electron beam welding of superduplex stainless steel.pdf

7/29/2019 Electron beam welding of superduplex stainless steel.pdf

1/24

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

Electron Beam Welding of SuperduplexStainless Steel S32750

Presenter:

Parvan ChavdarovAcademic Education and DegreesPost-graduate student Weldability of duplex stainless steels Technical University of Sofia2001 Master degree on technology of metals, TechnicalUniversity of Sofia, BulgariaPresent professional position:Expert in Industrial Services in TUV Rheinland Bulgaria


2/24


Electron Beam Welding of SuperduplexStainless Steel S32750

Authors: Read.Dr Serafim Serafimov Technical University of SofiaDipl.eng.Parvan Chavdarov post-graduate student TUV Rheinland SofiaAdress: Zona B-19, Dr Kalinkov str.18, app.10, Sofia, BulgariaPost Code 1309, tel. +359 898529051

e-mail: [email protected]

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 theshape of the weld, the microstructure and respectively the precipitation of secondphases in superduplex stainless steel S32750. Plates with 10mm thickness have beenbutt 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 beenmeasured. The availability of intermetallic phases is examined.


3/24


1. Introduction

Duplex stainless steels are a family of grades combining good corrosion resistancewith high strength and ease of fabrication. Their physical properties are between thoseof 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 becauseof their advantages:

Very high corrosion resistance especially against pitting and crevice corrosion inaggressive media containing chlorides and fluorides;

Higher mechanical properties (because of the ferrite phase) than the austenitestainless steels which are predominantly used in practice;

Higher ductility (because of the austenitic phase) than the ferrite stainless steelswhich reflects on the weldability.

Better compatibility to C-Mn structural steels, because the coefficients of heatexpansion are comparatively close, which results in lower thermal stresses.

All these advantages result in extremely wider usage of duplex alloys in many industrybranches 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 whichincreases its cost-effectiveness considerably. The desire structure is obtained by heattreatment at approximately 1050 to 1100C (solution annealing). The optimum ratiobetween 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, whichconsists of a mixture of about 50% volume austenite in ferrite grains. Both cast andwrought products have roughly equivalent volume fractions of ferrite and austenite,which in the case of wrought components, contain a rolling texture obtained by hotworking, followed by a solution annealing and quench. The optimum phase balance formodern 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. Thephase proportions and their respective compositions are indicated for a given alloyanalysis and annealing temperature, with the high temperature stability of the duplexstructure being influenced more by nitrogen content, than by Cr or Mo. The addition of0,25%N to a 25%Cr alloy produces a ferrite volume fraction of about 50% at 1250C,compared to nearly 80% ferrite with 0,18%N. Nevertheless, it is difficult to predict themicrostructure of a duplex alloy from simplified diagrams, due to the effects of otheralloy elements, which modify the phase fields.


4/24


Fig.1: Concentration profiles inthe ternary Fe-Cr-Ni constitutiondiagram at 70% and 60%Fe. Theeffect of nitrogen additions isshown in the first figure.

2.2 Heat treatment

Numerous structural changes can occur in the duplex stainless steels duringisothermal and anisothermal heat treatments. Most of these transformations areconcerned with the ferrite, as element diffusion rates are approximately 100 timesfaster than in austenite. This is principally a consequence of the less compact lattice ofthe BCC crystal structure. Moreover, the ferrite is enriched in Cr and Mo, whichpromote the formation of Intermetallic phases. Besides, element solubility in the ferritefalls with a decrease in temperature, increasing the possibility of precipitation duringheat 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 affectprecipitation kinetics and the stability of phases formed [1].

2.2.1 Temperatures above 1050C

Duplex stainless steels solidify completely in the ferrite field for standard grades andnormal 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 examplefrom 1050C to 1300C, leads to an increase in ferrite content. Further, as hetemperature increases, there is a reduction in the portioning of substantial elementsbetween the phases. In addition, the ferrite becomes enriched in interstitial elementssuch as carbon and nitrogen.

Heat treatment in the temperature range 1100-1200C can have a dramatic influenceon the microstructure of a wrought product. The grains can be made equiaxed byprolonged treatment at high temperature or can be rendered acicular, with aWidmannstaetten type structure by cooling an intermediate rate. A dual structure,consisting of both coarse and fine austenite grains, can be obtained by stepquenching, with or without simultaneous mechanical strain. These acicular structuresare also encountered in weld deposits.

2.2.2 The 600-1050C nose


5/24


The alloy lean grades are the least prone to Intermetallic phase precipitation andrequires exposure of at least 10 to 20 hours to initiate formation at temperatures below900C. For that reason, a solution annealing temperature below 1000C can be chosenfor this material.

More alloyed grades are more sensitive to precipitation due to the molybdenumcontent. This element not only increases the rate of Intermetallic precipitation, but alsoextents the stability range to higher temperatures. That is why higher solutionannealing temperatures are needed above 1000C.

The superduplex alloys show the greatest propensity for precipitations, due to theirhigher Cr, Mo and W contents. However, it should be emphasized that the precipitationkinetics in these high alloy grades are, at worst, equivalent to the superaustenitic orsuperferritic stainless steels [2]. And still, by taking precautions during heat treatment,including rapid removal from the furnace followed by water quenching, the superduplexalloys can be used satisfactory in industrial applications. However, especial care isrequired for heavy section components during all stages of production.

Precipitates re-dissolve during a solution anneal, which for the superduplex gradesmust be performed at 1050C or above. A few minutes at 1050-1070C are sufficientfor grade S32750. Similar high temperatures are necessary for welds as consumablestend to contain higher Ni, Si and Mn contents than base materials. The higher Ni-content encourages high austenite contents when annealed and results in enrichmentof 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(1040C compared to 1100C) can be used for weldments made with matchingconsumables [3].

2.2.3 The 300-600C 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 400C. A muchshorter incubation time is found for more alloyed grades, containing molybdenum,which would appear to accelerate hardening [1]. The 25%Cr and superduplex alloysshow the widest temperature range for hardening and shorter incubation timesaccordingly. 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 thetransformation curve. Therefore, it is difficult to avoid phase transformations, such as precipitation, during the reheating of heavy section products (for example ingots,castings, thick plate, etc.), and so a solution treatment should be performed at asufficiently high temperature to re-dissolve any such phases. On the other hand, duringcooling, the slow nucleation rate at high temperature and the sluggish growth rate atlower temperatures make it relatively easy to avoid the formation of phase, even inthe case of air cooling of certain castings of heavy plate.

2.3 Characteristics and morphology of precipitates


6/24


In welds of duplex stainless steels some detrimental second phases can occur in thetemperature range of 300-1100C. The tendency to precipitation depends mainly onthe 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 ofsuperduplex steels because of the higher content of such elements [4]. The followingparagraphs describe the various phases which have been observed in duplex alloys.Their character and morphology of these phases vary considerably, as do the time forthem to form and their influence over the properties.

On fig.2 is given the diagram for the precipitation of the various phases for gradeS32750. It can be easily seen what is the incubation time for the formation of a certainphase. Details for each phase can be observed in table1

Fig.2: Time-temperature

transformation diagram for alloyS32750.

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

Fe-Cr-Mo 30 4 7 600-1000 tetragonal Fe36Cr12Mo10

25 3 14 700-900 BCC- Mn 65 2,5 13 300-525 BCCR Fe2Mo 25 6 35 550-650 trigonal

Fe7Mo13N4 35 3 34 550-600 cubic

550-650 orthorhombic

Cu-rich Not definedType 1 same as ferrite 650

Type 2 24,3 11 3,4 650-800 FCC2

Type 3 700-900

Cr2N 72 6 15 700-950CrN cubic

M7C3 950-1050M23C6 58 2,5 12 650-950 FCC

Table 1: Crystallographic characteristics of particles observed in duplex stainlesssteels

2.3.1 Intermetallic phases

Sigma ()


7/24


The deleterious Cr, Mo rich -phase is a hard embrittling precipitate, which formsbetween 650 and 1000C, often associated with a reduction in both impact propertiesand corrosion resistance. At the peak temperature of around 900C, ferritedecomposition to sigma takes as little as two minutes in superduplex alloys. Accordingto some authors, the formation of sigma phase should be concerned with therequirement for pre-existing M23C6 particles [5]. Certainly, this phase has been found tonucleate at temperatures above 750C in association with such particles with thefollowing order of preference: / phase boundaries, austenitised / sub-grain

boundaries and high energy / grain boundaries (table 1). These nuclei can grow intocoarse plates, lamellar eutectoid + 2 (fig.3), or + lamellar aggregates. In the lastcase, the interlamellar ferritic region has a high dislocation density attributed to thevolumetric expansion from to . Further, in the case of phase boundaries, forinstance when transforms to or 2, the remaining becomes enriched in Cr andMo, and denuded in Ni, enhancing -formation and, for the same reason, growthprogresses into the destabilized ferrite.

Fig.3: SEM micrograph of + 2eutectoid. Steel S32750 after 72hours at 700C.

The formation of sigma is encouraged by the presence of Cr, Mo, Si and Mn. Ni is alsofound to enhance -formation, but reduce the equilibrium volume fraction [6]. Thisoccurs as Ni induces -formation and so concentrates the -promoting elements in theremaining ferrite.

Chi () phase

Like -phase, -phase forms between 700 and 900C, although in much smallerquantities. However, enrichment of ferrite with intermetallic forming elements during a

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

phase (.2). Like sigma, -phase often forms on the / boundary and grows intothe ferrite. The cube-cube orientation relationship (table 1) ensures continuity between

and the -matrix. This phase has similar influence on corrosion and toughnessproperties as sigma, but, as both phases often co-exist, it is difficult to study theireffects individually.

Alpha prime ()

The lowest temperature decomposition within duplex steel is that of alpha prime (),which occurs between 300 and 525C, and is the main cause of hardening and 475embrittlement in ferritic stainless steels. It is suggested that -formation is aconsequence of the miscibility gap in the Fe-Cr system, whereby ferrite undergoesspinodal decomposition into Fe-rich -ferrite (table 1) and a Cr-rich , or, just outsidethe spinodal but still within the gap, classical nucleation and growth of occurs. Alphaprime is often associated with the co-precipitation of Cr2N in the form of sub-grainnetworks of Cr2N needles interspersed within a film of .


8/24


R, , phases

R phase, also known as Laves, (Fe2Mo) precipitate in small quantities between 550and 650C after several hours exposure (table 1) [1]. They form at both intra- andintergranular 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 -nitride has been identified at intragranular sites in duplex weld metal afterisothermal heat treatment at 600 C for several hours [7]. It is Cr Mo rich and so hasbeen previously confused with -phase.

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

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

Cu-rich epsilon () 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 atlower temperature leads to the precipitation of extremely fine Cu-rich -phase particlesafter 100 hours at 500C, which significantly extend the low temperature hardeningrange of the duplex grades. Although the reported temperature range for theirformation varies, it would seem that they all form in the same temperature regime as2.

2.3.2 Secondary austenite (2)

Secondary austenite can form relatively quickly and by various mechanisms dependingon the temperature (table 1). About 650C, 2 has similar composition as the

surrounding ferrite, suggesting a diffusionless transformation, with characteristicssimilar to martensite formation [9].

At temperature range 650 and 800C where diffusion is faster, many Widmanstaettenaustenite forms can precipitate (fig.4). In this range, 2 obeys the Kurdjumov-Sachsrelationship, its formation involves diffusion as it is enriched in Ni compared to theferrite matrix (table 1). Even though there is some enrichment of N in 2 compared tothe matrix, both Cr and N contents of 2 are substantially below that of primaryaustenite.

Fig.4: Optical micrograph of 2 insuperduplex welds, x1000. Etch:electrolytic sulphuric acid.


9/24


In the 700-900C range, a eutectoid of 2 + can form (fig.3), as 2 absorbs andrejects Cr and Mo, encouraging Cr, Mo-rich precipitates, such as sigma phase.Similarly, one form of 2 which forms at / 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 M23C6 and M7C3

M7C3 forms between 950 and 1050C (fig.2) at the / grain boundaries. However, asits 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 everseen [10].

In duplex grades with moderately high carbon levels, like S32750, of about 0,03%, thecarbide M23C6 rapidly precipitates between 650 and 950C (fig.2), requiring less that 1

minute to form at 800C. Precipitation predominantly occurs at / boundaries, whereCr-rich ferrite intersects with carbon rich austenite. This type of carbide can be foundalso at the / and / boundaries and to a lesser degree inside he ferrite andaustenite grains. Several precipitate morphologies have been recorded includingcuboidal and aciculat particles, as well as a cellular form, although each type will havean associated Cr depleted zone in its vicinity.

2.3.4 Nitrides Cr2N and CrN

Nitrogen is added to duplex alloys to stabilize the austenite, and to improve strengthand pitting resistance. The solubility of N is considerably higher in austenite than inferrite, and has been shown to partition to the former phase. Above the solution

annealing temperature (about 1040C), the volume fraction of ferrite increases, untiljust below the solidus a completely ferritic microstructure can be present, though in thein the higher alloy grades some austenite may remain. At these temperatures, the Nstability in ferrite is high, but on cooling the solubility drops and the ferrite becomessupersaturated in N, leading to the intragranular precipitation of needle-like Cr2N. In asimilar manner, Cr2N is most likely to form after higher solution heat treatmenttemperatures and forms rapidly even if quenched from such temperatures [11].Welding process favors the formation of another nitride in the heat-affected zone: thecubic CrN, table.1.

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

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

triple points and inclusions. The latter form of Cr2N has been stated to affect pittingcorrosion [10].

2.4 Electron beam welding of duplex steels

Normally duplex steels are weldable using welding procedures generally used for highalloyed steels. The experience with such comparatively new welding method likeelectron beam welding is still limited. However, there have been a few successful


10/24


applications and there is every reason to expect that procedures will be developedmore fully. Electron beam welding is especially suited to produce joints of heavysection materials in one or two passes. Unfortunately, it tends to produce rapid coolingrates 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 lowoxygen content in the weld. Still the qualification of the procedure must be alert to thepossibility of excessive ferrite in the HAZ and even in the weld when the high speedwelding 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-steamchannel in the welded metal is created. The electron beam continues to penetratethrough this channel to bigger depth and thus it is possible thicknesses up to 200 mmto be fully penetrated with minimum width of the weld. This is a prerequisite for theformation 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 penetrationdepth and the weld width is 3,6:1 (10:2,75), but it canmuch bigger (50:1).

3. Experimental results

3.1 Welding machine

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


11/24


This electron beam machine is a conventional one. It is composed of an electron beamgun, 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 needfor shielding gases and fluxes.

The electron beam gun has a tungsten filament which is heated, freeing electrons. Theelectrons are accelerated from the source with high voltage potential between acathode 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 isdirected out of the gun column and strikes the workpiece.

The potential energy of the electrons is transferred to heat upon impact of theworkpiece and cuts a perfect hole at the weld joint. Molten metal fills in behind thebeam, 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 0,5m3, and thuseliminating oxidationand contamination.

3.2 Welding details

Plates from superduplex stainless steel S32750 with thickness of 10mm have beenbutt welded with electron beam without filler material and in one pass. As a result, sixwelds are accomplished and for each of them the heat input is different. For the firstthree 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 theferrite phase in the weld and the heat-affected zone. On the contrary, the other platesare welded with one and the same welding speed, but with different current. The ideais to be found out what is the minimum value of the welding current that secures fullpenetration. 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 0,185 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 ofthe welds

For each weld macroscopic analysis has been done in order to be studied he influenceof the welding parameters over the shape of the welds (penetration depth,


12/24


reinforcement height and the weld width). The etching is done by a mixture of 10 mlHNO3, 20 ml HCl and 30 ml H2O for 3-4 minutes depending on the temperature of thesolution.

n fig.7 is shown how the width of the weld (B) changes with the increase/decrease ofthe welding speed (Vw) and the current (I) accordingly with constant voltage (U=const)in both cases.

8 10 12 14 16 18 20 22 24 26

1.0

1.2

1.4

1.6

1.8

2.0

2.2

2.4

2.6

2.8

3.0

B(mm)

V(mm/s) Fig.7: Influence of Vw on the weld width

50 55 60 65 70 75

1.6

1.7

1.8

1.9

2.0

2.1

2.2

2.3

2.4

2.5

B(m

m

)

I(mA)

Fig.8: Influence of Ion the weld width

In the same way, the effect of changing the welding speed (Vw) and current (I) over thereinforcement height (Hr) and the penetration depth (Hpen) is studied and depicted onthe following four fig.9-fig.12. The voltage again is constant U=53 kV.

U=const, I=const

B(mm) Vw(mm/s)

1,1 25

1,55 15

2,75 10

U=const, Vw=const

B(mm) I(mA)

1.75 52

1,8 65

2,0 75

U=const, I=const

Hr(mm) Vw(mm/s)

1.2 25

1,3 15

1,8 10


13/24


8 10 12 14 16 18 20 22 24 26

1.2

1.3

1.4

1.5

1.6

1.7

1.8

H(mm)

V(mm/s)

Fig.9: Influence of Vw on the reinforcement height

50 55 60 65 70 75

1.45

1.50

1.55

1.60

1.65

1.70

H(mm)

I(mA)

Fig.10: Influence of Ion the reinforcement height

8 10 12 14 16 18 20 22 24 26

5

6

7

8

9

10

H(mm)

V(mm/s)

Fig.11: Influence of Vw on the penetration depth

U=const, Vw=const

Hr(mm) I(mA)

1,45 521,65 65

1,70 75

U=const, I=const

Hpen(mm) Vw(mm/s)

5,1 25

6,3 15

10,0 10

U=const, Vw=const

Hpen(mm) I(mA)

7,15 52

7,60 65

8,45 75


14/24


50 55 60 65 70 75

7.0

7.2

7.4

7.6

7.8

8.0

8.2

8.4

8.6

H(mm)

I(mA)

Fig.12: Influence of Ion the penetration depth

The experiments have proved something which is known for a long time, namely thatwith the increase of the welding speed (it varies from 10 to 25 mm/s) under constantother circumstances, the values of all parameters pertaining to the shape of the welddiminish. The biggest influence has been exerted on the width of the weld from2,75mm to 1,1mm. Certainly this is not a surprise, because the higher the speed is, theless the heat input is. At the same time the penetration depth also becomes muchsmaller from full penetration to 5,1mm, i.e. two times decrease. The smallestmodification refers to the reinforcement height 40%.

The effect of growing the welding current on the shape of the weld is just the oppositefrom what was said for the welding speed. With the increase of the current (from 52 to75mA), all above mentioned parameters grow bigger, as the ratio is almost the samefor 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 themtransforms to austenite as the temperature falls to about 1000C depending on alloycomposition. There is little further change in the ferrite-austenite balance at lowertemperatures. In case of welding the metallurgical processes are less equilibrium thanduring heat treatment, because the heating, respectively the cooling rate are very highand there is no enough time for the diffusion to make the composition uniformthroughout the whole volume.

Bearing in mind how decisive the ferrite content is for the mechanical and corrosionproperties of the duplex grades, it was measured for all six welds. Conclusions aremade about the influence of the welding regime (welding parameters) over thiscriterion, therefore over the properties of the steels in question.

For the measurement of the ferrite content a ferritoscope Foerster-1.053 has beenused. An additional graduation is made in order to be increased the measuring scopeof the apparatus. After changing the distance between the measuring inductive drilland the surface of the examined sample by means of a non-metallic folio with 0,5mmthickness, the power of the measured signal has been altered and the possibility for


15/24


the measurement of the ferrite content has been increased to 100%. Thus the scope israised, but the sensibility of the apparatus becomes smaller.

The results from the measurements of the ferrite phase in the welds are ordered intable 3.

Ferrite content in the welds, %Weld NoMeasurement results Mean value

1 63, 65, 57, 60 63,752 65, 60, 66, 62 63,253 50, 55 52,504 58, 60, 58, 53 57,255 57, 55, 60, 58 57,50

6 58, 50 54,00Table 3: Ferrite content in the welds

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

Ferrite content at the fusion boundaries, %Fusion boundaryfor Weld No Measurement results Mean value

1 43, 48, 47, 45, 46, 48, 43, 45, 44 45,442 45, 44, 46, 48, 43, 45, 44, 43, 38 44,003 50, 48, 50, 53, 49, 48, 45 49,00

4 47, 47, 43, 45, 43, 45, 48, 48, 44 45,555 44, 44, 38, 45, 46, 53, 45, 45 45,006 44, 47, 46 45,67Table 4: Ferrite content at the fusion boundaries

The relevant values of this phase in the base metal are presented in the following table5.

Ferrite content in the base metal, %Measurement results Mean value

38, 37, 40, 45, 42, 44, 38, 41 40,63Table 5: Ferrite content in the base metal

The difference in the ferrite content pertaining to the welds, the fusion lines and thebase 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 lowtemperatures. As a result, more quantities of high-temperature ferrite can be observedin the structure. On the other hand the initial balance between both phases is disruptedalso in the heat affected zones (which for the electron beam welding are very thinzones, like a line), but not to this extent like in the welds. That is why the mean value ofthe ferrite content in the HAZ is smaller than in the weld, but still is bigger incomparison with the base metal.


16/24


3.4 Metallographic analysis

Metallographic analysis on preliminary prepared five sections from duplex steelS32750 has been performed. For that purpose an optical microscope Leica DM6000Mwhich allows optical magnification x2000 has been used. The sections are prepared(coarse and fine grinding with subsequent polishing) with Struers. The testing hasbeen 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 thebase 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


17/24


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

Fig.16: Microstructure at the boundary between the weld and the base metal underdifferent magnifications: a) x500; b-c) x1000

Fig.17: Microstructure in the base metal x1000

Weld 3


18/24


a b c

Fig.18: Microstructure in the weld under different magnifications: a) x750; b) x1000; c)x2000

a b c

Fig.19: Microstructure at the boundary between the weld and the base metal underdifferent magnifications: a) x200; b) x1000; c) x2000


Weld 4

a b c



19/24


a b c



Weld 5

a b c


a b c


20/24



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

- Zone 1 coarse grains in the middlepart of the weld;

- Zone 2 strongly prolonged throughthe front of the grains;

- Zone 3 dispersive structure at theboundary with the base metal;

- Zone 4 fusion boundary;- Zone 5 base metal.

Zone1 Zone2 Zone3 Zone4 Zone5

Fig.28:Microstructure at theboundary weld base metal, 200


21/24


The base metal has ferrite-austenite structure with ratio between both phases closeto 50%-50% and with a texture of the grains (fig.29). Non-metallic inclusions arenot outlined in the section field.

a b

Fig.29: Microstructure of base metal:a) x200 (weld 3); b) 1000 (weld 4)

The weld structure is non-equal and consists of grains with different sizes, which isa result of the different speed of heat-conduction in the material. In all welds theformation of skeleton-like structure can be seen (fig.30). Under big metallographicmagnifications (1000, 2000)the structures which set up this skeleton are clearlydiscernable. On the periphery of the grains (1) white zones can be observed (2)and at the boundaries needle-shaped precipitations (3).

b

c d

Fig.30: Microstructure in weld, 1000:

1

2

3

1

2

3


22/24


) weld 5; b) weld 4; c), d) weld 1

Main phase component in the welds is -ferrite. Inside the -ferrite grainsprecipitations of 2 are registered and at the boundaries austenite, which hasparticularly a Widmannstaeten structure. Intermetallic phases like -type are notregistered upon using metallographic analysis.

The formed heat affected zone (zone 4) is very thin (10-35m) and is visible as aboundary between the weld and the base metal. The microstructure in this zone isidentical 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 theroot are visible porosity, non-uniform surface, etc.


23/24


4. Conclusions

With the increase of the welding speed all parameters which pertain to theshape of the weld, namely penetration depth, weld width and reinforcementheight, 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 toohigh.

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 theweld.

The ferrite content in the weld is bigger that the relevant one in the heat-affected zone and the base metal, because the structure is not equilibrium inthe weld. In comparison with the heat treatment, the welding processes arefaster, the cooling rate is higher and the time for diffusion is shorter. As aresult 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 examplewhen 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 m. This is an advantage for this weldingmethod, because very small part of the base metal is affected by the weldingprocess.

The measurement of the ferrite in the base metal has shown that the volumeof this phase is the range which is given in the literature between 40% and60% in annealed and quenched condition.

The initial position of the plates is annealed to 1050C and quenched in orderto be re-dissolved all second phases in the structure. In comparison with theferritic stainless steels, duplex alloys are famous with the slower diffusion rate,i.e. the incubation time for the phase precipitation is longer. This means thatthe nose of the curve is translated to the right.

The review of the metallographic photos has not shown the precipitation ofintermetallic phases. The explanation of this fact can be found in the time-temperature transformation diagram. According to this diagram the incubationtime is enough long to secure that no phases will precipitate. Thereforeelectron beam welding is proper for welding of duplex stainless steel S32750for the lack of second phases in the microstructure which could affectembrittlement and reduction in the corrosion resistance.

The more alloyed is one duplex grade, the more susceptible it is to theprecipitation of second phases. That is to say the standard duplex stainlessgrades are more susceptible to the precipitation than the lean ones. The

chance these phases to form in the high alloyed grades are much bigger thanin the standard steels. The superduplex alloys show the greatest propensityfor precipitations, due to their higher Cr, Mo and W contents.


24/24

5. References

[1] Charles J: Proc confDuplex 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 andWeld 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 Technol7, 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 Technology8, 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 StainlessSteels 91, Beaune, France, Oct. 1991, Vol.1, 347-362;

electron beam welding of superduplex stainless steel.pdf

Documents