Cranfield University

Mark Keeler

The effects of heat input and interpass temperature on the

microstructure of duplex stainless steel during welding

School of Applied Sciences

MSc

Welding Engineering

Cranfield University

School of applied sciences

MSc Thesis

Academic year 2010-2011

Mark Keeler

The effects of heat input and interpass temperature on the

microstructure of duplex stainless steel during welding

Supervisor: Paul Colegrove

Academic Year 2010 to 2011

This thesis is submitted in partial fulfilment of the requirements for

the degree of MSc Welding Engineering

© Cranfield University, 2011. All rights reserved. No part of this publication may be reproduced

without the written permission of the copyright holder

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Abstract

Intermetallic precipitates are known to have an extremely detrimental effect on both the mechanical

and corrosive properties of duplex stainless steel. The aim of the experiment was to find out if it

was interpass temperature or heat input that the most important factors in the formation of these

precipitates during welding.

Two sets of 5 butt welds were made on Duplex 2205 stainless steel using the Manual Metal Arc

Welding process. The first set of 5 welds with 1.0 KJ/mm heat input and the second 5 at 1.5 KJ/mm.

All plates started off with a pre heat of 150:C. The inter pass temperature for the first plate in each

set was also 150:C, this was increased in 50:C increments for each successive plate till the fifth one

had an inter pass temperature of 350:C.

It is known that the formation of intermetallic precipitates depends on time and diffusion at

elevated temperatures, these conditions can be found during the welding thermal cycle. The

complex welding cycles in multi-run welding can take the weldments to temperatures above that in

which these precipitates, and austenite, can exist and therefore they can revert back to the original

ferrite phase. At these elevated temperatures large ferrite grain growth can occur. If the

temperature cools too quickly then these other phases will not have time to reform leaving the

material with large ferrite grains. Both the presence of the precipitates and large ferrite grain

growth will have an equally detrimental effect on the toughness and the mechanical properties of

the weldment.

It was found that at the extremely high interpass temperature of 350:C and 1.5KJ heat input that

intermetallic precipitates become visible through micrographic examination, whereas the plate with

1.0KJ heat input at the same inter pass temperature did not. Although at lower interpass

temperatures of 200:C the phases were showing signs of having an affected structure which would

also have some detrimental effect on the materials properties.

From the results obtained it would seem that the formation of intermetallic phases is normally an in

service problem and only through poor welding practice did we find it attributable to welding.

However excessive heat input's and interpass temperatures do lead to an altered phase composition

and thermo mechanical history for the weldment which can affect the intended life cycle.

Keywords:

Duplex stainless steel, heat affected zone, micrographic examination, intermetallic precipitates,

interpass temperature, heat input, welding thermal cycle.

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Acknowledgements

I would like to thank Tom Donnellan (Sheffield and Perth) for his suggestion for the topic for this

thesis.

Because my funds were tight I'd like to thank Paul Sawyer from RCR Heat Treatment (Perth) and

Colin Lorrimar from Metallurgical Testing Services (Perth) and their staff for helping me set up my

experiments and with the testing afterwards.

Dr Supriyo Ganguly gave excellent advice and I would like to thank him for this.

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Table of contents

Abstract Acknowledgements Table of contents Table of figures Table of tables Table of graphs 1 Introduction 1 1.1 background 1 1.2 objective 2 Literature review 2 2.1 background 2 2.1.1 What is stainless steel? 2 2.1.2 Austenitic and Ferritic phase structures 2 2.2 Austenitic Stainless steels 7 2.2.1 Welding problems of austenitic stainless steels 7 2.3 Ferritic Stainless steels 8 2.2.1 Welding problems of ferritic S/S 8 2.3 Duplex stainless steel 8 2.3.1 Effects of alloying elements 9 2.3.2 Physical properties 12 2.3.3 Welding Duplex stainless steel 14 2.3.4 Corrosion resistance 15 2.3.5 Intermetallic compounds-precipitates 15 2.4 Maximum temperatures 18 2.4.1 475 Embrittlement 18 2.4.2 Recommended temperature 18 2.4.3 t12/8 cooling rate 19

3 Experiments and methodology 20

3.1 Experiment 20

3.1.1 Equipment 21

3.2 Material and weld preparation 22

3.2.1 Material composition 22

3.2.2 Plate preparation 23

3.2.3 Thermocouple 25

3.3 Welding 26

3.3.1 Preheat 26

3.3.2 Consumables 26

3.3.3 Welding 26

3.3.4 Heat input 27

3.4 Testing 29

3.4.1 Material preparation for testing 29

3.4.2 Etching 29

3.4.3 Metallographic examination 30

3.4.4 Ferrite count 32

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4 Results 34

4.1 Materials prior to experiment 34

4.2 Time at peak temperature 35

4.3 Macrographs 37

4.3.1 Parent plate 37

4.3.2 Heat affected zone 39

4.4 Ferrite count 40

5 Discussion 42

5.1 Welding Temperatures 42

5.2 Microstructure 42

5.2.1 Parent plate microstructure 42

5.2.2 HAZ microstructure 43

5.2.3 % Ferrite 43

6 Conclusion 45

7 References 47

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Table of figures

Figure 1 Transformation from Ferritic to Austenitic grain structure. 2

Figure 2 Crystallographic changes due to addition of nickel. 3

Figure 3 Fe- Cr phase diagram showing gamma (y) loop. 3

Figure 4 The Schaeffler diagram. 4

Figure 5 Delong Diagram. 5

Figure 6 WRC-1992 diagram for predicting weld ferrite content and solidification mode. 6

Figure 7 Relationship between Chromium and Nickel contents for basic types of stainless steels. 6 Figure 8 Schematic illustration of the mechanism of solidification cracking. 7 Figure 9 Mechanisation of sensitisation. 8 Figure 10 Thermodynamic equilibrium system Fe-Cr-Ni for 68% Fe 9 Figure 11 Transformation from Ferritic to Duplex to Austenitic grain structure if too much

nickel is added. 10

Figure 12 Thermodynamic equilibrium system Fe-Cr-Ni for 68% Fe with the

addition of nitrogen. 11

Figure 13 Typical stress-strain curves for carbon and stainless steels. 13 Figure 14 Differences in impact strengths between austenitic, ferritic, ferritic austenitic (duplex) and martensitic stainless steels. 13 Figure 15 Impact energy, charpy V, of UNS S31803 (2205) parent metal and welded joint. 14 Figure 16 Time temperature transformation diagram showing effect of alloying elements on

precipitation reactions in duplex stainless steels. 15

Figure 17 Isothermal precipitation diagram for 2205 duplex stainless steel, annealed

at 1050:C. 16

Figure 18 Formation of Sigma phase on the grain boundaries of ferrite. 17

Figure 19 Photograph of the 16mm plate showing the kerf left from the plasma cutting process which needed to be removed. 23 Figure 20 Photograph showing monitoring of excess heat generated during grinding 24 Figure 21 Plate prepared for welding. 24 Figure 22 Thermocouple attachment unit 25 Figure 23 Rear of plate showing thermocouples at 10mm spacing's from the

weld centre line. 25 Figure 24 MMAW electrode baking oven and MMAW electrode hotbox used for preheating the plates and electrodes prior to welding. 26 Figure 25 A completed weld showing the plate at a 10 degree angle. 27 Figure 26 Lapping block and emery paper used to prepare samples for etching. 29

Figure 27 S31803 showing an unaffected structure. 30

Figure 28 Possibly affected structure of S31803 and J92205. 31

Figure 29 An affected structure of S31803 and J92205 31

Figure 30 showing a centre line phase structure in J92205. 32

Figure 31 Graphical representation of the magnetic field produced by the probe. 33

Figure 32 Areas of the sample tested with the ferrite scope. 33

Figure 33 Macrographs of the Duplex plate in the supplied un-welded condition. 34

Figure 34 Graph of peak temperatures for 1KJ/mm and 150:C interpass temperature. 35

Figure 35 Graph of peak temperatures for 1KJ/mm and 200:C interpass temperature. 35

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Figure 36 Graph of peak temperatures for 1KJ/mm and 250:C interpass temperature. 35

Figure 37 Graph of peak temperatures for 1KJ/mm and 300:C interpass temperature. 35

Figure 38 Graph of peak temperatures for 1KJ/mm and 350:C interpass temperature. 35

Figure 39 Graph of peak temperatures for 1.5KJ/mm and 150:C interpass temperature. 36

Figure 40 Graph of peak temperatures for 1.5KJ/mm and 200:C interpass temperature. 36

Figure 41 Graph of peak temperatures for 1.5KJ/mm and 250:C interpass temperature. 36

Figure 42 Graph of peak temperatures for 1.5KJ/mm and 300:C interpass temperature. 36

Figure 43 Graph of peak temperatures for 1.5KJ/mm and 350:C interpass temperature. 36

Figure 44 Duplex 2205 parent plate (A) 1KJ heat input 150:C interpass;

(B) 1.5KJ heat input 150:C interpass. 37

Figure 45 Duplex 2205 parent plate (C) 1KJ/mm heat input 200:C interpass;

(D) 1.5KJ heat input 200:C interpass. 38

Figure 46 Duplex 2205 parent plate (E) 1KJ heat input 250:C interpass ;

(F) 1.5KJ heat input 250:C interpass. 38

Figure 47 Duplex 2205 parent plate (A) 1KJ heat input 350:C interpass. 38

Figure 48 Duplex 2205 HAZ (A) 1KJ heat input 350:C interpass;

(B) 1.5KJ heat input 350:C interpass. 39

Figure 49 Duplex 2205 HAZ (A) 1.5KJ heat input 350:C interpass showing the weld zone,

fusion zone and parent material;

(B) 1.5KJ heat input 350:C interpass. Showing the HAZ near the cap. 39

Figure 50 Duplex 2205 HAZ (A) 1.5KJ heat input 350:C interpass. 39

Figure 51 Graph of % Ferrite count 1KJ/mm heat input. 40

Figure 52 Graph of % Ferrite count 1.5KJ/mm heat input. 40

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Table of tables

Table1 Alloy additions that promote ferrite and austenite. 4

Table 2 Comparative properties of the stainless steel alloy families. 12

Table 3 Precipitation reaction (and others) temperatures for duplex S/S. 16

Table 4 Upper temperature limits for duplex S/S from pressure vessel design

codes for working out maximum allowable stress values. 18

Table 5 Sanvik steels recommended welding variables. 20

Table 6 Heat inputs and interpass temperatures used during the experiment. 21

Table 7 Equipment list from experiment. 21

Table 8 Parent plate composition. 22

Table 9 Consumable composition. 22

Table 10 1KJ/mm Heat input welding variables and observations. 28

Table 11 1.5KJ/mm Heat input welding variables and observations. 28

Table 12 % ferrite count in the supplied un-welded plate. 34

Table 13 Cumulative time at temperatures above 350:C (1KJ heat input). 37

Table 14 Cumulative time at temperatures above 350:C (1.5KJ heat input). 37 Table 15 Average % Ferrite count 1KJ/mm heat input 41

Table 15 Average % Ferrite count 1.5KJ/mm heat input 41

Table of Equations

Eq 1 Nickel equivalent. 4

Eq 2 Chromium equivalent. 4

Eq 3 WRC-1992 Nickel equivalent. 6

Eq 4 WRC-1992 Chromium equivalent. 6

Eq 5 Ferrite formation during cooling. 11

Eq 6 Heat input KJ/mm 27

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1. Introduction

1.1 Background Duplex stainless steel has a two phase structure austenite and ferrite in roughly equal proportions.

As such it combines the positive features of both these types of stainless steel. Such as excellent

resistance to corrosion, high strength, toughness and overall good weldability. In fact it is readily

welded by all the most common welding process such as SAW, MMAW, GTAW, FCAW etc.

Because of these properties duplex stainless steel is increasingly being used instead of austenitic

stainless steels in a wide range of industries such as the pulp and paper industry, desalination plants,

tankage, oil and gas industry.

Ideally the microstructure of the heat affected zone would have the same phase composition as the

original parent plate. Due to the repeated thermal cycles experienced in multi-pass welding this is

extremely unlikely resulting in the formation of unwelcome intermetallic precipitates which reduce

the corrosion and toughness values around the welded joint. If the thermal cycle is tightly controlled

then the properties of this zone can still be kept within acceptable limits. However these properties

can be cast aside through lack of control of the welding process.

1.2 Objective The aim of this study was to see at what temperatures intermetallic phases and other harmful micro

structural changes take place during and to see if this indeed a problem that is associated with

welding.

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2. Literature review

2.1 background

2.1.1 What is stainless steel?

A steel that resists corrosion is known as stainless steel. This resistance to corrosion is due to the

Chromium forming a Chromium oxide film on the surface of the metal that becomes a passive layer

that insulates and protects the surface. The Chromium must be in concentrations of at least 11-12%

in solid solution before the oxide (1-nm thick) can cover the surface sufficiently to make it stainless.

(1)

The following are the groups of steels that make up the stainless steel (S/S) family.

Austenitic stainless steels: AISI 200 and 300 series

Ferritic stainless steels: AISI 400 series

Martensitic stainless steels: AISI 400 series

Duplex (ferritic-austenitic) Stainless steels: AISI 300 series

Precipitation hardening stainless steels

2.1.2 Austenitic and Ferritic phase structures

Figure 1 Transformation from Ferritic to Austenitic grain structure (2)

Figure 1 shows the grain structure of the ferrite and austenite are completely different.

3

Figure 2 Crystallographic changes due

to addition of nickel (3)

Figure 2 shows Ferrite has a body centred cubic structure and Austenite has a face centred cubic

structure. If nickel is added, in sufficient quantities, then the structure can be changed from body

centred to the face centred as above.

Figure 3 Fe- Cr phase diagram showing gamma (y) loop (4)

The above iron-chromium phase diagram Austenite can only be found between 800C-1400C and

with less than 17% chromium. In order to achieve an austenitic structure at room temperature

alloying elements must be added to stabilize the austenite outside these ranges.

By far the most important alloying element for promoting austenite is Nickel but it is not the only

alloy added to stainless steel to alter its structure. Chromium when added will move the structure

back into a ferritic phase. Other alloys have a similar effect but to a lesser extent these are then

grouped as per the table below.

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Table 1 Alloy additions that promote ferrite and austenite (5)

In order to predict what the approximate composition of the material will be the Schaeffler diagram

was developed in 1949 to graphically show the phases in a weld deposit (also a general guide to the

parent microstructure). (6)

Figure 4. The Schaeffler diagram (7)

Each of the elements that were found to be austenite formers are given a weighting factor in

relation to nickel (the most powerful austenite former) to form the nickel equivalent equation. The

Chromium equivalent equation works in the same way for the ferrite formers.

Ni eq (Nickel equivalent) = % Ni + 30%C + 0.5% Mn (1)

Cr eq (Chromium equivalent) = % Cr + %Mo + 1.5%Si + 0.5% Cb (2)

5

By plotting the materials composition using the above equations the structure of the material can

be determined at room temperature.

The Delong constitution diagram was developed in 1956 that also introduces nitrogen as a strong

austenite former in the chromium equivalent area that lies between 18 and 27 (8)

Figure 5 Delong Diagram (9)

This diagram concentrates on a smaller area of the original Schaeffler diagram and introduces the

ferrite numbering system instead of the ferrite percentage. The ferrite number can be measured

using a ferrite meter and works by measuring the ferrite using the non destructive magnetic

induction method. This diagram works well for the 300 series of austenitic steels but not for duplex.

(10)

The WRC ferrite numbering system starts at 0 just like the % ferrite but the slowly diverges (see

above).

Further research by McCowan et al (11) led to the WRC (Welding research Council) 1988 diagram.

This was further modified by Kotecki and Siewert (12), for copper, till we have the WRC 1992

diagram used today.

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Figure 6 WRC-1992 diagram for predicting weld ferrite content and solidification mode. (13)

WRC-1992 Ni eq (Nickel equivalent) = % Ni + 35%C + 20%N + 0.25%Cu (3)

WRC-1992 Cr eq (Chromium equivalent) = % Cr + %Mo + 0.7% Nb (4)

Figure 7 Relationship between Chromium and Nickel contents for basic types of stainless steels. (14)

The above diagram shows were the various groups sit in relation to each through their Nickel and Chromium contents in a more graphical interpretation of the Schaeffler diagram.

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2.2 Austenitic Stainless steels

Austenitic S/S have the best overall resistance to corrosion of all the stainless steels with those with

increased molybdenum (Type 316) and nickel being better at resistance to pitting corrosion in high

chloride environments.

The minimum requirement for an austenitic microstructure is generally 16.5% chromium with

further additions of nickel, manganese, nitrogen and carbon. Other elements can be added to

improve the corrosion, oxidation properties etc.

In order to eliminate the delta ferrite that can form in austenitic S/S the material is normally

supplied in the solution annealed condition. It is heated to approximately 1050:C then rapidly

cooled to room temperature. In conditions such as autogenous welding delta ferrite can reform if

welded without suitably alloyed consumables. Unlike ferritic S/S austenitic S/S does not experience

large grain growth and as such can be welded in thick sections.

2.2.1 Welding problems of austenitic stainless steels

Because austenitic S/S have a 50% greater efficiency of thermal expansion than carbon steels and

30% lower co-efficient of thermal conductivity this can lead to problems of distortion during

welding. But the main problems encountered with welding austenitic S/S are hot cracking (Figure 8),

sensitisation/ intermetallic precipitation (Figure 9).

Figure 8 Schematic illustration of the mechanism of solidification cracking (15)

8

Figure 9 Mechanisation of sensitisation. M stands for mixed metals (16)

2.2 Ferritic Stainless steels

Ferritic stainless steels normally have greater than 17% chromium to around 30% but virtually no

nickel. Many consider ferritic S/S to be inferior to austenitic S/S but this is not the case, they are a

cheaper alternative. They have good resistance to high temperature oxidisation, liquid corrosion,

stress corrosion cracking and pitting. (17)

They are just as "Stainless" as austenitic S/S as it is not the nickel but the Cr content that makes

them stainless. They are highly resistant to stress corrosion cracking and have similar yield strength

to carbon steels. They have a similar co-efficient of thermal expansion as carbon steels which makes

for easier fabrication due to less heat distortion than the austenitic grades.

2.2.1 Welding problems of ferritic S/S

Large grain growth in the HAZ leads to poor toughness values (see Figure 14) and embrittlement

occurs when held at elevated temperatures which can decrease corrosion resistance (18) . The TWI

(19) mentions that the loss of toughness is significantly less when welding less than 6mm sections

whereas Atlas steels (20) points out that for this reason most ferritic S/S are only supplied up to

about 3mm, EN 1011-3:2000 recommends a maximum of 2.5mm (21). For heavier sections ferritic

stainless steels such as 3CR12 (ASTM A240 UNS S41003) give better toughness values through

micro alloying which prevents grain growth in the heat affected zone.

2.3 Duplex stainless steel

As the name Duplex suggests this is a stainless steel that contains 2 different phases, austenitic and

ferritic. In order to achieve the advantages of duplex S/S strict control over the microstructure's

balance of approximately 50/50 ferrite/austenite must be maintained. However this 50/50 is not

always possible to maintain throughout. So long as the smaller of the two phases is present in at a

minimum of 30% by volume and that two phases maintain large but separate structures and not as a

scattering of precipitates then the duplex properties will be maintained. (22)

9

Figure 10 Thermodynamic

equilibrium system Fe-Cr-Ni

for 68% Fe (23)

Figure 10 shows that depending upon the proportion of Cr and Ni present the material can form into

3 different phases.

Primary ferrite (α)

Primary austenite (y)

A mixture of ferrite and austenite (α+y) or duplex

2.3.1 Effects of alloying elements

Further changes to the boundaries of where these phases form can also be influenced by the

addition of the alloys noted in Table 1.

Chromium

As with the other stainless steels chromium in sufficient quantities forms a protective chromium

oxide layer that gives duplex S/S it excellent corrosion resistance. Equation 2 shows chromium to be

the strongest ferrite former. The amount of chromium in duplex S/S is higher than for the austenitic

S/S, at least 20%. Unfortunately intermetallic phases are also promoted at these higher level of

chromium.

Molybdenum

Molybdenum is the second ferrite former found in equation 2 and like chromium helps protect

against corrosion, especially pitting corrosion. Again as with chromium large amounts of

10

molybdenum can be detrimental due to the formation of intermetallic phases. For this reason it is

normally kept below 4% in duplex S/S.

Nitrogen

To counter the tendency of the chromium and molybdenum to form intermetallics nitrogen is added

to delay their precipitation. Due to the lowering of the intermetallics the toughness of the duplex

S/S increases. Also nitrogen is an austenite former and promotes y phase. It is also the strongest

solid solution strengthening element in duplex S/S so greatly increases strength. Like molybdenum

it increases the pitting corrosion resistance as well as crevice corrosion. If the nitrogen is present in

too large a quantity then it can form chromium nitrides (see figure 16) which can form brittle areas

in the material.

Nickel

As can be seen in Figure’s 1 and 2 it is the addition of Nickel that gives the austenite it's face centred

cubic structure. Equation 1 shows it to be the strongest of the austenite formers. Like nitrogen it

delays the formation of harmful intermetallic phases but to a much lesser degree. Figure 14 shows

that the toughness of austenitic S/S is superior to that of ferritic S/S. This is due to the face centred

cubic structure of the y phase.

Figure 11 Transformation from Ferritic to Duplex to Austenitic grain structure if too much nickel is

added. (24)

With the austenitic phase forming approximately 50% of the duplex microstructure this gives it

toughness values much higher than ferritic S/S.

Figure 10 shows for a 68% Fe-Cr-Ni composition, the material will pass from the liquid (L) state to

liquid and ferrite (L+α) then to the ferrite (α) phase above the duplex forming area (α+y). In the

case of this example this area corresponds to approximately 4%Ni-29%Cr to 8%Ni-24%Cr. The

material will then cool further and some of the ferrite will transform into austenite. The ferrite-

austenite equilibrium is normally reached at approximately 1000:C with little change thereafter. If

on the other hand the material is quenched just below its solidification temperature the diffusion

process by which the austenite forms will not have time to take place and consequently the material

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will be almost 100% ferrite at room temperature (25). If the material is not quenched then an

equilibrium structure of ferrite, austenite and intermetallic phases will be the equilibrium structure

(26 and Figure 18) as it cools without annealing.

If a line is followed the through the eutectic point (approx 12% Ni) it can be seen that the material

will pass through the liquid - austenitic - ferrite region and then becomes duplex for a short period

before passing through and into the austenitic region. No more transformation can take place and

the material will stay as an austenitic stainless steel.

Figure 12 Thermodynamic equilibrium system

Fe-Cr-Ni for 68% Fe with the addition of

nitrogen. (27)

The addition of nitrogen is illustrated above (Figure 12) it clearly shows that it raises the

transformation temperature of ferrite to austenite dramatically. As diffusion is more rapid at these

higher temperatures this also gives the austenite more time to form. The material is then annealed

at approximately 1050⁰C-1150⁰C to leave a ferrite/austenitic structure that is not in equilibrium.

The % of ferrite can be calculated if the annealing temperature is known through the following

equation.

% ferrite = -20.93 + 4.01 - 5.6 + 0.016T (5)

Where T = annealing temperature in :C

= % Cr + 1.73% Si + 0.88% Mo

= % Ni + 24.55% C + 21.75% + 0.4% Cu

The final phase balance of the duplex S/S is therefore influenced not just by the chemical

composition but also the rate of cooling.

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2.3.2 Physical properties

Table 2 Comparative properties of the stainless steel alloy families. (28)

As can be seen in Table 2 and Figure 14 Duplex S/S offer many of the beneficial properties of its

constituent phases. From the Austenitic phase it gets corrosion resistance and toughness. From the

Ferritic phase it gets its chloride stress corrosion resistance and strength.

Even down to 80:C they have good toughness values but this is not quite cryogenic conditions.

Compared to austenitic S/S duplex S/S are more difficult to machine or form. The weldability is only

a little less than austenitic but much greater than ferritic stainless steels.

13

Figure 13 Typical stress-strain curves for carbon and stainless steels. (29)

Graph 1 shows the higher tensile strength of duplex S/S compared to austenitic S/S. Because of this

thinner material, or down gauging, can be used over austenitic S/S. This gives a saving in materials

and weight but has a higher initial cost (30). In order to decrease the likelihood of sensitisation

austenitic stainless steels now come in a lower carbon form EG 316L this lowering of the carbon

decreases the mechanical properties and yield strength of the original 316 making the gap between

duplex and austenitic s/s even greater. This saving of weight is very important for industries such as

offshore oil and gas where weight savings are at a premium.

Figure 14 Differences in impact

strengths between austenitic, ferritic,

ferritic austenitic (duplex) and

martensitic stainless steels. (31)

As can be seen from Figure 14, above, there is a clear transition zone for duplex S/S which lies

between the values seen with austenitic and ferritic S/S. This moving of the transition zone from

that of the ferritic to much colder temperatures makes the duplex much more useful in industrial

settings.

14

Figure 15 Impact energy, charpy V, of UNS S31803 (2205) parent metal and welded joint. (32)

The choice of welding process employed also alters the toughness values significantly. The lower

impact values seen for MMAW welding, in Figure 15, over that of the TIG welding, and especially the

un-welded, is probably due to the growth of the ferrite grains in the HAZ. As the MMAW has a lower

heat input than TIG the weldment would cool quicker so leading to increasing the likelihood of

ferrite formation.

2.3.3 Welding Duplex stainless steel

Duplex S/S has good weldability. When welding duplex S/S some problems can be encountered

during the rapid heating and cooling of the welding cycle, it can be difficult ensuring the weld metal

and heat-affected zone (HAZ) retain the corrosion resistance, strength, and toughness of the base

metal. During welding of duplex S/S it is necessary to design welding procedures that lead to a

favourable phase balance after welding to avoid precipitation of detrimental intermetallic phases.

As noted in section 2.3.1 if the weld metal cools too rapidly then the austenite has little time to form

so leading to excessive ferrite being retained. This problem is magnified even more during welding

as the weld metal cool relatively rapidly which can lead to poor mechanical properties and a loss of

corrosion resistance.

Early generations of duplex S/S were prone to poor toughness and corrosion resistance in the HAZ

due to the formation of excessive ferrite during welding. To counter this modern duplex S/S have

higher levels of nitrogen (Figure 12) then was the case in the past leading to improved mechanical

properties (Figure 13). This and the over alloying of the electrode consumable with nitrogen should

ensure the austenite phase will form a favourable phase balance in the weld metal and HAZ in the as

welded condition. If this was not the case then a post weld heat treatment would be needed which

can lead to undesirable intermetallic phases forming.

Even with the over alloying of nitrogen if the welding heat input is too low the weldment will

experience rapid cooling which again will lead to excessive ferrite retention.

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2.3.4 Corrosion resistance

Duplex S/S is equally as good, and in some cases better, as other stainless steels in resisting

corrosion and stress-corrosion cracking in harsh environments. This is especially so in high chloride

environments as long as the ferrite stays above 30%.

Duplex S/S there are evaluated for two types of corrosion, pitting and crevice corrosion. The ways of

measuring these are.

Critical pitting temperature (CPT)

Critical crevice corrosion temperature (CCT)

Both of these are determined by the temperature above which corrosion will start to become

detectable within a 24 hour period in a ferric chloride solution. The ferric chloride simulates real life

environments such as sea water etc. Typically the CCT starts at 15:C to 20:C lower than the CPT.

The ASTM G48 standard is followed for testing of these critical temperatures (33)

2.3.5 Intermetallic compounds-precipitates

These are compounds that contain at least 2 metals and can form hard brittle phases which lowers

toughness values. As can be seen in Figure 16 the formation of these is time temperature

dependant and as they form on the boundaries between ferrite and austenite they strip away some

of the constituent elements through diffusion . In the case of those containing chromium this leads

to localised areas around the intermetallic phase where the chromium content will fall below the

percentage necessary to ensure corrosion resistance.

Figure 16 Time temperature transformation diagram showing effect of alloying elements on

precipitation reactions in duplex stainless steels. (34)

16

Due to the increased amounts of chromium and molybdenum duplex S/S has a greater likelihood to

form detrimental precipitates than other stainless steels this is due to the high diffusivity of

chromium and molybdenum into the ferrite phase.

Figure 17 Isothermal precipitation diagram for 2205 duplex stainless steel, annealed at 1050:C. (35)

2205 2507

:C :C

Solidification range 1470 to 1380 1450 to 1350

Scaling temperature in air 1000 1000

Sigma phase formation 700 to 950 700 to 1000

Carbide precipitation 450 to 800 450 to 800

475:C embrittlement 300 to 525 300 to 525

Table 3 Precipitation reaction (and others) temperatures for duplex S/S. (36)

17

Figure 17 shows how little time is needed in order for some of these precipitates to form. In the

case of Sigma, Chi, Carbide and Alpha prime this can occur approximately between 3 and 15

minutes. Table 3 shows the temperature ranges required for these precipitation reactions to take

place. As can be seen in the case of sigma at the lower temperature of 700:C it takes about an hour

for its formation it is only at the higher temperature of 900:C that this can form in 3 minutes.

Sigma phase (σ)

Due to its hard brittle nature it can drastically reduce the toughness values of the material. As can be

seen in Figures 18 and 9 molybdenum and chromium precipitate into the sigma phase leaving a

depleted area that is prone to corrosion. As these elements have a higher solubility in ferrite than

austenite the sigma forms on the outside of the ferrite grain boundaries. Because molybdenum and

chromium are ferrite formers (section 2.3.1) the lowering of these elements in this area of the ferrite

causes austenite formation on these boundaries. Figure 18 shows the sigma then growing into the

ferrite grains. Due to the welding cycle it is more prevalent in the Heat affected zone (HAZ). (37)

Figure 18. Formation of Sigma phase on the grain boundaries of ferrite. (38)

Wong, K.W. et all (39) claim that if the material is heated to temperatures above the sigma forming

temperatures the sigma phase has a tendency to transform back to ferrite and at these types of

temperatures ferrite and austenite are the thermodynamic stable phases. They did not indicate how

long it took for the ferrite to nucleate back from the sigma phase.

The other precipitates in Figure 17 can also form but sigma phase is the most detrimental to the

mechanical properties.

Secondary austenite In Figure 18 as well as showing σ phase forming on the boundary it also shows secondary austenite

also forming. Because molybdenum and chromium are ferrite formers (section 2.3.1) the lowering

of these elements in this area (due to the formation of σ phase) in the ferrite causes austenite

formation on these boundaries. This austenite is depleted of molybdenum and chromium and as

such the corrosion resistance of the material is further reduced (26).

18

2.4 Maximum temperatures

2.4.1 475 Embrittlement

Figure 16 shows there are two groupings of precipitation the lower group are centred around 475:C

at the nose of the curve and as such there are known to cause "475:C embrittlement". One of the

most important is alpha(a') phase which will lead to hardening, loss of corrosion resistance and

embrittlement. As can be seen if Figure 17 this takes at least 20 minute to start at the 475:C

temperature. Some precipitation can take place at temperatures as low as 300:C, although at these

low temperatures this can take months if not years to take effect.

2.4.2 Recommended temperature

For the reasons given above it is recommended the there are maximum temperatures in which it is

suitable to use duplex S/S.

It is recommended by Davison and Redmond (40) that duplex S/S should not be considered for

applications were the in service temperature is above 315:C. W.T.I.A. Technical note 13 (41) gives a

lower figure of 300:C for the embrittlement to commence. Even lower temperatures of 280:C for

un-welded and 250:C for welded structures were given by Van Nassau, L et all (42)

Grade Condition ASME TüV

:C :C

2304 Unwelded 315 300

2304 Welded, matching filler 315 300

2304 Welded with 2205/2209 315 250

2205 Unwelded 315 280

2205 Welded 315 250

2507 Seamless tubes 315 250

Alloy 255 Welded or unwelded 315

Table 4 Upper temperature limits for duplex S/S from pressure vessel design codes for working out

maximum allowable stress values. (43)

The German TüV welding code, Table 4, anticipates that temperature toughness suffers with welded

structures due to the effects of precipitation and large grain growth in the ferrite phase which

causes a hardening and embrittlement. For 2205 material the temperature of the maximum

allowable stress is reduced by 30:C to 250:C for this reason.

19

2.4.3 t12/8 cooling rate

The rate of cooling between 1200:C and 800:C can be critical in achieving a the correct balance of

phases in duplex S/S. If the cooling is to rapid through this t12/8 zone then chromium nitrates can

form which will favour the formation of ferrite (25), however chromium nitrates can lower the

corrosion resistance. Whereas slower cooling through this range will favour austenite and in some

cases intermetallic precipitates (44). A cooling rate of 10 seconds through this critical zone was

proven to give an acceptable phase balance by Geipl, H. (45) This cooling rate is very difficult to

calculate as it involves taking direct temperature measurements in the weld pool as it solidifies.

Sieuri and Sandstrom (37) found that a cooling rate of

would be best to give the correct ferrite/austenite balance whilst at the same time limit the

formation of sigma phase during the critical t12/8 range.

20

3 Experiments and methodology

3.1 Experiment

The formation of intermetallic compounds rely on:

Diffusion

Time

Composition and thermal/thermo mechanical history (46)

The experiments were structured to see how much heat input and interpass temperature has a

bearing on the final microstructure of 2205 Duplex S/S. To see which has the most influence on the

changes in microstructure the experiments used a varying interpass temperature and two different

heat inputs. EN 1011-3:2000 (47) recommends for a heat input range of between 0.5KJ/mm and

2.5KJ/mm which is the same as the guidelines from Sandvik (48) but they also give a maximum

interpass temperature of <250:C. The Welding and Technology Institute of Australia (49)

recommends no more than <1.5KJ/mm and <150:C interpass temperature.

Table 5 Sanvik steels recommended welding variables. (50)

Because of these variances in recommended inputs for these experiments two different heat inputs

were used, 1KJmm and 1.5KJ/mm. The interpass temperature was also increased in increments of

50:C from a base of 150:C. As diffusion is time temperature dependant it was hoped that when any

formation of intermetallics was noted in the micrographs the minimum formation temperature and

time could be deduced from the time temperature graphs recorded during the welding. In total 10

test plates were welded 5 at 1KJ/mm and 5 at 1.5KJ/mm (see table below).

21

Inter Pass temperature 1.0KJ/mm Heat Input 1.5KJ/mm Heat Input

150:C Test plate 1 Test plate 6

200:C Test plate 2 Test plate 7

250:C Test plate 3 Test plate 8

300:C Test plate 4 Test plate 9

350:C Test plate 5 Test plate 10

Table 6 Heat inputs and interpass temperatures used during the experiment.

3.1.1 Equipment

Cig Weld Transarc MMAW welder Data logger: Pico USB TC-08

Smithweld S-9F electrode hot box Generic desktop computer

Smithweld S-150H Electrode baking oven Digital temperature gauge

Bosch 125mm grinder Amp/Volt meter

Iron free hand tools Thermocouple attachment unit

Iron free cutting discs Wet and dry abrasive paper 40 grit to 1200 grit

Powertech DC supply Ferrite scope FMP30 (Fischer)

Sodium Hydroxide reagent (caustic soda) Olympus GX71 metallurgical inverted microscope

Leica DFC 295 3MP digital microscope camera

Table 7 List of equipment used in the experiment.

Variables for this experiment were:

Interpass temperature

Heat Input

Settings maintained for the experiments were:

Amps

Volts

Travel speed

Note: The welder varied the travel speed to achieve the required heat input but once the necessary

heat input was achieved it was maintained throughout (See notes/observations from tables 10 and

11).

22

3.2 Material and weld preparation

3.2.1 Material composition

Parent plate

The duplex S/S material was supplied by ArcelorMittal and complied with UNS

S32205(2205)/UNS31803 specification.

Chemical composition (%) of 2205 as supplied (heat number 72631)

C Mn P S Si Ni Cr Mo N

0.024 1.830 0.029 0.0005 0.324 5.467 22.574 3.023 0.165

Ferrite content 50.08% (Determined through automatic image analysis of volume fraction of ferrite)

Table 8 Parent plate composition (taken from the material test certificate).

The material was solution annealed at 1050-1100:C with a 2min/mm holding time.

Consumables

The welding consumables were Avesta 3.25 x 350mm (ASME SFA 5.4 E2209) MMAW electrodes.

Chemical composition (%) E2209 (Cert 9342 lot 09116)

C Mn P S Si Ni Cr Mo N Cu

0.024 0.6 0.026 0.017 0.94 9.01 23.07 3.07 0.13 0.16

Ferrite content 38.5% (see calculation below)

Table 9 Consumable composition.

The supplier did not stipulate the ferrite content but using EQ 3 and EQ4 and plotting this with the

WRC-1992 diagram (fig 6) we can calculate this.

Ni equivalent = % Ni + 35%C + 20%N + 0.25%Cu (3)

9.01 + (35 x 0.024) + (20 x 0.13) + (0.25 x 0.16) = 12.49

Cr equivalent = % Cr + %Mo + 0.7% Nb (4)

23.07 + 3.07 + (0.7 x 0) = 26.14

WRC-1992 = 55FN

In order to convert a FN to a ferrite % (the parent plate was measured as a % ferrite) the FN is

multiplied by 0.7 to give an approximation of ferrite by %. (51)

55FN x 0.7 = 38.5%

Note: No dilution has been accounted for in the above equations.

23

3.2.2 Plate preparation

A relatively thick 16mm plate was selected for the experiment. It was hoped that the thick material

would allow enough welding runs to be deposited and allow the plate to "hold" the temperatures

achieved. For this reason the plates were kept to a small overall dimension of 150mm long x 75mm

wide to reduce the dissipation effect by conduction as seen in large plates with a large surface area.

The plates were plasma cut under water. The bevels were prepared using a nine inch grinder and

iron free disks. All signs of the plasma "as cut" (Figure 19) were removed by at least 2mm as per the

recommendation of TWI Job knowledge 106 (52) to remove the HAZ caused by the cutting process.

Figure 19 Photograph of the 16mm plate showing the kerf left from the plasma cutting process which needed to be removed

As the aim of the experiment was to see the effect of heat input on the complete weldment the

temperature was monitored even during the bevel preparation during grinding.

As can been seen in Figure 20 the temperature quickly rose to 200:C. In order to eliminate the

possibility of the grinding heat input affecting the results the grinding was stopped at regular

intervals and the plates cooled. They were cooled in distilled water to eliminate contamination from

the local water supply.

24

Figure 20 Photograph showing monitoring of excess heat generated during grinding.

As the root run is the most difficult part of the welding procedure a 6mm backing bar was used. This

would save time as any delays in the first run would make the plates preheat temperature to fall

below the 150:C minimum selected for the experiment. This would not be a problem with the lower

inter run temperature test plates but with the higher interpass temperatures of 300:C and 350:C

time would be critical and it would be essential to maintain a uniformity throughout all the plates. A

gap in the backing strip of about 10mm was left and a small root run was welded with a 2mm throat

thickness (Figure 22 and 23) onto which the 1st thermocouple would be attached. In this way it was

hoped to get a temperature reading through the centreline of the weldment and quite close to the

first weld run. The backing bar overlapped the parent plate by about 10mm to act as run on and run

off strips.

Figure 21 Plate prepared for welding.

25

3.2.3 Thermocouple

To measure and record the thermal cycle during welding 4 sets of thermocouples were attached to

the rear of the plates to be welded. A thermocouple attachment unit (TCAU) was used for this task

(Figure 22). The TCAU is portable unit with a capacitor, it works by discharging the stored energy of

the capacitor through the thermocouple wires causing them to weld to the plates.

Figure 22 Thermocouple attachment unit

Four sets of thermocouple wires were attached to the rear of the plates. The first one was attached

on the weld centreline and the subsequent ones at 10mm intervals.

Figure 23 Rear of plate showing thermocouples at 10mm spacing's from the weld centre line.

26

3.3 Welding

3.3.1 Preheat

The plates were pre heated in an MMAW electrode baking oven to 180:C. (Figure 24) This was in

excess of the required 150:C preheat but this allowed time for attaching the thermocouple wires to

the data logging unit and welding the plate on to the work bench prior to commencing the

experiment.

3.3.2 Consumables

The welding electrodes were Baked at 250:C for 3 hours and held in a hot box (Figure 24) prior to

use as per the manufacturers recommendations.

Figure 24 MMAW electrode baking oven and MMAW electrode hotbox used for preheating the plates and electrodes prior to welding.

3.3.3 Welding

The welding process chosen for the experiment was Manual Metal Arc Welding (MMAW). MMAW is

widely known used and accepted in the welding industry. It can be used in most site conditions even

in relatively strong winds. Because of this it is often used for repair welding were it is favoured for

filler and capping passes. It can be used for root welds but this can be prone to operator error. Its

simplicity and ease of set up made it more favourable over fully automated welding due to the fast

turnaround time for commencing the subsequent welding runs. This would not be an issue in the

plates welded with the lower inter-pass temperatures, 150:C to 250:C, but the rapid cooling of the

higher inter-pass temperatures meant that welding had to recommence rapidly after the removal of

the slag and cleaning of the weld. Semi-automatic and fully automatic welding such as Metal Inert

Gas Welding (MIG) can be prone to stoppages which can take time to rectify during which the plates

could have cooled below the desired minimum inter-pass temperature.

27

3.3.4 Heat input

To measure the heat input the following equation was used:

Heat input KJ/mm =

x process efficiency (6)

Note: The travel speed was calculated using the formula: Length/seconds x 60

A test plate was used first and it was found that to achieve the necessary travel speed required for

1.5Kj heat input problems were found as the thick slag from the 2209 electrodes had a tendency to

get in front of the weld pool and form slag traps. To counter this all the plates thereafter were set

with a 10 degree incline this seemed to eliminate the problem. (Figure 25)

Figure 25 A completed weld showing the plate at a 10 degree angle.

Before the plates were pre heated each pair of thermocouples were visually inspected and tested for

electrical continuity. Unfortunately some of the thermocouples became dislodged either whilst

connecting to the data logger or during welding. The thermocouple that was closest to the weld

centre line was chosen as the indicator of the temperature measurements. If this thermocouple

became dislodged then the next thermocouple measurements were used.

The ambient temperature during the welding was 18 to 20:C.

28

1KJ/mm test plate. 150⁰C preheat

Test plate number Inter-run temperature Number of runs to complete the

weld

Notes/ Observations

1 150:C 9 Some difficulty in achieving the desired 1KJmm with some runs only at 0.7KJ/mm and some at 1.2KJ/mm. The average achieved throughout was approximately 0.85KJ/mm

2 200:C 10 Similar problems with varying heat input as plate 1. Travel speed was varied and 1KJ/mm was achieved.

3 250:C 11

4 300:C 10

5 350:C 10

Table 10 1KJ/mm Heat input welding variables and observations.

1.5Kjmm test plate. 150⁰C preheat

Test plate number Inter-run temperature Number of runs to complete the

weld

Notes/ Observations

6 150:C 9 Thermocouple 3 became dislodged during welding

7 200:C 8

8 250:C 8 Thermocouple 3 became dislodged during welding

9 300:C 8

10 350:C 9 Thermocouple 1 caused problems which delayed the start of the welding so the preheat was 135:C prior to commencement of welding

Table 11 1.5KJ/mm heat input welding variables and observations.

Note: The smaller number of runs needed for plates 6 to 10 can be attributed to the greater

distortion and metal deposition due to the higher heat input causing the plates to close up more.

As can be seen above some problems were encountered getting the correct travel speed, normally

in the first and second runs, the travel speed in the consequent runs was speeded up or slowed

down to compensate. These problems were greater with the smaller heat inputs and could have

been caused by the quite wide 5mm root gap. In order to keep the heat input low a fast travel

speed was used with only a small weave on the 1.0KJ/mm welds. For the 1.5KJ/mm plates a slower

travel speed was used and a larger weave this led to more control over the heat input.

29

3.4 Testing

3.4.1 Material preparation for testing

On completion of the welding the plates were cut up using an automatic hacksaw to give samples

approximately 70mm in length. The samples were taken from the centre of the test piece where the

thermocouples had been attached. The hacksaw cut left a very rough finish which was removed

with a belt grinder. Care was taken to keep the heat down just as in the plate preparation in section

3.2.2. The samples were then "lapped" on a granite block (Figure 26) using wet and dry abrasive

papers. The grit sizes were 40, 80, 120, 240, 320, 400, 600, 800, 1000 and finally 1200. Each paper

had a finer and finer abrasive content which left a mirror like finish to the samples.

A final polishing using a diamond suspension solution of 6 microns and 1 micron grit gave the

samples a satisfactory finish prior to etching.

Figure 26 Lapping block and emery paper used to prepare samples for etching.

3.4.2 Etching

The etchant used was Sodium Hydroxide mixed with distilled water this is commonly called caustic

soda. The sodium Hydroxide crystals were added to the water a small amount at a time till a ratio of

40% Sodium Hydroxide to 60% water was reached. Only small amounts were added at a time as the

Sodium Hydroxide produces an exothermic reaction when in contact with water and can be

dangerous.

The samples to be etched were carefully lowered into the etchant solution to avoid splashing. The

samples were etched using an electrochemical etching technique. A Powertech D.C. power supply

was connected to a positive and negative electrode. The cathode electrode was placed in the

solution and the anode electrode placed on the polished surface of the sample. The samples were

etched using 2.5 volts for 3 to 5 seconds. This gave the surface a very light etch. If the etching was

too heavy then the intermetallic phases would be masked by the austenite on the grain boundaries

which would make them much harder to observe. The samples were finally rinsed with acetone and

inspected visually to make sure the etching had been successful.

30

The etchant reacts more so with the ferrite than the austenite so leaving the austenite slightly

elevated to the ferrite in the micrographs.

3.4.3 Metallographic examination

ASTM A 923 (46) sets out testing methods to be employed for detecting detrimental intermetallic

phases in duplex S/S. The three tests are:

Method A- Sodium hydroxide etch and metallographic examination.

Method B- Charpy impact testing.

Method C- G48 corrosion test.

For the purpose of this study only method A has been followed. The etching and micrograph do not

indicate the exact type of detrimental phases found but the absence of them would indicate an

unaffected structure. If intermetallics are observed then this is an indication that there will have a

detrimental effect on toughness and corrosion resistance. To quantify the detrimental effect would

require further testing using charpy impact testing to measure the reduction in the materials

toughness and a G48 corrosion test to measure the lowering of the materials corrosion resistance.

A general survey was carried out using an Olympus GX71 metallurgical inverted microscope with a

Leica DFC 295 3MP digital microscope camera to capture the images. This was started on the centre

line of the weld to the root. This was followed by moving the sample 2mm and repeating till one

half of the weldment was surveyed to the HAZ zone plus 10mm.

As the formation of intermetallic phases is dependent on temperature and cooling rate a more

thorough metallographic examination concentrated on the region's most likely to contain these

phases. This area is in the centre of the weldment as this is the area that cools most slowly.

ASTM A923-03 categorises the microstructures as either Acceptable or Non acceptable structures.

Acceptable structures

Figure 27 UNS S31803 showing an unaffected structure. (500 x magnification) (46)

Figure 27 shows no signs of intermetallic formations and the boundaries between the phases are

smooth.

31

Non acceptable structures

Figure 28 Possibly affected structure of S31803 (left x 500) and J92205 (right x400). (46)

Figure 28 shows possibly affected structures with possible intermetallic phases, black dots and with

a wavy appearance of the phase boundaries.

Figure 29 An affected structure of S31803 (left x 500) and J92205 (right x400). (46)

Figure 29 clearly showing an affected structure with the formation of intermetallic phases, which

show up as black dots, on the boundaries.

32

Figure 30 showing a centre line phase structure in J92205. (500 x magnification) (46)

Figure 30 Shows intermetallic phases as near continuous phase normally in the mid thickness region

of the material.

In the case of acceptable structures (Figure 27) no further testing would be required. For those

samples showing unacceptable structures (Figures 28 to 30) then further testing to test methods B

and C would normally be carried out.

3.4.4 Ferrite count

The Metallographic examination in ASTM A 923 requires the weldment to be destructively tested.

Another testing method is through ferrite measurement with a ferrite scope. ASTM A799 (53) and

ISO 8249:2000(E) (54) describes this testing method. Because the austenite is non magnetic and the

ferrite is magnetic then the amount of magnetic permeability of the sample can be measured. The

ferrite scope can give either % ferrite or a ferrite number of the ferro-magnetic phase. The scope

gives a measurement of the amount of force required to "tear off" the measuring device which in

this case is a Ferrite scope FMP30 (Fischer).

The ferrite scope is usually used to measure the surface of the weldment but because the samples

had already been prepared for metallographic examination then the areas inspected under the

microscope were also tested with the ferrite scope. No further preparation was required for this

testing at the points 1 to 9 shown on Figure 32 as the surface preparation for the micrographs was

more than sufficient. The prepared surface is important as it ensures the probe sits perpendicularly

to the surface being measured. When the probe comes into contact with weldment a closed

magnetic circuit is created which is energized by a low frequency magnetic field. A voltage is

induced by this magnetic field in the probe coil, the magnetic permeability is calculated by this

induced voltage Figure 31 shows this graphically. (53)

33

Figure 31 Graphical representation of the magnetic field produced by the probe.

The weld was tested at the points 1 to 9 shown on Figure 32. The cap of the samples were also

prepared to test the surface. The cap, point 10, and part of the parent plate, point 11, were draw

filed flat, this filed area was a minimum of 10mm wide till all weld ripples had been removed.

Emery paper was attached to the file and filed to give a finer finish till a 600 grit finish was achieved

this is in accordance with ISO 8249:2000(E) 5.3.1 Surface finishing prior to ferrite testing. (54).

Figure 32 Areas of the sample tested with the ferrite scope.

Each of the areas of Figure 32 where tested 7 times each, the highest and lowest reading were

discounted, the other 5 readings were used to give an average % ferrite count for that area. As can

be seen in Figure 32 the areas chosen were down the centre line of the weld, the fusion zone and

the fusion zone + 5mm as well as the areas on the cap as shown.

34

4 Results

4.1 Materials prior to experiment

Figure 33 Macrographs (500 X magnification) of the Duplex plate in the supplied un-welded

condition. The dark colour shows the ferrite and the light blue the austenite.

Results of fraction of ferrite Micrograph Fraction of ferrite in %

Measurements Requiremements 1 51.38 2 49.68 3 50.88 4 49.02 5 51.28

Average 51.45 35-55 Std Deviation 1.05

95% Cl 1.30 % RA 2.57

Table 12 % ferrite count in the supplied un-welded plate.

The Macrographs (Figure 33) and ferrite count (Table 12) were taken from the certificate of

compliance provided by the manufacturer.

35

4.2 Time at peak temperature

Figure 34 Figure 35

Figure 36 Figure 37

Figure 38

0

100

200

300

400

500

600

0

44

5

89

0

13

35

17

80

22

25

26

70

31

15

35

60

40

05

44

50

48

95

53

40

57

85

62

30

66

75

71

20

75

65

80

10

84

55

Tem

per

atu

re in

°C

Time in Seconds

1KJ Heat input 150°C interpass temperature

0

200

400

600

800

1000

1200

0

35

5

71

0

10

65

14

20

17

75

21

30

24

85

28

40

31

95

35

50

39

05

42

60

46

15

49

70

53

25

56

80

60

35

Tem

per

atu

re in

°C

Time in seconds

1KJ Heat input 200°C interpass temperature

0 100 200 300 400 500 600 700 800 900

1000

0

220

440

660

880

1100

1320

1540

1760

1980

2200

2420

2640

2860

3080

3300

3520

3740

3960

4180

Tem

per

atu

re in

°C

Time in seconds

1KJ Heat input 250°C interpass temperature

0

200

400

600

800

1000

1200

0

195

390

585

780

975

1170

1365

1560

1755

1950

2145

2340

2535

2730

2925

3120

3315

3510

3705

Tem

per

atu

re in

°C

Time in seconds

1KJ Heat input 300°C interpass temperature

0

200

400

600

800

1000

1200

1400

0 1

80

36

0

54

0 7

20

90

0 1

08

0

12

60

14

40

16

20

18

00

19

80

21

60

23

40

25

20

27

00

28

80

30

60

32

40

Tem

per

atu

re in

°C

Time in seconds

1KJ Heat input 350°C interpass temperature

36

Figure 39 Figure 40

Figure 41 Figure 42

Figure 43

As can be seen in Figure 23 four pairs of thermocouples were attached originally before welding.

The heat inputs from the thermocouples recording the highest temperatures were used to populate

the graphs in Figures 34 to 43. This was usually the one on the weld centre line.

0

100

200

300

400

500

600

700

800

900

1000

0

51

0

10

20

15

30

20

40

25

50

30

60

35

70

40

80

45

90

51

00

56

10

61

20

66

30

71

40

76

50

81

60

86

70

91

80

Tem

per

atu

re in

°C

Time in Seconds

1.5KJ Heat input 150°C interpass temperature

0

100

200

300

400

500

600

700

800

900

1000

29

5 6

05

91

5 1

22

5 1

53

5 1

84

5 2

15

5 2

46

5 2

77

5 3

08

5 3

39

5 3

70

5 4

01

5 4

32

5 4

63

5 4

94

5 5

25

5

Tem

per

atu

re in

°C

Time in seconds

1.5KJ Heat input 200°C interpass temperature

0

200

400

600

800

1000

1200

1400

0 21

5

430

64

5

860

10

75

1290

15

05

1720

19

35

2150

23

65

2580

27

95

3010

32

25

3440

36

55

3870

Tem

per

atu

re in

°C

Time in seconds

1.5KJ Heat input 250°C interpass temperature

0

200

400

600

800

1000

1200

1400

0

190

380

570

760

950

1140

1330

1520

1710

1900

2090

2280

2470

2660

2850

3040

3230

Tem

per

atu

re in

°C

Time in seconds

1.5KJ Heat input 300°C interpass temperature

0

100

200

300

400

500

600

700

800

0 1

65

330

4

95

660

8

25

990

1

155

1

320

1

485

1

650

1

815

1

980

2

145

2

310

2

475

2

640

2

805

2

970

3

135

Tem

pe

ratu

re in

°C

Time in seconds

1.5KJ Heat input 350°C interpass temperature

37

Heat input and inter-run

temperature

Time above 350:C Minutes

Time above 400:C Minutes

Time above 450:C Minutes

Time above 500:C Minutes

1KJ/mm 150:C 4.66 1.75 0.41 .08

1KJ/mm 200:C 14.66 7.66 3.66 1.83

1KJ/mm 250:C 19.08 11.83 6.33 3.08

1KJ/mm 300:C 24.16 15.25 9.0 4.66

1KJ/mm 350:C 32.58 19.55 12.5 7.0

Table 13 Cumulative time at temperatures above 350:C (1KJ/mm heat input).

Heat input and inter-run

temperature

Time above 350:C Minutes

Time above 400:C Minutes

Time above 450:C Minutes

Time above 500:C Minutes

1.5KJ/mm 150:C 9.66 5.08 2.41 1.42

1.5KJ/mm 200:C 9.33 4.58 3.16 2.08

1.5KJ/mm 250:C 20.16 13 6.58 4.42

1.5KJ/mm 300:C 22.5 14.08 9.33 6.5

1.5KJ/mm 350:C 28.66 17.25 10.25 7.0

Table 14 Cumulative time at temperatures above 350:C (1.5KJ/mm heat input).

The temperatures that were recorded with the data logger during welding are shown in Tables 13

and 14 (above). The times are the total amount of time that each sample stayed above the

temperatures of 350:C, 400:C, 450:C and 500:C during all the weld runs.

4.3 Micrographs

4.3.1 Parent plate

Figure 44 Duplex 2205 parent plate (A) 1KJ/mm heat input 150:C interpass; (B) 1.5KJ/mm heat input

150:C interpass.

38

Figure 45 Duplex 2205 parent plate (C) 1KJ/mm heat input 200:C interpass; (D) 1.5KJ/mm heat input

200:C interpass.

Figure 46 Duplex 2205 parent plate (E) 1KJ/mm heat input 250:C interpass ; (F) 1.5KJ/mm heat

input 250:C interpass.

Figure 47 Duplex 2205 parent plate (A) 1KJ/mm heat input 350:C interpass.

39

4.3.2 Heat affected zone

Figure 48 Duplex 2205 HAZ (A) 1KJ/mm heat input 350:C interpass; (B) 1.5KJ/mm heat input 350:C

interpass.

Figure 49 Duplex 2205 HAZ (A) 1.5KJ/mm heat input 350:C interpass showing the weld zone, fusion

zone and parent material; (B) 1.5KJ/mm heat input 350:C interpass. Showing the HAZ near the cap.

Figure 50 Duplex 2205 HAZ (A) 1.5KJ/mm heat input 350:C interpass

40

4.4 Ferrite count

Figure 51 Points 1 to 11 from Figure 32

Figure 52 Points 1 to 11 from Figure 32

30

35

40

45

50

55

60

0 1 2 3 4 5 6 7 8 9 10 11 12

% F

err

ite

co

un

t

Sample point

% Ferrite count 1KJ/mm heat input

1Kj 150:C 1Kj 200:C 1Kj 250:C 1Kj 300:C 1Kj 350:C

30

35

40

45

50

55

60

0 1 2 3 4 5 6 7 8 9 10 11 12

% F

err

ite

co

un

t

Sample point

% Ferrite count 1.5KJ/mm heat input

1.5Kj 150:C

1.5Kj 200:C

1.5Kj 250:C

1.5Kj 300:C

1.5Kj 350:C

41

Average % Ferrite count 1KJ/mm heat input

Heat input

and inter-

run

temperature

Point

1

Point

2

Point

3

Point

4

Point

5

Point

6

Point

7

Point

8

Point

9

Point

10

Point

11

150:C 32.08 35.56 36.24 39.76 39.88 41.58 50.82 55.06 55.68 41.98 42.88

200:C 37.02 36.10 33.18 41.34 39.8 38.56 50.82 47.72 47.14 44.58 43.72

250:C 37.74 34.84 37.00 41.42 41.22 39.10 53.00 49.96 46.58 37.64 43.72

300:C 42.84 36.24 34.74 48.7 41.86 39.76 54.76 52.08 49.46 47.14 41.58

350:C 43.68 37.22 35.62 45.24 42.2 37.84 52.58 55.74 49.04 46.26 42.62

Average 38.67 35.99 35.35 43.29 40.99 39.36 52.39 52.11 49.58 43.52 42.90

Table 15 Points 1 to 11 taken from Figure 32

Average % Ferrite count 1.5KJ/mm heat input

Heat input

and inter-

run

temperature

Point

1

Point

2

Point

3

Point

4

Point

5

Point

6

Point

7

Point

8

Point

9

Point

10

Point

11

150:C 41.24 39.12 35.84 41.78 45.62 33.50 53.66 46.74 47.64 48.33 41.50

200:C 39.42 37.46 36.98 43.98 42.68 37.68 49.62 46.42 49.88 55.34 40.46

250:C 38.46 37.28 33.40 40.94 39.74 37.74 55.36 52.82 54.16 43.38 40.74

300:C 39.00 41.04 37.64 40.96 40.38 39.62 53.54 55.60 48.84 39.38 41.26

350:C 37.92 36.86 33.84 41.14 39.56 38.7 54.44 53.30 53.46 45.36 41.80

Average 39.21 38.35 35.54 41.76 41.60 37.45 53.32 50.98 50.80 46.36 41.15

Table 16 Points 1 to 11 taken from Figure 32

42

5 Discussion

5.1 Welding Temperatures The highest temperature recorded was for the 1.5KJ/mm HI 250:C interpass temperature where the

thermocouple recorded a temperature of 1321.:C (Figure 41) on the first welding run. This is not

surprising as this run was closest to the thermocouple and was consistent with most samples. As the

subsequent welds were placed further away from the point of recording of the temperature the

peak value diminishes. As all thermocouples were attached to the underside it was not possible

record the temperatures in the areas where the highest "holding" temperatures were expected i.e.

around the centreline of the plate and weld.

Even with the thermocouples placed on the underside both the weldments with the highest

interpass temperatures (350:C) stayed over 500:C for 7 minutes. At the lower interpass

temperatures the temperature rarely went above 500:C for the 1KJ/mm heat input (Figure 34) and

less than 2 minutes for the 1.5KJ/mm heat input (Figure 39).

The graphs show that the welds with the lower interpass temperatures, Figure's 34 and 39, were

held at temperatures above 150:C for approximately two to three times longer than Figure's 38 and

43. This was due to the long waiting times for the interpass temperature to drop back down to

150:C. The intervals between those welds with the higher interpass temperatures was very short.

In the case of the interpass temperatures of 350:C once the slag had been removed and the weld

wire brushed then the weld had to recommence almost immediately. Because of this the weldment

with 1.5KJ/mm and an interpass temperature of 150:C stayed above 150:C for a total 144 minutes

which was more than twice as long as long as the weldment with an interpass temperature of 350:C

which only stayed above 150:C for a total of 66 minutes during the experiment. This shows that

even though the plates with the higher interpass temperatures reached much higher temperatures,

the total time spent above the minimum interpass temperature (150:C) was actually much shorter.

Obviously the temperatures that were recorded during these times would have been much higher in

the areas of most interest i.e. around the centre of the weldment.

If the plates were of a thinner material then this much shorter welding duration could possibly not

provide sufficient time for the changes in microstructure that were observed.

5.2 Microstructure

5.2.1 Parent plate microstructure

The micrographs in Figure 44 show no signs of change in the parent plate as these plates were with

the lowest heat input and interpass temperatures of the two sets of test plates so this was not

surprising. In Figure 45 "C" showed no signs of change whereas "D" started to show signs of

"affected structure" (45) in the parent plate. All the micrographs exhibit this trend but more so with

Figure 47 with the phase boundaries showing signs of the wavy appearance seen in Figure 28. If

Figure 47 is compared to the original structure seen in Figure 33 then the changes are even more

noticeable. This is due to the austenite growing into the ferrite through diffusion with the increasing

heat inputs and temperatures as was noted in section 3.1.

43

5.2.2 HAZ microstructure

The only micrograph that showed any definite intermetallic precipitation was for the weldment of

1.5KJ/mm 350:C interpass temperature (Figure 48 "B"). The intermetallics were shown up as black

dots on the boundaries between the phases. The weldment with the lower heat input, 1.0KJ but the

same interpass temperature (Figure 48 "A"). exhibited no signs of intermetallics present.

The micrograph's in Figure 49 show the area of the HAZ with large areas of ferrite, in fact the

austenite is found in thin bands in a sea of ferrite. Figure 49 "A" was taken in the centre of the

weldment so was held at an elevated temperature for an extended period but still there is more

ferrite than austenite. This is due to the welding thermal cycle in which any austenite forming will

revert back to ferrite when the next weld run is placed over the top of the previous run. This can be

seen in Figure 12 where the HAZ temperatures will rise taking the material back into the ferrite

forming range. As the temperature starts to fall back into the austenite forming range the next weld

will be deposited and the cycle will be repeated.

In Figure 49 "B" the micrograph was taken close to the surface near the cap of the weld and the lack

of austenite is even greater. This is due not to the welding thermal cycle reverting the austenite

back to ferrite as in Figure 40 "A" and Figure 40 "B" but to the rapid cooling of the weldment at the

cap so very little austenite can form from the ferrite as it passes through t12/8 cooling zone too

quickly.

In areas that have not been subjected to the extremes of the welding thermal cycle, i.e. on the

boundary between the HAZ and parent material, the temperatures have not reached that in which

the austenite has reverted back to ferrite. However this area has sat at temperatures where the

diffusion of the austenite from the ferrite can continue. Figure 50 appears to show the austenite

"growing" as Widmanst tten plates into the ferrite.

5.2.3 % Ferrite

Using the WRC 1992 diagram (Figure 6) it was estimated that the weld deposit should have been

38.5% ferrite.

Looking at Tables 15 - 16 and Figures 51 - 52 a pattern can be seen. The first 3 readings, points 1 to

3, were from the weld metal with point 1 being near the surface this had the highest ferrite, the

ferrite % decreased towards the base of the plate. This was consistent with points 4 to 6 for the HAZ

and 7 to 9 for the parent plate. This would indicate that because the base had a longer time at an

elevated temperature then the austenite had more opportunity to form.

Points 10 (cap) and 11 (plate surface) on Figure 32 also showed difference in ferrite readings over

the corresponding readings (points 1 and 7) taken just 3mm below the surface. The average for 10

44

was 44.36% against 38.62% for point 1, the average for 11 was 41.3% against 54.1% for point 7. This

can be explained for the weld metal as the cooling rate for point 10 through the t12/8 zone faster

than the 10 seconds required for the transformation of austenite from the ferrite (45). The areas of

11 and 7 would have passed through these temperature ranges much more slowly but this does not

explain why the surface has a much lower ferrite count than the sub-surface.

Point 3 had the lowest average ferrite count of all the points tested with an average of 35.35% for

1KJ/mm and 35.54 for the 1.5JK/mm samples. This low reading cannot be explained as an effect of

the cooling effect of the root being open to cooling in air If this was the case then the ferrite count

would actually be higher than what was measured. The low ferrite count must have come about

due to this area having the greatest amount of dilution between the weld and parent metal.

Even so the lowest average ferrite count of 35.35% was just above the 30% needed to retain the

duplex S/S properties so in this case was acceptable.

45

6 Conclusion

Although there didn't seem to be a huge difference between 1KJ/mm and 1.5KJ/mm heat inputs at

the various temperatures it was noticed that the first indication of an affected structure took place

at 1.5KJ/mm with only 200:C interpass temperature and yet the corresponding plate with same

temperature but a lower heat input there was no sign of this. When the 1KJ/mm plate had an

interpass temperature of 250:C the micrographs also showed these signs. All the following test

pieces exhibited these changes. It was also noticed that the only sample in which intermetallics

formed was of the higher heat input at 350:C and yet again the sample with 1KJ/mm showed no sign

but as noted in section 5.2.1 the microstructure had already shown signs of change. It may well be if

an interpass of 400:C had been applied to the 1KJ/mm sample then this would have started to

exhibit similar signs of intermetallic phases forming. From this it can be concluded that the increase

in heat input does have a direct affect on the microstructures final phase structure.

The increasing interpass temperature did have a marked effect on the structure of the HAZ and the

parent plate near the HAZ. The recommended interpass temperatures given by Sandvik (48 and

Table 5) of 250:C if used in conjunction with the maximum heat input of between

1.5KJmm/2.5KJmm would lead to a HAZ with an affected structure. The recommendations from the

Welding and Technology Institute of Australia (49) of <1.5KJ/mm and <150:C interpass temperature

would seem to be correct.

The welding procedures carried out in the experiments whilst acceptable for the lower interpass

temperatures were unrealistic above 250:C. Minimum interpass temperatures of 350:C using

MMAW welding are almost impossible to physically carry out in anything except small weldments.

The complex welding cycles encouraged the austenite in the HAZ to transform back to ferrite as

predicted by Wong, K.W. et all (39). This will lead to poor toughness values and reduced corrosion

resistance due to large ferrite grain growth.

As can be seen from tables 13 and 14 the precipitates need more time at elevated temperatures

than is normally seen in welding to form harmful precipitates.

From these results it would seem that intermetallic precipitation is not a welding problem but is an

in-service phenomenon where the weldment stays at an elevated temperature for long periods of

time and intermetallic precipitates can form slowly.

Once a satisfactory weld procedure has been ran and passed all the relevant testing necessary for

the particular welding code being used then the welding variables must be strictly adhered to. For

the MMAW welding used during the experiment the maximum interpass temperature of 150:C was

the only temperature where a satisfactory microstructure was produced for both the of the heat

inputs used. The heat input did seem to have an effect on the microstructure but so long as the

weld procedure has passed the relevant testing then this should be acceptable if the interpass

temperature is kept low enough.

The detrimental effects of heat input and interpass temperature on the microstructure of duplex

stainless steel during welding are therefore controllable. It is up to the design engineer to take into

consideration the thermo mechanical history imposed during welding to ensure that the service

46

temperature stays within acceptable limits. If this is adhered to then the weldment should free

from metallurgical problems throughout its life cycle.

47

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