THE MICROSTRUCTURE AND MECHANICAL

THE MICROSTRUCTURE AND MECHANICAL PROPERTIES OF A Ni-Cr WELD METAL

by

AMEER KHELIL IBRAHEEM, B.Sc. Eng.

A dissertation submitted to the Victoria University of Manchester in partial fulfilment of the requirements for the degree of MASTER OF SCIENCE

D epartment of Metallurgy and Materials Science University of Manchester/UMIST Grosvenor Street, Manchester, Ml 7HS

June 1990

i

DECLARATIO N

Unless stated otherwise the work is original and n o part of it has been submitted in support of an application for another degree or qualifi cation of this or any other university or other institute of learning.

iii

ACKNOWLED GEMENTS

I would like to express my thanks to: •

Dr. G.W. Lorimer and Dr. R. Pilkington for their invaluable help and guidance throughout the work and for their careful reviewing of the entire manuscript and making many helpful suggestions and criticisms.

•

The Ministry of Higher Education and Scientifi c Research in Iraq for their financial support .

•

Professor K.M. Entwistle for the provision of the research facilities.

•

E. l 4 research group for their continuous help and suppm:t; especially D. Tyrer, M. Ahmed, H. Arabi and Z. Larouk.

•

Members of the technical group of the Materials Science Centre for their invaluable help.

•

Ruth Greenbank for typing the manuscript .

iv

ABSTRACT

A st udy has been made of the effect of heat treatment on the microstructure and sho rt term mechanical properties of welded Inconel 182. Optical, X-ray and transmission electron microscopy were carried out to study the effects of stress relieving, solution treatment and ageing at temperatures between 400· and 7oo· C on the microstructure of the welded material. Solution treatment was observed to remove the inhomogeneities in the welded and stress rel!eved materials;

also it affected the grain size. The precipitation density ( observed optically ) was increased when the alloy was aged above 6oo· C. Neither long range ordering nor short. range ordering were observed.

Creep tests have been carried out under constant load conditions in air over a

stress range of 380-450 MP a at a temperature of soo· c: Minimum creep rates were calculated for each applied stress. Optical, scanning and transmission electron microscopy were used to examine the failed specimens. The heat treatment prior to creep, especially the solution treatment, was found to have an important effect on the mechanical properties of welded Inconel 182. The heat treatment affected the strain hardening. It was found that the applied stress determined the amount of deformation

prior to creep ( during loading ) and this had an effect on all the subsequent stages of

creep . Most of the creep failures exhibited ductile and shear fracture, although limi ted areas of intergranular fracture were also observed.

V

LIST OF ABBREVIATIONS AND SYMBOLS Unless otherwise stated, all the abbreviations and symbols used in this dissertation have the following meaning. HA Z HP IP CEGB AC WQ PFZ LRO Tc SRO ll'i

TEM T

(J

t k

tp

m A (J

n A' A"

c (J'

Heat Affected Zone High Pressure Intermediate Pressure Central Electricity Generating Board Air Cool Water Quench Precipitate Free Zone Long Range Ordering order-disorder transition temperature (critical temperature) Short Range Ordering short range ordering parameter Transmission Electron Microscopy absolute temperature absolute melting point strain instantaneous strain constant time constant transient strain time constant secondary creep rate constant stress creep index constant constant constant constant temperature dependent parameter activation energy for creep gas constant activation energy for self diffusion lattice diffusion strain rate

vi

H,H' V a

Dv

d K cgb Dgb w

b p N C' d' B p e

r h

(3*

V z

B'

nu merical constants which take the values "' 1 0 and"' 1 50/ 1r, respectively atomi c vo lume lat ti ce diffusion coefficient grain size B oltzman's constant grai n boundary diffusion strain rate grain boundary diffusion coefficient grain goundary width B urger's vector the total area swept out by the dislocation on its new and old glide plane number of dislocation sources climb rate of dislocations distance between dislocation pile-ups numerical constant "' 20 density of mobile dislocations distance between jogs rate of thermal recovery rate of strain hardening constant mobility of climbing dislocations impedance factor constant shear modulus friction stress cumulative stress reduction time to failure Monkman-Grant constant \-

vii

CONTENTS

PAGE

1

CHAP TER 1 : INTRODUCTION C HAPTER 2: AUSTENITIC-FERRITIC TRANSITION JOINTS 2.1 2.2 2.3 2. 4 2.5

3 4 5 5 6

Introduction Austenitic Steel Transition Joints Ferritic Steel Transition Joints Ni-Base Transition Joints Inconel Welding Electrode 182

CHAPTER 3: MICROSTRUCTURE 3. 1 3.2

3.3 3.4 3.5 3.6

8 9 9 10 10 11 13 16 18

Introduction Strengthening in Ni-Base Alloys 3.2.1 Solid Solution Hardening 3.2.2 Precipitation Hardening 3.2.3 Carbides Compositional Effects Ageing Characteristics in Ni-Cr-Fe Alloys The Order-Disorder Transformation Ordering in Ni-Base Alloys ,,

CHAPTER 4: CREEP 4.1 4.2 4.3 4.4

4.5 4.6 4.7 4.8

Introduction Stress Dependence of Creep Temperature Dependence of Creep Creep Mechanisms 4.4 . 1 Diffusion Creep Mechanisms 4.4.2 Dislocation Creep Mechanisms 4.4.2 . 1 Dislocation Climb Models 4.4.2.2 Dislocation Jog Models 4.4.2.3 Harper-Darn Creep Models 4.4.2.4 Recovery-Creep Theories Creep Fracture Creep of Particle-Hardened Alloys Creep in Ni-Cr-Fe Systems Effects of Ordering on Creep

20 21 22 22 23 23 24 24 24 25 26 27 29 32

CHAPTERS : EXPERIMENTAL PROCEDURE 5 .1 5. 2

Material and Heat Treatments Metallographic Techniques 5 .2.1 Optical Examination 5.2.2 Transmission Electron Microscopy 5.2.3 X-ray Diffraction 5.3 Creep Testing

35 35 35 36 36 37

viii

CHAPTER 6: RESULTS 6.1

6.2 6.3

Material Characterization 6. 1. 1 As-Welded Material 6 . 1 .2 Stress-Relieved Material 6 . 1 .3 Solution Treated Material 6.1.4 Solution Treated and Aged Material 6. 1 .4. 1 Hardness Measurements 6. 1.4.2 Lattice Parameter Measurements 6 . 1 .5 Metallography of Solution Treated and Aged Material 6 . 1 .5. 1 Optical Metallography 6 . 1. 5. 2 Transmission Electron Microscopy 6.1 .5.3 X-ray Diffraction Creep Results Metallography of Failed Creep Specimens 6.3 . 1 Optical Metallography 6.3.2 Transmission Electron Microscopy 6 .3.3 Scanning Electron Microscopy

39 39 39 39 40 40 40 40 40 41 42 42 44 44 45 45

CHAPTER 7: DISCUSSION 7.1

7.2

Material Characterization 7 . 1 . 1 As-Welded Material 7 . 1 .2 Stress-Relieved Material 7 . 1 .3 Solution Treated Material 7 . 1 .4 Solution Treated and Aged Material 7. 1 .5 Ordering in Inconel 182 Creep Behaviour 7.2. 1 Effect of the Heat Treatment Prior to Creep 7.2.2 Effect of the Applied Stress 7.2.2. 1 Loading Strain 7.2.2.2 Creep Rate 7.2.2.3 Creep Ductility 7.2.2.4 Creep Failure

47 47 47 48 48 49 51 51 52 52 53 54 55

CHAPTER 8: CONCLUSIONS

58

C HAPTER 9 : FUTURE \VORK

59

References

60

Tables

66-72

Figures

73-1 18

1

CHAPTER 1 : INTRODUCTION

Nickel--chromium alloys are used in nuclear power plant for welding dissimilar metal components which are designed to operate at temperatures up to

600"

C. One

type of such nuclear power plant is the Fast Breeder Reactor where t.he reactor core, contained within a primary loop, is cooled by unpressurized liquid-metal (a sodium-potassium alloy) . Some of the heat is removed from the primary loop to a secondary loop of liquid sodium through an intermediate heat exchanger; this heat is then used to produce superheated steam in a secondary heat exchanger, and hence, drive a turbogenerator. The piping in the heat exchanger is manufactured from creep resistant type 316 austenitic steel and the steam generator of low alloy 2t% Cr-1 % Mo steel. The coefficient of thermal expansion of the austenitic steel is greater than that of the low alloy steel. This will create potential cracking problems if the weld filler does not have an appropriate coefficient of thermal expansion. Inconel 1 82* is used for such dissimilar metal welding; it has the'nominal composition shown in Table 1 1. Inconel is resistant to a great variety of corrosive media. Its chromium content makes it superior to pure nickel under oxidizing conditions, while its high nickel content enables it to retain a considerable corrosion-resistance under reducing conditions.

The requirements of a structural material used in a nuclear reactor will vary with its specific purpose in the reactor. There are some general characteristics which may be outlined. Mechanical properties, such as tensile strength and creep rupture strength must be adequate for the operating conditions in order to help design the components for use within the reactor.

In the present work, the inter-relationship between the creep properties and micro structure have been studied for the Inconel weld metal. Microstructural

*

Inconel is a registered trade mark of the INCO group of companies

2

examination using X-rays, optical and electron microscopy have been performed to determine the effect of initial heat treatment and ageing treatment on the material.

,_

3

CHAPTER 2: AUSTENITIC-FERRITIC TRANSITION JOINTS

2.1 In trod uction

The greatest proportion of steels which are used in modern electrical generating

plant are low alloy steels. These are less expensive than the austenitic steels, have low coeffici ents of thermal expansion and can be heat treated to give a high proof stress.

At low temperatures ( up to 454 C ) 2 and high pressures it is possible to use carbon •

and carbon-manganese steels. However, as the temperatures are increased the use of low alloy creep resistant steels is more appropriate because of the increased resistance to oxidation due to their chromium content. At temperatures between 550· and 650" C, austenitic steels can be used because it is easier to obtai n the necessary creep resistance in an austenitic matrix than in a ferritic matrix2 . The high strength of the austenitic steels allows the use of thinner sections; this leads to a lower weight of piping, a lower cost of bending, easier welding and a reduction i n the total system '

weight. Pipe of low wall thickness has high heat tzansfer efficiency which may permit a reduction in overall length. Expansion loops may be elimin,ated and considerable cost saving may result from a closer spacing of the major components.

The use of austenitic steels in the higher temperature regions requires techniques of joining them to the ferritic steels used at the lower temperatures. Asutenitic/ferritic transition joints have been developed for this purpose.

In the 194 0's, transition joints for boiler service were made with austenitic stai nless steel filler metals. However after approximately ten years service, failures occurred in the austenitic filler metals and in the 1950 's steps were taken to develop superior transition joints. Lundin 3 has reviewed the literature on the use and failure of transition joints since 1 935. His findings can be summarized as follows:

4

i)

Fai lures invariably occur in the HAZ of the ferritic material adjacent to the w eld int erface, which is becoming weak and hard with time, due to carbon mi gratio n from the ferritic material into the weld metal. Increasing the Ni content of the filler metal alters carbon solubility, makes carbides less stable, changes diffusivity and in general retards carbon migration from the ferritic material.

ii)

Cracking in transition joints arises from: the difference in thermal expansion coefficients of the base and filler metals, internal pressures in transition joint piping, external bending stresses and vibrational loadings.

iii)

Failure mechanisms begin with the formation of a carbon-depleted soft ferritic zone whkh is constrained by the surrounding harder and stronger material; the strain accumulation is relieved by creep at elevated temperature. Creep damage in the form of cavitation, grain boundary sliding and tearing results in cracking in the carbon-depleted soft zone along and adjacent to the weld interface.

2.2 Austenitic Steel Transition Joints Austenitic filler metals were first used for dissimilar metal welding in Germany in 1935. It is only since 1949 that they have been given serious consideration in the welding of power plant piping, particularly for joining austenitic to ferritic super­ heater tubes and the joining of 2t% Cr-1 % Mo steam piping to austenitic turbine piping4. However, several major problems occurred due to 5: i)

the effect of differential expansion between the austenitic and ferritic materials and the possibility of fatiguing action at the austenitic-ferritic interface resulting from any possible cyclic temperature conditions,

ii)

the effect of carbon migration from the ferritic to the austenitic material resulting from either or both post-weld heat treatment and service temperature, and

iii)

the preferential stress-oxidation at the interface.

5

One type of this group is AISI 316 austenitic steel (Table 2 ) .

2 . 3 Ferritic Steel Transition Joints Ferritic steels containing chromium and molybdenum have been used in steam generating plant6. 2t% Cr-1 % Mo low carbon ferritic steel (Table 2) t1as been used in aggressive conditions such as liquid metal cooling circuits in nuclear reactors7, where its chromium content gives good oxidation resistance at high temperature and molybdenum improves creep strength. However, this steel fails to resist decarburiz­ ation by molten sodium. Additions of niobium and titanium were found to improve resistance to decarburization and stress rupture propertiess.

The ferritic materials can be categorized in two groups; non-heat treatable and heat treatable. Non-heat treatable ferritic steels behave similarly to the austenitic steels. One of this group is the completely ferritic type 446 stainless steel (Table 2) . Non-heat treatable steels are susceptible to embrittlement near 4 75° C 9. Also, some austenite may be formed during processing and then transfon11 to martensite upon cooling during welding. These alloys are given a post processing heat treatment to relieve stresses and temper any martensite that may exist.

The heat treatable ferritic steels have a wide range of structures and properties. They include the low alloy steels and the high alloy martensitic steels that posses unique high temperature and/or corrosion resistant properties9. One of this group is 2t% Cr-1 % Mo steel. The mechanical properties can be changed due to the phase transformations which occur during processing. A pre-heating and post-weld heat treatment are recommended to eliminate the under-bead cracking and to equalize the hardness values over the weldment9.

2.4 Ni-Base Transition Joints Nickel-chromium alloys containing little or no iron have been used in industrial appli cations where corrosion resistance and high temperature oxidation resistance are

6

required 10. Investigationsll' 12' 13 of ferritic-to-austenitic steel joints using iron and nickel-b ase filler metals have shown that nickel-base filler metal is superior to iron­ b ase filler metal . In addition, nickel-base alloys greatly reduce carbon migration from the ferritic material, have coefficients of thermal expansion close to those of the ferritic steel, and achieve relatively good success in cyclic high temperature applications. However, some failures have occurred14.

Ni-Cr-Fe welding products such as Inconel 82 and 92, Inconel welding

electrode 182, and Inco-weld 'A' (Table 2 ) are widely used for welding of low alloy steels to austenitic steelst5. These products are also used for many other combinations of materials. Table 3 shows a selection of Ni-based welding products for di ssimilar metal weldst6.

2.5 Inconel Welding Electrode 182 Inconel has been known and used in high tE;mperature applications for many years17. Recently, Inconel has been used for heat exchanger t.u bing .and tube-sheet overlays in the pressurized water reactor. Some of the physical properties at different temperatures, such as thermal expansion coefficients, stress rupture properties and elevated-temperature tensile properties for Inconel 182 are shown in Figures 1 ,2,3 respectivelyts. Table (4) shows the other specification of Inconel 182.

Inconel 182 is also used for shielded-metal arc welding of Inconel alloys 600,

6 0 1 , and Incoloy 800 ( Table 2 ) . The weld metal has excellent high temperature strength

and

oxidation

resistance

and

can

meet

stringent

radiographic

requirementst9.

Early investigations 20 concerned with Inconel 182 showed tl�at it had good resistance to stress-corrosion cracking. It was also selected for joining the ferritic-to­ austenitic steel in the primary coolant pipes of the Experimental Gas-Cooled Reactor at O ak Ridge2t .

7

The coefficient of thermal expansion of Inconel is similar to that of ferritic steels. The use of Inconel weld metal also reduces the carbon migration from the ferritic steel to the weld metal, while extensive carbon migration into the austenitic steel weld metal has been reported22.

Slaughter and Housley21 showed that Inconel 1 82 appears to be suitable for joining ASTM A-212, grade B carbon steel to type 304 stainless steel and ASTM A-387, grade C and D low alloy steels to type 304 stainless steel ( Table 2 ) .

Kent23 has reported that failures were initiated from defects, in the form of

sharp cracks , located at or near Inconel 182 interface with Rex 500 ( Table 2 ) in joints · of HP and IP turbine pipework in the 500 MW turbo-generators for the CEGB Ferry bridge Power Station. '

Nicholson24 tested cross-weld tensile specimens which were machined from welds made in 2t% Cr-1% Mo steel plate using either AISI � 16 austenitic steel or

Inconel 1 82 filler metals. Low ductility failures close to the ferritic steel /weld metal interface were produced by iso-stress temperature-acceleration tests at either 62 or 100 MN/m2 over the temperature range 575.-640. C. The endurance of joints using

Inconel 1 82 was greater than those made using AISI 316 austenitic steel by a. factor of "' 2. 5 and 3 for tests at 62 and 100 MN/m2, respectively.

8

CHAPTER 3 : MICROSTRUCTURE

3 . 1 Introduction Nickel-base alloys comprise a complex range of alloys. They are precipitation hardened alloys and t heir high temperature strength depends on t he precipitates w hich form during heat treatment. At high temperatures t he precipitates are taken into solid solution with a corresponding loss of strengt h. Attempts have been made to produce alloys containing particles which are stable and maintain strength to a high temperature. Most of these are dispersion-hardened alloys which are weaker than conventional alloys at low temperatures25. The dispersion strengthening is obtained by the inclusion of a small percentage of inert particles, usually refractory oxides, t hrough powder metallurgical processes. The usual dispersant is thoria and it is introduced by one of t he following routes :

i)

Direct addition of thoria t o t he loose metal powder and milling.

ii)

eo-precipitation of nickel and thorium oxides and selec, tive reduction of the former to metal.

iii )

Wetting nickel powder with a solution of a thorium salt, drying and thermally decomposing t he salt to oxide2 6 .

T he earliest alloy of t his type was TD Nickel containing 2% thoria; later TD Ni Cr was produced containing 20% chromium and 2% thoria 2 6 .

Dispersion hardening can be combined with precipitation hardening in the Mechanical Alloying process. This process produces a material with good cree p rupture properties and high hot--corrosion resistance2 7. It combines the high­ temperature strengt h of dispersion hardening with the low-temperature strength of precipitation hardening. Mixed elemental nickel and chromium powders, a master­ alloy powder containing the titanium and aluminium, and a dispersion of yttria, are milled together by a dry technique in a high-energy ball mill. Ytt ria is preferre d to

9

thoria because of t he radioactive nature of t he latter. T his milling process promotes particle welding and attrition; t he produce is t hen consolidated b y extrusion in evacuated, sealed containers 2 6 .

3.2 Strengthen ing in Ni-Base Alloys T he deformation of pol ycrystalline material occurs b y eit her t he deformation of t he individual grains or b y t he relative movement of t he grains. The deformation mechanism depends on t he temperature of deformation and t he strain rate. In nickel-base allo ys t he mechanisms which contribute to t he strengt h of the grains are solid solut ion hardening and precipitation hardening. The properties of t he grain boundaries are determined by t he carbides formed in t hese regions.

3.2. 1 Solid Solution Hardening T he t heories of solid solution hardening can be roug hl y categor ised into t wo '

groups28. Firstly, t hose w hich depend on various models of dislocation locking where t he dislocation is at rest, like Cot trell locking2 9, and t he chemical interaction 3o t heories. Secondly, t hose w hich are concerned wit h t he frictional resistance of solute atoms to t he movement of dislocations. Many experimental results indicate that addition of solute raises t he level of t he stress -strain curve as a whole31'32 and t his increases wit h t he increase of solute concentrat ion. T his s trengt hening comes from t he resistance of moving dislocations to various configurations of solute atoms. The magnitude of t his frictional force is sensitive bot h to t he atomic size differences and differences in elastic properties between solute and solvent atoms. Varency differences are also likely to be significant 2 8.

At high temperatures, diffusion and dislocation cross-slip determine the strength. Atoms which diffuse slowly such as mol ybdenum and tungsten are effective solid solution additions , as are elements which decrease t he stacking fault energy, such as cobalt 33.

10

3.2.2

t

Precipitation Hardening

Pu r e n ick el r etains a face-centred -cub ic crystal structure throughout the whol e emp erature range u p t o its melting p oint , and this is convent ionally t ermed /34. The

high t emp erature strength of most n ickel -bas e all oys r el ies on the 1' phase. Th is

phase is bas ed on the formula N i3( Ti,Al ) and it can hav e a rang e of comp os it ions d ep ending on the t itanium and alum inium cont ents of the all oy . Th e1:e is a fl exib il ity in the composit ion s ince the n ick el can b e r eplaced t o s om e ext ent by cobalt , molybd enum, chromium and iron and th e alumin ium and t itan ium can b e r eplaced by n iobium , tantalum and vanad ium 35.

7

'

has an ord ered face-centred-cubic crystal

stru ctur e which is coher ent w ith the face-centr ed -cub ic mat rix. Th er efor e, the precip itat e has a l ow surface energy w hich r esults in good l ong -t im e stab il ity at el evat ed t emperature, i. e. a slow coars en ing rat e. The r es istance t o d isl ocat ion movement is relat ed t o the s ize of the 7' part icl es , t o t he v olume fract ion pres ent and t o t h e energy ass ociat ed w ith t h e int erphas e b oundary. Th e s iz e of th e 1' particl e is d et ermin ed by the .cool ing rat e aft er the solution treatment and the ageing temperatur e. The volume fract ion of 1' is proport ional t o the amount of t otal hard ening addition , e.g. Al, T i and Nb 36 , and it forms at early stage of the ag eing treat ment , further ageing involves coarsening of the 7'

part icles 37.

A l ess frequ ently us ed hard ening phas e in n ickel -bas e mat erials is 1" which is bas ed on the formula Ni3Nb . It has ord ered body -centr ed-t et ragonal crystal stru cture and it is coh erent with the matrix35.

3.2. 3 Carb id es M etall ic carbid es can form in n ickel -bas e all oys at the grain b oundaries and w ith in the grains. Carb ides ar e generally harder and more b rittl e than the all oy mat rix. Their dist ribut ion al ong grain b oundaries is imp ortant in d et ermining th e mechan ical prop ert ies . Mon o carbides of th e gen eral form ula MC, where M is

11

t it anium, tantalum, niob ium or tungsten , form during melting. The y are stable and us uall y appear as isolated cube-s haped part icles 38. T he y are di fficult to dissolve in t he solid p hase and t hey restr ict grain growt h during solu t ion treatment. Complex carb ides w it h t he formula M2 3C6 , M 7 C3 or M 6 C can be d issolved by solut ion treatment in t he range 1050· to 1 200" C 35, T hey are preciP.itated from t he solid p hase 39, Precipitat ion in t his case is heterogeneous w it h preferred nucleation sites at t he grain boundaries and dislocations. At grain boundaries t he carbides form as block y precipitates, often retaining co herency w it h one of t he boundar y matrix grains 4o.

In M 2 3C6 , M is usuall y c hromium, iron, to a lesser extent nickel , cobalt or mol ybdenum depending on t he base composi t ion of t he allo y. M 7C3 , w here M is usuall y chromium, is stable at higher temperatures t han M 23C6 . If t he molybdenum \.

and tungsten contents of t he alloy are s ignifica11t , M6 C has t he tendency to form instead of M 2 3C6 . Dreshfield41 has studied t he format ion of t his, carbide in a range of nickel-base allo ys and has establ is hed a relat ions hip between chromium content , mol ybdenum and tungsten content and carbide i dentit y. T he results are summar ised in Figure 4 38. Generally, M 23C6 , M 7 C3 and M C are observed to be stable (in some Ni-based alloys) up to about 1050 " ' uso· and 1 200" c, respectivel y4 2 .

3 . 3 Compositional Effects Chromium is added to increase t he corrosion res istance of nickle-base alloys by t he formation of an ad herent oxide, Cr 2 03. Chromium also forms chromium carbides which give strengthen ing at high temperature 35.

Alum inium, t itanium and niobium are added to strengt hen t he n ickle-base

alloys b y t he format ion of t he 1' phase, Ni 3(Ti,Al) . N iob ium can be substituted for

titanium and alumin ium in 1', or if the n iobium content is hig h enough (above about

4%)35, it can form a separate p hase N i3Nb which is des ignated 1* (in USA 111). /11

12

forms with difficulty in the binary alloys under normal conditions of quenching and ageing43. Small additions of aluminium and iron43 or chromium35 promote the formation of 1" phase, but if the aluminium concentration is high enough, only 1' forms38. Both titanium and niobium also form carbides. Aluminum forms the stable oxide Al2 03 which has a useful effect on the o xidation resistance. Al 2 03 can compensate for the loss in Cr2 03 when the chromium content is reduced to improve the stress-rupture properties of the nickel-base alloys 35.

Iron is added to replace nickel in order to make the alloy less expensive. Increasing the iron content will decrease the oxidation resistance, since iron forms a non-adherent oxi de35.

Cobalt is added in place of nickel since it reduces the solubility of aluminium and titanium in t he nickel chromium matri x and therefore helps to maintain the \-

strength at high temperature. Cobalt also increases the workability of nickel-base alloys 35.

Molybdenum, tungsten and tantalum are used to provide solid solution strengthening at high temperatures. They dissolve to a certain e xtent in 1' and affect the lattice parameters of this phase which determine its strengthening characteristics. In addition they form complex carbides with each other and with other elements , such as chromium and iron35.

Boron and Zirconium are added to improve the creep strength and ductility. The mechanism by which this improvement occurs is not clear44.

Carbon can form mono carbides M C, such as Ti C, TaC, Nb C, ZrC and V C, also it can form either M 2 3C6 , where M is chromium or M 6 C if there is more than 6% molybdenum plus tungsten45.

13

Manganese i s added up to about 5% to improve the resistance to sulphur attack at elevated temperature, since i t forms MnS in preference to Ni3S 2 4 6, but it does not hav e significant changes on the general properties of the alloy. High concentrations of silicon (13.8%)35 in nickel-base alloys result in 1' phase (Ni3Si ) . In the early stages of ageing r' particles appear to be nearly spherical, then they tend to change to a cube morphology. The reason for this change is not clear; it has been suggested47 that it may be due to a change i n the degree of order i n the

' 1,

(Ni3Si ) .

3.4 Ageing Characteristics in Ni-Cr-Fe Alloys The ternary constitutional phase diagrams for Ni-Cr-Fe at 650. , 14oo·

noo·

and

C are shown in Figure 5 48, these illustrate that Inconel 182 has a face­

centred-cubic ( r) structure at temperatures of 650. and

noo·

c.

Few i nvestigations have been made of the ageing qharacteristics of the Ni-Cr-Fe systems. Nicholson49 studied the interfacial structures between 2t% Cr-1 % Mo steel and Inconel 1 82 weld metal in the as -welded condition and after ageing at temperatures between 630" and 700" C for times up to 6000 hr. Examination of the as-welded specimens revealed a band of martensite 0.5 to 50 pm wide between the steel and weld metal. During ageing between 630• and

7oo·

C precipitates developed

in this band and led to arrays of interfacial carbides , extended ageing at temperatures greater than 675. C led to the dissolution of the precipitates. The p i·ecipitates were predominantly M23 C6 and M 6C and lay in the ferritic steel, about 1 f.J.m from the interface. Dix and Savage so studied the structural changes observed in Inconel x-750 weld metal (Table 2) as a result of isothermal ageing for (15 sec - 30 min) in the temperature range

7oo·

to

goo·

C. This alloy has higher Ni , Ti and Al contents an d

lo wer Mn and Nb contents than Inconel 182. The hardening p hase was coherent r',

r 14

Ni3(Al,Ti) 51. Carbides were formed at the grain boundaries and their morphology was a function of the ageing temperature. Cellular carbides were observed after ageing at 700" and 760" C while globular carbides were observed after ageing at 815 " c.

Raymond52 s tudied the effect of heat treatmen t on rup ture proper ties and corrosion of Inconel x-750. The alloy s tructure was dependent o � the treatment applied. The typical heat treatment given (for this alloy , which i s used as a turbine blade material) was 1 150" C/4h/ A C, 850" C/24h/ AC fol lowed by 7 oo· C/20h/ A C. It was mentioned that the hardening phase was

' 1,

Ni3(Al ,Ti) and the carbides were

Ti( CN) wi th Cr23 C6 in the grain boundaries. Susukida et al5 3 s tudied the changes in properties of Inconel x ( the o ld name of Inconel x-750) which occurred during long term heating. The same heat treatment '

speci fied by Raymond52 was given to the alloy. Sus ukida et al5 3 found that i t was di fficul t to observe changes in the micros tructure by long t�rm heating. Electron microscopy revealed that the

' 1

particles , Ni3(A l,Ti) , had a diameter of 0.05 to 0.1

J.Lm and that there were carbides in the grain boundaries , which precipitated finely and intermi ttently before heating, became larger and coarser with an increase in heating time. It was observed after heating at 7 oo· C for 1 0 ,000 h that

' 1

par ticles

grew in diameter to between 0.1 and 0.2 J.Lm and carbides in grain boundaries formed a continuous distribu tion.

Dix and Savage54 found that 1' , Ni3(Al ,Ti ) , was formed in Inconel x-750 in t he temperature range 704 -871 " C as a coherent precipi tate. The shape of •

1

'

depended

on the previous thermal his tory; i t was spherical if formed at low ageing temperatures and cubic when formed by long exposures to hi gh ageing temperatures . The matri x contained Cr2 3C6 , ( Cb ,Ta) C and Ti( CN) carbides.

15

Venkiteswaran et al55 stud ied t he effect of quenc hing rate on t he micro­ s tru cture of Inconel x-750 at t he different stages of t he heat treatment sequence. The res ults are s hown in Table 5. T he material had a 1 1 3 I-'m grain s ize. It was observed t hat water quenching (WQ ) tended to result in greater cont inuit y of precipitat ion along grain boundaries. Furthermore, WQ after intermediate ageing tended to reduce t he s ize and spacing of

' 1

part icles w it hin t he grains and reduce t he s ize of the

precip itate free zone (PFZ) adjacent to grain boundaries. In situat ions

1-4

(Table 5)

where WQ was used following solut ion treatment , air cool ing (AC) after the intermediate stage produced changes in t he distribut ion and t he s ize of

' 1,

also

increased t he s ize of t he PFZ. In samples No. l and No.5 (Table 5 ) , examined after intermediate ageing, it was found t hat t he intragranular precipitates were cr ystallographically orientated. In bot h cases the PFZ's at t his stage were narrower t han in full y heat treated samples and it was not poss ible to resolve clearl y t he precip itates along t he grain boundaries. W hen air cooling was used follow ing t he intermediate heat treatment, t here was a d iscontinpous dis tribut ion of precip itates at t he grain boundar y. In N imonic SOA 2 7'5 6 (Table 2 ) , most of t he carbides at t he grain bo undaries were completel y d issolved at temperatures above 1150 " C and were precipitated in a way depending on t he cooling rate and the temperature of ageing. Only titanium carbonitr ide was present in t he matrix at temperatures above 1 150" C, but on slow cooling to around 1000" C, a Cr7C3 carbide was precipitated . Cr7C3 can be transformed to Cr2 3C6 by ageing at 950· C for 24 h, while 14,000 h was needed to complete t his transformat ion at 700" C.

1

'

phase ( N i3(Al,Ti) ) was completel y dissolved at temperatures above sso · C.

Water quenching from above t his temperature retained t he phase in solut ion. Coarse precipitat ion of

1

'

can be achieved b y slow cool ing ( in a furnace) from above t he

solu t ion temperature to t he ageing temperature, and maintained at t he latter temperature for a s uitable period.

16

The temperature of the solution treatment determined the softening of the mat rix and controlled the resultant grain size. Double ageing treatments were developed either to obtain duplex 1' structures, or to precipitate car �i des retained in s olution.

3.5 The Order-Disorder Transformation If dissimi lar atoms in an alloy attract each other in preference to like atoms and if the alloy exists at or near a composition that can be expressed by a s imple formula such as A B, A 2 B, A3B, etc. an ordered solid solution or super lattice may be formed in which speci fic atom types occupy specific lattice sites. At high temperatures, above the ordering temperature, thermal energy will overcome the tendency to order and a random solid solution wi ll result.

Most of the order-disorder transformations in Ni-base alloys occur by a nucleation and growth type process5 7 . Nucleation, and growth refers to a model in which small highly ordered regions nucleate and then grow UI).til they contact each other within a disordered matrix. Long Range Ordering (LRO ) is established when the ordered arrangement extends over many adjacent unit cells5 8. Just above a critical temperature (Tc) regular arrays of atomic species of LRO do not exist, but the statistical probabi lity of finding ordered arrangements as near neighbours may be higher than that in a random solution. This is known as Short Range Ordering (SRO ) . There are four principal models of SR0:59

i)

The statistical model in which the SRO parameter

ll'i

represents the probabi lity

of finding B-type atoms in the ith coordination sphere around an A-type atom. ii)

The microdomain mo del in which small, long-range ordered domains , with dimensions which can reach some tens of nanometers , are separated by antiphase domain boundaries or twin boundaries or other discontinuities.

17

iii)

iv )

The lattice defec t model 6 0 in which dis locations and lattice defects are responsib le for the formation and stabil ization of inhomogeneities of concen tration and degree of order. The dispersed order model6 1 -6 3 in which a dispersion of small highly ordered particles are presen t in a d isordered matrix, wh ich are resis tan t to coarsen ing.

The cho ice between these models is sometimes very d ifficult, because the models d iffer in homogeneity or heterogeneity and in the SRO s tab ility. The ex is tence of the d ispersed ordered s tructure has been experimentally confirmed in the b inary Ni-Cr alloys5 8' 62'6 3. The existence of LRO phase can be detected by electron d iffraction techn iques, i.e. cer tain reflections in the reciprocal lattice occur when the ordered phase is presen tso. It should be no ted that failure to detect such superlattice reflections does ,_

no t necessarily rule out the presence of the ordered,phase. I t might be present but not identifiab le if either the amoun t of the ordered phase is b �low the threshold of detection, or if the s tructure factor causes the diffraction from cer ta in planes to be weak5 0. Superlattices can also be detected by X-ray or neutron d iffraction . Neutrons are scattered from the nucleus rather than the electronic shells as in X-ray diffracti ons.

Some physical changes

in

the alloy accompany

the order-disorder

transformation through the critical temperature. These are changes in specifi c heat, elec tr ical res istivi ty, magne tic proper ties, lattice cons tants, diffusion and hardness. These proper ties change very rap idly near the critical temperature. Also the order-d isorder transformation influences the mechan ical proper ties .

The resis tiv ity behav iour of some commercial Ni-Cr alloys has been s tudied in order to ob tain a better unders tand ing of the order ing reaction, through its dependence on compos ition and its effects on the physical proper ties of the alloys 64.

18

g_.6 Orde ring in N i-Base A lloys

Studies 65 '66 on the formation of SRO and LRO based on the Ni 2 Cr super ­

s tructure have been carried out on the binary alloys Ni 2 Cr and Ni 3Cr. Electron diffraction patterns showed diffuse scattering and streaking of fundamental reflections in directions. Also, TEM images exhibited mottled cont rast. These showed that a finite degree of SRO was formed in both alloys. They also exhibited lattice contraction and changes in electrical resistivity, as shown in Figures 6 and 7 58. In Ni 3Cr alloys the structure remained in the SRO state even after 25,000h when exposed to temperatures below 525" C. No evidence of ordering had been observed after 10,000h ageing, but small coherent precipitates were imaged in dark field micrographs by matrix strain contrast. They were dispersed through the sample. While i n N hCr alloys SRO began to transform to LRO structure after a few hundred hours of ageing58. The lattice contractions of the Ni 3Cr and Ni 2 Cr alloys were"' 0.05% and"' 0 .25%, respectively after ageing at 475" C fo � 25,0 00h 58. Klein et al 67 , did not report superlattice reflections in N hCr until after 730h at 4 79" 9· Goman'kov et al 6 8, studied the formation of N hCr superlattice in polycrystalline of Ni-(20 to 35) % Cr and single crystal of Ni-32.9% Cr (approximately stoichiometric) specimens. The heat treatment was quenching from 1200" C in water and annealing at 500" C for

1036h. No superlattice lines were found by neutron diffraction from poly- crystalline specimens , but superlattice lines were detected from single crystals. Neutron diffraction studied65 indicated that T was about 590" C for Ni 2 Cr alloys. c

In Ni-Cr-Fe systems, Dix and Savage5 o reported that the hardness of Inconel x-750 was increased after 1 minute ageing in the temperature range from 704 to •

887" C due to the formation of the ordered phase I' · Dix and Savage50 mentioned that failure to detect superlattice reflections in electron diffraction patterns does not necessarily rule out the presence of I'·

Marucco and Nath69 studied the order-disorder transformation in Ni-Cr alloys containing Fe. The Ni 2 Cr phase field in the system extended to 5 % Fe. It was found that within the limits of the Ni 2 Cr phase field, the magnitude of contraction ass ociated with ordering depended on the chromium content of the alloy. The kinetics of LRO also depended upon the Ni: Cr ratio, the higher ratio resulted in higher kinetics of LRO . A reduction in the critical temperature resulted from increasing the Fe content.

Metcalfe and Nath62 studied the ordering reaction of the alloy Nimonic 80A (Table 2), this alloy has higher Ni, Cr, Ti and Al contents and lower Fe content than Inconel 1 82. Electron diffraction patterns along the < 100> matrix direction showed diffuse scattering and the spots appeared streaked in the < 1 1 0 > direction . After 12,300h ageing at 450 " C, discrete superlattice diffraction spots appeared which were consistent with NhCr. After 30,000h the lattice parameter contraction was 0. 1 1 %, ' while it was 0.03% after 100h at 450 " C 62 . Measurements of the (31 1) .diffraction peak position for Nimonic 80A with temperature by X-ray �iffraction showed

a.

c

discontinuity at about 570" C which was an indication of T for an ordering reaction.
0.4 T The creep process starts with an instantaneous strain on the first loading of a specimen, followed by three sections: oprimary creep, secondary creep and tertiary creep (Figure 8).

Primary creep shows a transient period of rapid strain flow eventually giving way to a linear steady state which is a secondary creepo This is then followed by tertiary creep and fracture. Dislocation movement can occur by processes such as cross-slip and climb. At elevated temperatures, dislocation motion is no longer necessarily bound to the slip planes. The increased freedom of dislocation movement permits recovery to occur at dislocation networks and tangles. A dynamic equilibrium of strain hardening and recovery by dislocation motion can enable a steady state to be reached with a constant creep rate.

The transient stage obeys a time dependent relationship of the form

( 4.1)

where secondary creep is subsequently observed, a linear term is added to equation (4 . 1 )

o

f = to

+ j]t1/3 + kt

( 4.2)

This is known a s the Andra.de (time) 1/3 law.7 0 This relationship has two limitations; first, an infinitely high initial creep rate is required , and secondly, a true steady state is not produced. A creep relationship that avoids these limitations has been found71 to

f = fo + fp(l - exp(-mt) ) + Est

(4 .3)

4 .2 Stress Dependence of Creep

At elevated temperatures, the stress dependence of the secondary creep rate is

generally described using the power law, often called Norton' s law7 2

(4.4)

At low stresses and very high temperatures (T

"'

0 . 9 Tm) , the creep rate varies

linearly with stress ( n= 1 ) and creep in this range is controlled by stress directed vacancy diffusion without dislocation movement.

At intermediate stresses, equation (4.4) is not linear and n is greater than unity.

At high stresses the power law will break down and the s�ress dependence of the creep rate is better expressed by the expression7 3

Es = A ' exp(f]CJ)

( 4.5)

Alternatively, a more unified description which covers a very wide stress range is produced. It reduces to a power law at low stresses and an exponential at high stress levels with respect to creep rate74

E s = A" (sinhf]CJ) n

(4 .6)

The above equation can be modified by including a temperature dependent parameter,

a,

in order to overcome a discrepancy which exists between the activation

energy for creep and that for self diffusion75

22

fs = C [sinh(,B' ( O" - a))] n

(4.7)

The accelerating creep rate in the tertiary region, results from two mechanisms ; namely structural instability leading t o neck formation and grain boundary damage leading to intergranular fracture at low strains.

4.3 Temperature Dependence of Creep The temperature dependence of the secondary creep rate may be described by an Arrhenius expression of the form: fs = A exp(-Qc/RT)

(4.8)

where Q c is the activation energy for the rate controlling mechanism. Generally, the activation energy for creep is equal to that for self diffusion Q576 and is usually found to be independent of stress and strain77. At low temperature Qc is a function of temperature78 ' 79, while above 0.5 T

m,

Qc is independent of temp erature7 3 .

The activation energy for creep may be determined experiment.ally, by making temperature changes during a creep test and recording the change in secondary creep rate and the activation energy will be

( 4 .9) where fst and fs2 are the secondary creep rate corresponding to the temperatures T1 and T2 respectively. 4.4 Creep Mechanisms There are two dominant creep mechanisms; dislocation ereep and diffusion creep. The operating mechanism will depend upon the temperature and stress . Deformation mechanism mapsso have been developed from a knowledge of the stress

23

and temperature dependence of the creep rate for the various mechanisms , an example is given in Figure 9 for Inconel x-750 81. Generally, at low stresses and at m,

temperatures below 0.3 T the creep curve exhibits only a transient flow, which can m

be described by a logarithmic law82. In the temperature range 0 . � T to 0.9 T

m,

thermally activated cross-slip and climb become increasingly i mportant . However, at low stresses, creep results from diffusional transport processes. A brief descri ption for some diffusion and dislocation creep models will be outlined.

4.4. 1 Diffusion Creep Mechanisms Diffusion creep is a high temperature, low stress phenomenon. The mechanisms of diffusion creep can be described in the way that , vacancy concentration develops due to the applied stress at the boundaries under tension and a depletion of the boundaries under compression. As a result, a flow of vacancies will take place, either through the lattice giving rise to Nabarro-Herring creep83 ' 84, or along the grain

boundaries giving coble creep85 ; and the creep ra� e is controlled by the diffusion of vacancies. The lattice and grain boundary diffusion creep mechani sms can be expressed respectively by:

tL tgb

=

=

H a Va Dv/d2 KT

H' a Va D g b W j d 3 KT

( 4. 1 0 ) ( 4. 1 1)

4.4.2 Dislocation Creep Mechanisms Power law creep is dislocation controlled. Recovery and thermally activated glide of dislocations are the rate controlling processes. A balance between increased strain hardening cause by glide processes and a reduction in strain hardening due to recovery processes is thought to produce a steady-state creep rate. Several theories have been developed to account for this phenomenon and some are briefly outlined below.

1

24

4.4 .2 . 1 Dislocation Climb Models Weertman's86 ' 87 model is the first physical model of creep and nearly all the other models can be derived from it88. In this model, dislocations are produced by sources in parallel planes and their edge portions are mutually trapped, forming relaxed multipoles; when the head dislocations of parallel planes annihilate by cli mbing toward each other, the blocked loops can then glide over the average distance L/2 (L radius of the head loop) and another loop is emitted by the source and creep continues. The emission and expansion of dislocation loops are the strain­ hardening phenomenon, the climb and annihilation of dislocations are the recovery process. The overall creep rate is

fs = b P N C'/d'

( 4. 1 2)

4.4.2.2 Dislocation Jog Models Jogs on a screw dislocation are always edge �ogs; and if the scr ew dislocation is

forced to move, then such jogs can only maintain their positim� on the dislocation by climb. These jogs may be classified as ei ther vacancy-emitting jogs or vacancy­ absorbing jogs.

The thermally activated motion of jogged screw dislocations as a possible mechanism controlling high temperature steady state creep has been considered by many workers89 ' 9 0 ' 91. When a screw dislocation, containing vacancy-emitting or vacancy-absorbing jogs, moves in response to an applied stress, the jogs can accompany the dislocation only by emission and absorption of vacancies for the dislocation motion to be maintained. An expression for the creep rate91 is:

fs

=

B Dv p sinh( CJ b2f/KT)

( 4 . 13)

4.4.2.3 Harper-Darn Creep Models Harper and Dorn9 2 suggested that a polycrystal material which was deformed in

25

tensile creep at high temperature exhibited a Newtonian viscous behaviour at very low stresses. Examination93 by etchpits and TEM, of an Al-5% Mg alloy, revealed that the dislocation density was quite low (mostly edge dislocations) and was independent of stress, also no subgrains were seen. These results led Langdon and Yavari94 to propose that Harper-Darn creep arises from the climb of edge dislocations under saturated conditions, i .e. when the concentration of jogs is high enough for climb to be controlled by vacancy diffusion. It seems that Langdon and Yavari94 considered that strain was produced by climb of the edge dislocations, i .e. creep by exchange of vacancies between dislocations rather than grain boundaries. The Harper-Darn model may be applied for creep at high stresses by assuming that the dislocation density is stress-dependent88.

4.4.2.4 Recovery-Creep Theories A recovery controlled steady state creep model has been suggested by Bailey95 and Orowan96. They produced the following expression for the steady state creep .

r

f s = li

( 4. 14)

In this model, the strength is provided by the attractive and repulsive junctions of the dislocation networks. Some of these dislocations break due to thermal fluctuations and when the released dislocation moves some way through the network, it will produce slip and then it will increase the dislocation density. This process produces the work hardening stage of the creep . Recovery of the dislocation networks takes place by the gradual (climb controlled) growth of larger meshes and the contraction of smaller meshes. This tends to increase the average mesh size and to decrease the dislocation density97. Recovery rate can be calculated by a Friedel98 analysis and the hardening rate is inversely proportional to the local stress99. Hence the creep rate, after substituting in equation (4. 14), is:

ts

=

j]* V 0"4

( 4.15)

26

Dislocation mechanisms in creep have been reviewed by LagneborglOO, by considering the deformation mechanisms of the primary and secondary stages of creep through recovery-dependent models.

4.5 Creep Fracture Creep fracture in many commercial alloys are i ntergranular or ductile in nature. There are two types of intergranular separation, depending on the applied load and temperature tO l:

(1)

At high loads and low temperatures in the creep range, the fracture tend to originate at grain boundary junctions (triple points) , rather than on the boundaries.

(2)

At lower loads and higher temperature, fracture results more from the formation of cavities along grain boundaries , especially those boundaries perpendicular to the loading direction.

Creep fracture is a complex process involving the interaction between matrix creep deformation and the accumulation of creep damage at grain boundaries. Discontinuities may exist on grain boundary ledges due to the damage which can act as sites for cavity nucleation. This nucleation occurs when two grains slide past each other as a shear process to produce initial fracture t o2 . Another possible nucleation site is at the intersection of a sub-boundary with a grain boundary t03.

Cavity growth can occur by either the diffusion of vacancies to cavities along the grain boundary and the plating of atoms from cavities on to grain boundaries10 4, or by deformation. Cavity growth may also occur by grain boundary slidingt 05 , either in the plane of the grain boundary or in the direction of the principle applied stress1 06 .

When cavities nucleate at triple points, they propagate as wedge-shaped cr acks along the grain boundary. This is distinguished from cavitation, and it is known

a.s

27

wedge cracking. Crack growth may be due to the increased wedging action of the asso ciated grain boundaries lying at more i nclined angles to the tensile axis . Wedge cr acks are different from coalesced cavities which they nucleate at the triple points1 07. These cavities inter link to form grain boundary facet sized cracks, and the number of such grains increases with time and results in a crack of sufficient length for final fr act ure108.

The ductile fracture starts when voids nucleate on precipitates. A hard precipitate disturbs both the elastic and plastic displacement field in a deforming body. The disturbance concentrates stress at the precipitate, this stress building up as the plastic strain increases until the local stress reaches a critical value, where it either fractures the precipitate or tears open the interface with the matrix, thereby nucleating a void. Further plasti city makes voids growth and when they are large enough or when the specimen itself becomes mechanically unstable, they coalesce and fracture the material81.

4.6 Creep of Particle-Hardened Alloys The creep behaviour for many alloys can be expressed over a wide range of stress and temperature by:

Es = A

O'n

exp (-Q c /RT )

(4. 16)

where A and n are constants for a given material and set of creep conditions. For pure metals the stress exponent n, usually takes a value between 3 and 5 I09, larger indices have been observed in a second phase-hardened materials. For example values between 5 and 13 have been measured for 1'-hardened Ni-based alloysH0-11 2 .

In an attempt to rationalize the high stress dependence of the creep rate of second phase-hardened materials, a number of theories have been developed.

28

Lagneborg113 has proposed a model which is based on a model for recovery of creep in pure metals (the network growth model) proposed by McLean and co-workers114-l16. It is well established117 that precipitates and inclusions retard grain growth by exerting a back stress on the grain boundary. Then the driving force (for growth) is decreased by the amount of the back stress. Applying this treatment to the growth of a dislocation network which is impeded by second phase particles, gives an expression for the stress dependence of the steady state creep113 : ( 4.17) where Z is an impedence factor which is dependent on the number and size of impeding particles and on the nature of the particle-dislocation interaction.

A similar approach has been proposed by Wilshire and co-workersl l 0-112 ' 118 ' 119 . It suggests that the rate controlling process is the _ growth of the dislocation network to generate links in the network, long enough to act as dislocation sources. Once a link of sufficient length is formed, a burst of slip takes place, increasing the local dislocation density so that the next slip event will take place elsewhere. The particle distribution will impede the generation of sources by network growth. At high stresses, once a source is formed, the dislocations generated may cut through or bow between the particles, depending on their size and spacing. At low stresses, deformation may be confined to grain boundary regions with creep taking place by grain boundary sliding. However, Parker and Wilshire119 indicated that the change in stress exponent from low to high stresses is a consequence of a change in the stress dependence of and < 1 1 1 > matrix direction with a

slightly convergent and parallel electron beams are shown in Figs. 46 ( a-f) for

specimens aged at 400° , 450 " and 500° C: Figs. ( a) and ( b) show diffraction patterns

along the < 100> with a slightly convergent beam for specimens aged at 400° and

42

450" C respectively. The diffraction patter n along the - matrix direction with a slightly convergent beam for a sp ecimmen aged at 50 0· c is sh own in Fig. (c) . Diffraction patterns along the < 100> matrix direction with parallel elect ron beam for

specimens aged at 400" , 450" and 500 " C are sh own in Figs. (d) , (e) and (f) respectively.

The diffraction patterns taken, with the electron beam alon g the < 100> matrix direction, (with parallel electron beam) do not reveal either superl attice spots characteristic of LRO or streaked reflections characteristic of SRO . A slightly convergent beam was used in an attempt to increase any diffuse scattering due to SRO; no significant diffuse scattering is obtained. Diffraction patterns taken, when the electron beam along the < 1 1 1 > matrix direction, yielded similar results.

6 . 1 .5.3 X-ray Diffractions

X-ray diffraction patterns using either a pow�er camera or diffractometer do

not reveal direct evidence of ordering. However the lattice parameter measurements which are carried out in the diffractometer may show a lattice contraction, section 6 . 1 .4.2.

6.2 Creep Results The results of creep tests carried out at 500" C in the stress range of 380-450 MPa are given in Table 8 and the creep curves are shown in Figs. 47-5 1 . Except for the specimen tested at 420 MP a, the creep curves are characterized by a high initial creep rate which is followed by a rapid transient during which the creep rate decreases to the secondary stage. During the secondary stage, the creep rate is almost constant . There is then a rapid transient to tertiary creep is followed, which leads to failure. The tertiary creep strain is high for most of the tests. The time spent in tertiary creep is very short compared to the total creep life. The primary stage for the specimen tested at 420 MP a i s not clearly distinguishable from the secondary stage, Fig. 49.

43

The variation of minimum creep rate with the applied stress is shown logarithmically in Fig. 52. A possible straight line relationship may exist between the secondary creep rate and the applied stress. Assuming a power law of the form: fs = A

(J

n

N

then the value of the creep index, n, is found to be 34.

The variation between the applied stress and rupture life is shown logarithmically in Fig. 53. The points distribution shows the relati onship which is represented by the solid curve, which is shown in Fig. 53.

The variation between the percentage loading strain and the applied stress is shown i n Fig. 54. Except those of the 430 and 450 MP a creep tests, the loading strain \ -

seems t o vary linearly with the applied stress. The loading strains o f the 430 and 450 MPa tests appear to be lower and higher (respectively) than their values which are predicted by the extended linearity which is shown i n Fig. 54. Except that of the 430 MPa test, the loading strain represents 84-93% of the total strai n to failure; for the 430 MP a test, the loading strain is 65% of the total strain to failure. The variation between the percentage loading strain and the applied stress for the 400 MPa creep tests is sho\vn logarithmically in Fig. 55. In Fig. 55, the percentag� loading strain shows a linear relationship ·for applied stresses below

"'

150 MP a with loading strains

less than 1 %. Above 150 MP a the linear relationship breaks down and non-linear relationship is started with a high rate (of increase in loading strain) .

The variation between the percentage creep strain and the percentage loading strain is shown logarithmically in Fig. 56. No systemati c correlation is apparent. However, when the percentage strain to failure is plotted against the applied stress, Fig. 57, the strain to failure increases systematically with the increase in applied stress. The values of the percentage strain to failure may represent the ductility of the

44

material which is quite high. The ductility values ranging from

rv

33 to

rv

55% in the

stress range between 380--450 MP a.

The reduction in the uniform cross-sectional area of the creep specimens increases from 25 to N 36% when the applied stress increases from 380 to 450 MP a. The Monkman-Grant constant ( C M-G ) which can be expressed as

has been calculated for the creep tests. The values of C M--G are found to vary from 2.33 x I0-3 to 13.2 x lQ-3. The relationship between the Monkman-Grant constant and the applied stress is shown logarithmically in Fig. 58.

6.3 Metallography of Failed Creep Specimens 6.3.1 Optical Metallography Optical micrographs which illustrate the general features of dal)'lage within the specimens are shown in Figs. 59-65. In general, the fracture path is transgranular, and some regions fractured by shear, Fig. 59-6 1 . Wedge like cracks have occurred as a result of creep damage in a specimen tested at 400 MP a, Fig. 60. Cracks are initiated along the gauge length of the specimens and some have grown to a large size

( 0.4 mm long ) , Fig. 6 1 . Also, large cracks located far from the fracture surface are rv

observed in the specimens, like the one ( 0.8 mm long ) which is observed in a rv

specimen tested at 420 MP a at a distance of

rv

2.8 mm from the fracture surface,

Fig. 62. Secondary cracks from the fracture surfaces of the specimens are observed, an example is given in Fig. 63 for a specimen tested at 420 MP a.

Transverse sections of creep specimens showing damage in the form of wedge cracking, Figs. 64 and 65. Subgrain decoration and a high precipitate density a,re observed, Fig. 65. A large amount of grain growth seems to have occurred in the failed

45

creep specimens,Figs. 59 and 6 1 . The grain size measurement is difficult , since only the widths of the grains are shown and their lengths are longer than what are shown by the mi crograph of the whole specimen, Fig. 59. However, the few complete grains show that they are elongated in the direction of the stress axis, Figs. 59 and 60.

6.3.2 Transmission Electron Microscopy Transmission electron micrographs of typical areas of failed creep specimens are shown in Figs. 66--{)8. The microstructures show a marked difference from those of the heat treated material before creep testing. The dislocation density has increased significantly. Parallel traces of slip planes can be observed, Fig. 66, and high dislocation density adjacent to a grain boundary are shown in Fig. 67. An example of the dislocation pile-ups are shown in Fig. 68. Ti-rich precipitates are observed in the microstructure of the failed creep specimens.

6.3.3 Scanning Electron Microscopy Scanning electron micrographs show different features ,of fracture, most are ductile in character with some regions of shearing, an example of a fr?-Cture surface is shown in Fig. 69 for a specimen tested at 420 MP a. However, limited areas of inter­ granular fracture are observed in a specimen tested at 400 MP a, Fig. 70. The fracture planes across the specimens are oblique. The angle between the normal to the fracture plane and the stress axis is about 45" in all specimens. An example of sheared region is shown in Fig. 7 1 . Dimples can be seen in most regions, an example is shown in Fig. 72. Both conical and shallow dimples with an open end are observed. The shallow dimples resulted from the joining of voids by shear.

Qualitative information showing that the amount of sheared regions present on the fracture surfaces of failed specimens increase with the increase of applied stress, on the other hand the amount of regions which have dimples decrease with the increase of applied stress.

46

Precipitates are observed in the roots of some dimples, an example i s shown in Fig. 73. Polished and un-etched specimen of failed creep specimen tested at 400 MP a show intragranular voids as shown in Fig. 74.

47

CHAPTER 7: DISCUSSION

7.1 Material Characterization 7 . 1 . 1 As-Welded Material Inconel 1 82 is used for the shielded-metal-arc welding of dissimilar alloys when the temperature of the application does not exceed 540" C16. One welding method used is the multiple-pass process which results in the appearance shown in Fig. 1 0. The microstructure of the welded material reveals a dendritic morphology, which is a typical microstructure of welded material.

Ti-rich precipitates are found in the as-welded material; these are believed to

be either titanium carbonitrides , Ti ( CN ) , or titanium carbides, Ti C, which form 4

during melting2 7 38. Titanium carbides are stable at high temperatures and have been observed in the as--cast and solution treated condition of many of Ni-based alloys13 o.

The hardness of the as-welded material is higher than any other value which is obtained after solution treatment and ageing. This reflects the complex annealing treatments and thermally induced stresses which are produced during the multiple­ pass welding process.

7 . 1 . 2 Stress-Relieved Material Stress relieving at 700 " C for 3h results in little, if any, softening. There may be some slight polygonisation of some of the dislocations, produced as a result of the stresses produced during solidification, and some limited removal of solute from solid solution. There is some evidence of precipitation at and adjacent to the grain

boundaries, Fig. 24. These precipitates may be complex carbides ( M5C and M 23 C5 ) which are stable at low temperature and are precipitated on heterogeneous sites4o.

Carbides of the form M 23 C 6 may be precipitated more than M5C, Fig. 4 41, since Inconel 1 82 contains approximately 19.5 atomic % Cr and less than 0.02 atomic %

48

( Mo

+

0 . 4 W ) . Further TEM i s needed t o confirm this proposal. M 23 C 6 has been

observed at the grain boundaries in Inconel x-750 and Nimonic 80A ( which have compositions close to that of Inconel 182) after ageing at 7oo· C for different times 27 ' 52 ' 54' 56.

7. 1.3 Solution Treated Material Solution treatment at 1 250" C for 1h followed by water quenching removes most of the inhomogeneities associated with the welded microstructure: the coring disappears and the dendrites are replaced by equiaxed grains. A large decrease in hardness from 187

'f

8 Hv / 30 Kg to 139

'f

5 Hv/ 30 Kg accompanies the change in

microstructure. The material solution treated contains a duplex size distribution of large and small ( equiaxed) grains. At high temperatures, a grain with fewer sides tends to become smaller under the action of grain boundary tension forces, while one with more sides tends to grow. A close examination of the shape of the grains of the specimen solution treated and the specimens solution treated and aged, Figs. 29 and 33-39, reveals that small grains usually have less than six sides ( i .e. are shrinking ) while the big grains have more than six sides ( i .e. are growing ) .

Water quenching from 1 250" C to room temperature interrupts the grain growth and retains most of the precipitates in solution, except Ti-rich precipitates which are stable at higher temperatures ( formed during melting ) 27 ' 3 8.

7. 1 .4 Solution Treated and Aged Material The grain size of the material solution treated and aged is similar to that of the solution treated material. There is a negligible grain growth, since the ageing treatments were carried out at low temperatures ( 400 " -7oo· C ) and for short times ( 3-24h ) . Grain growth in Inconel occurs at higher temperatures and longer times. The

grain size of the chromium-containing alloys is restricted up to 1 0 1 0 ° C because of the (finely dispersed ) chromium carbides which block the grain growth. Above 1 0 1 0 ° C,

these carbides begin to coalesce and dissolve and the rate of grain growth increa.ses131.

49

The precipitation density ( as observed with optical microscopy ) within the

grains in specimens aged at 400 " ' 450° ' soo· or sso· c is similar to that of the specimen solution treated at 1 250· C for l h and water quenched. However, the increase i n hardness which is observed at these temperatures may indicate that sub-micron precipitates have probably formed. At high ageing temperatures ( above

6oo· C ) , extensive precipitation is observed at the grain boundaries. Most of the analysed precipitates are Ti-rich; a few are Ti-Nb-rich and Si-rich precipitates are also observed. Ti-rich precipitates are also observed in the as-welded and solution treated materials, so these precipitates are thought to be TiC or Ti{ CN ) which are stable at high temperature ( may be formed during melting ) 27 ' 3 8. Such an assumption

gives the impression that these precipitates may not be formed during ageing. Then,

the extensive precipitation which is observed above 6oo · C might be Ti-Nb-rich,

Si-rich and / or another kind of precipitation which has not been identified. More transmission electron microscopy should be carried out to confirm this proposal.

7 . 1 .5 Ordering in Inconel 182 The formation of the ordered phase,

' 1,

occurs at an early stage of ageing with a

volume fraction dependent on the total amount of additions of Al, Ti and Nb36. Inconel 1 82 has a very low Al content and low Ti and Nb contents, so that the volume fraction of

1

'

is expected to be lower than other Ni-base alloys, such as Ni monic SOA

and Inconel x-750.

In many superalloys the formation of

1

'

during ageing is associ ated with

the

disappearance of the carbide (MC) and with an increase in complex carbides of the form (M 2 3 C6)130, through the reaction

which may also be described as

50

(Ti,Mo)C

+

(Ni ,Cr,Al,Ti)

�

(Cr,Mo) 2 3C5

+

Ni3(Al,Ti )

If such a transformation is applicable for Inconel 182,carbides of the form Cr 2 3C 6 (since Inconel 1 82 contains < 0.01 wt% Mo) should be observed. However, the presence of these carbides are not proved. Such a transformation is either not applicable for Inconel 1 82, or it is applicable only for extended periods of ageing.

There is no direct evidence of LRO or SRO obtained by TEM diffraction patterns along the < 1 00 > and < 1 1 1 > matrix directions with either a slightly convergent or parallel electron beams, from specimens aged for 24h in the temperature range between 400 " - 650" C or 3h at 700" C, following solution treatment at 1 250" C for 1h. Also, X-ray diffraction patterns using either powder camera or diffractometer do not reveal direct evidence of ordering. However, the lattice parameter measurements which were carried out in the diffractometer may show a lattice contraction (Section 6 . 1 .4.2) .

Short range ordering has been observed at an early stage of ageing at temperatures below 525" C in the binary Ni 2 Cr and Ni3Cr alloys58. The transformation from SRO to LRO took a few hundred hours to start in Ni 2 Cr while such transformation was not observed even after 25,000h in a Ni3Cr alloy5S. Klein et al67 also observed superlattice reflections after 730h at 4 79" C in Ni 2 Cr. While Goman'kov et al6S did not observe superlattice spots after ageing (the binary Ni-Cr alloys) at 500" C for 1 036h, but superlattice lines were detected from single crystals of N

about a stoichiometric composition (which is Ni - 32.9% Cr) 6S. Marucco and Nath69 determined that the kinetics of LRO depended upon the Ni : Cr ratio: the higher the ratio the higher the kinetics of LRO , but the ordering reaction was delayed by increased Fe content. It is probable that LRO is not observed in Inconel 182 because it has a different composition than the other alloys investigated (or it is far from the stoichiometric composition). SRO may not have been observed because of the short ageing times which were used (3-24h) in the present work .

51

Other workers5B ' 62 ' 65 ' 66 ' 69 have reported that the formation of /1 was associated with a lattice contraction and the magnitude of the contraction depended on the chromium content and the ageing period69. 0 . 1 1 % lattice contraction was observed in Nimonic 80A after ageing at 450" C for 30,000h, it was only 0 . 03% after 100h6 2 . The lattice contractions of the Ni 3 Cr and Ni 2 Cr alloys were "' 0 .05% and "' 0.25% respectively after ageing at 4 75" C for 25,000h 58. In the present work a maximum contraction of 0.3

'f

0 . 1 % is observed after ageing at 500" C for 24h. This

value is higher than those observed in Nimonic 80A, Ni 3 Cr and NhCr, despite the difference in Cr content between Inconel 182 and these alloys. Further , in the present work, the differences between ( lattice parameter) readings made at various

temperatures are very small and the error associated with the measurement overlap for most of the readings. It is questionable if these differences are significant.

7.2 Creep Behaviour Limited results of the short term mechanical properties of Inconel 182 have been obtained from this work and more are required to fully, understand the creep behaviour of this alloy. However, from the available results, the effect of initial heat treatment prior to creep as well as the effect of the applied stress on the mechanical properties is outlined below.

7.2. 1 Effect of Lhe Heat Treatment Prior to Creep The effect of initial heat treatment prior to creep can be illustrated when the

results of this work are compared with the ( unpublished ) work of Tyrer 132 . Creep specimens of as-welded material were stress-relieved at 700" C for 3h and water quenched; these specimens failed after about 120 and 3060h when tested at 450 and 430 MP a, respectively and at 500" C132 . However the creep specimens of the present

work ( solution treated at 1250" C for 1h, water quenched, stress relieved at 700 " C for 3h and water quenched) failed after 3 and 40h when tested at 450 and 430 MPa.,

respectively. Such results may reflect the importance of the heat treatment prior to creep in determining the mechanical properties of Inconel 182. The heat treatment

52

may affect the strain hardening of the material, especially if a solution treatment is used. The solution treatment of 1 250 " C for 1h softens the welded material. It reduces the hardness of the welded material from 187

'F

8 to 1 39

'F

5 Hv /30 Kg. The followi ng

stress relieving at 700" C for 3h increases the hardness to 149

'F

5 Hv/30 Kg, however,

this value is less than the hardness of the welded material after stress relieving at 700" C for 3h only ( 185

'F

7 Hv/30 Kg) . This effect (of the heat treatment) on the

hardness may reflect a similar effect on the creep properties of Inconel 1 82. Raymond52 mentioned that the heat treatment affected the rupture properties of Inconel x-750. A typical heat treatment was used, 1 150" C/4h/ A C , ·sso· C/24h/ AC

followed by 700" C/20h/ AC. Such treatment provides precipitation of the carbides (chromium carbides) on the grain boundaries as well as 7' . Air cooling allows these precipitates to grow. It was also mentioned that air cooling followed the heat treatment resulted in better creep rupture properties of Inconel x-750

55.

So the heat

treatment prior to creep and the rapid cooling rates which were used in this work may have affected the rupture properties of Inconel 182.

7.2.2 Effect of the Applied Stress 7.2.2. 1 Loading Strain The variation between the

to t a l

loading strain and the applied stress for

the 400 MP a creep test, Fig.55, shows a linear and non-li near relatimiships of low and high loading strai n values, respectively. Creep tests of this work were carried out with high loading strains (between

N

28-52% ) These tests resulted in ductile and shear .

fracture, however another mode of fracture is observed in a specimen tested with

N

35% loading strain, in which, limi ted areas may have failed by intergranular fracture, Fig. 70. So creep tests of intermediate and low loading strains are required to be done to reveal their effects on the subsequent creep and failure mechanism.

It

is thought that some insight into this behaviour may be gai ned by

considering t he activity of dislocations in the material . Loading above the flow stress will generate dislocations throughout the material . These dislocations will

move

53

through the material until they reach an obstacle to motion such as the grain boundaries or precipitates. These may restrict the dislocations movement , giving rise to pile-ups, until a critical value is reached. An increase in the applied stress may allow a greater freedom of dislocation movement and thereby allow greater deformation to occur on loading. In this work a high precipitate density at the grain boundaries (with a moderate density of intragranular prepcipitates) is observed in the specimens prior to creep, Fig.39. Such precipitates may restrict the deformation which arises from loading (up to 1 50 MP a) , in which less than 1 % strains are . observed. Above 150 MPa, large and rapid elongation with a 11 tearing11 sound was observed (and heard) during loading and the creep machine was unloaded and adjusted many times to compensate for this large elongation. The loading strains of the present work represent 65-93% of total strain to failure, and they may have important effect on the subsequent creep and rupture properties.

7.2.2.2 Creep Rate

The minimum creep rates (except that of the 420 MPa test) are observed to be correlated reasonably well with the applied stress, Fig. 52, but a higher creep rate (6 73 x 1Q-4h-1) is observed for the 420 MPa test than that of the 430 MPa (3.29 x .

1 Q-4h-1). This may be due to the higher loading strain of the 420 MPa test which resulted in higher deformation at the start of the creep.

The creep index, n, of 34 is representative of the high stresses which were applied during creep tests. Such a value considerably exceeds those for pure metals which are between 3 and 5 109. Indices of values between 5 and 13 have been observed for 1'-hardened Ni-base alloysll0 -11 2 . Larger values were observed for Inconel 82 12 1. Creep tests at 510" C and at stresses above 430 MP a, resulted in creep index of 70, while at lower stresses (lower than 430 MP a) the creep index was 28 1 2 1, and hence the present results are reasonably compatible with the previous work.

54

The high stress dependence of the creep rate o f second phase hardened materials has been explained by Lagneborg113. The creep process is described

as

consecutive

events of strain hardening and recovery. Precipitates retard the growth of the dislocation networks by exerting a back stress on the network. A similar approach has been proposed by Wilshire and co-workersll0-112 ' 118 ' 1 19. They suggest that the rate­ controlling process is the growth of the dislocation network to generate links of sufficient lengths which allow burst of slip to take place. The distribution of the second phase particles impedes the generation of sources by network growth . At high stresses the dislocations generated may cut through or bow between the particles , depending on the particle size and their spacing. The dislocations then are affected by a frictional stress created by the parti cles. It is suggested that the applied stress operates on all dislocations, but only some of them are capable of becoming mobile, and then determining the creep rate.

The grain size has an effect on the creep rate, since the grain boundaries retard dislocation movement. An intermediate grain size of 1 13 J-lm was found to have the lowest minimum creep rate and resulted in maximum rupture life in Inconel x-750 55. In this work it seems that a large amount of grain growth may have occurred, Fig.59. It is difficult to give an explanation based on the present results, further study is required on this subject.

7.2.2.3 Creep Ductility There is no clear relationship between the creep strain and the percentage loading strain, Fig.56. However, except for that of the 430 MP a test, the creep strain appears to decrease slightly with the increase of loading strain. The creep strain of the

430 MPa test is higher than the other tests. However, the lower loading strain ( than

that which is predicted by the solid line i n Fig. 54 ) of the 430 MP a test may have had an effect. Less deformation may be present in the specimen after loading and more deformation is needed to reach the cri tical deformation to failure during creep. Creep

55

tests with lower loading strains (lower stresses) than those of the present work are required to reveal any relationship with the creep strain.

The strain to failure may reflect the ductility of the material. It increases with the increase of applied stress, Fig . 57, since the amount of deformation in the specimen is proportional to the applied stress. High values of strain to failure (ductility) are observed in this work, because of the large applied stresses. They range between

IV

33 to 55% in the stress range between 380-450 MP a. The large elongations

may be one of the characteristics of Ni-Cr-Fe alloys. Elongation of

IV

50% was

observed in Inconel 82 at stresses above 430 MPa and 5 1 0 · c 12 1. Also, a ductility of 56% was observed in Inconel x-750 when tested at 537• C and with a strain rate of 6 .72 x 1 0-2;s 54.

7.2.2.4 Creeo Failure '

The high stresses which were used in this work are reflected by the time to . failure; at 450 MP a the life time was only 3h. The value of 45Q MP a may be close to the value of the tensile strength of Inconel 182 at 500· C. The approximate tensile strength of Inconel 1 82 at 500" C (which is determined by Fig. 3 18) is 510 MPa. However the real value is not specified. Higher life times were observed when the creep specimens were tested at lower stresses.

The fracture mode is shear and ductile in character; at a stress of 400 MPa, there was a limited amount of intergranular fracture, Fig. 70 . A change in the fracture mechanism may have occurred at lower stresses. However, creep tests at lower stresses are required to reveal the validity of such a proposal.

A distinction has been made between the terms ' cavity' and 'voi d ' . The term ' cavity' is used to describe intergranular holes, typically 1 to 4 ttm in diameter, the growth of which is probably diffusion-control led. The term 'void' is used to describe any other holes (intragranular in location) with a diameter of about 1 0 ttm and above.

56

In this work, no cavities were observed ( although this does not necessarily mean that

they do not exist at some stage during creep ) ; most observations were of

intragranular voids, usually associated with a ductile failure mode, Fig. 74. Voi ds usually nucleate on precipitates ( and inclusions ) . A hard precipitate ( or

inclusion ) disturbs both the elastic and plastic displacement field in a deforming body. The disturbance concentrates stress at the precipitate, and the stress builds up as the plastic strain increases, until the local stress reaches a critical value. The stress then either fractures the precipitate or tears open the interface with the matrix, thereby nucleating a void. Further plastic deformation makes the voids grow and when they are large enough, or when the specimen itself becomes mechanically unstable, they coalesce and fracture the material81. A typical dimple structure on the fracture surface of a sample which contains large precipitates is shown in Fig. 73. It is clear from this micrograph that fracture may have occurred by the growth of i nternal voids and that void nucleation occurred around the precipitates .

The fracture mode which is observed i n this work is a combination of shear and ductile failure, Fig. 72. In high strength material, the strengthening is often attained

through dispersion of precipitates ( or added particles ) . The stress present may crack the interfaces between precipitates and a notch-like defect, either o riginally present

or formed during deformation, may concentrate the stress sufficiently to link the notch with a void. The crack then spreads by this process of tearing between the crack and the next void. The stresses to cause such tearing are relatively high, limiting this kind of fracture to materials of high strength133 . Large cracks are observed in specimens tested at 400 and 420 MPa, Figs.60 and 62 respectively, these cracks may be produced during the large loading strain.

Shear fracture was also observed when Inconel 82 was tested at 430 MPa and 5 1 0 " C, and the fracture surface made a 45° angle with the stress axisl 21 .

57

Transverse sections of creep specimens show wedge-like cracking, Figs . 64 and 65. Such cracks may be detected when the cross section passes through an intragranular crack (like that which is shown in Fig.62) .

High precipitation density with subgrain decorations are observed in the creep specimens, Fig.65. Intensive study is required to understand this behaviour. However, heterogeneous precipitation may have enhanced due to the high dislocation density which is observed by transmission electron microscope, Figs.66-68.

The Monkman-Grant relationship can be used to predict component endurance and its general form is given as

This relationship is empirical and will allow an· estimate of rupture life once the minimum creep rate has been determined. The values of from 2.33

x I Q -3

to 13.2

xi 0-3.

CM�

are found to vary

The reduction of creep rate and increase

in

rupture

time leads to an overall reduction in the Monkman-Grant constant.

At the 450 MPa test, in spite of the higher minimum creep rate than those of the specimens tested at lower stresses, a short lifetime (3h) gives lo.wer Monkman­ Grant constant than predicted from Fig.58. The specimen tested at 450 MPa may have failed prematurely due to the high deformation which was produced during the loading.

Further work on this project should be devoted to fully understand properties of Inconel 182.

the creep

58

CHAPTER 8: CONCLUSIONS ' The following conclusions concerning the microstructure and the mechanical properties of welded Inconel 182 have been established.

1.

Most of the precipitates which were observed i n the welded, solution treated and solution treated and aged materials were Ti-rich. A few Ti-Nb-rich and Si-rich precipitates were observed in the solution treated and aged materials.

2.

The inhomogeneities which were observed in the welded and stress relieved materials were removed by solution treatment at 1 2so· C for 1h.

3.

The precipitation density (observed optically) increased when the solution treated alloy was aged above 600" C.

4.

The hardness of the welded material was higher than after solution treatment and ageing. This reflected the complex annealing treatments' and thermally induced stresses which were produced during the multiple-pass welding process.

5.

Neither LRO nor SRO were observed after ageing at temperatures between 4oo·

6.

and 700" c.

A homogenization treatment of 1250· C for 1h, followed by ageing at 7oo· C for 3h produced much poorer creep properties than in a sample given only the 700" C treatment.

7.

The applied stress determined the amount o f deformation induced during loading prior to creep, the loading strain is relatively small for stresses up to "' 150 MPa above this value high loading strains (up to N 52%) were observed.

8.

Most of the failures observed exhibited ductile and shear fracture.

59

CHAPTER 9 : FUTURE WORK

There is a need for further work to be carried out on the characterization and creep behaviour of Inconel 1 82 which includes the following.

1.

To determine the effect of long term ageing at temperatures b �tween 4oo· and 7oo· C on the microstructure and creep properties.

2.

To determine the effect of complex heat treatment, for example 3-stage treatments as developed for Inconel x-750 and Nimonic 80A,

52 ' 53 ' 55 ' 5 6

on the

creep properties. 3.

To carry out creep tests with lower stresses than those of this work for a full understanding of creep behaviour.

4.

Further investigation of grain growth and the precipitation behaviour during creep.

60

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66

Table (1):

Composition limits of Inconel 182 welding electrode

Element

c

Mn Fe

Cr

Nb Si Ni Ti

Cu s p

Weight per cent Maximum Minimum

0.1 9.5 10 17 2.5 1 1 0.5 0.015 0 .03

5.0 6 13 1 .0 59 •·

Chemical compositions ( weight percent ) of the alloys which are mentioned in this dissertation

TABLE (2) :

Ni

c

Mn

Fe

s

Si

Cu

Cr

AISI 3 1 6 A ustenitic steel

14

.08

2

62

.03

1

.07

18

2t% Cr-1 % Mo Ferritic steel

.23

.12

.55

95

.018

.4

.27

.2

1 .5

72

.03

1

-

2.25

Alloy

Al

-

Ti

Nb

.01

-

3 Mo - . 045 P

.01

.012

.02

1 Mo - . 014 P - .02 V

446 Ferritic steel

-

Inconel 82

72*

.02

3

1

.007

.2

.04

20

Inconel 92

71*

.03

2.3

6.6

.007

.1

.04

16.4

70

.03

2

9

.008

.3

.06

15

Inconel 600

76*

.08

.5

8

.008

.2

.2

15.5

Inconel 6 0 1

60.5

.05

.5

14.1

.007

.2

.5

23

1 .4

Incoloy 800

32.5

.05

.8

46

.008.

.5

.4

21

.4

.23

. 74

98.7

.026

.2

.08

2

68.5

.03

1

.11

.57

96.7

.02

. 78

,_

.. 1 2

.55

95.8

.035

.25

Rex 500

.5

.07

2

90

-

.4

-

6

Inconel x-750

73*

. 04

.5

7

.005

.2

.2

Nimonic 80A

69.5

.1

1

5

1

In eo-weld A

1

ASTM A-212, grade B

-

304 Stainless steel

9.3

ASTM A-387, grade C ASTM A-387, grade D

(* Cobalt included)

-

-

1

-

-

-

-

-

25

-

19

-

-

-

-

-

-

-

.6

2.5

.04 P - .25 N 2

3.2

-

.4

-

-

1.8

-

1 .5 Mo

.015 p . 045 p

-

-

-

-

-

.015 P - .99 Mo

.25

-

1 Mo - .2 V - .5 W

15.5

.7

2.5

1

18

1 . 15

2.25

-

1 .33 2.2

-

-

Others

-

. 0 1 1 P - .5 Mo

2 Co

O':l -..J

Table (3)

Ni-based welding products for dissimilar metaJ welding. The appropriate filler metals and electrodes are selected for specific welding processes. The table shows filler metals for tungsten inert gas (TIG), metaJ inert gas (MIG) and submerged-arc welding processes and electrodes for the metallic-arc welding process. .

. Filler Metals for T.I.G . , M.I.G. and Submerged-Arc Welding �v,,o

��

Nic:hi 20D

·-

.. � �� MONEL

I N C O N EL allay 600

N I M O NI C allay 75

I N COLOY

800 BOOH

al lays 1nd

= c

"'0

�

� M6 C 16 12 8

MGC

': M 23 CG

4

· MOlybdenum

Fig .

(4 )

+

6

8

10

0.4 -tung sten, atomic Ofc,

The rel ati ons hip bet we en chr om i um con ten t, mo lyb den u m t un gst en con ten t and car bid e iden t i ty and rela t i v e qu ant i ty

+ 0. 4

' e

Ct

' 0 �

�

Cr

,(

l. ;X

,( /\,