METALLURGICAL AND MECHANICAL PROPERTIES OF

Ni-BASED SUPERALLOY FRICTION WELDS

by

Sujith Sathian

Department of Metdugy and Materials Science University of Toronto

Toronto, Canada

A thesis submitted in confonnity with the requirements for the degree of Master of Applied Science

Graduate Department of Metaiiurgy and MateriaIs Science University of Toronto

Toronto

Copyright by S. Sathian, 1999

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METALLURGICAL AND MECHANICAL PROPERTIES OF

NCBASED SUPERALLOY FRICTION W L D S

by

Sujith Sathian

Master of Applied Science

Graduate Department of Metallurgy and Materials Science

University of Toronto

1999

ABSTRACT

Friction welding is a solid state joining process and can be used to successfully join

similar and dissimila. Ni-based superalloy base materials. The objective of this project

involves the optimization of welciing parameters and post weld heat treatment procedures

dunng sirnilar Ni-base wrought alloy and dissimilar (Ni-base wrought alloy/Ni-base cast

alloy) friction welding. Extensive microstructural analysis was carried using a combination

of optical, scanning electron microscopy, transmission eIectron microscopy and atomic force

rnicroscopy. The application of a Ire-solution + stabilization + precipitation] post weld

treatment procedure produced the optimum balance of joint t ende strength and ductility

properties in Ni-base wrought alloy fiction joints. TEM analysis provided an insight

concerning the various mechanisms that occur in friction welded joints, particularly with

regard to changes in precipitate chemistry and in particle distributions.

ACKNOWLEDGMENTS

1 would like to express my sincere gratitude to my supervisor, Prof Tom North for

his persistent encouragement and inspiration during the course of this study.

I a h like to acknowledge Prof. G. Benâzsak and Prof' Z, Wang and Prof G.

Weatherly of McMaster University for their participation in useM discussions.

I would like to thank Pratt & Whitney Canada, Longueuil for fiinding this project.

My sincere thanks to Mt. Dave Thomas, Ms. Isabelle Bacon, Mr. Andy Weaver, Dr. Simon

Durham and Mr. Alain Bouthillier of Pratt & Whitney Canada for their usefùl suggestions

and technical discussions.

1 am grateful for the award of a University of Toronto Open Fellowship and a Pratt

& Whitney Canada Graduate Scholarship.

The technical support of Mr. Fred Neub and Mr. Sa1 Boccia and the administrative

efforts of the office st&s are deeply appreciated. I would like to thank the members of our

welding research group for their valuable discussions.

Finally, 1 would like to thank my wife, S h i , and my parents for their support and

encouragement.

TABLE OF CONTENTS

ABSTRACT

ACKNOWLEDGEMENTS

LIST OF FIGURES

LIST OF TABLES

CHAPTER 1: INTRODUCTION

CHAPTER 2: LITERATURE REVIEW

Metallurgy

Friction Welding

Problems Associated with Similar

Friction Welds

Problems Associated with Dissimilar

Friction Welds

CHAPTER 3: EXPERLMENTAL PROCEDURE

3.1 Ni-base Wrought AlloyMi-base

Wrought Alloy Friction Welds

xvii

Preliminary Tests to Generate the 34

Baseline Data

Opthkation of Friction Welding 36

Parameters

ûptimization of the Post Weld Heat 37

Treatment Procedure

3.2 Ni-base Cast Aiioy/Ni-base Wrought AUoy 38

Dissimilar Friction Welds

3.3 Experimental Testing 39

3.4 Transmission Electron Microscopy 41

3 -5 In-Situ Transmission Electron Microscopy 43

3.6 Atomic Force Microscopy 45

CaAPTER 4: RESULTS AND DISCUSSION OF 46 - 108

Ni-BASE WROUGHT ALLOY FRICTION W L D S

Generating the Baseline Data

4.1.1 Base Metal Microstructure

4.1.2 As- Welded Joint Microstructure

4.1.3 Combineci Effects of Friction

Pressure and Time

4.1.4 Rotational Speed and Weid quality

4.1.5 Forging Pressure and Weld Quality

4.1.6 Muence of Base Metai Condition

Optïmization of Welding Parameters

4.2.1 Weld Profiles

4.2.2 Hardness Profiles

4.2.3 Tensile S trength Properties

Optirnization of Post Weld Heat

Treatrnent Procedure

4.3.1 Microstructural Aspects

4.3.2 Hardness Profiles

4.3.3 Tensile Strength Properties

Tensile Strength of Ni-base Wrought Alloy

Base Material

CHAPTER 5: RESULTS AND DISCUSSION OF 109 - 121

Ni-BASE CAST ALLOY/Ni-BASE WROUGHT

ALLOY FRICTION WELDS

5.1 Ni-base Cast M o y Base Metal Microstructure 1 09

5.2 Welding Parameters and Post Weld 11 1

Heat Treatment

5.3 Hardness Profiles 117

CHAPTER 6: CONCLUSIONS 122- 123

FUTURE WORK 123

REFERENCES 124 - 128

LIST OF FIGUIWS

Figure 1

Figure 2

Figure 3

Figure 4

Figure 5

Figure 6

Figure 7

Figure 8

Figure 9

Figure 10

Figure t 1

Figure 12

Figure 13

Figure 14

Figure 15

Rupture strength behavior of three superalloy classes

(Fe-Ni, Ni- and Co-based superalloys)

Evolution of Ni-based superalloy microstructures

Microstructures of Ni-base cast ailoy and Ni-base

wrought alloy base materials

Schematic diagram indicating the relationship between the

percent volume hction of y' and the content of

hardeners (Al+TI) in superalloy base material

Schematic diagram showing the relationship between the

solution treatment temperature and the (AhTi)

content of the Ni-based superalloy

Pseudo binary diagram for Ni-based superalloys

Tirne-Temperature-Transformation diagram for Ni-base

wrought alloy base material

Primary and secondary y' particles in Ni-base cast

alloy superalloy base material

Basic steps during fiction welding of Ni-based superalloys

Direct - drive fiction welding machine

Different stages during direct dr ive fnction welding

Influence of fiction pressure and rotational speed on the

bum-off rate during welding

The relation between fnction pressure and forging

pressure on the notched tensile strengths of dissimilar

Al based MMC/ATSI 304 SS fiction welds

Reference chart indicating the weldability of different

material combinations

DBerent microstructural regions in friction welds:

Zpl is the Mly plasticised region,

vii

Figure 16

Figure 17

Figure 18

Figure 19

Figure 20

Figure 2 1

Figure 22

Figure 23

Figure 24

Figure 25

Figure 26

Figure 27

Figure 28

Figure 29

Figure 30

Zpd is the partly deformed region,

Zud is the undeformed region

Experimentally measured and calculated fluïd flow

region in aluminum aiioy ~ c t i o n welds

The relationship between the width of the flow region

and the fnction pressure applied during welding

operation ( m w s indicate the flow regions)

Predicted solute distribution within the HAZ in Al-Mg-Si

alloy fiction welds

Schematic representation of the HAZ hardness distribution

following fnction welding of Al-Mg-Si alloy base material

Spiral defect formation in aluminum alioy fiction welds

Particle agglomeration in yttcia containhg ODS alloy

fiîction welds

Intennetallic formed at the interface of a dissimilar

Al-based MMC/staidess steel fiction weId

Martensite formation in the stainless steel substrate in

a dissimilar Al-based MMC/stainless steel fiction weld

Materid transfer in a dissimilar Al based MMC/stainless

steel fiction weld

Microcrack formation in a dissimilar Al based

MMC/stainless steel fnction weld

Eutectic Formation at the bondline of a fiction joint

between Nickel and Magnesium

Sectionhg of Ni-base wrought dloy fiction welds prior

to metallography

Schematic showing the measurement of axial shortening

during fiction welding

Location of micro-hardness tests in Ni-base wrought

alloy fiction joints

Tensile sample design (ail dimensions are in impenal Units)

Figure 3 1

Figure 32

Figure 33

Figure 34

Figure 35

Figure 36

Figure 37

Figure 38

Figure 39

Figure 40

Figure 41

Figure 42

Figure 43

Figure 44

Figure 45

Figure 46

Figure 47

Figure 48

Extraction of TEM test sample following friction welding

Extraction of discs h m the foi1 using the EDM Process

Location of off-centered discs for TEM examination

Jet thinning of the discs

Heating chamber used during in-situ TEM microscopy

Working p ~ c i p l e in AFM Microscopy

SEM microstructure of solution-treated Ni-base

wrought d o y base material

Hïgher Magnification Mew of solution-treated Ni-base

wrought ailoy base material

Microstructure of hardened Ni-base wrought alloy

base material

TEM microstructure of solution-treated Ni-base

wrought alloy base material

XEDS pattern of the solution-treated Ni-base

wrought alloy y matrix

XEDS pattern of y' particles in the solution-treated

Ni-base wrought alloy base material

TEM microstructure of hardened Ni-base wrought

alloy base materid

XEDS Pattern fiom a grain boundary carbide in hardened

Ni-base wrought alloy base material

XEDS analysis of y matrix in hardened Ni-base

wrought alloy base material

XEDS pattern fiom a y' precipitate of the hardened

Ni-base wrought alloy base material

TEM micrograph fiom the dynamically recrystallised zone

(solution-treated Ni-base wrought alloy fiction weld

in the as-welded condition)

TEM micrograph showing the hcture of particles

in the partiy deformed region in a solution -treated

Figure 49

Figure 50

Figure 5 1

Figure 52

Figure 53

Figure 54

Figure 55

Figure 56

Figure 57

Ni-base wrought aiioy fnction weld

Particle segregation in the undeformed region of a

solution-treated Ni-base wrought ailoy fnction weld.

As-weldeà condition.

TEM rnicrograph showing the dynamically

recrystaïfised region in hardened Ni-base wrought

alloy fiction weld. As-welded condition.

Segregation of y' in undeformed region of hardened

Ni-base wrought alloy fnction weids.

As-welded condition.

TEM micrograph of a carbide contained in the y matrix of 56

hardeneà Ni-base wrought ailoy fiction welds.

AFM Image of Ni-base wrought alloy base materiai 59

(the arrow shows the grain boundaq)

AFM Image of the HAZ region in a fiction weld produced 60

using the conditions; Friction Pressure: 350 MPa

(tirne: 1 Os), Forging Pressure: 350 MPa (the: 2s),

Rotational Speed: 1000 rpm

Welded joint produced using a low fiction pressure

and fiction time (fnction pressure = 250 MPa,

fnction time = 2s). The other parameters were; forging

pressure = 250 MPa, forging t h e = 1 S.

Grain size variation in a traverse fkom the edge of

the partially recrystallised region into the as-received

Ni-base wrought alloy base matenal. These welds were

made using the conditions:

A. Friction Pressure: 250 MPa (time: 2s) + Forging Pressure:

250 MPa (time: ls), B. Friction Pressure: 300MPa (time: 6s)

t Forging Pressure: 300 MPa (time: ls), C. Fnction Pressure:

400 MPa (time: 14s) + Forging Pressure: 400 MPa (time: 2s).

Muence of rotational speed ( b m 1000 rpm to 64

2000 rpm) on grain size variation measured at the

component centerline. As-welded Ni-base wrought alloy

fiction joint.

Figure 58 Grain size changes on either side of the bondline

(in the stationary and rotating) components of the

fiction weld. As-welded Ni-base wrought alloy fiction joint.

Figure 59 Microhardness profiles in as-welded Ni-base 66

wrought aUoy/Ni-base wrought alloy joints in a traverse fkom

the bondline into the base matenal

(al1 measurernents at the centerline of the component).

The welding conditions comprised:

A. Forging Pressure: 200 MPa (tirne: 2s)

B. Forging Pressure: 350 MPa (time: 2s)

C. Forging Pressure: 500 MPa (time: 2s)

(Friction Pressure: 350 MPa, Rotational Speed: 1000 rpm)

Figure 60 Microhardness values before and &er post weld 68

heat - treatment in a traverse from the bondline into the

solution-treated Ni-base wrought alloy base metal.

The welding conditions compnsed:

Friction Pressure: 350 MPa (Friction Time: 10s)

Forging Pressure: 350 MPa (Forging Tirne: 2s)

Rotational Speed: 1000 rpm

Figure 61 Micro-hardness results before and &er post weld

heat- treatment in a traverse h m the bondline into the

hardened Ni-base wrought alloy base metal.

The welding conditions comprised:

Friction Pressure: 350 MPa (Friction Time: 10s)

Forging Pressure: 350 MPa (Forging Time: 2s)

Rotatiod Speed: 1000 rpm

Figure 62 Ni-base wrought alloymi-base wrought alloy welded

joint produced using a fiction pressure = 350 MPa and a

fiction tirne = 2s.

The other welding parameters were forging pressure

= 350 MPa, forging time = 2s and rotationai speed: 1000 rpm

Figure 63 Influence of fiction pressure on the amount of 73

axial shorteaing during fkiction welding of Ni-base

wrought atloy base material

Figure 64 Flash regions produced using different fiction pressures 73

(275,325 and 375 MPa) during Ni-base wrought

alloy/Ni-base wrought alloy fnction welding.

The other parameters were fnction time = 1 Os, forging

pressure = 350 MPa, forging tirne = 2s and

rotational speed = 1000 rpm.

Figure 65 Weld profiles produced using different fiction pressures 75

(275 and 325 MPa) during Ni-base wrought

aNoy/Ni-base wrought alloy fiction welding.

The other parameters were Friction Time = los, Forging

Pressure = 350 MPa, Forging Time = 2s and

Rotational Speed = 1000 rpm.

Figure 66 (a-c)Hardness profiles at the bondine in Ni-base

Figure 67

Figure 68

wrought alloyMi-base wrought alloy welds produced

using different fiction pressures

(275,300,325 and 375 M'Pa). The other welding parameters

were fiction tirne: 10s. forging pressure: 350 MPa,

forging tirne: 2s, rotational speed: 1000 rpm.

Typical stress -strain cuve produced during testing of 78

Ni-base wrought ailoy fiction joints.

The welding conditions comprised:

Friction Pressure: 375 MPa (Friction Time: los),

Forging Pressure: 350 MPa (Forging Time: 1 Os),

Rotational Speed: 1000 rpm

Tensile strength properties (yield strength and uitimate 79

tensile strength) of Ni-base wrougbt ailoy friction welds

Figure 69 Tende Ductility (total elongation) during tende testing of 79

Ni-base wrought alloy fiction joints

Figure 70a-b Fracture sufaces of broken tende sample fiom 81

Ni-base wrought alloy weld

Figures 7 1 a-b Magnined view of the above sample (see Figure 70a-b) 81

Figure 72

Figure 73

Figure 74

Figure 75

Figure 76

Figure 77

Figure 78

Figure 79

TEM micrograph showing massive precipitation in the 84

dynamicaily recrystailised region of a Ni-base

wrought alloy fiiction weld foilowing p s t weld

heat treatment ( P m

C-curve showing different precipitation mechanisrns in

Ni-based superalioys

XEDS pattern Erom the y matrix in post weld heat treated

PWHT] Ni-base wrought alloy friction weld

XEDS pattern fkom a y' precipitate (using

conventional TEM)

XEDS pattern fkom a y' particle (using STEM mode)

TEM microstructure following [RS + PWHT] treatment of a solution-treated Ni-base wrought

d o y friction weld

Variation in re-solution temperature with [Al+Ti] content 88

in the superalloy

Tirne-Temperature-Transition diagram for Ni-base 88

wrought alloy base material

Figures 80a-d Optical micrograph of the bondline region when an ST 90

heat treatment procedure is followed by PWHT

Figure 8 1 TEM micrograph f?om the dynarnically recrystallised 92

region (solution-treated Ni-base wrought alloy

base material prior to Wction welding)

Figure 82 TEM micrograph f?om the dynamically 92

recrystallised region (solution-treated Ni-base

Figure 83

Figure 84

Figure 85a-c

Figure 86

Figure 87

Figures 88

Figure 89

Figure 90

Figure 9 1

Figure 92

Figure 93a-f

wrought alloy base material pnor to fiction welding)

following in-situ heat treatment

Dynamically recrystallised region (hardened Ni-base 93

wrought d o y base material prior to Ection welding)

TEM micrograph fiom the dynamically recrystallised region 93

(hardened Ni-base wrought alioy base material prior to

fiction welding) foliowing in-situ heat-treatment

Hardness profiles produced using various heat treatment cycles 96

(a) After P WHT Treatment

(b) After RS + P m Treatment

(c) M e r ST + PWHT treatment

Schematic diagram showing the variation in hardness wi th 96

aging tirne and temperature

Schematic illustrating different precipitation effects 98

Typical stress-strain curve for w + P WHT] 98

Ni-base wrought alloy fiction welds

(solution-treated Ni-base wrought alloy base material

pnor to fiction welding)

Tensile strength properties (yield strength and ultimate 99

tensile strength) of solution-treated Ni-base wrought alloy

fiction welds given different heat treatments

Tensile Ductility (total elongation) of solution-treated 99

Ni-base wrought aLloy niction welds given different

heat treatments

Tensile strength properties (yield strength and ultimate

tensile strength) of hardened Ni-base wrought

alloy niction welds given different heat treatments

Ductiiity (total elongation) of hardened Ni-base

w u g h t dloy niction welds given different

heat treatments

Fracture surfaces of broken tensile test samples

extracteci h m heat treated Ni-base wmught alloyMi-base

Figure 94

Figure 95

Figure 96

Figure 97

Figure 98

Figure 99

Figure 100

Figure 101

Figure 102

wrought alloy welds

(a) Prior base metal condition: Solution treated Ni-base

wrought ailoy, Heat treatment: PWHT, (b) Prior base

metal condition: Hardened Ni-base wrought alloy,

Heat treatment: PWHT, (c) Prior base metal condition:

solution-treated Ni-base wrought doy ,

Heat treatment: RS + PWHT, (d) Prior base metal

condition: Hardened Ni-base wrought alloy,

Heat treatment: RS + PWHT, (e) Prior base metal condition:

solution-treated Ni-base wrought alloy,

Heat treatment: ST+PWHT, (f) Prior base metal

condition: Hardened Ni-base wrought alloy,

Heat - treatment: ST + PWHT Tensile strength properties weld strength and ultimate

tensile strength) of as-received Ni-base wrought

alloy base material given different heat treatments

Tensile strength properties weld strength and

ultimate tensile strength) of hardeneci Ni-base wrought

alloy base material given different heat treatments

SEM micrograph of Ni-base cast alloy base material

Magnified view of Figure 96

TEM micrograph of Ni-base cast alloy base material

XEDS pattern kom a cubical y' particle in Ni-base

cast ailoy base material

XEDS pattern h m a spherical y' particle in Ni-base

cast alloy base material

Joint interface in Ni-base cast alloy/Ni-base wrought

alloy fnction joint

As-Welded microstruchire of Ni-base cast

alloy/Ni-base wrought alloy dissimilar fnction weld

Figure 1 03 Crack formation foiiowing pst weld heat treatment 116

of a Ni-base cast ailoy /Ni-base wrought aiioy fiction weld

Figure 104 TEM micrograph of the bondhe region in a Ni-base IL8

cast aüoy/Ni-base wrought ailoy fiction weld

following RS (re-solution) + PWHT (stabilization

plus precipitation) heat treatment

Figure 105 XEDS spectra fkom a sphencal y' particle (hm

Ni-base cast alloy/Ni-base wrought alloy Ection welds)

Figure 106 XEDS pattern h m the y ma& (Ni-base cast

ailoy/Ni-base wrought d o y fiiction weld)

Figure 107 Hardness profiles in Ni-base wrought alloy / Ni-base

cast ailoy fiction welds

A. As- welded condition B. After Re-solution + PWHT

Table 1

Table 2

Table 3

Table 4

Table 5

Table 6

Table 7

Table 8

Ni-base superalloys and their composition

AUoying elements and their effects on superalloys

Typical heat treatment routes for Ni-base wrought aIloy

and Ni-base cast alloy bar materiai

Different types of carbides in Ni-based superalloys

Experhental testing rnatrix during preliminary

fiction welding trials

Experimental test matrix when examining the combined

effects of fiction pressure and time during Ni-base

wrought dloyMi-base wmught alloy friction welding

Experimentai test matrix when examining the

inauence of rotational speed on the Ni-base wrought

dioy Enction welds

Test matrix when exiunining the influence of

forging pressure during fiction welding of 19 mm

diameter Ni-base wrought alloy bars

CHAPTER 1

INTRODUCTION

Ni-based superalloys are dlo ys of Ni - Co - Cr containhg controlled additions of 1 0

or more trace elements. These materials are extensively used during the manufacture of

aircraft aero-engines. However, they are highly susceptible to cracking when conventional

fusion weldùig is used to fabricate them. In contrast, fiction welding is a solid state joining

process and can be used to successfiilly weld a range of Ni-based superalloy base materials.

The key joining parameters during fiction welding operation comprise friction

pressure and time, forging pressure and t h e and rotational speed. It bas been established

[North, 19991 that the fiction pressure is the key variable during the fiction joining

operation. The fiction time mut be sufncient to allow adequate fictional heaîing but not

so long that excess flash is produced during the welding operation. The influence of forging

pressure on nnal weld quality is unclear. However, some investigators have suggested that

longer forging time has a beneficial effect in reducing the residual stress in certain fiction

weld geometry's Pacon, 19991. For this reason, the influence of forging pressure and

joining time on weld quality constituted an important aspect of this program.

Rotational speed is determined by the type of direct drive fiction welding machine

used during fabrication. The torque increases with reduction in rotational speed. The torque

produced during the direct drive fiction joining process is also controlled by fiction

pressure since the range of rotational çpeeds that can be selected is relatively narrow- For

this reason, the present study focused on examining the effects of varying the rotational

speed in the first instance and then, having established an acceptable rotational speed value,

the influence of fiction pressure, fiction the, forging pressure and forging time were

inves tigated.

The fiction joining operation creates a My-plasticised region at the contact region

and heat affect4 zone regions on either side of the bondine. In duminum ailoy welds, for

example, the width of the hliy plasticised region is about 2 mm and the width of the heat

affected zone is about 4 mm on each side of the bondline (depending on the joining

parameters applied). The formation of HAZ regions is important when Ni-based superalloys

are fnction welded, since post weld heat treatment is mandatory. The required heat

treatment temperatures for Ni-base wrought alloy and Ni-base cast alloy base materials are

quite different. The corresponding heat - treatment temperatures are approx. 100-1 SOC

higher for Ni-base cast d o y compared to Ni-base wrought alîoys. When the solution-

treated Ni-base wrought alloy base material is Ection welded the joints are heat-treated

using a re-solution + stabihtion + precipitation treatment. Also in dissimila. Ni-base cast

alloy (available in the hardeneci condition)/Ni-base wrought alloy joints, the welds are heat-

treated using re-solution + PWHT involving a stabilization plus precipitation thermal cycle

corresponding with that for Ni-base wrought alloy base material.

The objective of this thesis is to examine the fnction welding characteristics of

similar Ni-base wrought ailoy/Ni-base wrought alloy joints and ushg this information

develop satisfactory parameter settings for dissimilar Ni-base wrought alloy/Ni-base cast

alloy fiction welding. For this reason the work effort was carrieci out in different stages,

narnel y,

O Investigating the propenies of Ni-base wrought alloymi-bare wrought alloy joints

0 Investigating the properties of Ni-base c m alloyNi-base wrought olloy joints

The factors influencing the mechanical and metallurgical properties of the cornpleted

welds were examineci in detail using a combination of optical, scanning electron

microscopy, fractography, microhardness and mechanical (tensile) testing. Transmission

Electron Microscopy was used to examine the changes in precipitate chemistry, particle

morphology and distribution resulting h m the fnction welding operation and fiorn

subsequent post weld heat-treatments.

CHAPTER 2

Ni-, Fe-,Co- based superaibys

Ni-, Fe- and Co-base s u p d l o y base materials are generally used at temperatures

above 550°C because of their excellent crap strength properties. Although some superalloy

material compositions can be forged and rolled into sheet, the more highly alioyed

chernistries are generaliy processeci as castings. Fabricated components are aiso made using

welding or brazing. However, many highly alloyed superalloy compositi011~ contain large

contents of hardening phases and are difncult to weld since they are highly susceptible to re-

heat cracking followiag post weld heat treatment [Duvall, 1966 and 19691. The mechanical

properties of s u p d o y base materials are controlled by composition and thermal

processing. For example Figure 1 compares the rupture strength properties of three alloy

classes (Fe-Ni, Ni- and Co-base superalloys) [MM, 19721.

Ni-based SuperaUoys

Ni-based superalloys are alloys of Ni, Co and Cr containhg controfled additions of

10 or more trace elements. The evolution of Nickel based superalloy is shown in Figure 2

[ASM Handbook, 19821 and a representative list of Nickel based superalloys and

compositions is presented in Table. 1.

Superalloy base materials derive their strengtb h m a combination of solid solution

hardenen and precipitating phases. Carbides may provide limited stmgthenuig directly or,

more commoniy, indirectly. In addition to those elements that produce solid solution

Figure 1 : Rupture strength behavior of three super alIoy classes (Fe-Ni, Ni- and Co-based superalloys) [MM, 19721.

Figure 2: Evolution of Ni-based superailoy microstructures [ASM Handbook, 19821 (Rm is the creep stress when samples are tested at 870C for 1 0 0 hours)

strengthening and promote carbide and gamma prime formation, other elements (B, Zr, Ce,

Hf) are added to enhance mechanical or chernical properties. Tabk 2 provides a List of

alloying elements and their effects in superailoys. Some elements produce readily

discemible changes in microstructure; while other elanents produce more subtle

microstructurai effects. The most obvious microstructurai effects involve precipitation of

geometrically close - packed phases such as y' and carbides.

Table 1 : Ni-base superdoys and their composition

Others

Elementd additions go into solid solution and provide one or more of the following

effects; improved strength (MoyT* W, Re) pathd, 19851, oxidation resistance (Cr,Al),

phase stability (in case of Ni) or increased volume fiactions (Vf) of favourable secondary

precipitates (Co) Ir\rathal, 19821. Other elements are added which fonn hardenuig

precipitates such as y' (Ni3(Al,Ti) and y" prime (Ni3Cb). y' phase is the key microstructural

feature responsible for the extraordinarily usefbl high temperature strength properties of Ni

and Fe-Ni base superalloys. Minor elernents (C&) are added to fonn carbides and bondes;

these elements plus others (CeJMg) are added for purposes of tramp element control. Some

elements (B, Zr, Hf) are also added to promote grain boundary effects Baldan, 19891.

Many elements (Co,Mo,W,Cr etc.), although added for their favourable alloying

capabilities, c m result, in some circumstances, in undesirable phase formation (sigma(a),

mu@), laves phases).

Table 2 : Alioying elements and their effects on superalloys.

- - E t r r a s n a m cn- Aluniinum, , , . , , , , , , , , , , . , . formr y* NWAî. TU; n t m r d m formrtion ork-gmmml .r

NtTi Titmnium .................. Fonnm r* NidAl. Ti) and MC urbidem Niobium. Tmntœlum, , , - , . , . . Formm body-œntœrœd tmtrrgonrl v" r d MC -bid- C m r b a n , . .................. Formm MC. Me. MaCa rnd &C arbidu; rtrbilitar fcc

matrcx Phorphorus. ............... Promotrr --ml pdpit i t ion of crrbidu Sitragen .................. Forma M t C m wbonitridar Chromium. ................ Oridition rriirtrrrœ; molid molutr'an mtrunehanine MolyW-num; Tua- - . . -. Solid .oIritioir atrœaqChrriis~ forum M& crrbid- Nickel . . -, . - , _ . _ . , _ , _ , _ _ - S t r b i l h C;cc mitrix; rom v*. i d i b i t inforinrtion o fdr leb

rio- p h n m Bor-=: Zi-nium, , . , . . , , , . . Impiovm erœœp propmrliœm; mtad formrtion of m i n bound-

Microstructure

The microstructure of superalloy base materiai is cornplex. It consists of an

austenitic face centred cubic m a t . (y) containhg a varïety of secondary phases. These

secondary phases are gamma prime ((Ni,Co),Al,Ti - intermetallic) and various carbides,

namely, MC &C, M G , and so on. Figure 3 shows optical microstructures of Ni-

base wrought ailoy and Ni-base cast ailoy base materials subjected to hardening treatments.

Cast superalloys generally have coarser grain &es, more aiioy segregation and

ïmproved creep and rupture strength characteristics Wmught superalloys are more uniform

and usuaily have 6aer grain sizes and improved tensie and fatigue strength properties.

Heat treatment routes

The optimum mimstnicture and mechanical properties are achieved using suitable

kat-treatment procedures. Table 3 details the typical heat treatment routes for Ni-base

wrought alloy and Ni-base cast aiioy base maîeriaIs peaver, 19981.

Table 3: Typical heat treatment routes for Ni-base wrought d o y and Ni-base cast alloy base material

Hardeners content

'Hardeners' are the total wt.% of (Ti+Al) in the superalloy. Figure 4 prager, 19681

shows the relationship between the percent volume fiaction of y' phase and the content of

hardeners. A linear relationship is observed. Other elements which are reported to be

present in y' are Nb, V, Hf, Ta, W and Mo.

Heat - Treatmcnt Solution - Tmtment

Stabilhtioh

Recipitation

Solution Treatment Temperatare

Solution treatment temperatures Vary from one alloy to another depending on the

AhTi contents (Le. the content of y') in the base material. Higher solutionizing

temperatures are required for ailoys containing increased Ai+Ti contents. The typical Ai+Ti

content in Ni-base wrought alioy is 4.5% and a solutionizing temperature of 1090°C for 4

hours completely dissolves the y' particles in the base material mi~l~~structure (see Figure 5)

peaver, 1 9981.

Ni-bars wrought .Uoy 109O0C for 4 ho- + Air Cool W C for 2 hours + Ai. Cool 76û°C for 16 hours + Air Cod

Ni-base cast alloy 1 190°C for 2 h o m + Air Cool

1

109o0C for 4 hours + Air Cool 870°C for 20 hours + 1

Air Cool

Figures 3 : Microstructure of Ni-base cast ailoy and Ni-base wrought ailoy base materials

Figure 4: Schematic diagram showing the relationship between the percent volume fiaction of the y' and the content of hardeners (Ai+Ti) in superalloy base material prager, 19681

Phase Diagram

A pseudo binary phase diagram for Ni-based superalloy is s h o w in Figure 6

[Schubert, 19801. The fouowing points should be noted:

The solution temperature varies k m one superalloy to another. For example,

alloy # 3 requires a higher solutionizing temperature thm alloy # 1. n ie

temperature difference is entirely dependent on the amount of y' prime in

superallo y.

Following the solution treatment, the alloys are heat treated to promote further

hardening. Ailoy #1 requires only one heat treatment cycle to precipitate al1 the

particles while alloys #2 and #3 require hvo different heat treatment cycles for

this to occur (see the figure 6). These subsequent heat treatments comprise a

stabilization treatment at 840°C for 2h and a precipitation treatment at 760°C for

16 h.

TTT Diagram

A typical TTT ( t h e - temperature - transformation) diagram for Ni-based superalloy

base material is shown in Figure 7 [Schubert, 19801 and is helpfid in understanding the

infiuence of t h e and temperature on the final microstructure. However, it must be noted

that for extended aging &es, diajgrams are only available for a few superalloy materials.

Stage 1 in this figure is the solution treatment. Stage II involves stabilization treatment.

Stage III is the precipitation treatment. The following effects occming during these stages:

Stage 1 : Complete dissolution of precipitates in the ma&.

Hot forming and working is carried out at this temperature and

high temperature carbides (MC) and gamma prime precipitate out during cooling to

room temperature.

S- 11: Coarsening of existing y' particles occurs [Baldan, 1989 and 19921.

Precipitation of low temperature carbides (e.g. S C & occurs dong the grain

bowidaries.

Figure 5 : Relation between the soiuuon ueatment temperature and the (-Ti) content of Ni-base wrought alloy peaver, 19981

o d u t i o n trmmorrtum for 7' :

- -------.- t---- 1140- t170°C

Seo -1010 OC

gure 6: Pseudo binary

sm Nucleation of secondary y' precipitation occurs within the existing y'

precipitates.

Precipitate growth is limited by low temperature khetics and the primary and secondary y'

particles in Ni-base cast ailoy (see Figure 8) [Antolovich, l982].

Carbides are important coastituents in superalloys. Carbides provide particle /

matrix strengthening- Different types of carbides are present Ui superalioys depending on

their composition and some of the important types are MC, M&, M,C, T C , (where M

stands for one or more metal atoms). Among the most cornmon are M& and MC

carbides. Table 4 lists various carbides found in different Ni-based superalloy base

materials.

Table 4: Different types of carbides in Ni-based superalloy

Allq cm-

ticooel718 MAR M-200

MAR M446 Nimmle 75 Rmb 41

Udimet 700-

The characteristics feahires of the different carbides comprise:

a) Carbide type - MC Tt is stable at elevated temperatures

Precipitation of MC is associated with the effective removal C from the matrix

and thus Little C is available for the forniaton of other carbides (e.g. K 3 C &

Time (s) Figure 7: Tirne - Temperature - TU-L~~U~UVU ~ ~ ~ q a m for Ni-base wrought alloy base material [Schubert, 1980)

Gamma Prime

>amma Prime

Figure 8: Primary and secondary Y' particles in Ni-base cast alloy base materid [Antolovich, 19821

Precipitation of MC carbides improves high temperature tende strength

properties.

Large amouuts of carbide precipitation result into poor weldability.

In Nb-containing superalloys (e.g. N i o n i c 80), formation of NbC during heat

treatment is sluggish compared to Tic. Addition of Hf in ailoy IN 792 replaces Ta in the

MC carbide particle paldan, 1989, Nathal, 19891. Further, Hf additions alter the

solidification sequence of the carbides in superalloys [Lecompte, 19881 during casting and

alter the carbide morphology and distribution during casring operatiom.

b) Carbide type - Mac

M6C (W, Mo&) has a cubic cubic lattice structure and forms, for example, in

Rene77 base material. M,C carbides are more stable than S C 6 carbides. However, at

elevated temperatures M,C carbides degeneration has been noted in Rene 41 and Mar-M-

252 base materials.

C) Carbide tjpe - Mg. M,C, forms at temperatures just below 1 050°C in Nimonic 80 base material. At

lower temperatures, M,C, carbides decompose to &C6 with a reverse transformation being

possible at temperatures exceeding 1060°C [Garosshen, 19851. Formation of M,C, has a

favourable infîuence on superalloy properties because this reduces the amount of &,C6

particles at grain boundaries and improves the creep Me of the material.

4 Carbide type - MJ3C6

%C6 has a cubic structure and forms preferentiaiiy at grain boundary regions.

Discrete carbides are the most favourable since they pin and limit grah boundary

movement. However, blocky and continuous carbides are highly deletenous [ASM

Handbook, 19721. Continuous carbide formation occurs in the temperature range of 870°C

to 980°C with the particles having an aspect ratio of 30 to 50 which favours crack nucleation

at the interface between the carbide and the matrix.

In H,- containing superalloys (e.g.DS 2Oûû He base material), Hf ties up C and

promotes the formation of a continuous network of carbides. The factors that determine

continuous carbide foxmation are the precipitation kinetics, the extent of prior precipitation,

the MC distribution and the yo' chemistry.

2.2. Friction Welding

Basic Mechanism

Friction welding is a solid-state joining process that produces coalescence due to the

heat developed between two surfaces when mechanicaiiy induced nibbuig motion occurs.

Basic steps in the fiction welding pmess are indicated in Figure 9. The welding process

can be characteriseci as a number of distinct stages. During Stage I, the substrates are

brought into contact at a low applied load (see Figure 9a) and the defornation process is

dominated by fictional Wear. In stage II, the applied load is substantially increased and

considerable fictional heating occurs at the bondline (see Figure 9b) and a balance of strain

hardening and softening processes is attained. During Stage III, fictional heat generation is

terminated, and the applied stress is substantiaüy increased to forge the heated material on

either side of the bondlule (see Figures 9c and 9d). An image of U of Toronto's direct-drive

fiiction welding machine is s h o w in Figure 10. The characteristics feahires of continuous

drive fiction welding are shown in Figure 11 [WRC Bulletin # 2041.

Weiding Variables

A number of parameters are important during the ltiction joining and comprise:

Friction pressure

Friction Time

Rotational speed

Forging pressure

Forging Time

Effects of Welding Parameters

Rotational speed produces the necessary relative velocity at the fayhg surfaces.

Lower rotation speeds produce higher torque values and necessarily involve senous work-

piece clamping problems and sometimes have been associatecl with non-uniform upsetting.

The equilibrium torque and the bum-off rates depend on the work-piece diameter, the

applied pressure and the speed of rotation. Equilibrium torque and bum-off rate are linearly

related with fiiction pressure and are inversely related with rotational speed (see Figure 12).

-1- CL.. .

Figure 9: Basic steps during fiction welding of Ni-based superalloys BulIetin # 204, 19821

Figure 10: Direct - drive fkiction welding machine

Figure 1 1 : Different stages during directdrive fiction weldhg Bulletin # 204, 19821

0.20 Sg.clman d m 0.75 In. db.

Pressure, psi

Figure 12: Influence of fiction pressure and rotational speed as the burn-off rate during welding pllis, 19921

[Ellis, 1992 and Hollander, 19631. A 1inear correlation exists between the equilibrium

torque and bum off rate over the range of Wction pressures comrnonly used, the slope being

dependent on the rotation speed selected.

The effects of fiction pressure and forging pressure on the notched tensile strengths

of dissimilar Al based MMCIAISI 304 Stainless Steel iiiction welds are illustrateci in Figure

13 Worth, 19951. It is apparent that the notched tensile strength increased when the fnction

pressure increased. However, it is worth noting that forging pressure had not a neghgible

effect on weld tensile strength,

Materials Welded

In p ~ c i p l e , fiction welding can be used to join alrnost any material combination. The

following matenal classes can be successfully readily joined;

Al1 types of steels

Aluminium and its alloys

Copper and its alloys

Titanium and its alloys

Nickel, cobalt and its aIloys

Dispersion strenghend alioys

However, base materials that contain weakening phases (e.g. graphite in cast irons

and manganese sulphide in steels) produce joints, which have poor strengths and

watisfactory microstnictures. A number of dissimilar metal combinations exhibit poor

weldability, e.g. aluminum alioydstainless steel and aluminum ailoylsteel, titanium

alloys/steel joints. hadequate joint strength properties result nom the formation of

intermetallic compounds at the bondlùie region. A reference weldability chart (indicates

various dissimilar metal combinations) is s h o w in Figure 14 W C Bulletin # 2041.

Figure 13 : The relation between fiiction pressure and forging pressure on the notched tensile strength of dissimilar Ai based MMC/AISI 304 SS Ection welds Worth, 19951

Figure 14: Reference chart indicating the weldabiiity of different matenal combinations [WRC Bulletin # 204, 19821

Microstructurai regions

The fiction welding opemtion markedly alters the microstructure of materiai on

either side of the bondine, i.e., in the heat-affecteci zone. The HAZ region can be divided

into a number of distinct regions, see Figure 15 wddling, 1994- 11. The three main regions

of specific interest are:

The fiilly plasticised region (Zpl) where the material is abIe to accommodate the

plastic strain by dynamic rezovery (or recrystallisation) of the microstructure.

The partly deformeci region (Zpd) where the plastic deformation is

accommodated by an increase in the dislocation density in the matrix grains. In

this region the temperature is sufnciently high to facilitate the dissolution of the

base metal hardening precipitates.

The undeformed region (Zud) that is characterized by partial reversion of the

base metai precipitates.

Modeling of the Welding Process

Bendzsak and North pendzsak, 19961 modeled the width of the flow region during

fiction welding operation by assuming that the materials in the contact region could be

treated as a highly viscous fluid. Assuming steady state conditions (see Figure 16) the width

of the flow regions is determind by the relation:

where p is the viscosity, r is the radius, a is the rotational velocity, p is the coefficient of

fiction, % is the outer radius of the test sample and Pa is the fiction pressure. Satisfactory

agreement was observed between the calculateci and the measured flow widths in Aluminum

alloy 606 1 friction joints. The width of the fully plasticised region in Al based MMC /

MMC friction welds was inversely pmportional to the fiction pressure applied during

welding (see Figures 17) Worth, 1998-11 -

Figure 15: Different microstructural regions in fnction welds wddling, 1994-11 Zpl is the M y plasticised region Zpd is the partly defomed region Zud is the undefonned region

-6CK) O 500 Axial Distance (pm)

Figure 16: Experimentally measured and calculated fluid flow regions in aluminum alloy fiction welds pendzsak, 19961

Figure 17: The relation between the width of the flow region and the Kction pressure applied during welding operation (arrows indicate the flow regious) Forth, 1998- 1 1. The welding conditions for welds A and B were as foiiows:

A. Friction Pressure : 60 MPa (tirne: 2s), Forging Pressure : 60 MPa (time: 2s) B. Friction Pressure : 180 MPa (time: 2s), Forging Pressure : 60 MPa (tirne: 2s) Rotational Speed : 2000 rpm

2.3. Problems Associated with Shilar Meta1 Friction Welds

a) Alterations in Precipitate Distniution and Morphology : Friction welding alters the

particle distribution wasiderably when age-strengthened base materials are fabricated.

Middlùig and Grong Wddhg, 1994 -11 developed a generic mode1 for the coupled

reversion for rod-shaped MhSi particla in Al-Mg-Si aluminium d o y s and applied this

during friction welding. Based on this approach, the volume hction (t) of precipitates

decreases fiom its initial value (fo), according to the relation:

where dt is the tirne increment at a temperature (T), t,' is the time required for complete

dissolution at the temperature concenieci, and a is the tirne exponent. The variation of t,'

with temperature is given by the relation :

where t,' is the time required for complete dissolution at the reference temperature Tr,., QS is

the metastable solws boundary enthalpy and Q, is the activation energy for diflkion of

particles in the undeformed matrix.

Using the above equations it is possible to calculate the solute distribution within the

of fiiction welds. The predicted particle dissolution cuves are shown in Figure 18.

b) Hardness Variations in the HAZ : Aiterations in the particle distribution, morphology and

microstnicture in the HAZ have a considerable influence on the mechanical properties of

completed welds (on the hardness profile, tende strength properties). The HA2 can be

divided into three different regions of interest according to the particle distribution and

general microstructure:

The fûlly plasticised region is where dynamic recovery of the microstructure

occurs. In this region the particles are cornpletely dissolveci.

The partiy deformed region whexe the plastic deformation is accommodateci by

an increase in dislocation density in the matrix grains and the temperature is

su££ïcient to facilitate dissolution of particles.

The undefomed region is characterized by the partial reversion of the base metai

precipitates.

Middiing and Gmng modeled the above regions when examining HAZ hardness

variations in Al-Mg-Si fiction welds wddling, 1994-23.

The net precipitation strength inmement, Amp , was caiculated using the relation :

Anp = k2f7f0 where (W) is the particle volume k t i o n and k, is the kinetic constant. This equation c m

be rewritten as

a = (~-o-/a--omin) = (Aq, /A%& = f7fo where cmin denotes the intrinsic matrix strength following complete particle dissolution and

a- is the original base metal strength.

The above equation provides a basis for obtaining information occurring reaction kinetics;

for example:

(i) Reversion Model: during particle dissolution the volume fiaction fell eom its initial

value, f,. The following equation delineated the dimensionless strength parameter within the

partly reverted region of the HAZ:

(ii) Naturd Aging Model: in fiction welding it is important to predict the final strength of

the HAZ following n a d aging. The precipitation strengthening increment was calculated

fiom the relation:

(iii) Work Hardening: dislocations are generated in the matruc matenal adjacent to bondline

as a result of straining. The strength contribution due to this factor a, was given by the

relation:

where y is a constant, which is a characteristic feature of the material selected. It proved

possible to calculate the HAZ strength distribution d e r friction welding by coupling the

above quations. Figure 19 shows schematic representaîion of the superimposed hardness

profiles produced using the above models. Partial dissolution of particles promp ted

softening in friction welds [MSu, 19911.

c) Spiral Defect Formation: A typical spiral defect obsenred on the fkacture surface of an Al

based MMC/A1 based MMC fiction joint is shown in Figure 20 Worth, 1998- 11. Spiral

defects are formed as a result of the formation of discontinuities in the flow of plasticised

materiai in the contact region. Spiral defects are fluid flow - induced defects formed when

material, magnesium nch segregates and oxide inclusion transfer to embedded regions close

to the stationary boundary of the joint. They remain trapped there for the reminder of the

welding cycle. The magnesium content of spiral defects ranged h m 15 wt % to 55 wt % in

different regions of spiral defects. Further, localized regions of a low melting point eutectic

material were observed and contained high magnesium contents. The spiral defects acted as

the sites for preferential failure sites during both fatigue testing and notch tensile testing of

MMC joints.

d) Particle Agglomeration and Particle Fragmentation: Particle agglomeration was observed

in MA 9561MA 956 fnction welds and was associated with the agglorneration of srna11

diameter (20 to 30 nm) yttria particles (see Figure 21). Oxidation early stage of the joining

operation also facilitated the formation of agglomerated particles. As expected, the creep

rupture properties of the fnction welds were detrimentaily affkcted by the particle

agglomeration process. Similar effects were also observed in FeA140 grade ODS

(containing yttria) alloy fiction welds [Inksoo, 19981. These particle agglomeration results

are quite different fkom those produced when Al based MMC material containhg much

larger (3 Fm to 45 diameter) reinforcing A$03 particles was fiction welded. In

MMC/MMC joints, the welding operation hgmented, not agglomenited, large 40,

particles Forth, 1996- 1,1997 and 1998-41.

1 2 3

Axiil distance . Z [mm]

Figure 18 Predicted solute distribution within the HAZ of Al-Mg-Si alloy friction welds Wddling, 1994- 11

Figure 19: Schematic representation of the HAZ hardness distribution welding of Al-Mg-Si alloy base matenal wyhr, 19911

Figure 20: Spiral defect formation in a l h u m alloy fiction welds FJorth, 1998-11

Figure 2 1 : Particle agglomeration in ytûia-containhg ODS alloy fiction welded joints NO*, 1996-11

2-4, Problems Associated with Dissimilar Metal Friction Welds

During the dissimilar metal welding process the microstructural changes are quite

different h m those in similar metal jcints. This is due to the diffient thermo-physical

properties (the melting points, thermal diftbivities, conductivity and diffusion coefficients)

and mechanical properties (elastic moâulus, yield strength) of the contacting substrates

[Sassani, 19981. Inadequate weld mechanical properties have been reported when dissimilar

substrates are fiction welded due to the formation of brittle intermetah phases at the

bondline- Further, brittle phases cm be formed in the HAZ, e.g. strain induced martensite in

dissimila. stainless steel /Ai based MMC friction welds. Since dissirnilar welding is an

important aspect in this thesis these topic will be discussed fùrther:

a) Intermetallic formation at the joint interface: Many metallic alloy combinations do not

produce satisfactory weld strength properties. This occurs since the joint strength is

detrimentaily affected by the foxmation of hard, brittle intermetallic phase at the bondline

region. During tensile loading, the intermetallic layer acts as a site for crack initiation and

propagation. For example, it has been reported that fnction joints between C u M contained

a thick intermetallic Iayer while Cu-70%WfAl welds contained thin intermetallic layers the

dissimilar joint interface. Similarly, FeAI and Fe2Al,, intermetaiiics have been observed in

fiiction welds between stainless steel and aluminium alloy base materials. Also, Fe,+, and

Fe$,, have been observed in fiction joints between Al and mild steel. The detrimental

influence of intermetallic layers on joint mechanical properties becomes apparent when a

critical intermetallic layer thickness is exceeded. A critical intermetailic layer thickness of 1

to 2 microns has been reported in AYC steel, Al / AIS1 3 16 stainless steel and in TUAISI 304

stainless steel fiction welds (see Figure 22) Forth, 1996-2 3.

b) Martemite Formation: When some dissimilar combinations are welded, local deformation

of the harder substrate occurs. For example, when welding Al based MMC to stainiess steel,

the differences in the flow strrngths of the harder and the softer substrates are considerable

so that almost al1 the deformation and plastic flow occurs in the softer substrate. However,

the harder substrate is still subjected to local deformation and TEM microscopy has

co-ed the formation of strain induced martemite in the stainless steel component of

MMC/AISI fiction welds (see Figure 23) [North, 1998-2 and 1998-31.

c) Matexial Transfer h m One Substrate to Another: The initial stage in the fiction welding

operation is characteriscd by a large number of adhesion/seizure/failure events. These

localiseci events transfer materiai h m one substrate to the other and vice-versa This

explains the high torque values generated at the start of the fiction welding operation. The

torque increases when in the flow stresses of the dissimilar combination increase. It has

been observed that generation of very high torque values can promotes hgmentation and

transfer of stainless steel material at the contact zone of stainless steeVMUC fnction welds

(sec Figure 24) Worth, 19951

d) Microcrack Formation at the Bondline: Microcrack fonnation commonly occurs in

dissimilar Al based MMC/AISI 304 stainless steel fiction welds when the fiction pressure,

fiction t h e and rotational speed are set low values (see Figure 25) Worth, 19951.

e) Trapping of Oxide Films: Trapping of oxides in the contact area is an inberent feature of

fiction welding. #en this occurs it can have a detrimental effect on the strength and

ductility of the cornpleted joints. This problem is severe when the oxide films are highl y

stable and very adherent at the contact surface. This problem is rnost apparent when

aluminium alloy substrates are fiiction welded. In Ti substrates however, the highly adherent

oxides dissolve in the tit;inium substrate at high temperature and their effect on weld tende

strength is negligible.

f ) Eutectic Formation at the Bondline: Eutectics are formed at the interface at low

temperatures (507OC for Ni and Mg) (see Figure 26) wazlett, 19621. Other alioy

combinations that are difficult to weld involve Zr to mild steel, Zr to stainless steel and Al

alloys to Cu.

Figure 22: Intermetaiiïc formed at the interface of a dissimilar Al- fiction weld Worth, 1996-21.

-bas4 MMC/stainless steel

Figure 23 : Martemite formation in the stades steel substrate of MMC/stainless steel friction weld Worth, 1998-21

a dissimilar Al-based

Figure 24 : Material transfer in a dissimilar Al based MMC/stahfess steel fiction weld Worth, 19951

Al based MMC/stainiess steel fiction weld

Figure 26 : Eutectic Formation at the bondhe of a joint between Nickel and Magnesium pazleît, 19621

Objective of the Thesis

Ni-base wrought aiioy and Ni-base cast alloy base materials were selected during

this study of the Mction welding characte&ics of Ni-based superdoys. The general

microstructure of these base materials is very cornplex, with the y matrix being hardened by

means of spherical a d o r cubical y' intermetallic particles and carbides. Ni-based

superalioys are highly susceptiile to cracking (iiquation cracking, re-heat cracking and

ductility dip cracking) during h i o n welding operatioris. Therefore, solid state joining

processes (fiiction welding) are recormnended for their fabrication. However, high

fictional Ioads and torques are rquired in order to produce defect fiee welds. For example,

it will be shown later that fiction welding of 19 mm diameter Ni-base wrought d o y bar

requires a fiction load of 12 T applied for a fictional time of 10s. This compares with

aluminum ailoy fiction welding where the fictional load is 2 T and the friction time is 0.5

S. During fiction welding of AIS1 304 stainless steel base material, the fiction load is 4 T

and the fiction time is 3 S. Moreover, fiiction welding operation generates considerable

microstructural changes at the bondline and in the HAZ regions on either side of the

bondline, and particularly with regard to the particle distribution and chemistry. These

microstructural changes have a significant effect on the tensile strength properties of

completed welds.

The initial work in this thesis focused on Ni-base wought alloy/Ni-base wrought

alloy friction welding. The influence of welding parameter selection and different post weld

heat treatment procedures on the final joint microstructure and tende strength propert-ies

were studied in detail. Transmission electron microscopy, in-situ transmission electron

microscopy and atomic force microscopy techniques were used to characterise the

microstructure. Microstructure - mechanical property correlations were established for Ni-

base wrought alIoy/Ni-base wrought alloy fiction welds. Finally, the results fkom Ni-base

wrought alloy/Ni-base wrought d o y fiction welds were used as the basis for the welding

parameter settings employed during dissimilar Ni-base wrought alloy/Ni-base cast alloy

friction welding operatiom.

CHAPTER 3

EXPERIMENTAL PROCEUDRE

3.1. Ni-BASE WROUGHT ALLOY FRICTION WELDS

Precipitation hardenable Ni-based superalloys retain theu strength at elevated

temperatures. For this reason it would be expected that very high fictionai loads (e.g.; 12

Tons) acting on 19 mm diameter Ni-base wrought alloy bar would be necessary in order to

obtain sound welds fiee of unbondeci regions. However, M t e d information is available in

the literature conceming fiction welding of Ni-based superalloy and for this reason the

initial welding trials were planned with the objective of geaerating baseline information.

The results fkom these preliminary trials are described in Section 4.1 of Chapter 4 in the

present thesis.

Following prelimuiary welding trials, the experimental work effort examined the

process envelope and the optimum welding parameter settings required during Ni-base

wrought alloy niction welding. The most important welding parameters (niction pressure,

Wction t h e , forging pressure and rotational speed) were varied systematically and their

effect on weld quality was investigated using a combination of optical, SEM and TEM

metallography, micro-hardness testhg and teasile tating. The results are discussed in

section 4.2 of Chapter 4.

Following friction welding the joints were pst weld heat- treated to restore their

mechanical properties and also to relieve the residual stresses generated during the fiction

welding operation. As expected, precipitate morphology and distribution were markedly

affecteci by the post weld heat treatment procedure selected. These fiction welds were heat-

treated using different post weld heat treatmcnt procedures and theù eff- on joint

microstructure and tende strength properties are described in Section 4.3 of Chapter 4.

3.1.1. Preliminary Tests To Genente the Bweliire Data

During these tests, the as-received Ni-base wrought dloy base materid was Ui the

solution-treated condition. The initial test joints were produceci using 25.4 mm diameter Ni-

base wrought alioy bar. It was quickiy leamed that joining of this base materiai required

high fiction pressures and long fiction times in order to produce satisfactory joints.

Testing of 25.4 mm diameter bars resulted in the University of Toronto's direct drive

fiction welding machine operating at, or above, its maximum capability. Consequently, it

was necessary to reduce the sample diameter in order to be able to d e l y operate within the

machine's capabiiities. For this reason, 19 mm diameter bar matenal was used during al1

subsequent trials (see Table 7).

During direct drive fiction welding, the weld is made between stationary and

Table 5 : Experimental testing ma& during preliminary fiction welding trials

rotating components. The rotating section is much shorter than the stationary section. In the

tests at U. of Toronto, the length of the stationary wmponent is 300 mm; it is 100 mm in the

case of the rotating component. It should be noted that these dimensions are the minimum

lengths for adequately safely holding the test sections during the fiction welding operation.

BAR A

BAR B

BAR C

' Ni-base wrought alloy bars A and B had different grain sizes ' ~ h e as-received base material was in solution-treated condition

Bar A' - Diameter

Average Grain Size

Tests Conducted

Bar B' - Diameter

Average Grain Size

Tests Conducted

Bar C - Diameter

Average Grain Size

Tests Conducted

25.4 mm

40-60 microns

Combined effects of fiction pressure and tirne

19 mm

120 microns

Effect of rotational speed (rpm)

19 mm

40-60 microns

Effect of forging pressure

Effect of base metal condition2

Methodology

The following welding parameters were examined:

a) Combined Effects Resulting Friction Pressure and Fricton Time Variations

The initial trials (Test # 1 and 2) were cSmed out usbg fiction pressures ranging

fiom 250 MPa (for a fiction thne of 2s) to 400 MPa (for a Eriction time of 14s). Tests # 3

and 4 were carried out using intexmediate fiction pressure and fiction time values (see

Table 8 below):

Table 6: Experimental test matrix when examinuig the combined effects of fiction pressure

TEST #4

400

14

400

2

and time during Ni-base wrought alloy fiction welding

b) Effects of Rotational Speed

in this test series, the rotational speed was varied £iom 500 rpm to 2000 rpm (ail

other joinuig parameters being held constant). The experimentai parameters are shown in

Table 9.

WELDING CONDITION AND

MEASURED VARLABLES

Friction Pressure (MPa)

Friction Tirne (s)

Forging Pressure @Pa)

Forging Time (s)

Table 7: Experimental test matrix when examining the influence of rotational speed on the Ni-base wought alloy fiction welds (using 25.4 mm diameter bar) TEST # 1 ROTATXONAL SPEED (rpm) 1 WELDING PARAMETERS

Rotational speed was kept constant at 2000 rpm Diameter of the bar used: 25.4 mm

TEST # 1

250

2

250

1

TEST # 2

300

6

300

1

Test # 5

Test # 6

Test#7

Test # 8

TEST #3

400

IO

400

2

2000

1500

1000

500

Friction pressure: 350 MPa

Friction thne: 10s

Forging pressure: 350 MPa

Forging the : 2s

c) Effects of Forging Pressure

The forghg pressure was v e e d h m 200 MPa to 500 MPa, all other joining

parameters being held constant (see Table 10).

Table 8: Test matrix when examining the influence of forging pressure during fiction - - - - welding of 19 mm diameter ~i-base-mught ailoy bars 1 TEST # s. 1 Forging Ressure (MPa) 1 Welding Parameters held constant I

d) ïnfluence of Base Metal Condition

Test # 9

Test # 10

Test # I l

In addition, the influence of initial Ni-base wrought alloy base material condition - whether it was in the hardened or solution-treated condition - on friction welding

performance was investigated. In these tests, the welding parameters compnsed:

Friction Pressure : 350 MPa Friction Time : 10s Forgiag Pressure : 350 MPa Forging T h e : 2s Rotational Speed : 1OOO rpm Bar Diameter : 19mm

200

350

500

3.1.2. Optimization of Welding Parameters

Friction Pressure: 3 50 MPa

Friction The: 10 s

Rotational Speed: 1 0 0 rpm

Forging Tirne: 2 s

The work carried out early in this study when generating the baseline data provided a

preliminary estimate of the fiction welding parameters needed to produced sound fiiction

Ni-base wrought alloy fiction welds using 19 mm diameter bar matexid. These parameters

comprised fiction pressure: 350 MPa, fiction time: los, forging t h e : 350 MPa, forging

time: 2s and rotational speed: 1000 rpm. These preliminary welding parameter settings

served as the basis for fiuther work aimed at determining the optimized welding parameters

during M e r Ni-base wrought alloy friction joining operations

The innuence of fiction pressure variation was examined in detaü, using the

conditions indicated below:

1 Solution Treated Ni-base wrought allorbars I v

Friction Welding Friction Pressure Variation: 275,300.325.350 and 375 MPa The aram met ers k e ~ t constant were:

Friction Time: 10s Foraina Pressure: 325 MPa Fornina Tirne: 2s Rotational S~eed: 1000 mm

Following the fiction welding operation al1 welded joints were subjected to a heat

treatment comprising re-solution + PWHT (stabilization plus precipitation). This heat-

treatment cycle comprised:

Re-solution 9 8 0 " ~ for 2 hours + Air Cool Stabilization 8 4 0 " ~ for 4 hours + Air Cool Precipitation 7 6 0 " ~ for 16 hours + Air Cool PWHT' Stabilization + Precipitation

3.1.3. Optimization of the Post Weld Heat Treatment Procedure

As mentioned earlier, as-received Ni-base wrought alloy base material was in

solution-treated condition. Hardened Ni-base wrought ailoy base materiai was produced by

heat treating the as-received base material using a [stabilization + precipitation] treatment

and this base material was employed d u d g a nurnber of welding trials. Tests were carried

out to examine the welding characteristics of solution-treated and hardened Ni-base wrought

ailoy base materials. The Ni-base wrought alloy bar materials were fiction welded using

the following fiction welding parameters:

- .. -- .

1 note that in the remainder of this thesis PWHT means a (stabilization + precipitation) heat treatment. It is not

a acronym for post weld heat treatment

Friction Pressure : 325 MPa Friction T h e : 10s Forging neJsure : 325 MPa Forging Time : 2s Rotational Speed : 1ûûû rpm Bar Diameter : 19 mm

M e r fiction welding the test joints were subjected to three heat treatments; (PWHT,

Solution-trcatad N i - b Friction 1. PWHT wmught d o y base materiai '

J Weldiag 2. RS+PWHT

* b 3. ST+PWHT

The heat-treatment temperatures and holding tirnes were:

Solution-treatment (ST) : 1090°C for 4 hours + Air Cool Re-solution (RS) : 980°C for 2 hours + Air Cool Stabilization : 840°C for 4 hours + Air Cool Precipitation : 760°C for 16 hours + Air Cool PWHT = [StabiIization + Precipitation]

1 PWHT 2 RS+PWH 3 ST+PWHT

H a r d e d Ni-base wrought alioy base materimaterial

3.2. Ni-BASE CAST ALLOYl Ni-BASE WROUGHT ALLOY

DISSIMLAR FRICTION WELDS

The as-received Ni-base wrought dloy base material was in the solution-treated

condition and Ni-base cast alloy was in the fully hardened condition (hardened Ni-base cast

alloy base material was produced by heat-treating the base material using a [stabilization + precipitation] treatment). The initial test welds were produced using a range of different

fiction pressures; 275,300,325,350 and 375 MPa The other welding parameter settings

were fiction tirne: 10 s, forging pressure: 350 MPa, forging tirne: 2s, rotational speed: 1 O00

Friction Welding

+

rpm. 19 mm diameter bars were used throughout. Foliowing fiction welding ali welded

joints were subjected to a heat treatment comprising re-solution + PWHT (stabiiization plus

precipitation), i.e. ;

Re-solution (RS) : 980°C for 2 hours + Air Cool S tabilization : 840°C for 4 hours + Air Cod Precipitatioa : 760°C for 16 hours + Air Cool PWHT = [Stabiiization + Precipitation]

3.3. EXPERIMENTAL TESTING OF SIMILAR AND DISSIMILAR Ni-

BASE ALLOY FRICTION WELDS

a) Metallography: Al1 test welds were cut and sectioned using a water-cooled &O3 cemnic

wheel (see Figure 27). The sectioned Ni-base wrought alioy welds were then rnolded and

polished to a 1 micron finish using diamond spray. The polished samples were then etched

using Marble's reagent (10 grams of CuSO4 + 10 ml HN03 in 100 ml of Water) with fkesh

etchant was being prepared in each case.

Cutting operaîion Slicing of the welded joint

Figure 27: Sectioning of Ni-base wrought alloy fiction welds prior to metallography

b) Measurement of Axial Deformation: The length of the rotating and stationary components

was rneasured pnor to the fiction welding operation. Following fiction welding the axial

length difference was recorded as the amount of axial shortening (see Figure 28 below).

Figure 28: Schematic showing the measurement of axial shortening during fiction welduig (Amount of axial shortenhg : X-Y)

C ) Micro-Hardness Profiles: A 200 g load was used during micro-hardness testing with

indentations being made at 200 microns intervals h m the joint centerline into the stationary

and rotating Ni-base wrought aUoy components (see Figure 29). AU hardness values were

made at the component centerline up to a distance of 10 mm h m the joint interface.

Figure 29: Location of micro-hardness tests in Ni-base wrought ailoy/Ni-base wrought aîloy joints

d) Tende Testing and Fractography: The tensile test specimen geometry is s h o w in Figure

30. Holding the externometer was particularly difncult due to the test specimen gauge

length. A strain rate of 0.05 in./iiJmin was employed during mechanical testing and al1

broken tensile samples examined using SEM hctography.

Figure 30: Tensile sample design (dl dimensions are in imperid units)

3.4. TRANSMISSION ELECTRON MICROSCOPY

TEM specimens were extracteci h m Ni-base wrought aUoy similar and Ni-base cast

alloyMi-base wrought aiioy dissimilm fiction welds which had been made using the

following parameters:

Friction Pressure 350 MPa Friction Tirne 10s Forging Pressure : 350 MPa Forging Time 2s Rotational Speed : 1 0 rpm Bar Diameter 19 mm

The Ni-base wrought aüoy fiction test joints were then given the following heat treatments:

Ni-base cast alloy/Ni-base wrought alloy fiction welds were given a re-solution + PWHT

(stabilization plus precipitation) heat treatment:

1. As-welded 2. PWHT 3. RS+PWHT

*

Solution-trcated Ni-base wrought alloy base ninttrial

i #

1. -4s-wtlded 2. P m 3. RS+PWHT

Hardened Ni-base wrought aüoy base materid

Friction -D Welding

wrought alloy base mterial

r a

- -

Friction Welding

1

1 NiOb- cast aUoy bue 1

*

-, Friction Weldhg b 1. kewtldcd

2. RS+PWHT

The niction welds were sectioned using a water-cooled ceramic wheel as shown in

Figure 3 1. The extracted components were mechanicaily thinned to a thickness of 75 - 150

microns. This was carrieci out m a n d y using Sic emery paper. A foil thickness of 75-100

microns produced the high quality images during TEM microscopy.

Figure 3 1 : Extraction of TEM test sample following fiction weld

An EDM process was used to extract the discs. This process is extremely slow and

time consuming. Discs were extracted at the bondline and also at off-centered locations in

the fiction weld (see Figure 32). The offlcentered locations correspondeci with the region

containing partly re-oriented grains in completed welds (see Figure 33).

Figure 32 Extraction of discs b m the foil using the EDM process

dise extracted .hm the dynamicall y recrystallised region

Figure 33: Location of offkentered discs for TEM examination discs

The jet polishing (Taiupol) technique was u s d to thin the TEM discs. Tub ,., conditions were apptied:

Solution used : 20 % Perchloric Acid + 80 % Ethanol Temperature o f the solution : - 40°C Potential between electrodes : 16 V Average current : 0.05 A

During the thinoing process, the temperature was maintaineci at - 40°C. Extreme

care was taken to maintain the temperature at -OC suice any rise in temperature during

thinning was extremely dangerous (see Figure 34). Adjacent regions (close to the central

hole produced during thinning) were transparent to transmitted electrons and were viewed in

TEM column.

Figure 34: Jet thinning of the discs

3.5 IN-SITU TRANSMISSION ELECTRON MICROSCOPY

The test specimen was placed on a specially designed fitniace containing a M o

resistance-heating coil. The maximum temperature attainable b ide the fiunace was 850°C.

When the specimen was heated it expauded markedly and this made imaging dinicult during

the initiai stages of the heating procas. AU images were coliected using a Video camera

attached to the TEM wlumn. A schematic representation of the equipment is shown in

Figure 35.

Figure 35: Heating chamber used during in-situ TEM microscopy

The as-welded joints were heat -treated inside the TEM column, using the procedure

hdicated below:

Hardened Ni-based wmught He&-treated inside the TEM ailoy in the as-welded condition column at 850°C for 10 minutes

--

Solution - tteatad N i - b d wrought alloy in the as-welded condition r Heat-treated inside the TEM

J

column at 850°C for 20 minutes

3.6. ATOMIC FORCE MICnOSCOPY

The working principle in AFM microscopy is simple and is illustrated in Figure 36.

The specimen holder vibrates due to a piezo-electric mechanism and the fiequency can be

easily varied. The cantilever is aàjusted such that it just touches the specimen and the laser

beam falls on the Sw cantilever and is then reflected back to a spi& diode. This generates a

potential in the split diode and an extemal potential amplifies the signal. The A-B potential

signai is nonnally kept at -2.2 V and the A+B signal is maintaineci at 7 . W . When the

specimen vibrates depending on the d a c e topography, the intensity of the reflected laser

beam on the split diode varies. During AFM examination, the electrical pulses fiom the split

diode are rnonitored continuously and the changes in (A-B)/(A+B) potential are plotted

versus the scanning distance.

Figure 36: Working principle in AFM Microscopy

The thickness of the foil is not very critical; however, great care must be taken not to

bend the foil during poüshing and etching. In the present study the foil was polished

manually using a polishhg wheel to 0.25-micron finish and was then etched using Marble's

(10 grams of CuS04 + 10 ml HN03 in 100 mi of Water) reagent.

RESULTS AND DISCUSSION

CHAPTER 4.1

Ni-BASE WROUGHT ALLOY FFUCTION WELDS

4.1. GENERATING THE BASE LINE DATA

Initial welding trials were carried out to generate the basehe data during Ni-base

wrought alfoy fiction welding. The welding parameters that a81ect the joint quality

comprise fiction pressure, fiction time, forging pressure, forging time, rotational speed and

the prior base material condition. In the present study, Ni-base wrought alloy base material

was weldeà in the solution-treated and hardened conditions. The welding parameters were

varied systematically and their effet on weld quality was investigated using a combination

of metallography and micro-hardtless testing. U of Toronto's Giction welding equipment

has a limitation regarding welding parameter settings. The maximum fictional load that can

be applied to a test piece is restricted to a maximum of 15 T. Further, the machine only

performs well using rotational speeds between 1000 to 2000 rotations per minute. The

initial welding tests provide a rough estimate of the operating envelope required during Ni-

base wrought d o y fnction welding.

4.1.1 Base Metal Microstructure

Ni-based superalloy material microstnictures are complex and comprise an austenitic

fcc matrix (y) contaking a variety of secondary phases. These secondary phases are the

gamma prime ((Ni,Co)&Ti intermetaliic phase and various carbides, namely, MC, MsC,

M7C3, Mac6. Figures 37 and 38 show SEM micrographs of Ni-base wrought alloy base

material in the solution-treated condition. An optical micrograph of hardened Ni-base

wrought alloy is shown in Figure 39. The grain boundary regions are etched preferentially

in this micrograph.

Figure 37: SEM microstructure of solution-treaîed Ni-base wrought alioy base materid

Figure 38: Higher magnification view of solution-treated Ni-base wrought alloy base material

Solution-treated Ni-base wrought alloy is an alloy supersaturatecl containhg different

elements and precipitation occurs when the ailoy is cooled to room temperature foilowing

solution treatment The most cornmon precipitates are y' particles and high temperature

carbides of the MC type. Figure 40 shows a TEM micrograph of solution treated Ni-base

wrought ailoy base material. The maîrix structure comprises unifody distributeci gamma

prime particles contained in a y ma&.

X-ray Energy Dispersive Spectrometry (XEDS) was used to examine the chemistry

of precipitates in the Ni-base wrought d o y base material. High resolution transmission

electron microscopy (HRTEM) was used instead a conventional TEM since the XEDS

pattern produced by the HRTEM microscope was more accurate and diable- Figure 41

shows an XEDS pattern produced by the marrix. The matrix compnsed an ailoy of Ni, Co

and Cr. However, trace amounts of Ti, Al and Mo were also indicated. The XEDS pattem

fiom a coherent precipitate is shown in Figure 42. This particle was low in Cr and Co and

contained mainly Ti, Al and Ni; its composition corresponded with the formdation

Ni3(Cr,Co)AiTi.

The microstructure of hardened base material was quite different fiom that of the

solution-treated Ni-base wrought alloy substrate. Hardened Ni-base superalloy contained

large volume fiaction of gamma prime particles and carbides compared to the alloy in the

solution-treated condition. A h y the hardening treatment promoted extensive carbide

precipitation dong grain boundaries, see Figure 43. An XEDS pattern nom a grain

boundary carbide is shown in Figure 44. No traceable contents of Al or Ti were present in

the carbide. The carbide was rich in Cr and cuntained smaller amounts of Co and Ni and

confirmed with the M d 6 type where M corresponds to Cr and Ni. Figure 45 shows an

XEDS pattern fiom the matrix and Figure 46 shows the XEDS pattem fiom the gamma

prime particles. The maîrix compnsed an alloy of Ni, Co and Cr. However, small amounts

of Ti, Al and Mo were indicated.

Figure 39: Microstructure of hardened Ni-base wrought alloy base material

XEDS of the mat&

-XEDS of the precipitate

Figure 40: TEM microstructure of solution-treated Ni-base wrought alloy base material

Client r Sugi Spthirn Job : Job numbor 387

CP* SpsetNrn 2(Sf12JQe f1:lT)

Energy (kew

Figure 4 1 : XEDS pattem of the solution-treated Ni-base wrought aiioy y matrix

Client : Suai Sathian ~ o b : Job numbor 387 Spectrum 1 < 5 1 l ~ l l : t l >

cps ?

Figure 42: XEDS pattern of y' particles in solution-treated Ni-base wrou&t alloy base material

XEDS spot

- XEDS analysis

Figure 43: TEM microstructure of the hardened Ni-base wrought alloy base material

Figure 44: XEDS Pattern h m a grain boundary carbide in hardened Ni-base wrought ailoy base material

Client : Sugi S m t h l r n Job : Job numbor 367

cps <5/12JB8 lP:OS>

Figure 45 : XEDS analysis of the y matrix in hardened Ni-base wrought aiioy base material

Figure 46: XEDS pattern fiom a y' precipitate of the hardened Ni-base wrought alloy base material

4.1.2. As-Welded Joint Microstructure

TEM discs were extracted h m the dynamically recrystaliïsed region and also nom

the partly defomed region in completed joints (see Figure 33). Figure 47 shows a TEM

micrograph of the dynamïcally recrystaliised grains in the as-welded joint. The grain size

was approx. 2 - 3 microns in diameter. This location corresponds with the M y plasticised

region (Zpl), where the material is able to accommodate plastic strain via dynamic recovery

(or recrystallisation) of the microstructure. Dissolution of y' is facilitated since these

particles are coherent with the rnaûix. Therefore, fiction welding creates a region of highly

alloy super-saturation at the bondline of as-welded joints.

Figures 48 shows the particle segregation and Figure 49 shows the particle fracture. The

TEM specimen was extracteci h m the off-centted location, Le. h m the partly deformed

region (Zpd), where the degree of plastic deformation is accommodated by an increase in the

dislocation density in the matrix gains. The base material grains were orientateci dong the

flow Iines due to fiction welding operation caused particle segregation while movement of

dislocation resulted the particle fhcturing.

Similar results were observed when hardened Ni-base wrought alloy base material

was fkiction welded. A TEM micrograph of the dynamically recrystallised grains in an as-

welded joint is show in Figure 50. No precipitates were observed in the dynamically

recrystallised region. Figure 5 1 shows a micrograph of the partly deformed region (Zpd).

Particle segregation was observed in this region. Figure 52 shows MC carbides, which are

easily differentiated fiom the gamma prime particles. The y' particles were coherent in

nature and were somewhat opaque to the transmitted electrons. However, the carbides

scatter the electrons and therefore appear dark.

TEM microscopy did not produce a complete p i c m of the particle

distribution over an area of 10 microns2. For this reason, Atomic Force Microscopy was

used to image the particle distribution over an area 15 micron x 15 micron. The scanned

area in an AFM cm be varied easily: 0.1 micron x O. 1-micron to 15 micron x 15 micron

(max.) regions can be investigated.

dynamicaily recrystallised grains

Figure 47: TEM mimgraph h m the dynamicaliy recrystallised zone (solution-treated Ni- base wrought alloy fiction weld in the as-welded condition)

Figure 48: TEM micrograph showhg the k t u r e of particles in the partly deformed region in a solution -treated Ni-base wrought alloy fkction weld

particle segregation

Figure 49: Particle segregation in the undefomied region of a solution-treated Ni-base wrought aiioy fkiction weId, As-welded condition

dynamicaiiy recrystalIised grains

Figure 50: TEM micmgraph showing the dynamically recrystallised region in hardeneci Ni- base wrought dloy fiction welds. As-welded condition.

particle segregation

Figure 5 1 : Segregation of y' in undeformed region of hardened Ni-base wrought alloy friction welds. As-welded condition

Figure 52: TEM micrograph of a carbide containeci in the y ma& of hardened Ni-base wrought alloy fiction welds

AFM is generally used for imaging d5u:e topographical features, e.g., the

topographical changes occurring when Si semiconductor material is doped with different

impurity elements. The limitation ofthe AFM technique is that the test specimen

preparation is very difncuit and it is difficuit to identifL the exact location that is being

imaged by the microscope.

Figure 53 shows an AFM image of Ni-base wrought alloy base metal. A grain

boundary region is clearly evident. Also, a topographical image of the dynamically

recrystallised region in the welded joint is shown in Figure 54. As rnentioned above it was

difficult to be precise conceming the area that was being scannai during AFM micmscopy.

4.1.3. Combined Effets of Friction Pressure and Friction T h e

The initial weiding trials were cmied out ushg fiction pressures ranging h m 250

MPa (for a fnction time of 2s) to 400 MPa (for a fiction time of 14s). Joints # 2 and 3

were made using intermediate fiction pressures (300 - 400 MPa) and fiction times (6 - 10s)

(see Table 8).

Unbondeci regions were fo& at the periphery of welded joints produced using low

fiction pressures (c 400 MPa) and friction times (40s). This is apparent in Figure 55,

which shows a micrograph of the welded joint produced using a fiction pressure of 250

MPa and a fiction time of 2s.

Figure 56 shows the grain sue variations in a traverse h m the bondiine into the as-

received base material. The bondline microstructure wmprised partially recrystallised

grains having diameters e 10 microns. Further h m the bondline, the microstructure

contained equiaxed dynamically recrystailised grains (having grain sizes in the range 10 - 20 microns). With increasing distance h m the bondline, the microstructure comprised te-

orienteci base metal grains. In efféct then was a transition h m a fine grain microstructure

at the bondihe to larger re-orieated base metai grains and fbally into the as-receivexi Ni-

base wrought alloy base material (where the grain sue was 60 microns). It would be

expected that the presence of very fine equiaxed grains at the joint centerline rnight have a

detrimental effect on elevated temperature properties (since grain boundary sliding

preferentially occurs at grain boundary regions that are aligned perpendicular to the

direction of tensile stress [Thamburaj, 19951).

The welding parameters applied during welding tests # 1 (friction pressure 250 MPa,

fnction time 2s) and # 2 (fnction pressure 300 MPa, fiction t h e 6s)) resulted in the

formation of unbonded regions near the joint periphery. The formation of these defects can

be explained as follows. When low fiiction pressure and iow fnction times are applied, the

filly plasticised region develops rapidly at the component centerline. However, formation

of the fully plasticised region at the periphery depends on transfer of plasticised material

fiom the component centerline towards the periphery. When the fnction welding time is

v e v short or the fiction pressure is low, this inhibits transfer of M y plasticised material

fiom the component centerline to the joint periphery and creates a region of weahess.

Satisfactory bonding across the whole joint interface and satisfactory weld profiles

were produced when higher fiction pressures (> 350 MPa) and longer fnction times (> 10 s)

were applied. However, since the fiction welding machine at University of Toronto had a

maximum rated capacity of 15 tons, a fiction pressure value around 400 MPa (on 25.4 mm

diameter bars) represented its lirnit of machine performance. For this reason, the peak

fiiction pressure used during subsequent fiction welding trials was necessarily limited to

350 MPa.

Based on the above commentary, it is concluded that the application of higher fiction

pressures (>350 M'a) and long Wction times (>los) are prime requirements during Ni-base

wrought alloy fiction welding and that this welding parameter combination produces welds

fiee of unbonded region located at the joint periphery.

Figure 53: AFM Image of Ni-base wmught aUoy base material (the arrow shows the grain boundary)

Figure 54: AFM Image of the HAZ region in a fiction weld produced using the conditions; Friction Pressure: 350 MPa (tirne: 1 Os), Forging Pressure: 3 50 MPa (the: 2s), Rotational Speed: 1000 rpm

Base me!

- xysîaiiised

iented base

ion

1 , - =

Figure 55: Welded joint produced using a low fiction pressure and fiction time (fiction pressure = 250 MPa, fiction time = 2s). The other parametm were; forging pressure = 250 MPa, forging time = 1 s, Diameter of the bar = 25.4 mm.

Distancm h m tho dge of the partklly nctyst8lli.d mgion (mm)

Figure 56: Grain size variation in a traverse h m the edge of the partially recrystallised region into the as-received Ni-base wrought alloy base material.

The welds were made using the conditions:

A. Friction Pressure: 250 MPa (time: 2s) + Forging Pressure: 250 MPa (time: 1s) B. Friction Pressure: 300MPa (time: 6s) + Forging Pressure: 300 MPa (tirne: 1 s) C. Friction Pressure: 400 MPa (the: 14s) + Forging Pressure: 400 MPa (tirne: 2s)

Rotational speed: 2000 rpm Bar Diameter: 25.4 mm Base metal grain size: 40 - 60 microns

4-1.4, Rotationaï Speed and Weld Qurlity

In this series of tests, the rotationai speed was varied h m 500 rpm to 2000 rpm (dl

other joining parameters k i n g held constant (f?ïction pressure : 350 MPa, fiction time :

los, forging pressure : 30 MPa, forging tirne : 2 s). Details of the experimental parameters

are shown in Table 9. It is weil known that the torque increases significantly when the

rotational speed decreases h m 2000 rpm to 500 rpm. As a result heat generation at the

contact interface markedy increases. Increased heat generation had a signincant infiuence

on the Gnal joint microstructure; e.g., when the rotational speed increased f?om 1000 to 2000

rpm the grain size increased h m 5 to 10 microns. This compares with an as-received Ni-

base wrought alloy base metal grain size of 120 microns (see Figure 57).

The torque produceci by the rubbing suffies substantially increased when the

rotational speed decreases. Consequently, although low rotational speeds inhibited the

formation of unbonded regions at joint peripheries when 25.4 mm diameter Ni-base wrought

alloy bars were welded, this methodology could not be generally applied because it

compromised the safe operation of the fiction welding machine (the safe operating range

being fiom 1000 to 2000 rpm).

Finally, there was a significant difference in the grain sizes measured on either side of

the bondline, i.e., in the stationary and rotating components (see Figure 58). Note that the

diameter of the Ni-base wmught alioy bar used prior to welding was 25.4 mm. This

diameter necessitates the application of high fiction ioads (loads>l ST). That means, the

welding machine was run at its maximum rated capacity and therefore, the results obtained

may not be the true representation. However, it can be presented for the cornparison

purposes. The grain size differences in the stationary and the rotating components resulted

fiom the different rates of heat absoption produceci by the clamping jaws in the case of the

shorter rotating section.

O 0.5 1 1.5 Dirbnce m m the cornponant t.ntreline(mm)

Figure 57: Influence of rotational speed (hm 1000 rpm to 2000 rpm) on grain size variation measured at the component centerline. As-welded Ni-base wrought alloy friction weld.

The measurements were carrieci out f3om the component centerhe

Welding conditions applied:

A. Rotational speed: 2000 rpm B. Rotationai speed: 1500 rpm C. Rotational speed : 1 O00 rpm

(Friction Pressure: 3 50 MPa (tirne: 1 0s) + Forging Pressure: 3 50 MPa (time: 2s) Grain size: 120 microns Bar Diameter: 25.4 mm

Distance from the wmponent centeriine (mm)

Figure 58: Grain size changes on either side of the bondline (in the stationary and rotating) components of the fiction weld As-welded Ni-base wrought alloy fiction joint.

The measurements were carried h m the joint centerline

Welding parameters applied:

Friction Pressure: 350 MPa (tirne: 10s) + Forging Pressure: 350 MPa (time: 2s) Rotational Speed: 1000 rpm Diameter of the bar: 25.4 mm Grain size: 120 microns

4.1.5. Forging Pressure and Weld Qurllty

In these t a c the forging pressure was varied fiom 200 MPa to 500 MPa, aii other

joining parameters being held constant (see Table 10). Forging pressure had no effect on the

grain size of recrystallised grains at the joint centerline. Figure 59 shows micro-hardness

profiles from the bondline into the as-received base metal (all measurments were carried

out at the centerline of the component). It is apparent that increased forging pressure had no

significant effect on hardness values (bearing in mind the typical scatter in resdts produced

during micro-hardness teshg).

4-1.6 Influence o f Base Metaï Condition

In general practice, solution-treated Ni-based superalloy material is fiction welded

and is then post weld heat-treated to produce the required joint strength and ductility

properties for any given application. In the case of Ni-base wrought alloy base material, this

involves the application of a two-stage thermal post treatment comprising stabilization at

8 4 0 ' ~ for 4 hours followed by a precipitation hardening treatment at 760'~ for 16 hours.

The initial stabilization heat-treatment promotes carbide precipitation at grain boundaries

and also coarsens the primary intennetallic phases. The low temperature precipitation

hardening treatment promotes precipitation of srnail secondary intermetallics, which

improve material strength. In the present study, fiction welding of both solution-treated and

hardened Ni-base wrought alloy base materials was investigated. Al1 fiction welded joints

produced using solution-treated and hardened Ni-base wrought alloy base materials were

post weld heat-treated using a stabilization + prezipitation thermal cycle.

Figures 60 and 61 show the micro-hardness results in a traverse h m the bondline into

the as-received Ni-base wrought alloy base metal. The indentation load was 200 g and 10

readuigs were taken on each test sample over a distance of 2 mm. When solution plus heat-

treated Ni-base wrought ailoy base material was fiction welded, this produced a sofiened

zone with a hardness of about 25 Hv lower than the adjacent base material. The hardness

trough was much deeper (about 150 Hv) in fiction welded joints produced ushg hardened

Distance from aie bondline (mm)

Figure 59: Microhardness profiles in as-welded Ni-base wrought alloy friction joints in a traverse fiom the bondline into the base material (al1 measureements at the centerline of the component).

The welding conditions comprised:

A. Forging Pressure: 200 MPa (the: 2s) B. Forging Pressure: 350 MPa (time: 2s) C. Forging Pressure: 500 MPa (the: 2s)

(Friction Pressure: 350 MPa, Rotational Speed: IO00 rpm) Bar Diameter: 19 mm

A - As-welded condition

Distance from the bondline (mm)

Figure 60: Microhardness values before and following post weld heat - treatment in a traverse fiom the bondline into the solution-treated Ni-base wrought alloy base matenal.

The measurernents were taken h m the bondline region

The welding conditions comprised:

Friction Pressure: 350 MPa (time: 1 Os) Forging Pressure: 350 MPa (time: 2s) Rotationl Speed: 1000 rpm Bar Diameter: 19 mm Base metal condition: Solution treated Ni-base wrought alloy

A. As- welded condition B. After PWHT

Distance from the bondline (mm)

Figure 61: Micro-hardness results before and following p s t weld heat treatment in a traverse fiom the bonche into the hardened Ni-base wrought alloy base material.

The measurements were carried out h m the bondline region

Welding conditions compnsed:

Friction Pressure: 350 MPa (time: 10s) + Forging Pressure: 350 MPa (time: 2s), Rotational Speed: 1000 rpm Bar Diameter: 19 mm Base metal condition: Hardened Ni-base wrought alloy

Ni-base wrought aiioy base material. The formation of softened zones on either side

of the bondline resulted h m solution and coarsening (over-aging) of intennetallic phases

and carbide pdc l e s during the thermal cycle in fiction welding. These softened regions

were completely removed foliowing p s t weld heat treatment. Similar hardness troughs

have been reported in the in the HAZ regions of 6Ml- T6 a1uminu.m fiction welds Malin,

19951.

To conclude, similar microstructural features were produced in post weld heat-treated

(stabiiization -t precipitation) joints when the starting Ni-base wrought ailoy base materid

was in the solution treated and hardened conditions. However, this may not mean that the

mechanical properties of the completed joints will necessady be similar. When using a

rotational speed of 1OûO rpm, microstnicturaUy sound, defeît fiee fiction welded joints

having excellent profiles were produced ushg the following welding parameters: a friction

pressure of approx. 350 MPa, a fiction t h e of 10 s and a forging pressure of 35OM.a (see

Figure 62).

Figure 62: Ni-bas fiction pressure =

M k i and a fiction time = 2s. The 0th- welding parameters were forging pressure = 350 MPa, forging time = 2s and rotational spccd: 1ûûû rpm

CHAPTER 4.2

RESULTS & DISCUSSION

4.2. OPTIMIZATION OF WELDING PARAMETERS

Initiai welding trids provideci a preliminary esthate of fiction welding parameter

settings needed to produce sound Ni-base wmught alloyMi-base wrought alloy friction

welds. These welduig parametm comprised a Wction pressure of 350 MPa, a fiiction tirne

of los, a forging pressure of 350 MPa, a forging time of 2s and a rotational speed of 1000

rpm. These parameter senings served as the basis for f.urther studia aimed at optiminng

welding parameter settings. In this work, the fiction pressure was varied systematically,

Le., 275,300,325,350 and 375 MPa The rotational speed was maintaineci constant at 1 O00

rpm for the foliowing reasons. When using rotational speeds exceedhg 1000 rpm, fiction

pressures >>350 MPa were required to produce sound Ni-base wrought alloy friction joints.

However, the application of such high fiction pressures pushed the U. of Toronto fiction

welding machine beyond its operating capability. Similady, the use of rotational speeds

c 1000 rpm had a detrimental influence on the performance of the fiction welding machine

(since the torque increased considerably when rotational speed was reduced). Ail test welds

were subjected to re-solution + PWHT (stabilization + precipitation) heat treatments and

were analyzed using a combination of optical, scanning microscopy, micro-hardness and

tende testing.

4.2.1. Weld Profiles

Figure 63 shows the effects of fiction parameters on the amount of axial shortening

produced during welding. The amount of axial shortening increased when the friction

pressure increased. For example, a fiction pressure of 275 MPa produced an axial

shortening value of 3 mm while a fiction piessure of 375 MPa produced an axial shortening

value of 8 mm.

300 360 Friction Prrwm (MP.)

Figure 63: Influence of fiction pressure on the amount of axial shortenhg during Wction welding of Ni-base wrought alloy base materiai

Figure 64: Flash regions produced using different friction pressures (275,325 and 375 MPa) during Ni-base wrought alloy fiiction welding. The other welding parameters were: fiction time = los, forging pressure = 350 MPa, forging time = 2s and rotational speed = 1000 rpm.

Figure 64 shows CCD images of the flash regions produced during welding using

different fiiction pressures. It is apparent that the amount of flash increased wheu the

niction pressure increased. For example, nibstantiai amounts of flash were produced when

a fiction pressure of 375 MPa was appl id When high niction pressure (>275 MPa) and

the friction times was los, the fbily-plasticised region developed rapidly at the component

centerline and transfer of plasticised material h m the component centeriine towards the

periphery occurred. When the fiction welding t h e was very short or the fiction pressure

was too low, this inhibitecl transfer of My-plasticised material fkom the component

centerline towards the joint periphery. Thus a region of weakness was created at the joint

penphery. Therefore, higher friction pressures (2275 MPa) are recommended based on a

flash formation perspective of Ni-base wrought alloy Ection welding.

The flow region in completed welds was deheated by the presence of dynamically

recrystallised grains. The width of this region decreased when the Wction pressure

increased. (see Figure 1 7). Bendzsak and North modeled pendzsak, 19971 the flow regions

produced during fiction welding and showed that the width of the flow region (h<,))

depended markedly on the welduig parameters selected, e.g.,

where, p is the viscosity, r is any radius of the sample, w is the rotational speed, q is the

coefficient of fiction, & is the outer radius of the sample and Pa is the applied pressure.

This relation indicates an inverse proportionality between fiction pressure and the width of

the flow region produced during the fiction welding operation. Ifit is assumed that the heat

generated at the contact region conducts equdy into the adjoining substrates, it would be

expected that the profiles of the hcat-affected-zone region on either side of the bondiine

would mirror that of the flow region This may explain the weld profiles shown in Figure

65.

4.2.2 Hardness Profiles

Figures 66a to c show hardness profiles in welds produced using different fiction

pressures. As noteâ earlier, a sofiened zone with a hardness of about 25 - 50 Hv lower than

the adjacent base material is produced when solution treated Ni-base wrought ailoy base

material is fiction welded, (see Section 4.1). The softened zones were completely removed

Friction Pressure: 325 MPa Friction Pressure: 275MPa Figure 65: Weld profiles produced using different fiction pressures (275 and 325 MPa) during Ni-base wrought aUoy friction welding. The other parameten were Enction time = los, forging pressure = 350 MPa, forging tirne = 2s and rotational speed = 1000 rpm

following the re-solution + PWHT (stabilization plus precipitation heat treatment).

However, when a fiction pressure of 275 MPa was applied sottened zones were still

apparent in Ni-base wrought alloy base materid immediately adjacent to the bondline.

Kowever, when a fiction pressure of 325 MPa or higher was applied the hardness profiles

were much more uniform (see Figure 66 c - d).

4.2.3. Tensile Strength Properties

Figures 67 shows a typical stress-strain curve produced when testing a Ni-base

wrought alloy/Ni-base wrought alloy weld. In ail the test samples the uitimate tensile

strength was the fiacture strength of the test samples. The ductility was evaluated as the

total elongation measued h m the stress- strain curve. In this thesis, the yield strength was

the stress corresponding to a strain of 0.02%.

Figure 68 compares the yield and ultimate tensile strengths while figure 69 shows the

ductility results. Ni-base wrought alloy fiction joints made using a fiiction pressure of 275

MPa had the poorest strength pmperiies (YS: 840 MPa, UTS: 1270 MPa). The highest

tensile strengths (YS : 925 - 1070 MPa, UTS : 1360 - 1475 MPa) were produced when

intemediate fiction pressures (300MPa to 350 MPa) were applied. In all the cases the

tensile sample failure occurred in base metal away fiom the weld region and the ductility

values exceeded 15% (Figure 69). It should be noted that the elongation measurements were

solely based on the extensometer rendings of broken tensile test samples and therefore

should be considered as approximate estimates only.

The ultimate tensile strength of as-received solution-treated Ni-base wrought alloy

base material was 1100 MPa and the ductilïty was 22%. Higher ultimate tensile strengths

(UTS values > 1275 MPa) and appreciable ductility levels (215%) were obtained when Ni-

base wrought alloy base material was welded using fiction pressures ranging fiom 325 to

375 MPa. Also, as indicated earlier, the tensile test specimen failure occmed in the base

metal away h m the bondline.

(a) Friction Pressure: 275 MPa

4 -2 ,,,je O 2 4

Distance frorn the bondllne (mm)

-4 2 O 2 4

Distance from the bondlino (mm)

(b) Friction Pressure: 325 MPa

-1 O -8 4 4 -2 #' O 2 4 6 8 10

Distance from the bondline (mm)

(c) Friction Pressure: 375 MPa

Figure 66 (a to c): Hardness profiles at the bondiine in Ni-base wrought aiioy friction welds produced using different fkiction pressures (275,325 and 375 MPa). The other welding parameters were Wction tirne: los, forging pressure: 350 MPa, forging tirne: 2s, rotational speed: 1000 rpm.

Figure 67: Typical Stress Strain c w e produceci during testing of Ni-base wrought ailoy fiction joints (Friction Pressure: 375 MPa Friction Time: 10s. Forging Pressure: 350 MPa, Forging Time: 1 Os, Rotational Speed: 1 Oûû rpm)

Yield Strength mtimate Tcosile Strength

Figure 68: Tensile strengths properties (yield strength and ultimate temile strength) of Ni- base wrought alioyMi-base wrought aiioy fiction welds

Figure 69: Tensile Ductility (total elongation) during tensile testing of Ni-base wrought alloy fiction joints

The fkcture surfaces of the broken teusile samples were examineci using scanning

electron mimscopy. Typical k t u r e d a c e s are shown in Figures 70 a and b and

rnagnified views of broken t d l e test specirnens are shown in Figures 71 a and b. Sample

failure occurred at grain boundary regions in an inter-granular fkture mode. In Ni- based

superalloys it is well documenteci thaî the grain ôoundaries are weak compared to the grain

interiors (since precipitation strengthens the grain interiors) [MM, 19821. Therefore, failure

is therefore more likely at grain bomdary regions, Le., the intergrandu failure mode is

more likely. Increasing the fiction pressure oniy had a minor influence, Le. secondary

cracks were observeci in fkactureâ tende test samples produced using high fiction

pressures.

Figure 70 a and b: Fracture SUff'es of broken tensle sample h m Ni-base wrought ailov/Ni-base wrounht allov weld

Figures 71(a-b): Magnified Mew of the above sample (see Figure 70 a-b)

CHAPTER 4.3

RESULTS & DISCUSSION

4.3. OPTIMIZATION OF POST WELD HEAT TREATMENT

PROCEDURE

Post weld heat treatment procedures Vary h m one application to another. For

example, it has been suggested that friction welding of hardened Ni-base wrought alloy bar

produces superior fatigue strength properties compareci to soiution-treated base material.

Following fiction welding the joints were heat-heated using dinerent procedures, namely:

O P m which involves stabilization + precipitation treatments:

S tabihtion : 860°C for 2 hours + Air Cool Precipitation : 740°C for 16 hours + Air Cool

RS (re-solution) + PWHT

Re-Solution : 980C for 4 hours + Air Cool S tabilization : 860°C for 2 hours + Air Cool Precipitation : 740°C for 16 hours + Air Cool

Solution - Treaîment : 1090°C for 4 hours + Air Cool Stabilization : 860°C for 2 hours + Air Cool Precipitation : 740°C for 16 hours + Air Cool

Precipitate morphology and distribution were markedly affectecl depending on the post weld

heat treatment procedure that is applied. In the present study fiction welds produced using

previously found welding parameter settings (fiction pressure: 350 MPa, fiction tirne: los,

forging pressure: 350 MPa, forging tirne: 2s and rotational speed: 1000 rpm) were subjected

to the above heat treatments. The effects of these heat - treatments were investigated using

a combination of transmission electron microscopy, micro-hardness and tende testing.

4.3.1. Microstructurai Aspects

a) Microstructure Following P m (Stabikation + Prenpiation) Treatment

Figure 72 shows massive (0.02-0.03 pm diameter) y' precipitation occurring in the

dynamically recrystalïised region of Ni-base wrought alloy friction welds given a PWHT

(stabilization + precipitation) heat treatment. The precipitation of particles follows a C-

curve relation, i.e., at temperatures exceeding T2 in figure 73, the kinetics of the

precipitaîion are controlled by diaision processes while at lower temperatures (temperatures

a,), they are controlled by nucleation phenornena. At very low temperatures, long tunes

are required for complete precipitation because the diaision rate is very slow. The rate of

precipitation is also very slow at temperatures just below the solvus line (see point 1, Figure

73). In this case nucleation is slow and precipitation is determined by the rate at which

nucleation can occur. Although high diffiision rates exist at temperatures just below the

solidus line, few nuclei are present. At intermediate temperatures, between these two

extremes, the precipitation rate increases to a maximum value so that the time required for

complete precipitation decreases. In this manner, the combination of moderate diffiision and

nucleation rates promote rapid precipitation. This explains why rapid precipitation occurs

during PWHT. In effect, nucleation during the PWHT (stabibtion + precipitation) heat

treatment will be favored thermodynamically rather than via d i h i o n and this will enhance

the precipitation of small particles within the ma&.

The superalloy matrix and the particles were analyzed using X-ray Energy

Dispersive Spectrometer (XEDS). This work was carried out using a conventional TEM

microscope and therefore, the values are not as accurate as in the case of high-resolution

transmission electron microscopy (HRTEM). However, the output is still usefid for

cornparison purposes. The XEDS result fiom the matrix is shown in Figure 74. X-ray

analysis of precipitates was evaluated using two différent techniques; (i) using a

conventional XEDS analyzer, (ii) using the scanning transmission electron microscope

(STEM) mode. In the normal TEM mode, condensing the bearn down to an appropriate size

massive precipitation of 7' particles

- XEDS of particle

XEDS of matrix

Figure 72: TEM micrograph showing massive precipitation in the dynarnically recrystallised grains of a Ni-base wrought alloy fiction welds following pst weld heat treatrnent (P WHT)

kinetics of precipitation: difkion controlled

Figure 73 : C<urve showing different precipitation mechanism in Ni-based superalloys

for micro-analysis can misalign the iiiumination system. Hence, the STEM mode is

preferred for XEDS anaiysis. The pattern produced using the nrst technique is shown in

figure 75 while the chart produced using the STEM technique is shown in figure 76. Both

approaches produced similar d t s , bearing in mind that only comparative elmental output

not absolute values are considemi. The XEDS results indicated that the matrix contained

approx. 3 wt% Ti and approx. 64 wt.% Ni. nie particles comprised 75% Ni and 10% Ti

and contained smaller amounts of Co and Cr. Similar results were produced when the

precipitates were examineci using the STEM mode. In this case, the values comprised 73

W.% Ni and 6.5 wt.% Ti and trace amounts of Cr and Co.

b) Microshrrciure Following RS me-Solution) + PWHT Treatment

Coagulation of y' particles occurred during the re-solution heat treatrnent. Also, fine

secondary precipitation occurred during the stabilization plus precipitation heat treatment

(see Figure 77). The coagulation priocess limited massive precipitation of the fine gamma

prime precipitates that were observe- in samples given the PWHT treatment (see Figure 72).

The re-solution temperature (980°C for 4 hours) exceeded T2 in Figure 73 and were

determined by diffûsion mechanisms and this enhanceci coagulation controlled the

precipitation kinetics.

Re-solution treatment temperatures vary fiom one superaiioy to another, depenciing

on the [Al+Ti] contents (on the y' content) in the Ni-based superailoy. Hïgher re-solution

temperatures are required for alloys containing high [Ai+Ti] contents. The heat-treatment at

980°C for 2 hours partiaiiy dissolves precipitates in the microstructure (see Figure 78) in the

Ni-base wrought alloy.

Figure 79 shows a typical 'IïT (the-temperature-transfomation) diagram that is

helpfbl in understanding the influence of time and temperature on the final microstructure

following heat treatrnent. Stage 1 in Figure 79 is the re-solution treatment while Stage II is

the stabilization treatment. Stage III is the precipitation heat-treatment. The following

processes occur during heat treatment:

X - R f i Y L t : 100s R t : 119s

P r s 1

Rem:

Figure 74: XEDS pattern h m the y matrix in p s t weld heat-treated (PWHT) Ni-base wrought dloy fiction welds.

Conventionai XEDS .:-RAY ~ t , : 100s P r s t : 100s Rem: 8s R t : 1215 17%Dt

Figure 75: XEDS pattem h m a y' precipitate (using conventional TEM)

X - R A Y L t : 1005 P r s t : R t :

100s 1225 18%Dt 1

STEM MODE

Rem: 0s

Figure 76: XEDS pattern fkom a y' particle (using STEM mode)

Figure 77: TEM miciostructure followkg RS + PWHT treatm base wrought alloy fiction weld

.ent of a solution-treated Ni-

Figure 78: Changes in re-solution tempcrature with [AhTi] content in the superalloy [weaver, 19981

Figure 79: Time-Temperature-Transition diagram for Ni-base wrought alloy base material [Schubert, 19801

a) Stage L-

0 There is partial dissolution of precipitates in the matrix, and high temperature

carbides (MC) and gamma prime precipitate during cwling to rmm temperature.

b) Stage lr:

0 Coarsening of existing y' phases occurs and also precipitation of low temperature

carbides (e-g-, MuCs) dong grain boundaries.

c) Stage iJlk

0 Nucleation of secondary fine y' precipitation occurs within the existing y'

precipitates. However, their growth is limited because of the low temperature

kinetics.

These different stages readily explain the coagulation of y' particles observed in welds that

were given a re-solution heat treatment (see figure 79).

c) Microstructure Following ST (Solution - Treatment) + PEUT Treatment

Figures 80 a to d show optical rnicrographs of the dynamicalfy recrystallised region

fomed at the bondline in completed welds. The right-hand micrograph in each set shows

the selected location at higher magnification. Considerable grain growth was observed

when the Ni-base wrought alloy fiction welds were heat-treated using a [ST + PWHT]

thermal cycle. Further, solution-treated Ni-base wmught alloy base material was more

susceptible to grain growth than hardened Ni-base wrought ailoy base material. The

hardened Ni-base wrought alloy base material contains high temperature carbides that are

very effective in pinning the grain boundaries and in reçtncting growth during post weld

heat treatment [Thamburaj, 19851.

d) Microstructure Following In-Situ Heat - Treatment in the TEM Microscope

The samples for in-situ microscopy were extracted hrn the dynamically

recrystallised region in as-welded joints. As pointed out eatlier this region is highly

supersaturateci with ailoying elements. When the test sample was heated in the TEM

column rapid precipitation occumd and the output was imaged. The test section was heated

inside the TEM column using a Mo resistance heating coi1 and the microstructural changes

Base Metal Condition: Hardened Ni-base wrought aiioy; Heat Treatment: PWHT

Base Metal Condition: Solution - Treated Ni-base wrought alloy; Heat Treatment: ST + PWHT

Figures 80 (a to d): Optical micmgraphs of the bondline region when ST heat treatment procedure is followed by PWHT

were imaged using a Video camem. The temperature ioside the fumace was graduaily

raised to 870°C (this was the maximum hanace temperature that could be used safély). This

temperature was marginaüy higher than the stabilisation temperature normaliy employed for

Ni-base wrought alloy base matenal. Figure 81 shows the r d t s produced when examining

as-welded joints made using solution-treated Ni-base wrought alloy base material. During

testing, the sample was heated for approx. 15 - 20 minutes and the ha1 microstructure was

photographed (see Figure 82). Massive nucleation of gamma prime occureci at 870°C and

since this temperature was higher than the stabilisation temperature for Ni-base wrought

alloy, base materid coagulation of the y' phase was apparent late in the heating cycle.

A similar approach was employed when examining as-welded fiction joints

produced using hardened Ni-base wrought alloy base material. The as-welded joint

microstructure is shown in Figure 83. The test specimen was held at 875°C for approx. 5

minutes. Since this holding temperature is higher than the stabilization temperature (860C)

normally used for Ni-base m u g h t alloy, considerable precipitation of y' was observed (see

Figure 84). Also it is worth noting that the jet-thinning process creates minute holes in the

specimen disc. When the test sample is heated these tiny holes expand considerably and

appear in the h a 1 microstnicture as small voids (see Figure 84).

4.3.3. Hardness Profiles

Figures 85 a to c show the hardness profiles in welds produced using different

welding procedures. The circular symbols refer to the use of solution treated Ni-base

wrought alloy base matenai prior to fiction welding; the square symbol indicate the use of

hardened Ni-base wrought alloy base materiai. The influence of using solution treated and

hardened Ni-base wrought alloy base materials prior to ection welding were examined:

a) Solution- Treated Ni-base wrought aihy

When a PWHT (stabiLization + precipitation) treatment was applied following fiction

welding, a hardened zone was fomed in Ni-base wrought ailoy base material adjacent to the

bondline. [Re-solution + P m ] andk [ST + PWHT] resulted in more uniforni hardness

Figure 8 1 : TEM micrograph h m the dynamically recrystailised region (solution-treated Ni- base wrought alloy base materiai prior to fiction welding)

y' precipitates

Figure 82: TEM micrograph h m the dynamicdy recrystallised region (solution-treated Ni- base wrought alioy base material prior to fiction welding) following in-situ heat treatment

Figure 83: Dynamicw tecrystallised region (hardened Ni-base wrought alloy base materials prior to fiction welding)

voids

y' particles

Figure 84: TEM micrograph fbn the dynamicaliy recrystallised region (hardened Ni-base wrought alloy base material prior to fiction welding) following in-situ heat treatment

profiles (see Figures 85 b and 85 c). The base metal hardness values were higher (approx.

4SOHv) when a PWHT treatment was applied. [PS + PWHT] treatment produced lower

hardness values (approx. 42SHv) in Ni-base wrought alloy base material. The most unifom

hardness profiles were obtained when an CRS + PWHT] thermal cycle was applied.

b) Hardened Ni-base woughf alloy

Hardened regions @eak batdlless: approx. 5OOHv) were formed in material adjacent

to the bondline when the fiction welded joints were heat-treated using a PWHT procedure.

When the fiction weids were heat-treated using ST or RS treatments prior to PWHT, the

hardness profile across the weld zone was uniform.

The typical aging (hardness) curve is a hc t ion of temperature (see Figure 86). The

significant feature is that extending the holding time and aging the test specimen too long at

a given temperature will decrease the sample hardness. It c a . be seen h m Figure 86 that

holding the superalloy at temperame Tl may attain a saturation value. However, holding at

temperature T2 wili produce a parabolic hardness relation. In-situ TEM micrographs

showed that the gamma prime particles had an extreme tendency for coarsening (over-

aging). Similar effects have been observed in creep samples of Ni-base cast alloy [Ahmet,

19941.

Temperature Tl corresponds with PWHT heat-treatment while T2 corresponds with

an RS post weld heaî treatmemt. Numerous tiny particles of gamma prime were precipitated

in the dynamically recrystallised region of Ni-base wrought alloy fiction welds following

the PWHT treatment. These particles were uniformly distributed and had an average

diameter of 0.02 microns. Similar precipitation behavior is likely in the Ni-base wrought

alloy base materiai. This would explain the peak hardness of 550 Hv observed in the

dynamically recrystallised zone and also the high hardness values in the base material

(approx. 435HV) foiiowing PWHT heat treatment.

Generally, an RS heat - treatment prior to PWHT (stabihtion + precipitation) is

quite uncornmon for Ni-base wrought d o y base matenal. However, the fiction welding

operation generates residual stress and non-uniform microstructures in the HAZ regions on

(a) After PWHT Treatment

4 -2 p O 2 4

Distance from the bondline (mm)

4 -2 Aje O 2

Distance from the bondline

(b) Mer RS + PWHT Treatment

4 -2 Ap O 2 4

Dîstrnca fmm the bondlino (mm)

(c) After ST + PWIfT treatment

Figure 85 a to c: Hardness profiles produced using various heat treatment cycles.

Figure 86: Schematic diagram showing variation in hardness with aging t h e and temperature.

either side of the bondline. Hence, the RS heat-treatment prior to PWHT is recommended to

relieve residuai stresses and aiso homogenise the weld microstructure. As seen in the TEM

"mgrapghs (see Figure 77), the particles in the dynamically recrystallised region were

0.05 pm in diameter following + PWHT] heat-treatment; this compares with a particle

diameter of 0.02 microns diameter foilowing the PWHT heat-treatment. Consequently, the

RS heat-treatment promotes over-aging and this reduces the haTdlless of the Ni-base

wrought alioy base material. This may explain why an overall reduction in hardness was

observeci in Ni-base wrought aiioy base material following the [PS+PWlWJ heat-treaîment.

However, the ST heat treatment was carried out at high temperature (1090C) for a long

duration (4 hours). This heat-treatment dissolves almost al1 the particles in the base

material. PWHT following the ST heat-treatment has the same effect as the PWHT

treatment with regard to the hardness profile. However, a ST heat-treatment pnor to PWHT

completely eliminates the residual stresses in the welded joint and HAZ regions and even

out hardness variations. In addition, extreme grain growth was observed following the ST

treatment (see Figure 80) and this will have a considerable infiuence on joint mechanical

properties. The different precipitation effects are shown schematicaily in Figure 87.

4.3.4. Tensile Strength Properties

Figure 88 shows typical stress-strain c w e s for Ni-base wrought dloy fiction welds

produced using different post weld heat treatment procedures. In al1 cases, the ultirnate

tensile strength was the fkacture strength of the test sample. The yield strength was the

stress corresponding to a strain of 0.02% and the ductility was the total elongation measured

using the stress- strain cuwe.

Figure 89 compares the yield and ultirnate tende strengths of joints made using

solution treated Ni-base wrought alloy base material while Figure 90 shows the

corresponding ductility values. Figure 9 1 compares the strengths of welded hardened Ni-

base wrought aiioy base material and Figure 92 shows the ductility values.

Foilowing PWHT Following RS+PWHT

Figure 87: Schematic illustrating different precipitation effects.

Figures 88: Typical Stress-Strain c m e s for w+PWHT] Ni-base wrought welds (solution -treated base materiai prior to fiction welding).

alloy fiiction

PWHT R S + M ST +PWHT

Varying heat treatment routes

ml- Yield Strength (MPa)

a Ultimate Tensile Strength (MPa)

Figure 89: Tensile strength properties (yield strength and ultimate tensile strength) of solution-treated Ni-base wrought ailoy fiction welds given different heat treatments.

br i s Mebl : M e c e i v e d Ni48rs wrought dloy

30

PWHT RS+PWHT ST+PWHT

Varying hart treaûnent mutes

Figure 90: Tensile Ductility (total elongation) of solution-treated Ni-base wrought alloy fiction welds following given different heat treatments.

Varying heat treatment routes

Yield Stieagth @Ga)

a t c Tensile Strength (MPa)

Figure 91 : Tensile strength pmperties (yield strength and ultimate tende strength) of hardened Ni-base wrought alloy fiction welds given different heat treatments.

Ba= Metal : Hardoned N i b r - wrought rlloy

PWHT RS + PWHT ST + PWHT

Figure 92: Ductility (total elongation) of hardened Ni-base wrought ailoy fiction welds given different heat treatments.

PWHT treatment produced the poorest weld strength properties (YS: 930 MPa, UTS:

1300 MPa) when solution- treated Ni-base wrought alloy base material was Ection welded.

The highest tensile strengths (YS: 925 MPa, UTS: 1380 MPa) were produced when a re-

solution treatment was applied prior to PWHT. In al1 the cases tende sample failure

occurred in Ni-base wrought dey base metal away h m the weld zone at appreciable

ductility values (elongations X5%) in ali joints. Similar tensile strengths and ductility

results were produced when hardeneci Ni-base wrought ailoy base materiai was employed

prior to welding. Re-solution prior to the stabilization plus precipitation treatment produced

the highest tensile strength properties (YS: 1080 MPa, WTS: 1428 MPa). However, Iower

strengths (YS: 870 MPa, UTS: 12ûûMPa) were found when the fiction welds were heat-

treated using an [ST + PWHT] thermal cycle.

The variation in tende strength properties can be explained via particle - dislocation

interactions. Dislocations are generated during tensile testing and an increase in tensile

strength is synonymous with incfea~ed difficulties in dislocation movement. A moving

dislocation must cut thorough the particles in its path or it must rnove around them.

Dislocations c m move around large diameter particles (Ormwan looping). However,

uniformly distributed srnall particles are cut as the dislocation moves. PWHT heat-treatment

prornoted the precipitation of small particles while the RS + PWHT treatment favored the

formation of both coarse and small particles. Further, RS + PWHT heat-treatment produced

large volume fraction of y' precipitates and therefore produced a matrix haviog high tensile

strength.

In general, solution treated Ni-base wrought alloy base material produced the highest

ductility values while the hardened Ni-base wrought alloy produced highest joint swengths.

In al1 the cases, sample failure occurred away h m the bondline in the adjoining base metal.

This is an indication that the variations in strength and ductility values were largely

determined by the Ni-base wrought alloy base matenal properties. For example, hardened

Ni-base wrought alloy base material containecl large volume fiactions of y' precipitates and

therefore had the highest strength values. High precipitate volume h t i o n in the matrix

decreases ductility.

In conclusion, application of a [RS+ PWHT] treatment produced the optimum

balance of strength and ductiiity properties and in all cases, tensile test specimen failure

occurred in base metal away h m the weld zone.

Typical fiacture surfaces are shown in Figures 93 a to f. The hcture mode was

intergranular since grain boundary regions in Ni- based superalloys regions are weak

compared to the grain interiors (since precipitation strengthens the grain interiors).

Consequently, failure is therefore more likely at grain boundary regions and therefore the

intergranular failure mode is more likely. It is worth noting that hardened Ni-base wrought

alloy base material contained large number of y' particles and carbides. Also, the hardening

treatment promoted extensive carbide precipitation dong grain bomdaries, see Figure 43.

An XEDS pattern h m a grain boundary carbide confimed that these carbides were rich in

Cr. This means that a depleted of Cr layer was created adjacent to the grain boundary

carbides. This depleted zone will produce a thin weak layer in which cracks can easily

nucleate and pmpagate during temile loading- This may explain why secondary cracks were

often observed in the failed hardened Ni-base wrought alloy tensile test samples (see Figures

93 a to 0.

4.4. Tensile Streugth Properties of Base Material

Friction welded joints comprise different microstructural regions, e.g., the region

containhg dynamically recrystallised grains, the heat afTected zone and the as-received base

material. During tensile loading for a given strain rate, these microstructural regions

respond differently. Weaker matenal preferentially necks and tensile failure occurs. In Ni-

base wrought alloy fnction welds the sample failure occurred wholly in the base matenal

indicating that it had the lowest tensile strength. In effect, the tensile strength of Ni-base

base material has a significant infiuence on mechanical properties of the Ni-base wrought

alloy fiction welds.

Ni-base wrought alloy base metai was subjected to different heat - treatment routes;

P m , RS+PWHT and ST + P m . Figures 94 and 95 compare the yield and ultimate

(a) Prior Base Metal Condition: Solution -Treated Ni-base wrought aiioy; Heat Treatment: PWHT

@) Prior Base Metal Condition: Hardened Ni-base wrought aiioy; Heat Treatment: PWHT

103

(c) Base Metal Condition: Solution -Treaîed Ni-1 base wrought doy; Heat Treatment: RS + PWHT

(d) Base Metal Condition: Hardened Ni-base wrc ~ught aiioy; Heat Treatment: RS + PWHT

Figure 93 a-E Fracture d a c e s of broken temile test samples extracted h m heat-treated Ni-base wought alloyMi-base mught dey welds

tensile strengths of solution treated and hardened Ni-base wrought alloy base materials. The

tende strengths of solution-treated Ni-base wrought alloy base material were YS: 590 MPa

and UTS: 970 MPa; these wrnpared with those in hardeneci Ni-base wrought alioy base

material (YS: 740 MPa and UTS: 1140 MPa). The highest tende strength values (YS: 762

MPa, UTS: 1136 MPa) were produced when PWHT hûit-treatment was carrieci out using

solution treated Ni-base wrought alloy base material. The tensile strength values following

the w+PWHT] heat-treatrnent were simila. to the tensile strengths of Ni-base wrought

alloy base material subjected to a PWHT heat treatment cycle: the values were YS: 700 MPa

and UTS: 1090 MPa for the solution treated Ni-base wrought alloy base material and YS:

760 MPa and UTS: 1 153 MPa for the hardend Ni-base wrought alloy base material.

However, the [ST+PWH'l'] heat-treatment produced the lowest tende strength properties; in

solution -treated Ni-base wrought alloy base material the vaiues were YS: 560 MPa, UTS:

864 MPa; in hardened Ni-base wrought d o y base matenal the values were YS: 600, UTS:

93 2MPa receptively .

PWHT heat-treatment pmmoted the precipitation of srnail (0.02 micron diameter) y'

particles while the CRS + PWHT] treatment favored the formation of both coarse (0.05

micron) and small diameter particles (0.01 micron) in the Ni-base wrought alloy base

material. Both these heat- treatments increased the tensile strength of the base material. An

ST heat-treatment pnor to PWHT dissolved ail the y' precipitates in the matrix and the

PWHT treatment promoted fiutber precipitation. However, it is possible that low volume

hctions of y' following [ST+PWHT] heat-treatment led to the lowest tensile strength

values: YS: S60MPa and UTS: 864 MPa for solution - treated Ni-base wrought ailoy base

materiai and YS: 600 MPa and UTS: 932 MPa for the hardened Ni-base wrought alloy base

matenal.

In general, solution treated Ni-base wrought alloy base matenal produced the highest

ductility values (>20%) while the hardened Ni-base wrought alloy produced highest

strengths (11 100MPa). In al1 the cases, specimen failure occurred close to the center of the

gauge length. This is au indication that the strain rate dong the gauge length was essentially

uniform. Hardened Ni-base wrought alloy base matenal had lowex ductility in temis of total

elongation values (approx. 8-100/0 lower) compareci to solution-treated Ni-base wmught

alloy base material.

ST + PWHT heat- tmatment improved the ductility @y 20-27%) in compatison the

to P m and = + P m heat-treatments. Extreme grain growth was observed when the

ST heat- treatment was applied. Moreova, carbides were completely dissolveci during the

heat-treatment cycle resulting in a matrix containing a lower volume h t i o n of the y' and

carbide precipitates. This explains high ductility and the low teasile strengths following the

ST+PWHT heat-treatrnent.

Yield Strength (MPa) m-+ t e Tensile Strength m a )

Figure 94: Tensile strength properties (yield strengh and ultimate tensile strength) of as- received Ni-base wrought alloy base material given different heat treatments.

Yield Strength (MPa) m-+ Ultimate Tensile Strength @Pa)

Figure 95: Tensile strength properties (yield strength and ultimate tensile strength) of hardened Ni-base wrought alloy base material given different heat treatments

CHAPTER 5

RESULTS AND DISCUSSION

Ni-BASE CAST ALLOYINI-BASE WROUGHT ALLOY

FRICTION W L D S

The results found during Ni-base wrought alloy/Ni-base wrought alloy friction

welding served as the basis for the selection of the welding parameters applied during

dissimilar Ni-base cast aiioyNi-base wrought alloy welding. Ni-base cast alloy material

was fiction welded to Ni-base wrought aUoy in the hardend condition and foilowing

welding operation the dissimilar joints were post weld heat - treated. It was previously

found that the optimum mechanical properties (ductility >15% and UTSX 300 MPa) were

produced when a re-solution + PWHT (stabilization plus precipitation) heat treatment was

applied to Ni-base wrought ailoy fiction welds. For this reason, it was recommended that a

RS + PWHT treatment should be applied in the case of dissimilar Ni-base cast alioy/Ni-base

wrought alloy fiction welds. Heat-treated welds were investigated using transmission

electron microscopy.

5.1 Ni-base Cast Ailoy Base Metal Microstructure

Ni-base cast alloy base material is normally available in the hardened condition, i.e.,

the alloy is subjected to a stabilization + precipitation heat treatment. Figure 96 shows the

typicai cast base metal microstructure and coÏnprises of incoherent and coherent y'particles

contained in a y matrix together with MC and M& type carbides. The mismatch between y'

precipitates and the y matriv increases when the volume fiaction of y' in the base material

increases. The shape of the y' particles changes from spherical to cubical when the volume

fiaction of y' in the y matrix exceeds 35 vol.%. In Ni-base wrought alloy base material the

y' particles were spherical; however, they were cuboidal shaped in Ni-base cast alloy base

1 Cast Sb

Figure 96: SEM micrograph of Ni-base cast d o y base mai

Figure 97: Magnified view of Figure 96

110

material (see Figure 97). Further, spherical y' particles have the lowest mismatch while the

cubical y' particles have the highest intemal strah (+ 0.0200% to - 0.0100%)). In spite of

this mismatch, base material containhg cubical y' particles exhibits excellent ductility, a

unique behavior. The inherent ductility of the y' particles prevents severe embrittlement,

uniike the strengthening that is developed by phases which have much higher hafdIless, cg.

alumina in an Al matrix.

Figure 98 shows a TEM mimgraph of Ni-base cast alloy base material, The

structure comprised u n i f o d y distributecl cubical y' particles dong with coherent sphencal

y' particles in a y matrix. The average size of cubic particles was approx. 0.5 microns; the

spherical particles were appox.0.05 microns in diameter. The chemistry of y' precipitates

was examined using X-ray Energy Dispersive Spectrometer (XEDS). High resolution

transmission electron microscopy (HRTEM) was used since the output was more accurate

and reliable. Figure 99 shows an XEDS pattern produced by the cubical particles. The

cubical particles were essentialiy an allay of Ni containhg Al, Ti and W. However, trace

contents of Cr and Co were also indicated. The XEDS pattern h m a coherent sphencal y'

particie is shown in Figure 100. These particles had the same chemistry as previously, but

were high in W and comprised maitily Ti, Al and Ni. These particles corresponded with the

formulation Ni3(Cr,Co)AlTiW.

5.2. Welding Parameters and Post Weld Heat Treatment

The initial friction welding tests were carried out by varying the fiction pressure

nom 275 MPa to 375MPa when using a fiction time of 105, a forging pressure of 350 MPa,

a forging t h e of 2s and a rotational speed of 1000 rpm. The Ni-base wrought alloy

substrate was severely deformed and Ni-base cast d o y base material was largely unaffected

(see Figure 101). However, the final weld interface profile was not perpendicular to the axis

of rotation of the test samples. The high temperature flow stress of Ni-base wrought alloy

base material is lower that that of Ni-base cast alioy base material and this is why the Ni-

base wrought alioy substrate deformed prefmtially and constitutes the principal

component in the flash. During friction welding the contact surface stresses are the highest

Figure 98: TEM micrograph of Ni-base cast d o y base material

Cliant : Sugl Smt t i l rn Job : Job numbœr 387 Spectrum 8 (8/5199 10:32>

Figure 99: XEDS pattern fiom a cubical 7' particle in Ni-base cast ailoy base material

Figure 100: XEDS pattern h m a sphericd y' particle in Ni-base cast alloy base material

Figure 101 : Joint interface in Ni-base cast ailoy/Ni-base wrought alloy fiction joint

at the weld centerline Worth, 19921 and the accumulation of these stretched elements

makes a major contriiution to the flash volume during welding. If deformation of the Ni-

base wrought alloy substrate was the only faetor controlling the shape of the Ni-base

wrought alloy/Ni-base cast aUoy joint interface; it would be curved pnor to the forging stage

in the fiction welding operation. However, since higher temperatures are produceci at the

outer periphery of the joint and this promotes preferential plastic flow in this location so that

the final joint interfce profile is that shown in Figure 101.

Figure 102 shows the as-welded microstructure at the joint centerline. An intennixed

zone was observed adjacent to the bondline. This is the region where elemental diaision

and mechanical mixing occurs as a resuit of the thermal and sWst ra in rate cycle in

friction welding. The width of this intermixeci zone depended on the welding parameters

selected, particularly on the fiction pressure selected. As expected, welds produced using

high friction pressures (375 MPa)-produced wider inter-mixed zones.

Al1 dissimilar Ni-base wrought alloy/Ni-base cast alloy joints were heat-treated using

a re-solution + PWHT involving a stabilization plus precipitation thermal cycle

correspondhg with that for Ni-base wrought alloy base material. Following heat treatment

cracks were observed at the bondline of welds produced using low fiction pressures (( 325

MPa). Figure 1 O3 shows micrographs of weided joints produced using a low fkiction

pressure (Wction pressure = 300 MPa, Wction tirne = 10s) and an intermediate fiction

pressure of 325 MPa (fiction tirne: 10s). Crack fiee dissimilar welds were produced using

friction pressure exceeding 325 MPa.

TEM discs were extracted h m the weld centerline. It is worth noting that Ni-base

cast alloy and Ni-base wrought alloy materials behaved differently while electrolytic

thinnllig process. The Ni-base wrought alloy substrate was thinned preferentially while Ni-

base cast alloy substrate stayed as it was. Thus, the transmission micrograph fiom a

dissimilar weld was of barely acceptable qualïty. For this reasoq a modified TEM

procedure will be required in friture* which will effectvely characterize dissimilar Ni-base

cast alloy/Ni-base wrought alloy welds. A TEM micrograph of the weld centerline

Figure 102: As-Welded microstructure of Ni-base cast aiioyMi-base wrought aiioy dissimilar fiction weld

Figure 103: Crack formation following post weld heat treatment of a Ni-base cast alloyMi- base wrought aiioy fiction weld

following re-solution + P m (stabiiization plus precipitation) is shown in Figure 104. The

location of this sample is close to Ni-base cast alloy substrate. As noted earlier, in Ni-base

cast alloy contains cohererit 0.05 pm diameter spherical particles and 0.5 pm diameter non-

coherent cuboidal y' particles in a y matrix. FoUowing fiction welding considerable

amounts of elemental a i o n take place across the weld interface (into the Ni-base cast

alloy componeni). It is worth noting that Ni-base wrought alloy base material is less rich in

Al, Ti and W than the Ni-base cast alloy alloy. Thus, diaision fiom Ni-base wrought alloy

material into the Ni-base cast alloys substrate w i U dilue the chemistry of the Ni-base cast

alloy alloy. As a result high temperature diffusion during the fiction welding operation

might result in partial dissolution of cuboidal y' particles and dissolve the smaller 0.3 pm

diameter spherical particles in the Ni-base cast dloy base matenal. Since the spherical

gamma prime particles are coherent with the gamma matrix they may be easily dissolved.

The XEDS pattern fkom a spherical type y' particle is shown in Figure 105. The

particle composition corresponds to Ni3 (Cr, Co) AlTiW. This 0.3 pm diameter sphencal y'

particle has approximately the same chemistry as that of the cuboidal particles (see figure

101). Figure 1 O6 shows an XEDS pattern h m the y rnatrix. As expected, the ma&

comprised an alloy of Ni, Cr and Co. However, smaller amounts of Ti, Al and W were aiso

indicated. Also, the matrix contains higher amounts of Cr than the y' particles.

It can be concluded that the application of a higher fiction pressure (>325 MPa) is a

prime requirement during Ni-base cast ailoymi-base wrought alloy fiction joining and

produces crack-fiee welds.

5.3. Hardness Profiles

Figure 107 shows the micro-hardness results produced in a traverse h m the bondline

into the as-received Ni-base wrought alloy base metal. The indentation load was 500 g and

25 readings were taken on each test weld. When solution-treated Ni-base wrought alloy

base material was fiction welded, a softened zone was produced in the Ni-base wrought

alloy substrate and had a hardness about 25 -50 H v lower than the surromding base

Figure 104: TEM micrograph of the bondhe region in a Ni-base cast alloy/Ni-base wrought ailoy friction weld following RS (re-sdution) + PWHT (stabiiization plus precipitation) heat treatment

Clknt : S u g I S m t h l r n Job : Job numhr 367 spmctrurn tS CWS/99 1 1 :46>

Figure 105: XEDS spectra h m a spherical y' particle (fiom Ni-base cast alloy/Ni-base wrought alloy fiction welds)

Figure 106: XEDS pattern ftom the y matrix (Ni-base cast alloyMi-base wrought d o y weld)

Nilase Cast Ailoy & { Ni-base Wrought Alloy

A 1 1

: O t

: 0. O C 0

O 0

E x

1

4 -3 -2 -1 ,#' O 1 2 3 4 5

Distance ftom the bondline (mm)

Figure 107: Hardness profile in Ni-base cast alloy/Ni-base wrought alloy fiction welds

A- As- welded condition B. Mer Re-solution + PWHT

4

r

The measurements were carrieci out at the center (see arrow)

Welding Condition :

Friction Pressure : 350 MPa Friction Time : 10s Forging Pressure : 350 MPa Forging Time : 2s Bar Diameter : 19-

material. This soAened zone r d t e d hm the solution and coarsening (over-aging) of

intemetallic phases during the thermal cycle in fiction welding. These softened zone

regions were completely removed following re-solution + PWHT (involving a stabilization

plus precipitation treatment).

CHAPTER 6

CONCLUSIONS 6.1. Ni-base Wrought AIIoy Friction Wel&

The optimum weldhg parameters were: Friction Pressure: 350 MPa, Friction Time:

los, Forging Pressure: 350 MPa, Forgïng Time: 2s, Rotationai Speed: 1000 rpm.

Highest ultimate tende strengths (UTS values: 1475 MPa) with appreciable ductility

levels (>15%) were obtained when Ni-base wrought alloy base material was welded

using an intermediate friction pressures (325 to 375 MPa). In al1 the cases the tensile

test specimens failure occwed in base metal away fkom the weld zone. In welds

produced using at higher fiction pressures, the hardness profile across the weld zone

was essentially unifonn.

Hardened regions were produced in matenal adjacent to the bondline when the

welded samples are directly heat-treated using a PWHT (Stabilization + Precipitation) procedure. When Wction welds are heat-treated using ST or RS

treatments prior to PWHT, the hardness profile across the weld zone is uniform.

Application of a [RS+ PWHT] heat-treatment produced the optimum balance of

strength and ductility properties. In d l the cases tensile test failure occurred in base

metal away fiom the weld zone.

Solution - treated Ni-base wrought alloy base matenal produced the highest ductility

vales (>20%) wwhi the hardened Ni-base wrought alloy produced the highest

strengths (>Il00 MPa). Hardened Ni-base wrought alloy base material had lower

ductility in texms of total elongation values (approximately 8-10% lower) compared

to solution - treated Ni-base wrought allay base material. An ST+PWHT heat - treatment improved the ductility @y 20-27%) in cornparison to alternative PWHT

and RS+P WHT heat-treatments.

5. Hardened Ni-base wrought alloy base material exhibiteci extensive Cr carbide

precipitation dong the grain boundaries. Coarse and elongated carbides having

aspect ratios > 15 were highly detrimental vis-à-vis mechanical properties since they

can acted as a stress raisers and promoted secondary crack propagation.

6. The base materials tensile strengths were @? WHT: 1 137 MPa; RS+PWHT: 1 O87

MPa; ST+PWHT: 862MPaI and compared with fiction welded joints PWHT: 1302

MPa, RS+PWHT: 1380 MPa, ST+PWHT: 1323 MPa]. The Ni-base wrought

alloy/Ni-base wrought ailoy joint strengths were 20-25% higher than the as-received

base matenal. Ail sarnple failures occurred in the adjohhg base material

6.2. Ni-base Cast AlloyNi-buse Wrought AlZoy Friction Wei&

1. The application of higher fiction pressures (>325MPa) is the prime requirements

during Ni-base cast alloy/Ni-base wrought alloy dissimilar fiction joining and that

this welding parameter combination produces crack-fkee welds following post weld

heat-treatment.

2. The fiction welding operation altered the original particle shape and morphology at

the bondline. Close to the bondline region in the Ni-base cast d o y substrate side,

the cubical gamma prime was altered to spherical in shape.

FUTURE WORK

It is recommended that future work should emphasise modelling of the

microstructural changes produced in Ection welded joints, particularly with regard to

changes in the particle distribution when différent p s t weld heat treatment procedures are

applied. A general mode1 correlating and predicting microstructural variations and their

effects on tensile strength properties is required for different dissimilar superalloy material

combinations.

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[ASM Haudbook, 19721

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[Nathal, 19821

[Nathal, 19821

M.V. Nathal and LJ. Elbert, "The Influence of Cobalt on the Microstructure of the Ni-base Superalloy Mar M-247, Met. Transactions A, Vol. l3A (1982), pp. 1775-1 783.

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