Electron Beam Welding of Superduplex Stainless Steel S32750

Electron Beam Welding of Superduplex 

Stainless Steel S32750 

Presenter: 

Parvan Chavdarov 

Academic Education and Degrees 

Post-graduate student – “Weldability of duplex stainless steels” 

– Technical University of Sofia 

2001 Master degree on technology of metals, Technical 

University of Sofia, Bulgaria 

Present professional position: 

Expert in Industrial Services in TUV Rheinland Bulgaria 

Originally presented at the Stainless Steel World 2007 Conference, Maastricht, the Netherlands

Electron Beam Welding of Superduplex 

Stainless Steel S32750 

Authors: Read.Dr Serafim Serafimov – Technical University of Sofia 

Dipl.eng.Parvan Chavdarov – post-graduate student – TUV Rheinland Sofia 

Adress: Zona B-19, “Dr Kalinkov” str.18, app.10, Sofia, Bulgaria 

Post Code 1309, tel. +359 898529051 

e-mail: P_Chavdarov@mail.bg 

Keywords: duplex stainless steel, welding parameters, ferrite content, microstructure, 

intermetallic phases, metallographic analysis 

Abstract: 

The aim of this paper is to present the influence of the welding parameters upon the 

shape of the weld, the microstructure and respectively the precipitation of second 

phases in superduplex stainless steel S32750. Plates with 10mm thickness have been 

butt welded with electron beam without filler by altering the heat input. Metallographic 

sections have been prepared and the ferrite content in the weld metal has been 

measured. The availability of intermetallic phases is examined. 


1. Introduction 

Duplex stainless steels are a family of grades combining good corrosion resistance 

with high strength and ease of fabrication. Their physical properties are between those 

of the austenitic and ferritic stainless steels and tend to be closer to those of the ferritic 

and to carbon steel. These steels are very attractive for various applications because 

of their advantages: 

• Very high corrosion resistance especially against pitting and crevice corrosion in 

aggressive media containing chlorides and fluorides; 

• Higher mechanical properties (because of the ferrite phase) than the austenite 

stainless steels which are predominantly used in practice; 

• Higher ductility (because of the austenitic phase) than the ferrite stainless steels 

which reflects on the weldability. 

• Better compatibility to C-Mn structural steels, because the coefficients of heat 

expansion are comparatively close, which results in lower thermal stresses. 

All these advantages result in extremely wider usage of duplex alloys in many industry 

branches like petrochemical, chemical, pulp and paper, off-shore, energy, gas fuel, 

mining, shipbuilding, power generations, marine transportations, food manufactures, 

etc. 

A balanced austenite/ferrite ratio of 50/50% is crucial for their high performance. 

Compared with normal austenitic steels, duplex steels contain less nickel which 

increases its cost-effectiveness considerably. The desire structure is obtained by heat 

treatment at approximately 1050 to 1100˚C (solution annealing). The optimum ratio 

between both phases can be influenced by the welding processes. 

2. Metallurgy of duplex stainless steels 

2.1 General 

Modern duplex stainless steels are characterized by a two phase structure, which 

consists of a mixture of about 50% volume austenite in ferrite grains. Both cast and 

wrought products have roughly equivalent volume fractions of ferrite and austenite, 

which in the case of wrought components, contain a rolling texture obtained by hot 

working, followed by a solution annealing and quench. The optimum phase balance for 

modern wrought products varies between manufacturers, but overall a range of 

between 45% and 60% austenite may be expected. 

Fig.1 shows a schematic section of the Fe-Cr-Ni diagram at the 70%Fe level. The 

phase proportions and their respective compositions are indicated for a given alloy 

analysis and annealing temperature, with the high temperature stability of the duplex 

structure being influenced more by nitrogen content, than by Cr or Mo. The addition of 

0,25%N to a 25%Cr alloy produces a ferrite volume fraction of about 50% at 1250˚C, 

compared to nearly 80% ferrite with 0,18%N. Nevertheless, it is difficult to predict the 

microstructure of a duplex alloy from simplified diagrams, due to the effects of other 

alloy elements, which modify the phase fields. 


Fig.1: Concentration profiles in 

the ternary Fe-Cr-Ni constitution 

diagram at 70% and 60%Fe. The 

effect of nitrogen additions is 

shown in the first figure. 

2.2 Heat treatment 

Numerous structural changes can occur in the duplex stainless steels during 

isothermal and anisothermal heat treatments. Most of these transformations are 

concerned with the ferrite, as element diffusion rates are approximately 100 times 

faster than in austenite. This is principally a consequence of the less compact lattice of 

the BCC crystal structure. Moreover, the ferrite is enriched in Cr and Mo, which 

promote the formation of Intermetallic phases. Besides, element solubility in the ferrite 

falls with a decrease in temperature, increasing the possibility of precipitation during 

heat treatment. 

Wrought and heat-treated products are considered to be segregation free, but in case 

of castings and welded joints the element segregation during cooling will affect 

precipitation kinetics and the stability of phases formed [1]. 

2.2.1 Temperatures above 1050˚C 

Duplex stainless steels solidify completely in the ferrite field for standard grades and 

normal cooling rates. This is followed by solid state transformation to austenite (fig.1), 

which is naturally reversible, so that any large increase in temperature, for example 

from 1050˚C to 1300˚C, leads to an increase in ferrite content. Further, as he 

temperature increases, there is a reduction in the portioning of substantial elements 

between the phases. In addition, the ferrite becomes enriched in interstitial elements 

such as carbon and nitrogen. 

Heat treatment in the temperature range 1100-1200˚C can have a dramatic influence 

on the microstructure of a wrought product. The grains can be made equiaxed by 

prolonged treatment at high temperature or can be rendered acicular, with a 

Widmannstaetten type structure by cooling an intermediate rate. A dual structure, 

consisting of both coarse and fine austenite grains, can be obtained by step 

quenching, with or without simultaneous mechanical strain. These acicular structures 

are also encountered in weld deposits. 

2.2.2 The 600-1050˚C nose 


The alloy lean grades are the least prone to Intermetallic phase precipitation and 

requires exposure of at least 10 to 20 hours to initiate formation at temperatures below 

900˚C. For that reason, a solution annealing temperature below 1000˚C can be chosen 

for this material. 

More alloyed grades are more sensitive to precipitation due to the molybdenum 

content. This element not only increases the rate of Intermetallic precipitation, but also 

extents the stability range to higher temperatures. That is why higher solution 

annealing temperatures are needed – above 1000˚C. 

The superduplex alloys show the greatest propensity for precipitations, due to their 

higher Cr, Mo and W contents. However, it should be emphasized that the precipitation 

kinetics in these high alloy grades are, at worst, equivalent to the superaustenitic or 

superferritic stainless steels [2]. And still, by taking precautions during heat treatment, 

including rapid removal from the furnace followed by water quenching, the superduplex 

alloys can be used satisfactory in industrial applications. However, especial care is 

required for heavy section components during all stages of production. 

Precipitates re-dissolve during a solution anneal, which for the superduplex grades 

must be performed at 1050˚C or above. A few minutes at 1050-1070˚C are sufficient 

for grade S32750. Similar high temperatures are necessary for welds as consumables 

tend to contain higher Ni, Si and Mn contents than base materials. The higher Nicontent 

encourages high austenite contents when annealed and results in enrichment 

of Cr and Mo in the remaining ferrite. This fact, combined with higher Si and Mn levels, 

increases the stability of Intermetallic phases. And yet, lower annealing temperatures 

(1040˚C compared to 1100˚C) can be used for weldments made with matching 

consumables [3]. 

2.2.3 The 300-600˚C nose 

The alloy lean grades are the least sensitive to hardening in this temperature range, 

and a significant effect is not recorded until about 3 hours exposure to 400˚C. A much 

shorter incubation time is found for more alloyed grades, containing molybdenum, 

which would appear to accelerate hardening [1]. The 25%Cr and superduplex alloys 

show the widest temperature range for hardening and shorter incubation times 

accordingly. This is the result of both the higher Cr and Mo contents and, if present, 

copper additions. 

2.2.4 Continuous cooling diagrams 

At temperatures near the solvus, the nucleation of precipitates is slow and their growth 

is fast, whereas the opposite is true at lower temperatures, near the “nose” of the 

transformation curve. Therefore, it is difficult to avoid phase transformations, such as σ 

precipitation, during the reheating of heavy section products (for example ingots, 

castings, thick plate, etc.), and so a solution treatment should be performed at a 

sufficiently high temperature to re-dissolve any such phases. On the other hand, during 

cooling, the slow nucleation rate at high temperature and the sluggish growth rate at 

lower temperatures make it relatively easy to avoid the formation of σ phase, even in 

the case of air cooling of certain castings of heavy plate. 

2.3 Characteristics and morphology of precipitates 


In welds of duplex stainless steels some detrimental second phases can occur in the 

temperature range of 300-1100˚C. The tendency to precipitation depends mainly on 

the content of alloying elements and therefore, it differs for the different duplex grades. 

The precipitation of phases is much more typical during heat treatment or welding of 

superduplex steels because of the higher content of such elements [4]. The following 

paragraphs describe the various phases which have been observed in duplex alloys. 

Their character and morphology of these phases vary considerably, as do the time for 

them to form and their influence over the properties. 

On fig.2 is given the diagram for the precipitation of the various phases for grade 

S32750. It can be easily seen what is the incubation time for the formation of a certain 

phase. Details for each phase can be observed in table1 

Fig.2: Time-temperature 

transformation diagram for alloy 

S32750. 

Particle Chemical formula Cr Ni Mo Formation range, ˚C Lattice type 

σ Fe-Cr-Mo 30 4 7 600-1000 tetragonal 

χ Fe 36 Cr 12 Mo 10 25 3 14 700-900 BCC- αMn 

α’ 65 2,5 13 300-525 BCC 

R Fe 2 Mo 25 6 35 550-650 trigonal 

π Fe 7 Mo 13 N 4 35 3 34 550-600 cubic 

τ 550-650 orthorhombic 

ε Cu-rich Not defined 

Type 1 same as ferrite ≤ 650 

γ 2 Type 2 24,3 11 3,4 650-800 FCC 

Type 3 700-900 

Cr 2 N 72 6 15 700-950 

CrN 

cubic 

M 7 C 3 950-1050 

M 23 C 6 58 2,5 12 650-950 FCC 

Table 1: Crystallographic characteristics of particles observed in duplex stainless 

steels 

2.3.1 Intermetallic phases 

• Sigma (σ) 


The deleterious Cr, Mo rich σ-phase is a hard embrittling precipitate, which forms 

between 650 and 1000˚C, often associated with a reduction in both impact properties 

and corrosion resistance. At the peak temperature of around 900˚C, ferrite 

decomposition to sigma takes as little as two minutes in superduplex alloys. According 

to some authors, the formation of sigma phase should be concerned with the 

requirement for pre-existing M 23 C 6 particles [5]. Certainly, this phase has been found to 

nucleate at temperatures above 750˚C in association with such particles with the 

following order of preference: δ/γ phase boundaries, austenitised δ/δ sub-grain 

boundaries and high energy δ/δ grain boundaries (table 1). These nuclei can grow into 

coarse plates, lamellar eutectoid σ + γ 2 (fig.3), or σ + δ lamellar aggregates. In the last 

case, the interlamellar ferritic region has a high dislocation density attributed to the 

volumetric expansion from δ to σ. Further, in the case of phase boundaries, for 

instance when δ transforms to γ or γ 2 , the remaining δ becomes enriched in Cr and 

Mo, and denuded in Ni, enhancing σ-formation and, for the same reason, growth 

progresses into the destabilized ferrite. 

Fig.3: SEM micrograph of σ + γ 2 

eutectoid. Steel S32750 after 72 

hours at 700˚C. 

The formation of sigma is encouraged by the presence of Cr, Mo, Si and Mn. Ni is also 

found to enhance σ-formation, but reduce the equilibrium volume fraction [6]. This 

occurs as Ni induces γ-formation and so concentrates the σ-promoting elements in the 

remaining ferrite. 

• Chi (χ) phase 

Like σ-phase, χ-phase forms between 700 and 900˚C, although in much smaller 

quantities. However, enrichment of ferrite with intermetallic forming elements during a 

long exposure to relatively low temperatures, i.e. 700˚C, favors the precipitation of χ- 

phase (фиг.2). Like sigma, χ-phase often forms on the δ/γ boundary and grows into 

the ferrite. The cube-cube orientation relationship (table 1) ensures continuity between 

χ and the δ-matrix. This phase has similar influence on corrosion and toughness 

properties as sigma, but, as both phases often co-exist, it is difficult to study their 

effects individually. 

• Alpha prime (α’) 

The lowest temperature decomposition within duplex steel is that of alpha prime (α’), 

which occurs between 300 and 525˚C, and is the main cause of hardening and “475 

embrittlement” in ferritic stainless steels. It is suggested that α’-formation is a 

consequence of the miscibility gap in the Fe-Cr system, whereby ferrite undergoes 

spinodal decomposition into Fe-rich δ-ferrite (table 1) and a Cr-rich α, or, just outside 

the spinodal but still within the gap, classical nucleation and growth of α’ occurs. Alpha 

prime is often associated with the co-precipitation of Cr 2 N in the form of sub-grain 

networks of Cr 2 N needles interspersed within a film of α’. 


• R, π, и τ phases 

R phase, also known as Laves, (Fe 2 Mo) precipitate in small quantities between 550 

and 650˚C after several hours exposure (table 1) [1]. They form at both intra- and 

intergranular sites, are molybdenum-rich and reduce pitting corrosion resistance. 

However, as those that precipitate at intergranular sites contain slightly more Mo (40% 

compared to 35%Mo), their influence on pitting resistance is more considerable. 

The π-nitride has been identified at intragranular sites in duplex weld metal after 

isothermal heat treatment at 600 ˚C for several hours [7]. It is Cr Mo rich and so has 

been previously confused with σ-phase. 

Heat treatment for several hours in the temperature range 550 to 650˚C (table 1) can 

lead to the formation of the heavily faulted needle-like τ-phase on δ/δ boundaries [8]. 

• Cu-rich epsilon (ε) phase 

In alloys containing copper and/or tungsten, other hardening mechanisms can occur. 

In the case of Cu, the super saturation of the ferrite due to the decrease in solubility at 

lower temperature leads to the precipitation of extremely fine Cu-rich ε-phase particles 

after 100 hours at 500˚C, which significantly extend the low temperature hardening 

range of the duplex grades. Although the reported temperature range for their 

formation varies, it would seem that they all form in the same temperature regime as 

γ 2 . 

2.3.2 Secondary austenite (γ 2 ) 

Secondary austenite can form relatively quickly and by various mechanisms depending 

on the temperature (table 1). About 650˚C, γ 2 has similar composition as the 

surrounding ferrite, suggesting a diffusionless transformation, with characteristics 

similar to martensite formation [9]. 

At temperature range 650 and 800˚C where diffusion is faster, many Widmanstaetten 

austenite forms can precipitate (fig.4). In this range, γ 2 obeys the Kurdjumov-Sachs 

relationship, its formation involves diffusion as it is enriched in Ni compared to the 

ferrite matrix (table 1). Even though there is some enrichment of N in γ 2 compared to 

the matrix, both Cr and N contents of γ 2 are substantially below that of primary 

austenite. 

Fig.4: Optical micrograph of γ 2 in 

superduplex welds, x1000. Etch: 

electrolytic sulphuric acid. 


In the 700-900˚C range, a eutectoid of γ 2 + σ can form (fig.3), as γ 2 absorbs and 

rejects Cr and Mo, encouraging Cr, Mo-rich precipitates, such as sigma phase. 

Similarly, one form of γ 2 which forms at δ/γ boundaries is found to be depleted in Cr. 

Either of these diffusion controlled reactions can render the area susceptible to pitting 

corrosion. 

2.3.3 Carbides M 23 C 6 and M 7 C 3 

M 7 C 3 forms between 950 and 1050˚C (fig.2) at the δ/γ grain boundaries. However, as 

its formation takes 10 minutes, it can be avoided by normal quenching techniques. 

Further, as modern duplex grades contain less that 0,02%C, carbides are rarely if ever 

seen [10]. 

In duplex grades with moderately high carbon levels, like S32750, of about 0,03%, the 

carbide M 23 C 6 rapidly precipitates between 650 and 950˚C (fig.2), requiring less that 1 

minute to form at 800˚C. Precipitation predominantly occurs at δ/γ boundaries, where 

Cr-rich ferrite intersects with carbon rich austenite. This type of carbide can be found 

also at the δ/δ and γ/γ boundaries and to a lesser degree inside he ferrite and 

austenite grains. Several precipitate morphologies have been recorded including 

cuboidal and aciculat particles, as well as a cellular form, although each type will have 

an associated Cr depleted zone in its vicinity. 

2.3.4 Nitrides Cr 2 N and CrN 

Nitrogen is added to duplex alloys to stabilize the austenite, and to improve strength 

and pitting resistance. The solubility of N is considerably higher in austenite than in 

ferrite, and has been shown to partition to the former phase. Above the solution 

annealing temperature (about 1040˚C), the volume fraction of ferrite increases, until 

just below the solidus a completely ferritic microstructure can be present, though in the 

in the higher alloy grades some austenite may remain. At these temperatures, the N 

stability in ferrite is high, but on cooling the solubility drops and the ferrite becomes 

supersaturated in N, leading to the intragranular precipitation of needle-like Cr 2 N. In a 

similar manner, Cr 2 N is most likely to form after higher solution heat treatment 

temperatures and forms rapidly even if quenched from such temperatures [11]. 

Welding process favors the formation of another nitride in the heat-affected zone: the 

cubic CrN, table.1. 

Isothermal exposure to the 750-900˚C temperature range (fig.2), produces 

intergranular Cr 2 N at δ/δ grain boundaries as thin plates on sub-grain boundaries, 

triple points and inclusions. The latter form of Cr 2 N has been stated to affect pitting 

corrosion [10]. 

2.4 Electron beam welding of duplex steels 

Normally duplex steels are weldable using welding procedures generally used for high 

alloyed steels. The experience with such comparatively new welding method like 

electron beam welding is still limited. However, there have been a few successful 


applications and there is every reason to expect that procedures will be developed 

more fully. Electron beam welding is especially suited to produce joints of heavy 

section materials in one or two passes. Unfortunately, it tends to produce rapid cooling 

rates and therefore highly ferrite in the melt zone, particularly in thin sections [12]. 

Nevertheless, the toughness remains high which can be attributed to the very low 

oxygen content in the weld. Still the qualification of the procedure must be alert to the 

possibility of excessive ferrite in the HAZ and even in the weld when the high speed 

welding capabilities of these methods are considered. 

Special feature of this welding method is the formation of the welding pool and bead. 

Because of the high temperature and concentration of energy, peculiar gas-steam 

channel in the welded metal is created. The electron beam continues to penetrate 

through this channel to bigger depth and thus it is possible thicknesses up to 200 mm 

to be fully penetrated with minimum width of the weld. This is a prerequisite for the 

formation of the so called “dagger” form of the weld (fig.5) 

Fig. 5: Formation of “dagger” form of the bead. 

In this case the ration between the penetration 

depth and the weld width is 3,6:1 (10:2,75), but it can 

much bigger (50:1). 

3. Experimental results 

3.1 Welding machine 

Installation LEYBOLD-HERAEUS EWS-1560 has been used for the welding. The 

sketch of the experimental configuration is shown on fig.6 

Fig.6 

1 – Devices for computer driven control; 

2 – Electron gun; 

3 – Source for acceleration; 

4 – Cathode feed; 

5, 6, 7 – Deflection/focus coils; 

8 – Leading; 

9 – Vacuum chamber. 


This electron beam machine is a conventional one. It is composed of an electron beam 

gun, a power supply, control system, motion equipment and vacuum welding chamber. 

The fusion of the base metals eliminates the need for filler material. Besides, the 

vacuum requirement for operation of the electron beam equipment eliminates the need 

for shielding gases and fluxes. 

The electron beam gun has a tungsten filament which is heated, freeing electrons. The 

electrons are accelerated from the source with high voltage potential between a 

cathode and anode. The stream of electrons then pass through a hole in the anode. 

The beam is directed by magnetic forces of focusing and deflecting coils. This beam is 

directed out of the gun column and strikes the workpiece. 

The potential energy of the electrons is transferred to heat upon impact of the 

workpiece and cuts a perfect hole at the weld joint. Molten metal fills in behind the 

beam, creating a deep finished weld. 

The electron beam stream and workpiece are manipulated by means of precise, 

computer driven controls, within a vacuum welding chamber of 0,5m 3 , and thus 

eliminating oxidation and contamination. 

3.2 Welding details 

Plates from superduplex stainless steel S32750 with thickness of 10mm have been 

butt welded with electron beam without filler material and in one pass. As a result, six 

welds are accomplished and for each of them the heat input is different. For the first 

three welds the parameter which values are being changed is the welding speed 

(mm/s) in order to observe its influence over the weld shape and the percentage of the 

ferrite phase in the weld and the heat-affected zone. On the contrary, the other plates 

are welded with one and the same welding speed, but with different current. The idea 

is to be found out what is the minimum value of the welding current that secures full 

penetration. The welding regimes are given in table 2. 

Weld No Voltage, kV Welding Current, mA Welding Speed, mm/s Heat Input, kJ/mm 

1 53 45 10 0,24 

2 53 45 15 0,16 

3 53 45 25 0,10 

4 53 52 15 0,18 

5 53 65 15 0,23 

6 53 75 15 0,27 

Table 2: Welding parameters 

3.3 Influence of the welding parameters over the shape of 

the welds 

For each weld macroscopic analysis has been done in order to be studied he influence 

of the welding parameters over the shape of the welds (penetration depth, 


einforcement height and the weld width). The etching is done by a mixture of 10 ml 

HNO 3 , 20 ml HCl and 30 ml H 2 O for 3-4 minutes depending on the temperature of the 

solution. 

Оn fig.7 is shown how the width of the weld (B) changes with the increase/decrease of 

the welding speed (V w ) and the current (I) accordingly with constant voltage (U=const) 

in both cases. 

U=const, I=const 

B(mm) 

V w(mm/s) 

1,1 25 

1,55 15 

2,75 10 

B(mm) 

3.0 

2.8 

2.6 

2.4 

2.2 

2.0 

1.8 

1.6 

1.4 

1.2 

1.0 

8 10 12 14 16 18 20 22 24 26 

V(mm/s) 

Fig.7: Influence of V w on the weld width 

U=const, V w=const 

B(mm) 

I(mA) 

1.75 52 

1,8 65 

2,0 75 

B (m m ) 

2.5 

2.4 

2.3 

2.2 

2.1 

2.0 

1.9 

1.8 

1.7 

1.6 

50 55 60 65 70 75 

I(mA) 

Fig.8: Influence of I on the weld width 

In the same way, the effect of changing the welding speed (V w ) and current (I) over the 

reinforcement height (H r ) and the penetration depth (H pen ) is studied and depicted on 

the following four fig.9-fig.12. The voltage again is constant – U=53 kV. 


H r (mm) 

V w(mm/s) 

1.2 25 

1,3 15 

1,8 10 


Hус(mm) 

1.8 

1.7 

1.6 

1.5 

1.4 

1.3 

1.2 

8 10 12 14 16 18 20 22 24 26 

Vз(mm/s) 

Fig.9: Influence of V w on the reinforcement height 


H r (mm) 

I(mA) 

1,45 52 

1,65 65 

1,70 75 

Hус(mm) 

1.70 

1.65 

1.60 

1.55 

1.50 

1.45 

50 55 60 65 70 75 

Iз(mA) 

Fig.10: Influence of I on the reinforcement height 


H pen (mm) 

V w(mm/s) 

5,1 25 

6,3 15 

10,0 10 

Hпр(mm) 

10 

9 

8 

7 

6 

5 

8 10 12 14 16 18 20 22 24 26 

Vз(mm/s) 

Fig.11: Influence of V w on the penetration depth 


H pen (mm) I(mA) 

7,15 52 

7,60 65 

Originally presented at the Stainless Steel World 2007 Conference, Maastricht, the Netherlands 

8,45 75

Hпр(mm) 

8.6 

8.4 

8.2 

8.0 

7.8 

7.6 

7.4 

7.2 

7.0 

50 55 60 65 70 75 

Iз(mA) 

Fig.12: Influence of I on the penetration depth 

The experiments have proved something which is known for a long time, namely that 

with the increase of the welding speed (it varies from 10 to 25 mm/s) under constant 

other circumstances, the values of all parameters pertaining to the shape of the weld 

diminish. The biggest influence has been exerted on the width of the weld – from 

2,75mm to 1,1mm. Certainly this is not a surprise, because the higher the speed is, the 

less the heat input is. At the same time the penetration depth also becomes much 

smaller – from full penetration to 5,1mm, i.e. two times decrease. The smallest 

modification refers to the reinforcement height – 40%. 

The effect of growing the welding current on the shape of the weld is just the opposite 

from what was said for the welding speed. With the increase of the current (from 52 to 

75mA), all above mentioned parameters grow bigger, as the ratio is almost the same 

for all of them. 

3.3 Ferrite content in the weld and the base metal 

It is well known that the duplex stainless steels solidify as ferrite and some of them 

transforms to austenite as the temperature falls to about 1000˚C depending on alloy 

composition. There is little further change in the ferrite-austenite balance at lower 

temperatures. In case of welding the metallurgical processes are less equilibrium than 

during heat treatment, because the heating, respectively the cooling rate are very high 

and there is no enough time for the diffusion to make the composition uniform 

throughout the whole volume. 

Bearing in mind how decisive the ferrite content is for the mechanical and corrosion 

properties of the duplex grades, it was measured for all six welds. Conclusions are 

made about the influence of the welding regime (welding parameters) over this 

criterion, therefore over the properties of the steels in question. 

For the measurement of the ferrite content a ferritoscope ‘Foerster-1.053’ has been 

used. An additional graduation is made in order to be increased the measuring scope 

of the apparatus. After changing the distance between the measuring inductive drill 

and the surface of the examined sample by means of a non-metallic folio with 0,5mm 

thickness, the power of the measured signal has been altered and the possibility for 


the measurement of the ferrite content has been increased to 100%. Thus the scope is 

raised, but the sensibility of the apparatus becomes smaller. 

The results from the measurements of the ferrite phase in the welds are ordered in 

table 3. 

Weld No 

Ferrite content in the welds, % 

Measurement results 

Mean value 

1 63, 65, 57, 60 63,75 

2 65, 60, 66, 62 63,25 

3 50, 55 52,50 

4 58, 60, 58, 53 57,25 

5 57, 55, 60, 58 57,50 

6 58, 50 54,00 

Table 3: Ferrite content in the welds 

The ferrite content at the fusion boundaries is given in table 4. 

Fusion boundary Ferrite content at the fusion boundaries, % 

for Weld No Measurement results Mean value 

1 43, 48, 47, 45, 46, 48, 43, 45, 44 45,44 

2 45, 44, 46, 48, 43, 45, 44, 43, 38 44,00 

3 50, 48, 50, 53, 49, 48, 45 49,00 

4 47, 47, 43, 45, 43, 45, 48, 48, 44 45,55 

5 44, 44, 38, 45, 46, 53, 45, 45 45,00 

6 44, 47, 46 45,67 

Table 4: Ferrite content at the fusion boundaries 

The relevant values of this phase in the base metal are presented in the following table 

5. 

Ferrite content in the base metal, % 

Measurement results 

Mean value 

38, 37, 40, 45, 42, 44, 38, 41 40,63 

Table 5: Ferrite content in the base metal 

The difference in the ferrite content pertaining to the welds, the fusion lines and the 

base metal is obvious. It is most in the welds, because the cooling rate is highest in 

these zones and the transformation from ferrite to austenite is impeded at low 

temperatures. As a result, more quantities of high-temperature ferrite can be observed 

in the structure. On the other hand the initial balance between both phases is disrupted 

also in the heat affected zones (which for the electron beam welding are very thin 

zones, like a line), but not to this extent like in the welds. That is why the mean value of 

the ferrite content in the HAZ is smaller than in the weld, but still is bigger in 

comparison with the base metal. 


3.4 Metallographic analysis 

Metallographic analysis on preliminary prepared five sections from duplex steel 

S32750 has been performed. For that purpose an optical microscope Leica DM6000M 

which allows optical magnification x2000 has been used. The sections are prepared 

(coarse and fine grinding with subsequent polishing) with Struers. The testing has 

been done in accordance with BDS EN 3690 (Visual-optical Methods) and BS EN 

1321 (Normal and Special Metallographic Analysis). 

Object of analysis for every section has been the weld, the fusion boundary and the 

base metal. 

3.4.1 Microstructures 

• Weld 1 

a b c 

d 

e 

Fig.13: Microstructure in the weld under different magnifications: a-c) x200; d) x500; e) 

x1000 

a 

b 


Fig.14: Microstructure at the boundary between the weld and the base metal under 

different magnifications: a) x200; b) x500 

• Weld 2 

a 

b 

Fig.15: Microstructure in the weld under different magnifications: a-c) x200; b) x1000; 

c) x1500 

a b c 


different magnifications: a) x500; b-c) x1000 

Fig.17: Microstructure in the base metal x1000 

• Weld 3 


a b c 

Fig.18: Microstructure in the weld under different magnifications: a) x750; b) x1000; c) 

x2000 

a b c 


different magnifications: a) x200; b) x1000; c) x2000 


• Weld 4 

a b c 


x2000 


a b c 




• Weld 5 

a b c 


x2000 

a b c 




• Weld 6 

a b c 

Fig.26: Microstructure in the weld under different magnifications: a) x500; b-c) x1000 

a b c 



3.4.2 Results 

• For all probes (welded with different regimes of electron beam welding) a typical 

structure with five distinguished zones can be observed (fig.28): 

Zone1 

1 

Zone2 Zone3 Zone4 Zone5 

- Zone 1 – coarse grains in the middle 

part of the weld; 

- Zone 2 – strongly prolonged through 

the front of the grains; 

- Zone 3 – dispersive structure at the 

boundary with the base metal; 

- Zone 4 – fusion boundary; 

- Zone 5 – base metal. 

Originally presented at the Stainless Steel World 2007 Conference, Maastricht, the Netherlands 

Fig.28: Microstructure at the 

boundary weld – base metal, х200

• The base metal has ferrite-austenite structure with ratio between both phases close 

to 50%-50% and with a texture of the grains (fig.29). Non-metallic inclusions are 

not outlined in the section field. 

a 

b 

Fig.29: Microstructure of base metal: a) x200 (weld 3); b) х1000 (weld 4) 

• The weld structure is non-equal and consists of grains with different sizes, which is 

a result of the different speed of heat-conduction in the material. In all welds the 

formation of skeleton-like structure can be seen (fig.30). Under big metallographic 

magnifications (х1000, х2000) the structures which set up this skeleton are clearly 

discernable. On the periphery of the grains (1) white zones can be observed (2) 

and at the boundaries – needle-shaped precipitations (3). 

1 

2 

3 

а 

b 

1 

2 

3 

c 

d 

Fig.30: Microstructure in weld, х1000: 


а) weld 5; b) weld 4; c), d) weld 1 

Main phase component in the welds is δ-ferrite. Inside the δ-ferrite grains 

precipitations of γ 2 are registered and at the boundaries – austenite, which has 

particularly a Widmannstaeten structure. Intermetallic phases like σ-type are not 

registered upon using metallographic analysis. 

• The formed heat affected zone (zone 4) is very thin (10-35µm) and is visible as a 

boundary between the weld and the base metal. The microstructure in this zone is 

identical as the one in the base metal. Intermetallic phases are not registered. 

• In most of the welds some typical for electron beam welding imperfections in the 

root are visible – porosity, non-uniform surface, etc. 


4. Conclusions 

• With the increase of the welding speed all parameters which pertain to the 

shape of the weld, namely penetration depth, weld width and reinforcement 

height, become smaller because of the low heat input. This speed should be 

chosen very carefully, because there is a risk of lack of penetration if it is too 

high. 

• The influence of the welding current over the shape of the weld is the following 

– the higher the current is, the bigger are the geometrical dimensions of the 

weld. 

• The ferrite content in the weld is bigger that the relevant one in the heataffected 

zone and the base metal, because the structure is not equilibrium in 

the weld. In comparison with the heat treatment, the welding processes are 

faster, the cooling rate is higher and the time for diffusion is shorter. As a 

result of this, high-temperature condition is fixed at lower temperatures. 

• The ferrite content in the weld is more when the heat input is less, for example 

when the welding speed is higher or the current is lower than some initial 

values. 

• The heat-affected zone in the materials, which are welded with electron beam 

is very thin, of the order of some µm. This is an advantage for this welding 

method, because very small part of the base metal is affected by the welding 

process. 

• The measurement of the ferrite in the base metal has shown that the volume 

of this phase is the range which is given in the literature – between 40% and 

60% in annealed and quenched condition. 

• The initial position of the plates is annealed to 1050˚C and quenched in order 

to be re-dissolved all second phases in the structure. In comparison with the 

ferritic stainless steels, duplex alloys are famous with the slower diffusion rate, 

i.e. the incubation time for the phase precipitation is longer. This means that 

the “nose” of the curve is translated to the right. 

• The review of the metallographic photos has not shown the precipitation of 

intermetallic phases. The explanation of this fact can be found in the timetemperature 

transformation diagram. According to this diagram the incubation 

time is enough long to secure that no phases will precipitate. Therefore 

electron beam welding is proper for welding of duplex stainless steel S32750 

for the lack of second phases in the microstructure which could affect 

embrittlement and reduction in the corrosion resistance. 

• The more alloyed is one duplex grade, the more susceptible it is to the 

precipitation of second phases. That is to say the standard duplex stainless 

grades are more susceptible to the precipitation than the lean ones. The 

chance these phases to form in the high alloyed grades are much bigger than 

in the standard steels. The superduplex alloys show the greatest propensity 

for precipitations, due to their higher Cr, Mo and W contents. 


5. References 

[1] Charles J: Proc conf Duplex Stainless Steels ’91, Beaune, 1991, Vol.1, 3-48. 

[2] Mancia F, Barteri M, Sasseth L, Tamba A, Lannaioli A: York ’87, vide ref.11, 

160-167. 

[3] Gunn R: Duplex Stainless Steels, Woodhead Publishing, 2003. 

[4] Leif Karlson: Intermetallic Phase Precipitation in Duplex Stainless Steels and 

Weld Metals. Metallurgy, Influence on Properties, Welding and Testing Aspects. Doc. 

IX-1920-98. 

[5] Goldsmith HJ: Interstitial Alloys, Plenum Press, 1967, 167. 

[6] Maehara Y, Ohmori Y, Murayama J, Fujino N: Metal Science 17, 1983, 541. 

[7] Nilsson J-O, Liu P: Mater Sci Technol 7, 1991, 853. 

[8] Redjaimia A, Metauer G, Gantois M: Beaune ’91, vide ref.2, Vol.1, 119-126. 

[9] Soulignac P, Dupoiron F: Stainless Steel Europe 2, 1990, 18-21. 

[10] Nilsson J-O: Materials Science and Technology 8, 1992, 685-700. 

[11] Herzman S, Roberts W, Lindenmo M: The Hague ’86, vide ref. 7, paper 30, 

257-267. 

[12] Bonnefois B, Charles J, Dupoiron F, Soulignac P: Proc conf Duplex Stainless 

Steels ’91, Beaune, France, Oct. 1991, Vol.1, 347-362;

Electron Beam Welding of Superduplex Stainless Steel S32750

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