi-
ciently long time, there will be only a single phase structure present. All alloy elements were
dissolved, Figure 8.8 between point P and Q.
When quenched to room
temperature in this condi-
tion, no precipitation will
take place. The alloy ele-
ments are forced to remain
dissolved, the crystal is out
of equilibrium. If such a
structure is subjected to an
age hardening at room or
elevated temperature, a
precipitation of a second
phase takes place in ac-
Classification of Aluminium Alloys
67
86
78
Mg
Si
Mn
Cu
ZnAl
Al Cu Mg
Al Mg Si
Al Zn Mg
Al Zn Mg Cu
Al Si Cu
Al Si
Al Mg
Al Mg Mn
Al Mn
non-h
ard
enable
allo
ys
hard
enable
allo
ys
© ISF 2002br-er08-05.cdr
Figure 8.5
Use and Welding Consumablesof Aluminium Alloys
Al - alloys Typical use W elding consumable
Al99,5electrical engineering
SG-Al 99,5Ti;
SG-Al 99,5
AlCuMg1 mechanical engineering, food
industriesSG-AlMg4,5Mn
AlMgSi0,5 architecture, electrical
engineering, anodizing quality
SG-AlMg5; SG-AlMg4,5Mn;
SG-AlSi5
AlSi5 architecture, anodizing quality SG-AlSi5
AlMg3 architecture, apparatus-, vehicle-,
shipbuilding engineering, furniture
industry
SG-AlMg3;
SG-AlMg4,5Mn
AlMg2Mn0,8 apparatus-, vehicle-, shipbuilding
engineering
SG-AlMg5; SG-AlMg3;
SG-AlMg4,5Mn
AlMn1 apparatus-, vehicle-engineering,
food industrySG-AlMn1;SG-Al99,5T
base material - aluminium
percentage of alloy elements without factor
© ISF 2002br-er-08-06.cdr
Figure 8.6
8. Welding of Aluminium 100
cordance with the binary system, the crystal tries to get back into thermodynamical equilib-
rium.
Depending on the level of
hardening temperature, the
precipitation takes place in
three possible forms: co-
herent particles (i.e. parti-
cles deviating from the
matrix in their chemical
composition but having the
same lattice structure),
partly coherent particles
(i.e. the lattice structure of
the matrix is partly re-
tained), and incoherent
particles (lattice structure completely different from the matrix), Figure 8.7. Coherent particles
formed at room temperature can be transformed into incoherent particles by increase of tem-
perature (i.e. enabling diffusion).
The precipitations cause a restriction to the
dislocation movement in the matrix lattice, thus
leading to an increase in strength. The finer the
precipitations, the stronger the effect.
At an increased temperature (heat ageing, Fig-
ure 8.7) a maximum of second phase has pre-
cipitated after elapse of a certain time.
Consequently a prolonged stop at this tem-
perature does not lead to an increased
strength, but to coarsening of particles due to
diffusion processes and to a decrease in
strength (less bigger particles in an extended
space).
Ageing Mechanism
solution heat treatment
quenching
ageing at slightly
increased temperature
coherent
precipitations,
cold aged
condition
temperature
rise
temperature
riselonger warm
ageing
longer warm
ageing
partly coherent
precipitations,
warm aged
condition
partly coherent
and incoherent
precipitations,
softening
stable incoherent
equilibrium phase
stable condition
stable condition
solidification of alloy elements
in solid solution
oversaturated solid solution,
metastable condition
coherent and partly coherent
precipitations, transition conditions
cold ageing -- warm ageing
repeated hardening
regeneration
cold ageing (RT ageing)
warm ageing
© ISF 2002br-er-08-07.cdr
Figure 8.7
Phase Diagram Al-Cu
700
600
500
400
300
200
100
0 1 2 3 4 5 mass-% 7
Q
P
liquid
liquid and solid
copper containingaluminium solid solution
aluminium solid solutionand copper aluminide(Al Cu)2
copper content ofAlCuMg
Copper
Te
mp
era
ture
© ISF 2002br-er08-08.cdr
Figure 8.8
8. Welding of Aluminium 101
After a very long heat ageing a stable condi-
tion is reached again with relatively large pre-
cipitations of the second phase in the matrix.
In Figure 8.7 is this stable final condition iden-
tical with the starting condition. A deteriorati-
on of mechanical properties only happens
during hot ageing, if the ageing time is exces-
sively long.
The complete process of hardening at room
temperature is metallographic also called age
hardening, at elevated temperature heat age-
ing. A decrease in strength at too long ageing
time is called over-ageing.
Figure 8.9 shows a schematic representation
of time-temperature curves during hardening
with age hardening and heat ageing.
Figure 8.10 shows the
strength increase of AlZnMg
1 in dependence of time.
The difference between age
hardening and heat ageing
is here very clear. Due to
improved diffusion condi-
tions is the strength increase
in the case of heat ageing
much faster than in the case
of age hardening. The
strength maximum is also
reached considerably ear-
lier. The curve of hot ageing shows clearly the begin of strength loss when held at a too long
stoppage time. This figure shows another specialty of the process of ageing. During ageing, a
Temperature - Time DistributionDuring Ageing
solution heat treatment
quenching
heat ageing
age hardening
2 4 6 8 10 12 14h
500
°C
400
300
200
100
0
Q
P
Te
mp
era
ture
Time
© ISF 2002br-er08-09.cdr
Figure 8.9
Increase of Yield Stress DuringAgeing of AlZnMg1
quenched Ageing time in h
0.2
% y
ield
str
ess
in N
/mm
²s
0.2
water quenching (~900°C/min)air cooling (~30°C/min)
10-1
100
101
10² 10³
380
320
260
200
140
80
120°C
RT
© ISF 2002br-er-08-10.cdr
Figure 8.10
8. Welding of Aluminium 102
second phase is precipitated from a single-phase structure. To initiate this process, the struc-
ture must contain nuclei of the second phase. However, a certain time is required to develop
such nuclei. Only after formation of nuclei can the increase in strength start. The period up to
this point is called incubation time.
Figure 8.11 shows the effect of the height of
ageing temperature level on both, mechanical
properties of a hardenable Al-alloy and on in-
cubation time. The lower the ageing tempera-
ture, the higher the resulting values of yield
stress and tensile strength. If a low ageing tem-
perature is selected, the ageing time as well as
the incubation time become extremely long.
Figure 8.11 shows that a the maximum yield
stress is reached after a period of about one
year under a temperature of 110°C. An in-
crease of the ageing temperature shortens the
duration of the complete precipitation process
by a certain value raised by 1 to a power. On
the other hand, such an acceleration of ageing
leads to a lowering of the
maximum strength. As the
lower part of the figure
shows, the fracture elonga-
tion is counter-proportional
to the strength values, i.e.
the strength increase
caused by ageing is ac-
companied by an embrit-
tlement of the material.
Influence of Ageing Temperatureand -Time on Ageing
260
0 10-2
10-1
100
101
102
103
h 104
190180
150 135
110°C
230260
30min
1day
30
20
10
Fra
ctu
re e
lon
ga
tio
nd
20
.2%
yie
ld s
tre
ss
s0.2
400
300
200
110
135
180
Te
nsile
str
en
gth
sB
110
135
150
180
230
500
400
300
200
Ageing time
%
N/mm²
N/mm²
205
260°C
190
150
190205°C
230
205
1week
1month
1year
© ISF 2002br-er08-11.cdr
Figure 8.11
Age Hardening of Al Alloys
0 30 70
100
200
300
400
N/mm²
% Strain
0
Te
nsile
str
en
gth
Rm
AlMg5
AlMg3
Al99,5
© ISF 2002br-er-08-12.cdr
Figure 8.12
8. Welding of Aluminium 103
Figure 8.12 shows a method of how to increase the strength of non-hardenable alloys. As no
precipitations are present to reduce the movement of dislocations, such alloys can only be
strengthened by cold working.
Figure 8.12 illustrates two essential mecha-
nisms of strength increase of such alloys. On
one hand, tensile strength increases with in-
creasing content of alloy elements (solid solu-
tion strengthening), on the other hand, this
increase is caused by a stronger deformation
of the lattice.
Figure 8.13 shows the effect of the welding
process on mechanical properties of a cold-
worked alloy. Due to the heat input during
welding, the blocked dislocations are released
(recovery), in addition, a grain coarsening will
start in the HAZ. This is followed by a strong
drop in yield point and tensile strength. This
strength loss cannot be overcome in the case
of a welding process.
Figure 8.14 illustrates the
mechanisms in the case of a
hardenable aluminium alloy.
As a consequence of the
welding heat, the precipita-
tions are solution heat treated
and the strength values de-
crease in the weld area. Due
to the age hardening, a re-
strengthening of the alloys
takes place with increasing
time.
Non-Hardenable Al Alloy
Distance from Seam Centre
HV
30
80 60 40 20 0 20 40 60 mm 100
0,7
0,6
0,5
0,4
0,3
0,20
50
100
150
200
250
300
N/mm²
R/R
p0
,2m
Ror
Rm
p0
,2
© ISF 2002br-er08-13.cdr
Figure 8.13
Hardenable Al Alloy
4 mm plates of: AlZnMg1F32start values: R =263N/mm²
R =363 N/mm²
welding method: WIG, both sides,simultaneously
welding consumable: S-AlMg5specimens with machinedweld bead
p0,2
m
Distance from seam centre
Str
ess
90 days RT
21 days RT
1 day RT
80
400
N/mm²
350
300
250
200
150
100
5080 60 40 20 0 20 40 60 100 mm 140
90 days RT
1 day RT
21 days RT
Rp0,2
Rm
© ISF 2002br-er-08-14.cdr
Figure 8.14
8. Welding of Aluminium 104
Figure 8.15 shows another
problematic nature of Al-
welding. Due to the high
thermal expansion of alu-
minium, high tensions de-
velop during solidification
of the weld pool in the
course of the welding cy-
cle. If the welded alloy indi-
cates a high melting
interval, cracks may easily
develop in the weld.
A relief can be afforded by
preheating of the material, Figure 8.16. With an increasing preheat temperature, the amount
of fractured welds decreases. The different behaviour of the three displayed alloys can be
explained using the right
part of the figure. One can
see that the manganese
content influences signifi-
cantly the hot crack suscep-
tibility. The maximum of this
hot crack susceptibility is
likely with about 1% Mg con-
tent (corresponds with alloy
1). With increasing MG con-
tent, hot crack susceptibility
decreases strongly (see also
alloy 2 and 3, left part).
To avoid hot cracking, partly very different preheat temperatures are recommended for the
alloys. Zschötge proposed a calculation method which compares the heat conductivity condi-
tions of the Al alloy with those of a carbon steel with 0.2% C. The formula is shown in Figure
Hot Cracks in a Al Weld
© ISF 2002br-er-08-15.cdr
Figure 8.15
Influence of Preheat Temperatureand Magnesium Content
1: AlMgMn 2: AlMg 2,5 3: AlMg 3,5
Preheat temperature
Weld
cra
ckin
g tendency
Cra
ckin
g s
usceptib
ility
Alloy content
X
X
X
X
100
80
60
40
20
%
0 100 200 300 400 °C 500
1
2
3
0 21 3 % 4
Si
Mg
© ISF 2002br-er-08-16.cdr
Figure 8.16
8. Welding of Aluminium 105
8.17, together with the related calculation result. These results are only to be regarded as
approximate, the individual application is subject to the information of the manufacturer.
Another major problem
during Al welding is the
strong porosity of the
welded joint. It is based on
the interplay of several
characteristics and hard to
suppress.
Pores in Al are mostly
formed by hydrogen,
which is driven out of the
weld pool during solidifica-
tion. Solubility of hydrogen
in aluminium changes
abruptly on the phase
transition melt-crystal, i.e.
the melt dissolves many
times more of the hydro-
gen than the just forming
crystal at the same tem-
perature. This leads to a
surplus of hydrogen in the
melt due to the crystallisa-
tion during solidification.
This surplus precipitates in
form of a gas bubble at the
solidifying front. As the
melting point of Al is very
low and Al has a very high heat conductivity, the solidification speed of Al is relatively high.
As a result, in the melt ousted gas bubbles have often no chance to rise all the way to the
surface. Instead, they are passed by the solidifying front and remain in the weld metal as
pores, Figure 8.18.
Recommendations for Preheating
Welding possible without preheating:AlMg5, AlMg7, AlMg4.5Mn,AlZnMg3, AlZnMg1
T in °C temperature of melt start (solidus temperature)
in J/cm*s*K heat conductivity
S
Al-Leg.
T in °C preheat temperaturevorw.
l
;
745TT
.LegAl
SVorw.
-l
-=
melting point pure aluminium
Increasing better weldability
Recom
mended p
reheat te
mpera
ture
Al 99,9
8R
Al9
9,9
Al9
9,8
Al 99,7
Al 99,5
Al 99
Al R
Mg0,5
Al M
g S
i 0,5
Al M
g S
i 0,8
Al M
g S
i 1
EA
l M
g S
i 1
Al M
g 1
Al S
i 5
Al C
u M
g 1
Al R
Mg 2
Al C
u M
g 0
,5A
l M
nA
l M
g 2
Al C
u M
g 2
Al M
g 3
Al M
g 3
Si
Al M
g M
n
Al Z
n M
g C
u 0
,5A
l Z
n M
g C
u 1
,5
mild
ste
el (0
.2%
C)
without pre
heating
660
600
°C
500
400
300
200
100
0
© ISF 2002br-er-08-17.cdr
Figure 8.17
Excessive Porosity in a Al Weld
© ISF 2002br-er-08-18.cdr
Figure 8.18
8. Welding of Aluminium 106
To suppress such pore
formation it is therefore
necessary to minimise the
hydrogen content in the
melt. Figure 8.19 shows
possible sources of hydro-
gen during MIG welding of
Al.
Figure 8.20 and 8.21 show
the effect of pure thermal
expansion during Al welding.
The large thermal expansion
of the aluminium along with
the relatively large heat af-
fected zones cause in com-
bination with a parallel gap
adjustment a strong distor-
tion of the welded parts. To
minimise this distortion, the
workpieces must be set at a
suitable angle before weld-
ing, Figure 8.21.
Ingress of Hydrogen Into the Weld
too thick and water containing oxyde layerby too long or open storagein non air-conditioned rooms
nozzle deposits and too steep inclinationof the torch cause turbulences
VS
too thick oxyde layer(condensed water)
dirt film(oil, grease)
H2
H2
Grundwerkstoff
Poren
festesSchweißgut
feuchte Luft
poorcurrent transition
irregularwireelectrodefeed
humid air
humid air(nitrogen, oxygen, water)
pores
solid weld metal
base material
© ISF 2002br-er-08-19.cdr
Figure 8.19
Weld Gap Adjustment
parallel gap
overlap
weld pool
weld pool
opening gap
© ISF 2002br-er-08-20.cdr
Figure 8.20
8. Welding of Aluminium 107
Examples to Minimise Distortion
wedge flame
© ISF 2002br-er08-21.cdr
Figure 8.21
9.
Welding Defects