Stellingen behorende bij bet proefschrift

Joining silicon carbide to au sten itic sta in less steel through diffusion w eld in g

Jan-Paul Krugers

1. Het voortdurend bezuinigen van universiteiten op onder- steunend technisch personeel leidt er onherroepelijk toe dat deze instellingen verworden van onderzoekinstel- lingen tot serviceverlenende instituten.

2. Het opstellen van een eenduidige norm voor de kwantita- tieve analyse van (siliciumhoudende) keramieken is meer dan wenselijk, teneinde resultaten van experimenteel onderzoek, verricht door verschillende onderzoekers, op een zinvolle manier met elkaar te kunnen vergelijken.

3. Gelet op een van de primaire taken van universiteiten, het overdragen van kennis aan studenten, dienen docen- ten verplicht te worden gesteld regelmatig een (opfris)- cursus 'didactiek en toepassing van audiovisuele midde- len in het onderwijs' te volgen.

4. Een menselijke samenleving kan alleen bestaan bij de gratie van verbale communicatie.

5. Het begrip 'nieuwe wereldorde' zoals dit door de intema- tionale gemeenschap wordt gehanteerd, is een utopie vanwege de inherente slechtheid van de mens.

6. Het toedienen van een vaccin tegen ziektes die tot epide- mieen kunnen leiden, dient van overheidswege verplicht te worden gesteld.

7. Onze leefomgeving is alleen te redden door een extreem doorgevoerde regeling op mondiaal niveau volgens het principe: de vervuiler betaalt.

8. In het onlangs verschenen rapport van de werkgroep ethische aspecten van neonatologie van de Nederlandse Vereniging voor Kindergeneeskunde ontbreekt ten onrechte een definitie voor de kwaliteit van het leven. Dit ondermijnt de rechtspositie van artsen die overgaan tot de levensbeeindiging van pasgeborenen.Doen o f laten? Grenzen van het m edisch handelen in de neonatologie, N V K , 1992.

9. In het artikel van lino, over vaste stof reacties tussen SiC en Ni, wordt ten onrechte met geen woord gerept van de aanwezigheid van grafiet in een reactiezone.Y. lino , J .M at.Sci., 26 (1991) 4399-4406.

10. De gevolgtrekking van Batfalsky et al., gebaseerd op het veranderen van de conushoek, dat een diffusielasverbin- ding zonder scheuren kan worden verkregen tussen koper en siliciumcarbide en niet tussen roestvast staal en siliciumcarbide, is aanvechtbaar.P. Batfalsky, J. G odziem ba-M aliszew ski, R. Lison, 5 th In t.C on f 'H igh Technology Joining' o f the B ritish Association fo r B razin g & Soldering, Brighton 3-5 N ovem ber 1987, p .13 /1- 13/7.

Joining silicon carbide to austenitic stainless steel through diffusion welding

Jan-Paul Krugers

Joining silicon carbide to austenitic stainless steel through diffusion w elding

proefschrift

ter verkrijging van de graad van doctor

aan de Technische U niversiteit Delft, op gezag van de Rector M agnificus, Prof.Drs. P.A. Schenck, het openbaar te verdedigen ten overstaan v an een com m issie

aangew ezen door het College van Decanen, op 19 januari 1993 te 16.00 uur

door

Johannes Paulus Hendrikus Maria Krugers,

geboren te Rotterdam , doctorandus in de scheikunde.

Dit proefschrift is goedgekeurd door de prom otor Prof. Dr. G. den O uden

Life’s good - B u t not fa ir a t a ll

Lou Reed

(from 'M agic and Loss')

C o n t e n t s

1 General introduction .......................................................................................................... 11.1 A pplication of ceramics 1

1.1.1 Historical review 11.1.2 Silicon carbide 21.1.3 Applications 31.1.4 Expectations for the fu ture 4

1.2 C eram ic/m etal joining techniques 61.2.1 Introduction 61.2.2 Joining in the solid state 61.2.3 Joining in the liquid state 9

1.3 Joining silicon carbide to stainless steel 121.4 O utline of this thesis 13References 15

2. Theory ............................................................................................................................... 162.1 In troduction 162.2 M eta l/m eta l diffusion bonding 162.3 M etal/ceram ic diffusion bonding 19

2.4 D iffusion processes and in terphase boundary m orphology 222.4.1 Introduction 222.4.2 General description o f the diffusion problem 232.4.3 Interphase boundary morphology 27

2.5 Diffusion in the solid state 292.5.1 Introduction 29

2.5.2 Multiphase diffusion in binary systems 302.5.3 Multiphase diffusion in ternary systems 32

2.6 Therm odynam ic aspects of M-Si-C system s 362.6.1 Introduction 362.6.2 Binary systems 362.6.3 Ternary systems 47

2.7 Metallic inserts 532.8 Active m etal brazing 562.9 M echanical testing 59

2.9.1 Test methods 59

2.9.2 Strength variations 60References 62

3. Experimental procedure .................................................................................................. 653.1 M aterials 653.2 Diffusion bonding equipm ent 68

3.3 M echanical testing 713.4 A nalytical m ethods 73

3.4.1 Optical microscopy 733.4.2 X-ray diffraction (XRD) 73

3.4.3 Electron probe micro analysis (EPMA) 743.4.4 Auger electron spectroscopy (AES) 75

References 77

4. Direct b o n d in g .................................................................................................................. 784.1 In troduction 784.2 H igh vacuum bonding experim ents 78

4.2.1 HPSiC/AIS1316 79

4.2.2 RBSiC/AISI316 924.3 Bonding experim ents in shielding gas 97

4.3.1 HPSiC/AIS1316 974.3.2 RBSiC/AISI316 98

4.4 Conclusions 99References 101

5. Interlayers.............................................................................................................................. 1025.1 In troduction 1025.2 Experim ents w ith interlayers 103

5.2.1 N i as interlayer material 1045.2.2 Cu as interlayer material 105

5.2.3 Cu-x%Ni as interlayer material 107

5.3 Conclusions 110References 111

6. Diffusion bonding experiments with copper-nickel interlayers.............................1126.1 In troduction 1126.2 C opper - 5 wt% nickel 113

6.2.1 Welding experiments 1136.2.2 Structure of the joint 113

6.2.3 Kinetics 1326.2.4 Mechanical behaviour 136

6.3 C opper -1 0 wt% nickel 1416.3.1 Introduction 141

6.3.2 Structure of the joint 141

6.3.3 Kinetics 1436.3.4 Mechanical behaviour 145

6.4 General m odel of the SiC/Cu-x% N i/A ISI316 joint 147

References 149

7. The role of residual s t r e s s e s ...............................................................................................1507.1 In troduction 1507.2 Stress state evaluation using linear elastic fracture m echanics 152

7.2.1 Computation of Kin for joints without an interlayer 1587.2.2 Computation of Ki n for joints with an interlayer 159

7.3 D iffusion bonding using interlayers of reduced size 1617.4 Diffusion bonding silicon carbide to silicon carbide 164

7.5 Conclusions 166References 167

Sum m ary ......................................................................................................................................... 168

Sam envatting ..................................................................................................................................171

D a n k b e tu ig in g ............................................................................................................................... 174

CHAPTER 1

General introduction

1.1 Application of ceramics

1.1.1 Historical review

A lthough the sanguine hopes concerning the in troduction and application of ceramic m aterials are not entirely redeem ed, it is not a too im puden t assertion to state that research and developm ent of ceramics are still expanding. Currently, in terest is m ainly focused on silicon based ceramics, especially silicon carbide and silicon nitride, because of their versatility. They are applied in autom otive energy system s, cutting tools, heat

exchangers, m echanical seals and so on.

The first serious efforts in the field of technical ceramics (also denoted as structural, special o r engineering ceramics) w ere m ade in 1953 in the U nited K ingdom . These attem pts concerned a program , com pletely financed by the governm ent (M inistry of Supply), w hose purpose it w as to find ou t w hether special ceramics could offer a solution to the existing technological m aterial problems. U p till then, the only w idespread application of technical ceramics involved alum ina (A120 3) w hich w as used in sparking plugs of com bustion engines. It soon appeared that alum ina w as inadequate for m any structural applications and silicon based ceramics w ere developed that were m ore suited. After 1973, the U nited K ingdom lost its leading position in the area of

technical ceramics w hereas o ther countries like Japan, the U nited States of Am erica and, to a lesser degree, the European C om m unity started to show interest in the research and developm ent of technical ceramics. A significant contrast betw een the research program s of the U nited K ingdom com pared w ith those of the o ther countries, is that the form er have a m ore fundam ental character w hereas the latter show a m ore application directed approach. This is connected w ith the fact that governm ent support in these countries is

m ore aim ed at the industry and concentrated on the practical u tilization of ceramic m aterials in gas turbines and heavy du ty diesel engines [1.1].N ow adays, it appears tha t m any prognoses have been too optim istic w ith respect to both industrial investm ents and applications of technical ceramics. This m ay largely be ascribed to the practical problem s connected w ith the in troduction of ceram ic m aterials in an "age of metals". A major obstacle in this respect is that m ost design engineers and constructors are no t fam iliar w ith ceramic materials.

1

The prim ary reasons w hy structural ceramics are of vital im portance to m aterial engineers are tha t they can replace m etals in high-tem perature engineering applications.

This provides tw o benefits:in contrast w ith the decreasing supply and accessibility of m etals such as chrom ium , cobalt, nickel, titanium and their alloys, the base m aterials for technical ceramics are am ply available;there is an increasing em phasis on energy conservation accom panied w ith the desire to use energy m ore efficiently; ceramic m aterials are m ore heat resistant and generally have a low er specific gravity than metals.

A severely lim iting factor of ceramics is tha t they are difficult to m anufacture in precise shape and that finishing processes are troublesom e and expensive. Furtherm ore, it has appeared up till now that ceramic products, even the uncom plicated shapes, are difficult

to fabricate reproducibly.

1.1.2 Silicon carbide

A prom ising candidate as high-tem perature structural ceramic for w ear-resistant applications is silicon carbide. In the following several characteristics w ill be enum erated together w ith some com m ents and com parisons w ith other ceramic m aterials [1.2].The elastic properties of silicon carbide are high com pared w ith those of alum ina and silicon nitride and also its hardness is superior to the hardness of these m aterials (it is only surpassed by diam ond, cubic boron nitride and boron carbide). This m akes silicon carbide very suitable for applications like mechanical seals and bearings w hich require resistance to abrasion. S trength values, however, are low er than those of zirconia and silicon nitride (still h igher than those of alum ina) and can be increased by reducing fracture origins like pores, inclusions and coarse particles th rough im proved m anufacturing processes. Typical values of the critical stress intensity factor KIC for SiC, Si3N 4 and Z r0 2 are 4.6, 5.5 and 8.0 MNm"3'72, respectively. A t h igh tem peratures, exceeding 1200°C, SiC dem onstrates greater strength than any other construction m aterial, including m etals, w hich m akes it pre-em inently suited for use in gas turbines o r heat exchangers. The therm al shock resistance is considered good w hen high-density m odifications are taken into account. Only silicon nitride appears to be m ore resistant to therm al shocks.

The therm al conductivity of SiC is relatively h igh (as w as already dem onstrated by the h igh value for the therm al shock resistance) w hich is considered beneficial for m aterials used for mechanical seals and bearings. The therm al expansion coefficient is ra ther low com pared w ith m etals and this also contributes to the high therm al shock resistance.

2

Silicon carbide possesses extrem ely high resistance to corrosion and is not attacked by either acids or alkalis. However, im purities a n d /o r sintering additives decrease the

resistance to corrosion.The m echanical reliability of ceramics as structural m aterials m akes or breaks the application in practice. As m entioned before, the fracture strength of ceramic m aterials is strongly dependen t on the dim ension and num ber of various defects and these are norm ally present in a b road statistical distribution (see C hapter 2). The am elioration of ceramics reliability can be acquired through design-related approaches (including avoidance of stress concentration) and through efforts to increase fracture toughness and

to reduce fracture-form ing defects.A m ost im portan t aspect of ceramic m aterials is their m icrostructure. The m icrostructure is strongly influenced by the m anufacturing process (raw m aterial properties like grain size, additive am ounts and properties, m oulding conditions, sintering conditions) which, therefore, should be controlled accurately.

1.1.3 Applications

The applications of silicon carbide com pacts can roughly be d iv ided into tw o categories:1 abrasion-resistant and corrosion-resistant com ponents;2 heat-resistant and refractory com ponents.

At present, m ost applications fall into the form er category, a lthough the developm ent of SiC com pacts as structural m aterials w ith high density w as originally started w ith the purpose of application to com ponents for gas turbine engines. A few exam ples of

applications of bo th categories w ill be given below.

A brasion-resistant and corrosion-resistant com ponents M echanical seals, pum p com ponents and abrasion-resistant liners fall in this first category.U ntil recently, m aterials like various m etals, carbon, WC and A120 3 w ere used for m echanical seals. As mechanical seals have to com ply w ith several requirem ents like ro tating at h igh speeds while m aintaining a seal against hot water, slurry, chemicals and oils, a w ide range of properties is needed, including resistance to abrasion, corrosion and therm al shock, a h igh elastic m odulus, a low friction coefficient and a low specific

gravity. It appears tha t silicon carbide is prom inent in each of these respects and can be applied as a replacem ent for conventional m aterials in severely corrosive a n d /o r erosive

environm ents.

Pum p com ponents like bearings, sleeves and pum p shafts are used in the sam e

3

environm ents as m echanical seals. In addition to the dem ands for m echanical seals m entioned before, these com ponents require h igh strength because they m ay experience

tensile stresses. Furtherm ore, the effect of dim ension on strength reduction is especially im portan t in this respect because m any of the com ponents are large.In earlier days, nozzles and liners for handling slurry and oil w ere m ade of metals because of their m echanical reliability. However, w ith the developm ent of high-strength SiC w ith its superior abrasion resistance, metals are being more and m ore replaced in

these applications.

H eat-resistant and refractory com ponents The goal of im proving therm al efficiency has led to the developm ent of gas turbine engines and gas turbines for pow er generation consisting of ceramics instead of m etals or superalloys. Potential applications for SiC involve com bustors, turbocharger rotor vanes and blades. N otw ithstanding the severe research efforts that have been conducted, the stage of practical application has yet to be reached.W hen therm al efficiency is an im portan t issue, it is of prim e concern to be able to perform heat exchange as well as to circulate and transport gas under pressure w ithout

reducing gas tem perature. In this respect, it should be noted that the use of heat exchangers and ceramic fans has already been m entioned in literature.

1.1.4 Expectations for the future

The expenditures on ceramics has increased significantly from 1970 till about 1985. After

this, bo th the investm ents in ceramics and research and developm ent on ceramic m aterials fell short of predictions m ade in the past.This stagnation m ay be partly attributed to the unrealistic faith in ceramic m aterials as

replacem ents for metals. It has become clear that ceramics cannot just substitu te metals b u t that also, in m any cases, joints betw een these m aterials have to be m ade. The latter appears to be a m uch harder task to accom plish than was expected at first. This resulted

in less enthusiasm and willingness to participate from the side of the industry than w as anticipated.Presenting figures of estim ates and prognoses w ith respect to the m arket size of structural ceramics is a difficult and risky task, especially w hen com parisons betw een different countries are involved. This is due to several reasons, such as m arket fluctuations th rough the years and the definition of structural ceramics w hich varies

from country to country. In view of this, separate estim ates and forecasts are given for

4

ind iv idual nations in U.S. dollars in Table 1.1 [1.3,1.4]. It is clear from this table that the

absolute values w ith respect to the m arket predictions are still increasing, b u t that the grow th of m arket size per year dim inishes for bo th the USA and Japan. Furtherm ore, it is obvious tha t bo th the estim ated m arket size and the average annual grow th rate for structural ceramics are low for W estern Europe in com parison w ith Japan and the USA.

Table 1.1 M arket estim ates and prognoses for advanced structural ceramics (millions of U.S. dollars; AAGR = average annual g row th rate) [1.3, 1.4].

USA AAGR (%) Japan AAGR(%)

W esternEurope

AAGR(%)

1982 276

49 (3 yrs)

1985 920

33 (2 yrs)

1987 171 1620

1989 36 (3 yrs) 920

1990 433

22 (5 yrs)

17 (13 yrs) 4 (6 yrs)

1995 1160

18 (5 yrs)

1150

2000 2645 13145

5

1.2 Ceramic/metal joining techniques

1.2.1 Introduction

W hen joining m etals to ceramics, tw o principal problem s have to be overcome.One problem concerns the accom plishm ent of well bonded m etal/ceram ic interfaces (that is, free from unbonded areas), the other problem concerns the effective relaxation of therm al stresses w hich originate from the therm al expansion m ism atch betw een metallic and ceramic com ponents to be joined a n d /o r the form ation of brittle reaction products w ith their ow n mechanical and physical properties.Various m ethods have been developed to join m etals to ceramics. Several classifications can be em ployed to discern one joining technique from another. In this section the

distinction is based on the state of aggregation of the m aterials to be joined. Obviously,

this categorization is rather arbitrary and sometimes one joining m ethod m ay be grouped in m ore than one division, b u t in that case the prevailing choice reported in

literature w ill be preferred.

1.2.2 Joining in the solid state

M echanical joiningThere are m any w ays in w hich metals can be attached m echanically to ceramics [1.5, 1.6]. These m echanical joining processes have in com m on that they are cheap and simple to perform . How ever, they also show draw backs like lack of continuity of the joint and the im practicability of dem ounting the bonded com ponents. Examples of mechanical joining techniques are crim ping, clamping, shrink fitting, bolting and screwing. It is obvious tha t various restrictions exist w ith respect to the applicability of mechanical joining. Im portan t aspects w hich have to be considered before perform ing mechanical

attachm ent of m etals to ceramics are:expenses m ay be high w hen complex diam ond m achining is required;gas tightness can be achieved by applying a soft m etal or a com pliant metallic

interlayer;

stress concentrations should be avoided;therm al m ism atch of the com ponents to be joined requires low -tem perature

applications.Practical applications of this joining technique are h igh-tem perature heat exchangers, high perform ance engines (keying of blades to the rotors of gas turbines and bolting of ceramic piston crowns in reciprocating engines). M ore com m on exam ples are furnace

6

refractories (using mechanical interlocking or metal brackets) and table lamps.

A dhesive bondingA dhesives m ay provide a simple and inexpensive w ay of joining m etals to ceramics.

However, there are a few constraints on the use of adhesives in joining m etals to

ceramics. The m ost prom inent are the low operating tem peratures of the joined p roduct (up to 200°C) and the low to m oderate strength (which can be increased by prem etallizing the ceramic [1.7]).M ore h igh-tem perature reliable are the ceramic adhesives based on inorganic com pounds. However, there is little know n about these adhesives w hich is the m ain reason w hy their application is still lim ited to non-critical areas w here exposure to mechanical or therm al shock is minimal.

Friction w elding

This w elding technique takes advantage of mechanical energy by converting it into heat at the joint to be w elded. In this process, the ceramic com ponent is kept stationary and is pressed against the rapidly spinning metallic com pound w hich flows plastically. The

time required to m ake a friction w eld lies in the range of a few seconds [1.7]. A serious

d isadvantage of the friction w elding equipm ent is that it is ra ther expensive.Friction w eld ing has been reported to be successful for several com binations of ceramics and m etals like alum inium , titanium and copper [1.8].

U ltrasonic bondingUltrasonic bonding is a m ethod in w hich tw o m aterials can be joined by im posing an

ultrasonic w ave upon the bonding interface. It has been dem onstrated that alum inium can be bonded ultrasonically to various ceramics in air [1.9]. The benefits of this joining

technique are the short process tim e (0.1 to 1 second) and the absence of both an external heating source and a flux. A serious draw back is tha t the surfaces to be bonded have to be very clean and free from oxide layers. For instance, titanium can only be bonded to silicon nitride in an ultra-high vacuum environm ent.

Pressure w eldingPressure w elding utilizes hot pressing or hot isostatic pressing of pow der com pacts consisting of m etals a n d /o r ceramics [1.10]. Metal pow ders w hich can be used in this processing m ethod, are refractory metals and superalloys. However, w hen preform ed sintered ceramic com ponents are applied, low er m elting m etals can be used in joining.

An im portan t prerequisite is that the m aterials to be used, have both m echanical and chemical compatibility.

7

A variant of this m ethod w hich is often used, is the graded pow der bonding technique

[1.6]. In this technique, a ceram ic/m etal bond is accom plished by pressing layers of ceramic pow der and m etal pow der together. G raded pow der seals are ceramic-metal com posite m aterials com posed of a series of such layers. The com position varies

continuously from the m etal to the ceramic and this type of seal can be produced by both hot pressing and pressing and sintering at elevated tem peratures. In this w ay it is

possible to keep properties like therm al expansion under control.

Diffusion bondingThe diffusion bonding process (also denoted as diffusion w elding, solid state bonding and therm ocom pression bonding) is a joining process w hich occurs in the solid state. It can be applied for bonding metals to ceramics. The properly prepared surfaces of the m aterials to be joined are brought into intim ate contact w ith each other under specific conditions of tem perature, time, pressure and environm ent. The process tem perature lies betw een 0.5 and 0.98 Tm (Tm is the m elting point in K of the lowest m elting material).

The process time m ay be varied from several m inutes till about 24 hours, depending on the process tem perature and the m aterials to be joined. The applied pressure is set above the level required to ensure uniform surface contact, b u t below the level that w ould cause macroscopic deform ation. In the initial period of diffusion bonding small-scale plastic deform ation of the metal com ponent is the m ost im portan t operating m echanism (metals generally do have a low yield stress at elevated tem peratures). The atm osphere usually is a vacuum or a shielding gas w ith a low oxygen activity; oxide ceramics som etim es are joined to m etals in air.Diffusion bonding m ay be executed directly or w ith the aid of a (usually) metallic interlayer. A n interlayer is often applied w hen undesired chemical reactions m ay occur betw een the m etal and the ceramic and form ation of hard and brittle products is to be

expected. Additionally, an interlayer is used w hen large residual stresses are expected to develop as a result of the therm al expansion m ism atch betw een the ceramic and the metal.

8

1.2.3 Joining in the liquid state

Brazing, braze w elding and soldering Roughly speaking, there are three techniques w hich closely resem ble each other bu t w hich differ significantly w hen looked at in m ore detail [1.11]. These are brazing, braze

w elding and soldering, respectively.Brazing represents a group of processes w hich produces coalescence of m aterials by heating to an appropriate tem perature and by using a filler m etal having a liquidus above 450°C and below the solidus of the base materials. The filler m etal is d ispersed betw een the closely fitted surfaces of the joint by capillary attraction.Braze w elding is a w elding m ethod in w hich the filler m etal is deposited in a groove or fillet exactly at the place w here it is to be used. Capillary action is no t a factor in d istributing the filler metal. In som e instances, lim ited base m aterial fusion m ay occur. Soldering involves coalescence of m aterials by heating them to an appropriate tem perature and by using a filler m etal (solder) having a liquidus not exceeding 450°C and below the solidus of the base m aterials. In this case, capillary action plays a param ount role. The essential aspect of a soldered bond is that a m etallic bond is produced by a m etal solvent action. In several instances a flux is applied. This is a liquid, solid or gas w hich, w hen heated, is capable of prom oting or accelerating the w etting of the m aterials to be joined by the solder. Its purpose is to rem ove and exclude sm all am ounts of oxides and other surface com pounds from the surfaces being soldered. Crucial in this respect is tha t the capillary attraction betw een the base m aterial and the

filler m etal is m uch higher than that betw een the base m aterial and the flux. The use of fluxes is also com m on in brazing and braze welding.A major problem in brazing and soldering ceramics to m etals is the inherent difficulty

to w et the ceramic m aterials w ith conventional brazing filler metals. W etting depends u p o n the surface tension of the braze and the base m aterials and also on physical and m etallurgical reactions betw een the braze and the base m aterials. The com m on criterion of w ettability can be expressed by the contact angle (9) w hich is given by Young's equation as follows:

Q Ifc ~ ^CM n ncos 6 = ------------- ( i .i)Ym

in w hich y is the surface or interfacial tension and the subscripts C, CM and M identify the ceramic surface, the ceramic-metal interface and the m etal surface, respectively. If the interfacial energy yCM is larger than the ceramic surface energy yc , then cos0<O

(m eaning that the contact angle is larger than 90°) and the liquid will not spread freely

9

or flow into capillary gaps (Fig. 1.1).Bonding can often be attributed to interlocking particles or penetration of liquid metal into surface pores and voids. W etting is therefore an im portant factor as far as bond form ation is concerned. The relatively poor w etting behaviour of ceramics can be rationalized by the fact that the ionic or the covalent structures of ceramics lack the delocalized bonding electrons present in metals and hence m etal/ceram ic interfaces can be considered as significant electronic discontinuities. In order to im prove the situation, the ceramic surface can be prem etallized so that the braze adheres to this metallic layer. W etting can also be prom oted by letting the braze react w ith the ceramic surface thereby changing the surface chemistry.

Fig. 1.1 W etting of a solid by a liquid.

The procedure of m etallizing which is m ost broadly used, is the m oly-m anganese process in w hich a pow der m ixture of a glass and m olybdenum or m olybdenum trioxide and m anganese or m anganese dioxide is applied to a ceramic surface by painting, screen

prin ting or spraying. The bonding is then realized by heating in a m oist hydrogen

atm osphere at 1500°C. Further im provem ent of the w ettability of the ceramic surfaces is gained by cladding the m etallizing coating w ith nickel. Finally, a bond is accom plished using a conventional braze which usually is the Ag-28Cu eutectic because of its m oderate m elting tem perature (780°C) and its superb ductility. Difficulties arise w hen the ceramic does not contain glassy binder phases or is not an oxidic ceramic. Furtherm ore, considerable skill is required in applying the m oly-m anganese m ixture, as problem s m ay be encountered due to poor w etting of the ceramic by the glass phase or

due to excessive w etting of the metallic particles [1.12],

A second m ethod of m etallizing, the so-called "active m etal brazing" m ethod, will be

Vapour

Liquid \ Vapour

9 Lr / Liquid

rsoua M9 > 90° - no wetting 0 < 90° - wetting

10

treated in more detail in C hapter 2, since this technique has been applied in brazing

silicon carbide to austenitic stainless steel. This technique has the great advantage of

being a one-stage process in contrast to the m oly-m anganese process.

Fusion w elding

Fusion w elding is a technique w hich comprises a broad range of w eld ing m ethods like arc w elding, laser beam w elding and electron beam welding. Fusion w eld ing of ceramics to m etals m ay be accom plished by filling the joint w ith m olten m aterial resulting from m elting the edges of both com ponents or w ith m elt from a filler of sim ilar m aterial. In the case of joining m etals and ceramics this technique is very lim ited for several reasons. In the first place there should be ideally a close m atch of m elting points and therm al contraction characteristics of not only the m etal and the ceramic, b u t also of the m aterial form ed in the w eld pool. Obviously, this sim ilarity is rarely encountered in practice. In addition, some ceramics (including SiC) sublim e or decom pose before m elting takes place, while others vaporize rapidly w hen molten. Nevertheless, some successful fusion w elds have been reported betw een refractory m etals (Mo, Ta, Nb) and oxide ceramics

[1 .6 ].

GlazingG lazing m ay be considered as the "mirror image" of brazing: m etallizing the ceramic surface prom otes its w etting by m olten m etal braze, w hereas oxidizing the m etal surface

prom otes the w etting by m olten oxide glass [1.5]. The surfaces of m ost m etals are covered by th in oxide films w hich can be dissolved in the glass so tha t adherence is lost. Therefore, it is com m on practice to thicken these films by preoxidation, as a result of

w hich they rem ain strongly attached to their m etal substrate. G lazing and brazing have a lot in com m on bu t a param ount difference is that glasses are m uch m ore brittle than brazes, w hich m eans that the contraction characteristics of the specim ens to be joined should closely m atch. However, the therm al contraction properties of m etals and ceramics generally are m arkedly different im plying that few com binations satisfy the criteria for successful glazing. Nevertheless, thriving results have been reported about

glazing m etal/ceram ic com binations by using glasses w hich devitrify w hen cooled, thereby show ing contraction properties w hich m atch those of metals.

11

1.3 Joining silicon carbide to stainless steel

As m entioned before, silicon carbide, used as a structural m aterial, has several favourable characteristics com pared to metals, superalloys and oxide ceramics, especially w hen placed in harsh environm ents. These advantegeous properties include hardness, heat resistance (decom position at approxim ately 2400°C), h igh therm al conductivity, low specific gravity, low therm al expansion coefficient, h igh corrosion resistance and m agnificent resistance against oxidation. These qualities m ake SiC exceptionally suited for use in gas turbine engines and heat exchangers [1.2],Serious draw backs of SiC are the low degree of reproducible m anufacturing and its inherent brittle character.

Austenitic stainless steel is often used in situations w here the tem perature does not exceed ±700°C and the environm ent is corrosive (for instance in petrochemical installations). A good example is austenitic stainless steel AISI316. It contains about 2.5 wt% m olybdenum (which enhances the corrosion resistance, especially p itting corrosion, and the creep resistance) and has the highest stress-rupture characteristics of all 300

series steels.The lim itations of the use of austenitic stainless steels are related to their poor corrosion resistance in chloride containing surroundings and strong inorganic acids, w ith the exception of nitric acid. Furtherm ore, their oxidation resistance to aggressive gaseous surroundings is inferior at elevated tem peratures.

Silicon carbide has often been (and still is) referred to as a prom ising candidate as high- tem perature structural m aterial in harsh environm ents. However, m onolithic ceramic constructions are seldom encountered. In m ost cases ceramic m aterials are used in com bination w ith metals. This has led to a situation in w hich the developm ent of joining techniques for ceramics to metallic materials has become decisive. The large-scale in troduction of ceramics will m ainly depend on this developm ent and the rate at which

this will occur.In the present study the attention is focused on the com bination silicon carbide and

austenitic stainless steel AISI316.

12

1.4 Outline of this thesis

In this d issertation the results are presented of a s tudy dealing w ith the joining of silicon carbide to austenitic stainless steel th rough diffusion welding.Diffusion w elding w as selected as joining technique for the follow ing reasons:

diffusion bonds generally are capable to w ithstand high process tem peratures; during the diffusion bonding process no liquid phase is form ed, only solid state diffusion processes take place (liquid phases are norm ally no t easy to handle, unless special precautions are taken);

the bonding process occurs in one single step in w hich several parts can be joined; in case com ponents of simple geom etry are to be joined, only a uniaxial pressure is needed to establish sufficient contact betw een the m ating surfaces; com ponents of m ore com plex geom etry require a m ore sophisticated and versatile (and therefore a far m ore expensive) variant of the diffusion bonding process, the so- called hot isostatic pressing technique;if the therm al contraction characteristics and the elastic m oduli of the com ponents to be joined do not m atch properly (as is usually the case) and the form ation of excessive reaction products w hich are detrim ental to the bond is possible, then (metallic) interlayers m ay be in troduced to accom m odate residual stresses that m ay occur a n d /o r to prevent unw anted diffusion of particular elem ents across the interface.

In C hapter 2 several theoretical aspects w ith respect to diffusion bonding are given, together w ith som e considerations concerning kinetics and therm odynam ics of binary and ternary systems. Moreover, the active m etal brazing technique is discussed and mechanical testing m ethods are review ed, together w ith statistical handling of m echanical strength data.In C hapter 3 a detailed description is presented concerning the m aterials and equipm ent w hich are used and of the analytical techniques w hich have been applied.

The direct diffusion bond ing experim ents, that is, w ithout m aking use of a m etallic insert betw een silicon carbide and austenitic stainless steel are dealt w ith in C hapter 4.After a brief evaluation of the results given in C hapter 4, diffusion bond ing experim ents betw een SiC and AISI316 w ith various interlayers (Ni, C u and binary Cu-N i alloys) are discussed in Chapter 5.In C hapter 6, diffusion w elding experim ents using two interlayer m aterials w hich are considered best su ited for joining silicon carbide to austenitic stainless steel, are further

exam ined. This exam ination includes the determ ination of optim al process conditions

w hile other aspects like interface m icrostructure (determ ined on the basis of detailed

13

analyses) and m echanical strength data and their m utual relation are taken into account.

Finally, the results of FEM calculations of the residual stress in the joint are presented in C hapter 7, together w ith tw o examples dem onstrating the effect of geom etry and

m aterial properties on the residual stress level.

14

References

1.1 J.T. Buma, De Ingenieur, 6 (1985) 32-37.1.2 K. Yamada and M. M ohri, Silicon Carbide Ceramics-1, S. Sdmiya (ed.), Elsevier

A pplied Science, London 1991, p .13-44.1.3 H igh Tech Ceramics (Part C), P. Vincenzini (ed.), Elsevier Science Publishers,

A m sterdam 1987.1.4 Ceram ics Today - Tom orrow 's Ceramics (Part D), P. Vincenzini (ed.), Elsevier

Science Publishers, A m sterdam 1991.1.5 M.G. N icholas and R.J. Lee, Metals and M aterials, 6 (1989) 348-351.1.6 M.M. Schwartz, Ceramic Joining, ASM International, Ohio 1990.

1.7 J. Fernie, W elding Institute Bulletin, S ep t/O ct 1990.1.8 A. Suzum ura, T. O nzaw a, S.K. Budhi, A. Ohm ori, Y. A rata, Proc.Int.M tg. on

Adv.Mats. (Vol.8), P ittsburgh 1989, p.269-274.1.9 K. M iyazaw a, S. M atsuoka, T. Fujii, T. Suga, Proc.Int.M tg. on Adv.M ats. (Vol.8),

P ittsburgh 1989, p.275-280.1.10 M. Erz and H.W. Hennicke, Ceramics in A dvanced Energy Technologies,

H. Krockel et al (eds.), D. Reidel Publishing Company, D ordrecht 1982, p .138-156.

1.11 W.H. Kearns (ed.), W elding H andbook Vol.2 (7th ed.), A m erican W elding Society, M iami 1978, p.370-458.

1.12 M.G. Nicholas, Designing Interfaces for Technological Applications, S.D. Peteves (ed.), Elsevier A pplied Science, London 1989, p.49-76.

15

CHAPTER 2

Theory

2.1 Introduction

In this chapter the diffusion bonding process will be dealt w ith in more detail. This will

include both m eta l/m eta l and m etal/ceram ic combinations.The m odelling of m eta l/m eta l diffusion bonding has already been treated extensively elsew here [2.1, 2.2], bu t will be sum m arized here for tw o reasons. Firstly, the m eta l/m eta l diffusion bonding m odel can serve as a proper basis for a m etal/ceram ic diffusion bonding m odel and, secondly, m eta l/m eta l diffusion bonding plays a role w hen applying a metallic inter layer betw een the ceramic and the m etal to be joined. Several a ttem pts have been m ade to m odel the m etal/ceram ic joining process in the solid state [2.3, 2.4]. Three aspects will be treated here in m ore detail, nam ely the

interface form ation of a m odel system and the diffusion processes w hich take place a t/ac ross the interface and their influence on the m orphology of the interphase boundary.Subsequently, some kinetic and therm odynam ic aspects w ill be discussed, thereby focusing on the silicon carb ide/m etal interactions. In addition, the role of metallic interlayers will be considered and the general reaction behaviour of various S iC /m etal couples will be presented. After that, the active metal brazing process is given a closer look at and, finally, some mechanical testing procedures and the statistical handling of

test results are reviewed.

2.2 Metal/metal diffusion bonding

Joining (dis)sim ilar m aterials to each other by m eans of diffusion bonding is a process w hich requires sm ooth and clean surfaces of the m aterials in order to ensure intim ate contact. This contact is facilitated by perform ing the diffusion bonding at elevated

tem peratures and by m aking use of a relatively small mechanical pressure.Though diffusion bonding is a joining technique which is know n since the sixties and w hich still rejoices at increasing interest and application, relatively little research has been carried o u t on the m echanism of bonding.In the case of conventional m eta l/m eta l diffusion bonding, that is w ithout m aking use of a diffusion aid (solid filler metal), a three-stage m echanistic m odel exists which properly describes the form ation of the weld [2.1].

16

In the first stage (Fig. 2.1) plastic deform ation of the contacting asperities occurs, principally by yielding and by creep deform ation. These m echanism s provide close contact betw een the tw o m ating surfaces over a large section of the interfacial area. W hen this first stage is com pleted, the joint consists of grain boundaries w ith voids in between.

c dFig. 2.1 Different stages during diffusion w elding show ing the course of events

leading to joint form ation [2.1].

D uring the second stage grain boundary diffusion plays a param ount role: the volum e

of the voids will be reduced and some of them will even disappear. Besides, the grain boundaries tend to shift to their equilibrium position w hich leads to the situation that

m any voids are located w ith in the grains.In the last stage the rem aining voids are elim inated by volum e diffusion of m aterial to

the void surface (equivalent to vacancy diffusion aw ay from the voids). This leads to the situation that the initial interface has (almost) com pletely disappeared and a continuous m aterial transition results.

It m ust be rem arked in this scope that the three stages often show some overlap and that one m echanism w hich governs one stage, can also be operative during o ther stages. A ccording to Derby and Wallach [2.2, 2.5] the course of events can be described by three

17

groups of m ass transfer processes analogous to pressure sintering (Fig. 2.2). These are: a bulk deform ation of the asperities by plastic yielding or creep; this process is

prim arily driven by the applied pressure du ring bonding (a m inor influence is to be expected from surface tension effects);

b m aterial transport from the void surface tow ards the neck, realized by surfacediffusion, volum e diffusion and through the vapour phase; the driv ing force

in this case is the difference in surface curvature across the surface of an interfacial void;

c m aterial transport from the interface tow ards the neck on the void surfaceaccom plished by diffusion along the bond interface and by bulk diffusion; the driv ing force is the gradient of the chemical potential along the bond line.

5 a n d 6

a surface so u rc e m echan ism s (1 an d 2 ) ; b bon d -lin e sou rce m echan ism s (3 and 4 ) , c bulk deform ation m echan ism s (5 and 6)

Fig. 2.2 Schematic view of different mass transfer paths during diffusion bonding [2 .2],

It is recognized now adays that achieving good interfacial contact is the rate determ ining

step in the joining of tw o m etals. The param eters determ ining the m echanism in each phase are process tem perature, pressure on the specimens and initial surface roughness.

In the specific case of diffusion bonding of copper it appeared tha t the dom inant m echanism is that of pow er law creep [2.2].

18

2.3 Metal/ceramic diffusion bonding

O nly few researchers have investigated the theoretical aspects of m eta l/ceram ic joining. In fact, those w ho d id study the theoretical features of bonding m etals to ceramics, generally based their view s on existing m odels of m e ta l/m eta l joining or sintering

m odels [2.6].In the case of m eta l/m eta l diffusion bonding the diffusion activated processes of local m ass transport at the interface are rate determ ining w hile for m eta l/ceram ic interfaces it is found that a m inim um process tem perature is necessary to achieve bonding

(intim ate contact betw een the tw o m ating surfaces is not sufficient).In general, the bonding process can be subdivided into tw o stages provided that no

(chemical) reaction takes place.The first stage is the deform ation and associated flattening of the initial asperities at the m etal side. The ceramic surface w hich also shows surface roughness, is assum ed to be

rigid and undeform able (Fig. 2.3). The em anating m echanism s are plastic deform ation,

creep and diffusion.The second stage (partly overlapping the first step), com prises adhesion. The driving force for interface adhesion is given by the Young-Dupre equation:

W a = 'I'mG + 'I'CG ~ ^CM ^2-1

in w hich Wa is the w ork of adhesion, y is the free surface or interface tension and the subscripts MG, CG an d CM stand for m eta l/gas, ceram ic/gas and ceram ic/m etal,

respectively.In this section the A1/A120 3 interface will be treated as a m odel for o ther m etal/ceram ic system s [2.4]. From other studies [2.7, 2.8, 2.9, 2.10] it appeared that grain boundary diffusion in A1 is the m ost probable bonding route w hen joining this m etal to A120 3. Thus the m odel w hich is considered, is bu ilt on the assum ption that contact grow th occurs by diffusion along the A1/A120 3 interface. In addition, the m odel has been m ade

complete th rough an adjustm ent of the sintering m odel of Johnson [2.6] and the theory

of m e ta l/m eta l diffusion bonding.The fact that A1 is m uch softer than A120 3 greatly facilitates the m odel because it may be expected that all plastic deform ation occurs in the m etal (Fig. 2.3). Furtherm ore, the ceramic surface is considered to consist of a series of sharp, relatively sm ooth sided asperities. These asperities have an idealized shape of conic sym m etry w hich splits the bonding process into tw o stages, thereby sim plifying the m odel once again (Fig. 2.3). The first stage involves the indentation of a stiff conical punch, w hile the second stage com prises the collapse of a flat symmetrical pore.

19

The m odelling of stage I is analogous to the solution for the self indentation of spheres according to Johnson [2.6] (w ith som e variation in the local geom etry), w hereas stage II can be described in the sam e w ay as done for pore collapse in sintering [2.11]. The volum e transfer rates v x and v2 for bo th stages are given by:

47t5Dbfi

kT Peff(2 .2)

27t5Db£2

kTb 2 - r 2

( b 2 + r 2) l n ( b / r ) - ( b 2 - r 2)Peff

(2-3)

w here 8Db represents the interface diffusion coefficient (interface thickness times diffusion coefficient), Q the atomic volum e of Al, peff the effective bonding pressure, k

the Boltzm ann constant, T the absolute tem perature and b and r are defined in Fig. 2.3. The effective bonding pressure is in fact the norm al interfacial pressure plus the contribution of Joule-Thom pson curvature terms. If it is assum ed tha t the contact surfaces are flat and the equilibrium void shape is elliptical, then the following equations are valid:

Stage I: p eff =a 2v y

(2.4)

Stage II: p eff =1-b" (2.5)

w here p is the external bonding pressure, y the Al surface energy, h the void height and a, b and r are defined in Fig. 2.3.If it is further assum ed that the geom etry of Fig. 2.3 is preserved then the contact grow th rate d a /d t and the void shrinkage rate d r / d t m ay be expressed by:

Stage I: — =6 d t

Vjb"

( b 2 - a 2 ) 7 t a 2 ( h / b )

(2 .6 )

20

Stage II:drd t

-v2b 2 (2.7)( b 2- r 2)7tr2( h / b )

U sing this m odel, it w as found tha t the calculated fraction of bonded area as a function of process tim e is in reasonable agreem ent w ith the m easured bond strength as a function of process tim e in the case of joining A1/A120 3. In this respect, it m ust be rem arked that the m odel is satisfactory for relatively rud im entary system s only. It considers only non-reactive system s w ith one phase present at either side of the interface. M oreover, it is assum ed that only in one m aterial deform ation occurs. In the case of joining silicon carbide to a m etal the situation is far m ore com plicated because phenom ena like interdiffusion and com pound form ation as a result of a chemical reaction are show n to occur [2.12].

a i2o 3

iX '^ S ^ ^ d i f f o s io n a iX ^| T mass transfer

1 1 Al 1 1 1

f r y

contact

contact

Stage II

Fig. 2.3 M odel of m eta l/ceram ic (A1/A120 3) diffusion bonding, proceeding in tw o stages:I. indentation of a conical punchII. collapse of a sym m etrical void [2.4].

21

2.4 Diffusion processes and interphase boundary morphology

2.4.1 Introduction

In this section a m odel is presented w hich describes the interaction betw een a m etal and a ceramic. This theoretical approach considers diffusion for m ulticom ponent and m ultiphase system s w ith in a given geometry, form ation of new phases and evolution of phase and interface morphology.M ost m odels for solid state reactions deal w ith binary or quasibinary system s in w hich one-dim ensional transport processes take place. This is not a very realistic and pragm atic approach because m ost technologically im portant solid state reactions involve ternary and higher order tw o-phase systems. This evokes a few crucial differences betw een the

tw o approaches w hich are sum m arized below.First of all there exists a difference in the num ber of therm odynam ic degrees of freedom betw een binary and m ulticom ponent systems. According to the Gibbs phase rule:

F = k - p +2 (2-8)

w here F stands for the num ber of degrees of freedom, k for the num ber of independent com ponents and p for the num ber of coexisting phases, a binary closed system

consisting of tw o phases at given p and T has no degree of freedom left (F=0) and, therefore, is com pletely therm odynam ically defined. This im plies that an interface in a binary closed system is stationary given the assum ptions of constant pressure and tem perature during a diffusional equilibration process.The situation is quite different w hen ternary (k=3) or higher order system s (k>3) are involved. In such cases the interface com positions are partly controlled by kinetics and non-planar interfaces m ay develop.A second distinction betw een binary and m ulticom ponent system s is that a phase

diagram does not give inform ation about the probability of occurrence and nature of products w hich m ay be form ed during a diffusional anneal betw een tw o starting com positions of tw o different phases. This is due to the fact that the diffusion path in the phase fields of the initial phases is governed by the diffusion coefficients of the different species. As a consequence different interface com positions m ay resu lt and yield different reaction paths depending on the ratio of the corresponding diffusivities.A final dissim ilarity betw een binary and m ulticom ponent system s is that diffusion w ith in the phases in m ulticom ponent systems is influenced by concentration gradients of all diffusing species and is not only determ ined by its ow n concentration gradient as

is the case w ith binary systems. This results in cross term s in the diffusion fluxes and

22

m ay lead to up hill diffusion.In the next sections a general description of the problem will be given, follow ed by a sim plified m athem atical m odel w hich predicts w hether or no t a given interface m orphology rem ains stable. The m odel w hich is developed by Backhaus-Ricoult and Schm alzried [2.13, 2.14, 2.15], is an elaboration of existing studies on solidification and

transform ation.

2.4.2 General description of the diffusion problem

For reasons of sim plicity the problem will be confined to a one-dim ensinal set-up w ith tw o hom ogeneous starting materials. In addition, it is p resum ed that it is know n w hich

phases form and that they are restricted to a banded geometry. O ther structural constraints are the electroneutrality condition and the lattice site balance. Further it is

supposed tha t reaction is restrictively controlled by bulk diffusion and that local equilibrium is established everyw here. W hen these assum ptions are taken in to account, then the individual interdiffusion fluxes of all diffusing species can be w ritten as a function of the driv ing force w hich is proportional to the electrochemical potential gradients of all species:

j. = E b . V p . (2.9)j 'i n

w here j, represents the flux of all diffusing species i, the transport coefficient and Vp| the chemical potential gradient of species j. In this special case the O nsager relation is

valid, that is, L ^ I y There are certain conditions w hich have to be fulfilled by the various fluxes w ith in one phase, like the site conservation in the sublattices, the electroneutrality condition, the point defect equilibria betw een different species and therm odynam ic equilibrium w ith respect to the form ation of ceramic or interm etallic com pounds. Supplem entary assum ptions can be taken into account concerning the different com ponents such as ideal behaviour in some of the phases or im m obility of certain species. The characterized set of prerequisites taken together w ith the num ber offluxes of diffusing species yields in every phase for an n-com ponent system n-1

interdependent com ponent fluxes of the sort:

j. = - E D. .VC. (2.10)Ji i 'I i

23

in w hich are the chemical diffusion coefficients and VCj the concentration gradients

of the com ponents.

The following boundary conditions are indispensable in order to describe the diffusion problem properly:

at tim e t=0, bo th metal and ceramic have hom ogeneous initial com positions

C . ( t = 0 ) = C.° , i = A . . . , N - l (2-11)

the initial com positions are conserved at a distance far aw ay from the interfaces during the process time

C , ( t , $ -> ± ~ ) = C." (2-12)

special interface com positions are established at coordinates i;*k

C .(t,i;‘k) = C.*k(t) , (k = l , ... different interfaces) (213)

the chemical potential of the com ponents is continuous at the interface

p k( ^ k, t ) = p<kt,)( ^ k't ) <2-14)

the local continuity equation has to be valid at every interface, m eaning that the interface velocity is represented by the difference of the com ponent fluxes at the

interface in the tw o phases d iv ided by the difference in interface concentrations

t*k =j k( r k, t ) - j < k̂ - k,t)

c >k(k) _ c .k(k*l)(2.15)

The set of equations (2.11)-(2.15) is complete for an n-com ponent system of k+1 phases

(m arked by su p e rsc rip t<k)) and k interfaces ^‘k, w ith k=l,..., N -l.A n analytical solution of the equations is impossible unless they are greatly simplified. A few assum ptions w hich can be m ade, are:

the chemical diffusion coefficients are constant;the m olar volum e of the phases is constant and rem ains unaltered w hen

24

com positions change;the bu ild -up of defect gradients during the diffusion process m ay be ignored; the interface energy term s are neglected (in m ost cases they are dom inated by chemical effects).

In [2.14] a solution for the special system A-B-O is given w here A and B are m etals and A is in contact w ith BOy (Fig. 2.4). Further assum ptions are that A does no t dissolve in the oxide BOy and oxygen and the tw o m etals diffuse on different sublattices w hich m eans that the nondiagonal elem ents of the transport m atrix in the m etal phase can be ignored.O ne practical exam ple of the foregoing is the system Al20 3/N b in w hich N b represents A and Al is B. O xygen diffuses interstitially in the m etal phase and bo th m etals diffuse via vacancies. In addition, the diffusion coefficients can be assum ed to be constant because the solubility of O and Al in Nb is small. All these sim plifications taken together yields the following solution of the differential equation system:

5*2 = 4 k ' t^ p(2.16)

x / ( U ) = 1 +

1 + erf\

kD7

erf ̂

2 ,/D 7(2.17)

w here i= A l,0 and 1 stands for the m etal phase and 11 sym bolizes the ceram ic phase. W hen equations (2.16) and (2.17) are solved numerically, the parabolic rate constant kp‘ and the interface com positions xA1* and xQ* are obtained.It m ay be clear from the foregoing that it is possible to calculate a parabolic rate constant

and interface concentrations for a special system but that the application of this approach is im practicable [2.17].

25

0

( a )

BA

/ / / / / / / / / / /A/ phase a' ' / / / / / / / / / A

c

h

Fig. 2.4 M odel of m etal/ceram ic system [2.14].a. phase diagramb. diffusion couplec. concentration profile after annealing.

26

2.4.3 Interphase boundary morphology

In some instances p lanar interfaces in higher order system s rem ain p lanar du ring a diffusional anneal, in others they become unstable and develop into non-planar interfaces. These events occur even in system s w here no new phases are form ed. In order to predict w hether or not a p lanar interface becom es unstable and in w hich w ay this unstable interface evolves w ith time, a m athem atical m odel has been developed [2.13]. W ith this m odel w hich has been adapted to diffusion controlled reactions in two- phase m ulticom ponent systems, the time evolution of perturbations of sm all am plitude can be calculated. Perturbations of small am plitude can occur by transport fluctuations

caused by changes in tem perature or m aterial defects. Three possible situations m ay develop, depending on the so-called critical param eter 3> w hich will be described below. The initial p lanar interface m ay either rem ain p lanar or develop into a m orphologically

unstable interface if all possible perturbations shrink w ith tim e or if they increase in am plitude and the interface adopts a non-planar morphology. This is depicted schem atically in Fig. 2.5.

* * *

($) > o <£=o ($> < o

t = 0 tj > 0 tj> t 2 tw

Fig. 2.5 D ependence of a m orphological evolution of a d isturbed interface on the critical param eter <I> [2.13].

27

In order to acquire an expression for the time evolution of a small deviation 8 from the average interface coordinate, the diffusion problem for a non-planar interface has to be solved. In the rem aining p a rt of this section only the solution will be given of the problem and this solution will be discussed in a qualitative way.Provided tha t a num ber of boundary conditions and assum ptions have been defined, an analytical expression is obtained w hich show s that the non-planar interface no longer represents an isoconcentration line. This expression symbolizes the time evolution of

small perturbations 5 and can be w ritten as:

S(q,t) = 8 ( q , t = 0 ) . exp (2 | q | $ s /T ) (218)

in w hich q, the Fourier transform of the y-coordinate, em bodies the w ave vector of the

perturbation, t the tim e and <£ the critical parameter.The equation show s that 8 increases w ith time for <£>0 and decreases for <t><0 (Fig. 2.5).

The critical param eter <£ is a function of the starting com positions of the tw o phases, the m obility of the constituents and the therm odynam ic properties of the system.It is evident that the solution given here greatly simplifies the actual situation. However, w hen higher precision in the m odel is pursued the m athem atics become too com plicated to handle. Therefore, a relevant postulate for the prediction of interface m orphologies w ould be very helpful. U nfortunately, such a postulate is not available at present.A n exam ple of problem s encountered in predicting the interface m orphology is the system SiC-Fe of w hich the macroscopical diffusion path and local m icrostructure have

been determ ined. In this case silicon containing a-Fe phase w ith carbon precipitates is form ed. A detailed analysis of the product scale show s that the carbon precipitates consist of small lamellae w hich are random ly distributed [2.15]. Experim ents have show n that only iron diffuses th rough the p roduct scale [2.16] and the in terpretation of the reaction m echanism is, therefore, relatively simple. However, the special fibrous structure of the carbon precipitates cannot be explained by any m odel, so far.

28

2.5 Diffusion in the solid state

2.5.1 Introduction

In this section a sim ple system will be considered in the form of an inhom ogeneous single-phase alloy. If this system is annealed, m aterial w ill flow in such a w ay that the concentration grad ien t w ill gradually disappear, the net flow of m atter being zero. This

situation is described by Tick's first law and the equation in its sim plest form is:

J. = -D.

r \ dc.

d x(2.19)

w here J, is the flux of com ponent i, D, the diffusion coefficient (or diffusivity) of com ponent i, c, the concentration of com ponent i and x the distance coordinate.Tick's first law can no longer be applied if steady state conditions are not established, tha t is, if the concentration varies w ith both distance and time. In that case, the variation of the concentration w ith tim e (t) is given by Tick's second law:

dc.

"dFd

dx

dc.

d 7D. (2 .20 )

Turtherm ore, it is experim entally determ ined that diffusion is dependen t of tem perature according to an A rrhenius rate equation:

D , = D .o e x P

/-QRT

V /

( 2 .21 )

in w hich Di0 represents the pre-exponential or frequency factor, Q the activation energy for diffusion, R the gas constant and T the absolute tem perature. The frequency factor is a function of the crystal structure, the atomic vibration frequency and the activation entropy.

Diffusion occurs, roughly speaking, in tw o ways:

th rough the crystal lattice; this w ay of diffusion is called lattice diffusion, volum e diffusion or bulk diffusion;

along dislocations, grain boundaries, interfaces or free surfaces (high-diffusivity

29

paths); this diffusion is often referred to as short-circuit diffusion.

Volume diffusion can be subdivided into two different types:a substitutional or vacancy diffusion, w hich is the dom inant form of diffusion in

face-centered cubic (fee) m etals and alloys and is operative in m any body-centered cubic (bcc) and hexagonal close packed (hep) metals as well as in ionic com pounds

and oxides;b interstitial diffusion, w hich is thought to occur in both m etals and non-m etallic

solids.

It is obvious that atom s m igrate easier along grain boundaries, interfaces and surfaces than through a crystal lattice. W hen the different types of diffusion are com pared w ith each other, it appears that:

Ds > Db > Dv (2.22)

in w hich the subscripts s, b and v stand for surface, grain boundary and vacancy, respectively. In general, it is found that the contribution of grain boundary diffusion to the total diffusion coefficient becomes im portant below tem peratures of about (0.75- 0.8)Tm, w here Tm is the absolute m elting tem perature, com pared to lattice diffusion. This is related to the fact that in m ost cases the activation energy for grain boundary diffusion

is low er than that for lattice diffusion.The diffusion along dislocations can also be significant at lower tem peratures (about 0.5 Tm) com pared to lattice diffusion. This is, once again, due to the fact that the activation energy for diffusion along dislocations is lower than that for lattice diffusion and,

therefore, the lattice diffusivity decreases m uch faster w ith T than the diffusion along dislocations.

2.5.2 Multiphase diffusion in binary systems

W hen tw o solid m aterials are in intim ate contact w ith each other at a relatively high tem perature, diffusion of either one or both com ponents across the interface m ay occur.

W hen it is assum ed that volum e diffusion is rate lim iting throughout the entire annealing process, then the diffusion layer thickness d and the diffusion tim e t are

related th rough the parabolic grow th constant as follows:

30

d 2 = 2 k tp

(2.23)

in w hich lq, represents the parabolic grow th constant.U nder certain conditions this relation is not appropriate and some anom alies are

discussed below [2.17].Often a thin layer of different com position than the tw o starting m aterials is presen t on one or both of the surfaces to be joined (for example, an oxide layer). This layer m ay act

as a diffusion barrier thereby decelerating the actual diffusion process and grow th of the diffusion layer. This phenom enon is m anifested in a d2 vs t p lot by the fact that the straight line will not go through the origin bu t crosses the tim e axis at a positive value, indicating that incubation plays a role.It is also possible tha t the rate lim iting step is a reaction betw een elem ents of the starting m aterials resulting in the form ation of a new com pound. In this case the reaction layer starts to grow proportionally to the process tim e bu t after a certain period of tim e the diffusion process becomes rate lim iting. This im plies that in the second stage of the process a d2 versus t plot yields a straight line. In practice, reaction-lim ited grow th is not often encountered. A well know n exam ple of reaction-lim ited linear grow th is

encountered w hen porous layers are involved, resulting from reaction betw een a solid substrate and a su rrounding gas or liquid.The third situation in w hich a deviation from the linear relation betw een d 2 and t occurs, is w hen layer grow th in the beginning is faster than in a later stage. The cause for this phenom enon is that initially small grains are present w hich results in fast grain boundary diffusion, w hereas during the annealing process grain grow th occurs so that

gradually the (slower) volum e diffusion becomes the rate lim iting process. The effect of these phenom ena on the d 2 versus t plot is that, although bo th processes are diffusion controlled and therefore yield a straight line, a change in the slope can be observed indicating that one diffusion process transform s into another.

In m any cases it is difficult to decide w hether a d 2 versus t plot or a d versus Vt plot should be used. A ccording to Pieraggi [2.18], a d versus Vt plot is to be preferred if an initial period of faster kinetics does not contribute to the steady state control in the later parabolic stage (for exam ple, in an oxidation reaction). H owever, if an initial period of faster kinetics does contribute to the steady state control or if an incubation tim e is involved w ithout a considerable transitory layer thickness, only a d 2 versus t p lo t will p rovide the correct value for the parabolic grow th constant. In this respect, it m ust be

m entioned tha t an appropriate choice of a graphical presentation is also difficult due to large scatter in layer thickness assessments. It is therefore m andatory to also consider

m orphological aspects like porosity and grain size w hen evaluating experim ental results.

31

2.5.3 Multiphase diffusion in ternary systems

W hen m ultiphase diffusion in ternary systems is considered, an extra degree of freedom in com parison w ith binary system s m ust be taken into account, w hich complicates m odelling of such system s considerably. In contrast w ith binary phase diagram s, in ternary system s tw o-phase regions are allowed. This m eans that predicting a penetration profile for elem ents in a ternary couple is very hard because of the large num ber of

hypothetical possibilities. The relation betw een either diffusion phenom ena or com position of a diffusion zone and the ternary phase diagram becomes less apparent.

In order to determ ine exactly a diffusion path (which is a route through an isotherm al section of a ternary phase diagram , m arking the place of the average com positions in planes parallel to the original interface throughout the diffusion zone) therm odynam ic data are no longer sufficient. A dditional inform ation is necessary about the diffusivities of all elem ents and their interdependence. It is evident tha t m odelling m ultiphase diffusion in ternary system s is strongly obstructed by these dem ands. Therefore, one has to resort to experim ental m ethods to determ ine the diffusion path. The know ledge experim entally obtained can thus be used in the developm ent of a diffusion m echanism. As already referred to, diffusion data are indispensable w hen investigating the joining

of m etals to ceramics. In Table 2.1 diffusion data are given for several m etals of interest. Inform ation about diffusion coefficients of metals is hard to get and, m ore im portantly, the values of such data are som etim es rather divergent and diffusivity m easurem ents are not alw ays accurately described.The availability of diffusion data for silicon carbide is even worse. In Table 2.2 some

data are given for Si and C, b u t these are data for a tem perature region w hich lies far above the tem perature region w hich has been applied in this investigation. If one

calculates the diffusion coefficients for a tem perature of 1300 K (incorrectly m aking use

of D0 and Q values from Table 2.2, assum ing D0 and Q to be independen t of tem perature), values in the order of 10'27 m 2/ s and 10'3° m 2/ s for respectively DC[IXKI1| and DSi|(xxni are obtained, w hich shows the relative low m obility of Si and C in silicon carbide, com pared w ith that of metals in the same tem perature range. In addition , Table 2.2 show s that im purities like alum inium (often used in its oxidic form as sintering additive) have great influence on the diffusivities of both carbon and silicon.

32

Table 2.1 Diffusion data of metals.

m atrix diffusingelem ent(wt%)

D0

(m2/s )

Q

(kj/m ole)

tem peratureregion

(K)

refs.

Ni Si (0-<l) 1.5*10"4 258.2 1393-1573 [2.19]

Ni Cu 4*10"5 257.7 1038-1339 [2.19]C u vol.diff. 5.2N0'12 145.7 573-777 [2.27]Cu 6.1N0'5 255 1080-1613 [2.28]

Cu Fe 2.7*10'4 265.7 1173-1323 [2.19]Cu 6.1*1 O'4 267.8 1173-1323

Fe Fe 8.9*10"4 313.8 1173-1323Cu 3.6*10"4 274.0 1173-1323

Cu Si0 3.7*106 167.4 973-1073 [2.19]4 4.1NCT5 201.7 973-10738 1.9*10'3 225.1 973-1073a-sol.soln. 1.1*10"3 200.2 938-1048

Cu Ni 1.4*10'4 228.0 1038-1339 [2.19]Ni 2.7*10"4 236.4 1016-1349 [2.24]Ni 3.8*10"4 237.6 1016-1349 [2.24]Ni 1.9*10'4 232.8 1128-1328 [2.23]N i g.b.diff. 8.2*10'5 142.8 573-777 [2.27]Ni vol.diff. 2.6*10'10 133.2 573-777 [2.27]Ni 1.9*10'4 236.3 1053-1310 [2.29]

Fe Si 7*1 O'6 243 1273-1463 [2.20]Cu 4.2-10"4 306.2 [2.25]Ni 1.F10"4 296.8 [2.25]

Fe Si (2) 2.9*10'4 229.1 1133-1463 [2.20]Cu 5.7*10'5 238.6 [2.25]Ni 1.3*10‘4 234.5 [2.25]

AISI316 Fe vol.diff. 1.2*10"6 228.5 1178-1483 [2.21]Fe g.b.diff. 6.1*10"4 177.2 1178-1483

AISI316 Cr 6.3*10'6 243 [2.22]M n 4.1*10-5 260Mo 1.7N0'8 143 1178-1497

FeCrl8N i8 Fe 5.8*10"5 280.5 [2.22]Cr 8.0*10"6 244.5

33

FeC rl7N il2 Fe 3.6*10'5 279.2 [2.22]Cr 1.3*10'5 263.8

Cu Fe 1.0*10'4 213.3 993-1333 [2.23]Fe 1.4*10'4 217.7 733-1343 [2.23]Fe 1.3*10'4 215.6 1005-1297 [2.23]Fe 1.6*10+2 389.2 1130-1320 [2.24]Fe 9.1*10"6 193.0 923-1073 [2.26]Fe 5.0*10‘5 208.1 1073-1323 [2.26]

C uN il Cu 1.9*10'4 246.6 1053-1163 [2.23]

C uN il.08 Ni 9.5*10'5 233.0 1053-1310 [2.29]

CuNi2.81 Ni 2.3*10'5 225.4 1053-1310 [2.29]

CuNi21.5 Cu 1.9*10'4 231.5 1136-1385 [2.23]Ni 6.3*10"6 208.1 1203-1386

Fe' C 8.1*1 O'7 82.5 1183 [2.25]

Fe C 7.4*10"5 159.0 1183 [2.25]

34

Table 2.2 Diffusion data of SiC.

m atrix diffusingelem ent

D„

(m2/s )

Q

(kj/m ole)

tem peratureregion

(K)

refs.

SiC(600 ppm Al)

C[0001] 3*10"2 952 2126-2301 [2.30]

SiC(100 ppm N)

Cjcooi] 2*10+13 1266 2126-2301 [2.30]

a-SiC c ,00011 8*io+1 715 2128-2433 [2.31]

a-SiC Si[oooi) 5*10‘2 697 2283-2553 [2.31]

a-SiC(620 ppm N)

C[00011 3*10+3 791 2128-2433 [2.31]

a-SiC(620 ppm N)

Si[0001] 1*10+1 789 2283-2553 [2.31]

6-SiC C vol.diff. 2*10+4 841 2128-2374 [2.32]

fi-SiC C g.b.diff. 4*10+3 564 2128-2374 [2.32]

B-SiC Si vol.diff. 8*10+3 912 2283-2547 [2.33]

35

2.6 Thermodynamic aspects of M-Si-C systems

2.6.1 Introduction

Silicon carbide is know n for its reactive nature at elevated tem peratures w hen brought in intim ate contact w ith metals. Phenom ena like dissociation, form ation of silicides,

carbides and carbon are frequently encountered w hen SiC is w elded or brazed to metals. In this study attem pts have been m ade to join silicon carbide to austenitic stainless steel AISI316 in a d irect m anner and by m aking use of metallic interlayers (copper, nickel and binary copper-nickel alloys). As m entioned before, solid state reactions betw een the ceramic m aterial and the metal(s) will take place during diffusion w elding. Therefore, therm odynam ic aspects of a num ber of M-Si-C system s will be treated in this section, w ith em phasis on the Cu-Si-C, Fe-Si-C and Ni-Si-C systems. Therm odynam ic considerations may, in com bination w ith kinetic aspects, reveal the form ation of reaction products w hen SiC is w elded to metals.

In section 2.6.2, binary system s will be treated w hereas in section 2.6.3, as far as data are available, ternary system s will be discussed.

2.6.2 Binary systems

Si-COnly a few thorough studies have been perform ed on the Si-C system because of difficult experim ental requirem ents [2.35]. Unfortunately, the results of the various

studies are som ew hat conflicting. Therm odynam ic calculations show ed that the cubic form of silicon carbide (6-SiC) is more stable than the hexagonal form (a-SiC) at any

tem perature below the peritectic point (2545 ± 40°C).The hexagonal form of SiC is well know n for its num erous polytypes w hich are modifications of the a-SiC structure. The a-SiC structure is m ade up of tetrahedra oriented in successive hexagonal layers in which the lattice param eters a and c are variable (Fig. 2.6). The stacking sequence of these units can be very diverse and the resulting phases are know n as poly types. According to the Ram sdell-notation, the four

basic (and m ost common) polytypes are denoted as: 3C, 4H, 6H and 15R (3C = fi-SiC). The num bers are the num ber of layers in the unit cell and the letter suffixes represent the crystal sym m etry: C=cubic, H =hexagonal, R=rhom bohedral. The d istribution of

polytypes in a-SiC is strongly influenced by the presence and quantity of im purities. Especially the am ount of alum inium has a dram atic effect (Fig. 2.7).

36

Fig. 2.6

Fig. 2.7

SicSchematic representation of the a-SiC structure [2.34].

100

80

60

|4 0

20

0 . 0 0 0 . 0 5 0.10 0.15 0 . 2 0 0 . 2 5 0 .3 0

Al c o n t e n t ( %)

The influence of the am ount of Al on the distribution of polytypes of SiC [2.34].

37

Fe-SiIn Fig. 2.8 the phase diagram for the binary iron-silicon system is depicted. In the iron- rich corner of the diagram (which is the m ost relevant region for the presen t study), three bcc alloys are present, designated as B2 and DOS (ordered modifications) and A2 (disordered modification); these are referred to as Oj, a , and a respectively. The iron-

silicon com pounds w hich are stable, are (w ith increasing silicon am ount) Fe2Si(13),

Fe5Si3(Ti), FeSi(e) and FeSi2, the latter occurring in a high (h) and low (1) tem perature modification, ^ and Ce respectively [2.36].The boundaries of the y-loop are of interest w ith regard to the diffusion bonding process betw een AISI316 and SiC. D uring the bonding process silicon carbide decom poses into Si and C and as a result an Fe-Si solid solution is form ed, the crystal structure of w hich is dependen t on process tem perature and silicon content (the (a+y) region at the vertex extends from 3.19 to 3.8 at.% Si at a tem perature of about 1423 K).

W eight P e rc e n t S ilico n70 BO 90 100504 030

1700

1500

o°c

1300uo

90812001

(aFe)

700

500100

Si90705030 40

A to m ic P e rc e n t S ilico nFe

Fig. 2.8 Fe-Si phase d iagram [2.35].

38

Fe-C

The Fe-C system is one of the m ost exhaustively investigated system s because of its im portance for steelmaking. O nly a few aspects of this system will be m entioned here [2.43]. The (meta)stable equilibrium phases of the Fe-C system at am bient pressure which are relevant to this thesis, are austenite or y-Fe (fee) w hich can dissolve 9.06 at.% C at

1153°C, ferrite or a-Fe (bcc) w hich can only dissolve 0.096 at.% C at 740°C and hexagonal carbon (C) or graphite. The y /a-tran sitio n occurs in a tem perature range depending on the C content (at 912°C for pure iron). This transition is accom panied by

a volum e change. W hen the y-Fe contains 2.97 at.% carbon or more, a eutectoid reaction (at 740°C) can occur from w hich both a-Fe and C result.

Ni-SiIn Fig. 2.9 the phase diagram of the binary nickel-silicon system is presented. The solid solubility of Si in N i is maximal at 1143°C (15.7 at.% Si) and decreases w ith decreasing tem perature. O n the o ther hand, the solubility of Ni in Si is negligibly small. In the

nickel-rich corner, N i3Si is present in three m odifications depending on the tem perature. These are 6, w hich has a com position range of 22.8-24.5 at.% Si and is stable till 1035°C,

and fi2 and fi3 w hich have narrow com position ranges from 990°C till 1170°C. The y- phase also show s a narrow com position range covering the stoichiom etric com position of 27.9 at.% Si. The existence of the y-phase (Ni3]Si]2, also designated as N i5Si2), N i2Si, N i3Si2 and NiSi are reported in bulk diffusion couples [2.35]. The N i2Si-NiSi eutectic m elts at 964°C and is therefore know n as the low est m elting eutectic in the Ni-Si binary system.

O nly few experim ental values for AGf are available and, for that reason, values for the enthalpies of form ation of nickel silicides at 298 K are show n in Table 2.3. The discrepancy betw een AG, at process tem perature and AH,, values at room tem perature is not dram atic and lies w ith in the experim ental error (a relative error of 6% is

in troduced in the case of NiSi and NiSi2). The enthalpy of form ation of nickel silicides as a function of the atomic fraction of silicon is depicted in Fig. 2.10.

39

W e i g h t P e r c e n t S i l i c o n

1500

1400

1300 -

126!

U 12151200 (Si)0111

3

2 1100-

Etb 1 0 0 0 - E~

(Ni)Lroo'

m .

900

825£820®C800

70090 10030 40 50 60 70 8010 200

Ni A t o m i c P e r c e n t S i l i co n Si

Fig. 2.9 Ni-Si phase diagram [2.35].

Table 2.3 Enthalpy of form ation AH^ (k j/m ole atom) of nickel silicides at 298 K.

com pound Barin [2.37] Elliot [2.38] M iedem a [2.39] M urarka [2.40]

N i5Si <-31

Ni,Si -37.1 -36 -35.1

N i5Si2 -43.0 -42

Ni2Si -43.9 -48 -49.3

N i3Si2 -44.8 -46

NiSi -42.7 -42.8 -45 -42.8

NiSi2 -29.3 -29.1 -31 -29.1

40

A H

f ( k

j/m

ole.

ato

m )

A M i e d e m a ( e x p . )

• M i e d e m a ( c a l c . )

o C h a r t ( e x p . )

v B a h r i n ( e x p . )

□ M u r a r k a ( e x p . )

v El l i ot t ( e x p . )

0 .00 0 .2 5 0.50 0 .75 1.00

Si / Si + Ni

Fig. 2.10 Diagram show ing the enthalpy of form ation at 298 K for several nickel silicides as a function of the silicon fraction of the com pound [2.37-2.41].

41

Ni-CThe phase diagram of the Ni-C system is relatively simple, see Fig. 2.11. It consists of a eutectic w ith lim ited term inal solid solubility of C in fee N i (m axim um of 2.7 at.% C at 1318°C) and lim ited term inal solid solubility of N i in C (graphite). There are no stable

carbides in the Ni-C system b u t a m etastable carbide, N i3C, has been produced by splat quenching a m olten Ni-C alloy [2.35].

W eight P ercen t C arbon4020

4000

3650

3300

2950

Uoy 8600L3f l 2250 ua£ 1900Q)e-

(C, g ra p h ite )— *•

'zn1200

(Ni)850

5004 0 50

A tom ic P ercen t60

Carbon70 90 100

C30

Ni

Fig. 2.11 Ni-C phase diagram [2.35].

42

Cu-NiThe phase diagram of the Cu-N i system is considered as a m odel system because copper and nickel are com pletely miscible over the entire com position range (Fig. 2.12). Phase separation into a , and occurs at 354.5°C and 67.3 at.% Ni. However, in som e instances

segregation of bo th Cu and Ni have been observed. C opper segregation in Cu-N i alloys

is a com m on phenom enon, observed by m any investigators (for instance [2.44-2.46]). C onspicuous, however, is the oscillatory behaviour of the copper concentration in the second and deeper atom layers. Furtherm ore, copper segregation is basically lim ited to the surface layer and decreases rapidly w ith increasing depth. In the case of copper- nickel alloys w ith a nickel content of 16 at.% or less, nickel segregation to the surface has been observed for the first tim e by Sakurai and co-workers [2.47]. They found, by using a focusing type time-of-flight atom probe in an ultra-high vacuum of less than 1040 Torr, tha t nickel segregation to the surface is extended to three or four atomic layers and that

the nickel content in the first layer is approxim ately 40 at.% and gradually decreases to its bulk value. U p till now, they have not been able to present a satisfactory theoretical explanation for this phenom enon.

U f • i a 11 1 I ’ f ; r t \ , c • k r i

(Cu.Ni

at + a2

o 10 30 40 60 70 60 90 10 0

Cu A t o m i c P e r c e n t N i c k e l Ni

Fig. 2.12 Cu-Ni phase diagram [2.43].

43

Cu-CThe phase d iagram of the Cu-C system consists of the following relevant equilibrium phases [2.43]: an fee term inal solid solution of copper w ith a m axim um solubility of about 0.04 at.% C and a graphite-type term inal solid solution of carbon w ith no solubility of copper (Fig. 2.13). Furtherm ore, carbon is know n for its low diffusivity into solid copper [2.42],

W e i g h t P e r c e n t C a r b o n0 0350.01 0.015

1200

1600

•0704"

(Cu)

300

0.180 16 0.220 0 1 0.20 14

C n A t o m i c P e r c e n t C a r b o n

Fig. 2.13 Cu-C phase diagram [2.43].

44

Cu-SiIn total there are nine equilibrium phases of the Cu-Si system (Fig. 2.14) of w hich only a few will be m entioned here [2.35]. O n the Si side of the binary phase d iagram a term inal solid solution exists w ith negligible solubility of Cu w hich dissolves interstitially in Si (m axim um of 0.002 at.% Cu betw een 1200 and 1300°C). O n the Cu side

of the diagram a term inal solid solution exists w ith a m axim um solubility of 11.25 at.% Si at the peritectoid tem perature of 842°C. Furtherm ore, there are three interm ediate phases q , rf and r f w hich are a high-tem perature phase m elting congruently at 859°C, a rhom bohedral phase rem aining stable betw een 467 and 620°C and an orthorhom bic phase rem aining stable below 570°C, respectively. These phases are also frequently denoted as Cu3Si. Finally, there are tw o m ore interm ediate phases of im portance, namely, a cubic interm ediate y-phase (also denoted Cu5Si) decom posing peritectoidally a t 729°C and a hep interm ediate phase k (also denoted Cu7Si) w hich is stable betw een 552 and 842°C.

W eig h t P e r c e n t C o p p e r20 40 50 60 100

1500

1414°C14 0 0 -

1300

1200

3 1000

<0900

£^ 600

( C u )802°C

700 710

55fl°C

500

V " -400

0 10 20 30 40 50 60 70 A0 90 :oo

Si A to m ic P e r c e n t C o p p e r Cu

Fig. 2.14 Cu-Si phase diagram [2.35],

45

Cu-FeThe Cu-Fe phase diagram shows several equilibrium phases w hich are m eaningful w hen

studying diffusion bonding copper-nickel alloys to AISI316 (Fig. 2.15). The following solid solutions are present in the phase diagram: a fee Cu-rich solid solution (often

referred to as the e-phase) w ith a m axim um iron solubility of 3.5 at.% Fe at the peritectic

tem perature of 1096°C, an interm ediate-tem perature iron-rich fee solid solution (y-Fe) and a low tem perature iron-rich bcc solid solution, a-Fe, in w hich the m axim um solubility is 1.88 at.% Cu at the eutectoid tem perature of 850°C. Below the eutectoid tem perature of 850°C, large coherent clusters of iron have been observed in the fee Cu matrix. Furtherm ore, it w as found that these precipitates could readily be transform ed to the stable bcc form by plastic deform ation [2.43].

W e i g h t P e r c e n t C o p p e r90

1600 - H

1485°C

11.5

95^( C u ) -

e

12.7

10090

CllVv

Fig. 2.15 Cu-Fe phase diagram [2.43].

46

2 . 6.3 Ternary systems

Si-C-FeIn Fig. 2.16 the ternary phase diagram of the Si-C-Fe system at 1123 K is given. In this diagram it is dem onstrated that all relevant iron silicides are in therm odynam ic

equilibrium w ith silicon carbide, including the solid solution of silicon in iron (a) w ith a m axim um solubility of 25 at.% Si. Schiepers [2.12] has show n that in the system Fe-SiC, the diffusion pa th traverses the three-phase region (a + SiC + C), crosses the tw o-phase area (a + C) and ends in the single phase sector a.

C

SiC

Fen-Fe ' S i , c - F e S i Ca - F e S i

Fig. 2.16 Isotherm al section of the ternary Si-C-Fe system at 1123 K [2.12],

47

Si-C-NiThe ternary phase diagram of the system Si-C-Ni has been determ ined by Schiepers [2.12] at 1123 K. In this case it is im m ediately apparent that not all nickel silicides are in therm odynam ic equilibrium w ith silicon carbide (Fig. 2.17). The phases N i3Si2, NiSi and NiSi2 are in therm odynam ic equilibrium w ith SiC w hereas N i2Si is in therm odynam ic equilibrium w ith bo th SiC and C. The phases N i3Si and N i5Si2 are in

therm odynam ic equilibrium w ith carbon. Also in this case the diffusion p a th has been

determ ined w hen placing SiC and Ni into intim ate contact w ith each o ther at 1123 K (the in terrup ted line in Fig. 2.17). Going from the SiC end to the three-phase field (5-Ni2Si +

C + SiC), the diffusion path crosses alternately the 8-Ni2Si field an d the (5-Ni2Si + C) regions. Subsequently, another alternating layer sequence is encountered, nam ely the (y- N i5Si2 + C) field and the y-Ni;Si2 field. After this the (y-Ni5Si2 + 6-Ni3Si) zone is penetrated and, along tie-lines, the S-Ni3Si region is reached. Finally, the N i end is reached along tie-lines.

c

6+C+SiCSiC

Ni Sii S i . N iS i N iSi

T -N ic S i - 6 - N i . S i

Isotherm al section of the Si-C-Ni phase diagram at 1123 K [2.12].

48

Fe-Ni-SiThis ternary system has been critically surveyed by Raynor and Rivlin [2.48] w ho constituted the phase diagram by m eans of in terpretation of data from literature and calculations perform ed by themselves. It w as established that a ternary phase (t,) is present consisting of approxim ately 54 wt% Fe, 35 wt% Ni and 11 wt% Si at 600°C.

Schiepers determ ined the Fe-Ni-Si phase diagram experim entally at 850°C and, departing from the inform ation obtained by [2.48], a major part of the isotherm al section of the Fe- Ni-Si phase diagram w as devised [2.12], This diagram is given in Fig. 2.18. A reasonably good consensus exists betw een the diagram proposed by Schiepers an d the diagram s presented by Raynor and Rivlin in their review.

'ummmmrn

Fig. 2.18 Isotherm al section of a part of the Fe-Ni-Si ternary phase diagram at 850°C [2 .12],

49

Cu-Ni-SiIsotherm al sections of the Cu-Ni-Si system w ere determ ined at 500 and 700°C by Sokolovskaya and co-workers [2.49]. Some relevant features of their findings will briefly

be dealt w ith. They found a ternary Ni-Cu-Si solid solution (a) w ith an fee structure. The solubility of silicon in this solution w as greatest near the pure m etal corners and lowest (only 1-2 at.% Si) in alloys containing 60-80 at.% Cu (for alloys containing 90 at.% Cu the silicon solubility was found to be 5 and 6 at.% at 500 and 700°C, respectively). An isotherm al section of the Cu-Ni-Si phase diagram is given in Fig. 2.19. In addition, it was found tha t solid solutions based on N i5Si2 and Ni2Si form quasi-binary system s w ith the

copper-rich a-solid solution. C opper stabilizes the high-tem perature m odification of N i2Si

w hich transform s at 825°C in the low -tem perature m ode (see also Fig. 2.9).Schiepers also com posed a ternary phase diagram of the Cu-Ni-Si system at 775°C and his results are consistent w ith those of Sokolovskaya et al. [2.12].

Fig. 2.19 Isotherm al section of the ternary Cu-Ni-Si phase diagram at 500°C [2.49].

tu ZO *0 60 80 Cu

50

Cu-Ni-FeThe ternary Cu-Ni-Fe phase diagram has been studied by G upta and co-workers [2.50]. They m aintained the phase designation as given by H ansen [2.42] for the binary systems. The m ost im portan t phases occurring in this system are the nickel-rich fee y-phase, the iron-rich fee y^phase, the copper-rich fee y2-phase, the bcc a-phase occurring in bo th Fe-

Ni and Fe-Cu binary system s and the cubic y’-FeNi3. In Fig. 2.20 the m iscibility gap in this ternary system is given for several tem peratures. This diagram is in reasonable agreem ent w ith the findings of Palm er et al. [2.51], w ho determ ined the solubility lim it of iron (the y2-phase solid solubility) at different tem peratures in copper-nickel alloys

containing 5 and 10 wt% Ni, respectively (Fig. 2.21).

30,

60 <

40

301050°

.950°

50°80 ,800 °

600 "

N 40 50 60a t . p e t . Cu

80 90 Cu

Fig. 2.20 The miscibility gap in the Cu-Ni-Fe system for several tem peratures [2.50].

51

cu

M * CENT HICKCL

Fig. 2.21 The y2-phase solid solubility limits for several tem peratures at the copper-rich corner of the Cu-Ni-Fe system [2.51].

52

2.7 Metallic inserts

W hen diffusion bonding m etals to ceramics, it is often desirable (and som etim es inevitable) to m ake use of (metallic) interm ediates. It is evident that metallic interlayers should m eet certain requirem ents in order to be considered suitable for increasing the

m echanical strength and dim inishing the residual stresses of m etal/ceram ic joints.The m ost im portan t requirem ents to be m et by the interlayer m aterial, are:

the therm al expansion coefficient (a) should have a value w hich lies betw een thatof the m etal and that of the ceramic to evade h igh residual stresses;the yield stress (Rp 0 2) should be low, so that therm al stresses w hich develop duringcooling from process tem perature to room tem perature can be accom m odated bythe metallic insert th rough yielding;the Young's m odulus should be as low as possible;the reactivity of the interm ediate should be lim ited in order to prevent excessive

form ation of detrim ental brittle phases w hich often have physical properties not m atching those of bo th the metal and the ceramic.

It is obvious that not a single m aterial exists w hich m eets all these requirem ents and, therefore, a com prom ise has to be found. In the case of silicon carbide m any studies

have dealt w ith the solid state reactivity of SiC w ith m etals a n d /o r alloys under different process conditions. A sum m ary of these studies is presented in Table 2.4.Some rem arks need to be m ade w ith respect to the contents of this table.First of all, it is very im portan t to know the type and the chemical com position of the SiC w hich has been used. The significance of this is clearly illustrated in the case of joining SiC to Nb using either PS-SiC or RS-SiC (the PS type contains no free silicon, in

contrast w ith the RS variant).Less obvious b u t just as m eaningful as the inform ation of the presence of free silicon is the inform ation about w hich kind(s) of sintering additives or im purities are present in the silicon carbide. This is particularly significant w hen kinetic aspects are considered.

53

Table 2.4 Chemical behaviour of SiC/metal couples at elevated temperatures.

type of SiC

metal processtem perature(K)

reaction products T (K) below w hich no reaction is observed

refs.

HIP Ni 973-1323 N i2Si, C, N i5Si2, N i3Si 873 [2.12]

HIP Fe 973-1323 Fe3Si, C 973 [2.12]

a Ti 1073 TiC 948 [2.52]1173 TiC, Ti5Si31273 TiC, Ti5Si31473 TisSi3, TiSij, Ti3SiC2, TiC1-x1523 Ti3SiC2, TiC,„y, Ti5Si3, G-Ti [2.53]1673 TiSi2, Ti3SiC2 [2-52]

PS Cr 1173-1373 Cr3Si, C r7C3, C r^ Q [2-59]H P 1273 Cr3Si, C r7C3, C r23C6 [2.56]Si/SiC 1423 (air) silicides [2.56]

a C u 1273 Cu87Si17 (K-phase), C 1273 [2.52]

a Pd 873 Pd;Si 873 [2.52]973 P d5Si, P d3Si1073 P d2Si, P d3Si, P d5Si1173 P d2Si, Pd3Si> 1273 Pd2Si

? Mo 1873 M o2C, Mo3Si2 [2.57]6 1473 Mo2C, Mo5Si3, M o5Si3C [2.58]

pow der V 1173-1373 Si, VC, V2C, V5Si3 [2.59]

RS Nb 1473 N b5Si3 [2.60]1673 N b5Si3, NbSi2

PS 1473-1773 N b5Si3 [2.61]

a H f 973<T<1273 HfC 973 [2.52]1573 HfC, H f2Si, Hf3Si2, H f5Si31673 HfC, H f5Si3, HfSi, HfSi2

6 (HP) Pt 1173-1273 Pt3Si, P t2Si, C [2.54]1373 Pt3Si, P t2Si, C, PtSi, P t12Si5

PS Zr 1773 ZrSi, ZrC [2.55]

HIP = ho t isostatically pressedPS = pressureless sinteredH P = h o t pressedRS = reaction sintered

54

A t first sight, refractory metals seem to be very suitable as in term ediate m aterial for joining stainless steel to silicon carbide because they have a therm al expansion coefficient w hich closely approaches that of silicon carbide. However, other physical properties of these metals are less advantageous w hen joining metals to ceramics. Both the yield stress and the Young's m odulus are rather high com pared w ith other m etals. Furtherm ore,

their affinity to oxygen, especially at h igh tem peratures and in oxidizing surroundings, m akes them less appropriate candidates for joining stainless steel to silicon carbide. A nother interesting feature that em erges from this table, is that the reaction behaviour of the various S iC /m etal couples can be grouped into three categories, namely:

M + SiC —> silicide + C (2.24)

M + SiC —> silicide + carbide (2.25)

M + SiC —» [ Si ] + carbide (2.26)

The first category is represented by metals w hich form stable silicides b u t do not formcarbides. In some instances the m etal m ay dissolve significant quantities of carbon bu t

w hen the solubility lim it is exceeded, carbon precipitates in the form of graphite. The S iC /Fe, S iC /N i, S iC /C u and S iC /P d couples correspond to this first category.The second category com prises m etals w hich form both silicides and carbides. This is characteristic for m ost of the refractory metals. Therefore, the SiC /C r, SiC/Ti, SiC/Ta, SiC /W , S iC /H f and S iC /Z r systems fit in the second category.In the last category, the m etals react w ith SiC w hich leads to the form ation of prim arily silicon and carbides. The silicon rem ains either in solid solution in the m etal or becom es incorporated in the carbide structure (provided that ternary carbides are stable).

N evertheless, in some instances small quantities of silicides have been observed. The

m etals Al, Nb and V react w ith SiC producing silicon and carbide(s) w hich is characteristic for the th ird category.

55

2.8 Active metal brazing

As has already been m entioned in C hapter 1, the m ost successful brazing technique w hen non-oxidic ceramics are concerned, is the active m etal brazing process. In fact, the active m etal braze is an alloy w hich can react w ith the material(s) to be brazed. It makes these m aterial(s) wettable by altering its /th e ir surface com position so tha t w etting is governed by chemical factors rather than physical factors [2.68]. M ost of the applied braze alloys are based on gold, copper, silver or nickel w ith additions of titanium or zirconium as active element. The m ost thoroughly described and longest established active m etal braze is the Ag-28Cu eutectic w ith 5-10% titanium added. The specific

am ount of titanium w hich has to be added , appears to be very im portant. Several studies have dem onstrated tha t a m inim um am ount of titanium is necessary for inducing w etting because there is a threshold level of the titanium activity needed for the form ation of m etal-rich products. The addition of other elem ents like tin and indium is associated w ith this activity concept: it decreases the titanium solubility in the braze alloy and, hence, increases its activity [2.69].The generally accepted use of a Ag-Cu alloy as base for an active m etal braze is logical for several reasons. Titanium is soluble in a Ag-Cu alloy and it m aintains its activity and hence the ease of form ing a w ettable reaction product. Furtherm ore, the preferred com position (Ag-28Cu) is a eutectic w ith a m elting point of 780°C. The alloy is ductile, enabling preform s to be m anufactured and accom m odation of therm al m ism atch stresses.

Unfortunately, the addition of Ti affects the ductility detrim entally, just like the in troduction of In or Sn stiffens and hardens the solvent alloy, b u t it has been dem onstrated that these complex alloys still have acceptable properties.

Several reports have been published in literature concerning the brazing of SiC to metals and in all of these, titanium is used as the active elem ent in the brazing filler metal. In Table 2.5 the results are presented of brazing experim ents perform ed by various

investigators.Some rem arks concerning Table 2.5 need to be made.Firstly, it is apparen t that there is a difference in w etting behaviour of active braze filler

metals w ith respect to the type of silicon carbide. It w as show n that only a small am ount of titanium is needed to w et PS-SiC w hereas RS-SiC is properly w etted w hen a titanium content of m ore than 10 wt% is p resent in the Ag-Cu matrix. This can be explained by the fact that RS-SiC contains some 10% free Si and it appears that titanium readily reacts w ith silicon at elevated tem peratures form ing silicides w hich obstructs w etting.

56

Table 2.5 Results of brazing experiments of SiC/m etal couples.

SiCtype

m aterialbeingjoined

matrix com position of filler metal

Ticontent

(wt%)

mechanicaltestprocedure

joint strength at RT (MPa)

refs.

PS stainlesssteel(AISI316)

Ag-Cu 1.525

shear 65503

[2.62]

RS 18 45

RS RS Cu

Ni

435071.5

shear 388060

[2.63]

C u Cu 43 67

PS PS Ag-Cu 24.58

shear702520

[2.64]

PS PS N i 016.933.345.4 56.262.5

shear 14405097100129

[2.65]

PS = pressureless sinteredRS = reaction sintered (w ith free silicon)

Furtherm ore, it appeared from the results that there exists a m axim um w hen the shear strength of a brazed joint is plotted against the titanium content in the braze filler metal. Small am ounts of Ti do not properly w et the ceramic m aterial, w hereas increasing the am ount of Ti results in increasing reaction layer thickness thereby low ering the joint strength.

A nother aspect w hich needs elucidation is the dependence of the m echanical strength

on the process tim e and process tem perature. In the table, values of the strength are given w hich are representative for the brazed S iC /m etal joints. How ever, it appears that bo th process tim e and process tem perature have a profound influence on the m echanical strength w hich is m easured. It has been show n in several cases that there exists an op tim um com bination of process time and process tem perature w ith respect to the strength of b razed joints. This feature will be faced again w hen diffusion bonding experim ents are evaluated.

Finally, a less com m on braze filler metal should be m entioned: a gold based filler m etal

57

(82% Au, 18% Ni) w hich has proved to be especially suitable for bonding silicon infiltrated SiC to itself. The strength of the bonded specimens at room tem perature (four- po in t bend test) appeared to be over 80% of that of the base m aterial [2.66].

58

2.9 Mechanical testing

2.9.1 Test methods

W hen the m echanical strength of a m etal/ceram ic bond is to be determ ined, tw o different strategies m ay be followed, nam ely the conventional m echanical testing and tests based on fracture mechanics. Both approaches w ill be discussed briefly in this section.

C onventional mechanical testing The conventional m echanical testing m ethods are used to determ ine bond strength by m easuring the stress needed to fracture the bonded surfaces. The m ost com m on m ethods to m easure this stress are the bend test, the shear test and the tensile test. Each of these test m ethods has its ow n advantages and drawbacks and, therefore, none of these test m ethods is pre-em inently suited for the estim ation of the fracture stress of ceram ic/m etal diffusion bonds. Two serious deficiencies of the conventional testing techniques are [2.70]:

the bond strength values determ ined by these m ethods are specim en size dependent because the failure strength of a m etal/ceram ic joint is affected by stress concentrations at flaws present at the interface or in the adjacent ceramic; therefore, altering the d istribution of flaw sizes by changing the dim ensions of the bonded area will result in a change in the m easured strength;the bond strength can only be determ ined provided that it does no t exceed the strength of the ceramic m aterial; if this is the case, the crack w hich has developed

will follow a p a th aw ay from the m eta l/ceram ic interface into the bulk of the ceramic.

N evertheless, conventional testing m ethods are often applied because they offer an easy and relatively inexpensive w ay to specify the role of process param eters and bond geom etry on the m echanical strength of m etal/ceram ic joints.

Testing based on fracture m echanics In fracture m echanics tests the interface fracture energy G c (or critical energy release rate) is taken as a m easure of bond strength. The value of Gc represents the energy w hich is needed to separate a joint by a single crack propagating from a sharp interface

notch or precrack into the m eta l/ceram ic interface at the beginning of fracture [2.70]. The fracture energy is obtained from the failure load using conventional m echanical testing or specially developed m ethods based on fracture mechanics, like the double cantilever beam test. In contrast w ith conventionally m easured bond strength values, the results

5 9

of fracture m echanics tests are independent of the size of the interface area. However, the interface fracture energy is influenced by the m icrostructure of the interface region. A n additional benefit of determ ining Gc-values is that they can be applied to derive the interface fracture resistance Kc of the joint. This offers the possibility to directly compare the Kc-value w ith the fracture resistance KIc of m onolithic materials.A n inherent d isadvantage of testing based on fracture mechanics, is that a pre-defined

notch or crack (w ith accurately know n dep th and height) has to be introduced into an interface w hich is not an easy task to perform and is relatively expensive.

2.9.2 Strength variations

W henever a physical param eter is m easured in a series of tests w hich are perform ed under sim ilar conditions, a statistical spread in the results will be encountered.

Evidently, this is also the case w hen the mechanical strength of m etal/ceram ic joints is tested. Two aspects should be considered in explaining this scatter.In the first place, there are intrinsic errors in the test m ethod. In a m echanical property test, for instance, the sensitivity of the load m easuring device and the tolerances in

specim en dim ensions will result in scatter of the m easured data.In the second place, the reproducibility plays an im portant role in m easuring a property of a series of specimens.In a m echanical strength test perform ed on ceramics, deviations from the m ean strengthvalue of about 25% are frequently encountered, w hereas the instrum ental error in astrength test on a single specim en is com m only lower than 1%.A statistical function w hich is used often and is applicable in m any cases (including the classification of mechanical strength data), is the Weibull cum ulative d istribution function. The sim plest form of this approach is based on a w eakest link m odel (com parable w ith the breaking of a length of chain) and is described in a rudim entary

w ay by D avidge [2.67].Starting from the w eakest link m odel it is proposed that under a stress level <y the probability PS(V) that a volum e V of a ceramic which is x times larger than a unit volum e V0 (having a survival probability PS(V0)), will endure this im posed stress, is:

Ps(V) = P^(V(|)X (2.27)

The probability that rup tu re takes place, is then defined as:

P R = - ln P s(V) = - x ln P s(V„) (2.28)

60

and the probability of rup tu re for an infinitesimal volum e dV is:

d P R = f(a )d V (2.32)

because In PS(V0) only depends on the stress.In addition, it is assum ed that:

a - o (2.33)

in w hich a u, ct0 and m represent m aterial constants; a u is the location param eter orthreshold stress, im plying an upper lim it to the flaw size (it is generally justified to setthis value equal to zero); a„ is the scale param eter or characteristic strength and m is the shape param eter or Weibull m odulus w hich reflects the variability in strength (the h igher the value of m , the less variable the strength).W hen equation (2.33) is substitu ted in equation (2.28) the following expression isobtained:

P J V ) = exp

/ \C - G

-VG„

- I u ; -

(2.34)

This equation can be w ritten in a m ore convenient form by taking twice the logarithm , w hich yields:

l n l n ( l / P J = InV + m l n ( c - o u) - m ln o (| (2.35)

This expression represents a linear relationship betw een ln ln ( l /P J and ln a w ith slope

m. W hen a linear regression analysis is perform ed on the m easured data, the Ps-values are assigned o n basis of the position i of the m easured strength am ong N ordered re­values.

61

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63

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2.61 M. Naka, T. Saito, I. Okam oto, J.Mat.Sci., 26 (1991) 1983-1987.2.62 T. Iseki, H. M atsuzaki, J.K. Boadi, Am.Ceram.Soc.Bull., 64 (1985) 322-324.2.63 M. Naka, T. Tanaka, I. Okam oto, Trans.JWRI, 15 (1986) 49-54.

2.64 J.K. Boadi, T. Yano, T. Iseki, J.Mat.Sci., 22 (1987) 2431-2434.2.65 M. Naka, H. Taniguchi, I. Okam oto, Trans.JWRI, 19 (1990) 25-31.2.66 R. Reichel and B. Wielage, Pract.Met., 25 (1988) 74-81.

2.67 R.W. D avidge, M echanical behaviour of ceramics, C am bridge U niversity Press,

C am bridge 1979, p. 132-156.2.68 M.G. N icholas, Designing Interfaces for Technological Applications, S.D. Peteves

(ed.), Elsevier A pplied Science, London 1989, p.49-76.2.69 M.G. Nicholas, Joining of Ceramics, M.G. Nicholas (ed.), C hapm an and Hall,

L ondon 1990, p .73-93.2.70 G. Elssner, Joining of Ceramics, M.G. Nicholas (ed.), C hapm an and Hall, London

1990, p .128-154.

64

CHAPTER 3

Experimental procedure

3.1 M aterials

The ceramic m aterials w hich w ere used for the bonding experim ents, are hot-pressed silicon carbide (HPSiC), reaction-bonded silicon carbide (RBSiC) and hot-isostatically- pressed silicon carbide (HIPSiC). The m ain difference betw een these types of SiC is the

w ay they are m anufactured, resulting in a varying chemical com position and a differing m echanical behaviour [3.1].H ot pressing is the oldest m ethod to density silicon carbide pow der w hich consolidates and sinters the m aterial in one step. The pow der is m ixed w ith a sm all am ount of

sintering additive and is placed in a graphite m ould, heated to 1900-2000°C and uniaxially pressed to near theoretical density. Sintering aids are necessary because silicon carbide is a covalent com pound (ionicity of only 12%), covalent com pounds being know n for their reluctance to sinter at m oderate pressure and tem perature (diam ond

synthesis occurs under a pressure of 3500 MPa and at 2300°C). A dditionally, silicon carbide tends to decom pose rather than m elt at high tem peratures. The m ost com m on sintering aids are B, A1 or com pounds containing these elements.

Reaction-bonded silicon carbide (also denoted as reaction-sintered or silicon-infiltrated SiC) is a com posite of SiC and metallic silicon. A pre-densified pow der (generally called "green body") consisting of silicon carbide, carbon and a carbon containing b inder is infiltrated w ith liquid or vaporous silicon w hich reacts w ith the carbon. This reaction

causes sintering of the silicon carbide pow der w hich w as already present. In order to obtain a dense m aterial, a surplus of silicon is used so that the final p roduct contains free silicon varying betw een 5 and 30 vol.% of silicon.H ot isostatic pressing is a process in which pressure in all directions is applied on a pow der preform in a vacuum sealed casing using a highly pressurized (up to 200 MPa)

gas atm osphere inside a specially constructed pressure vessel. S intering occurs at high

tem perature w ithout m aking use of sintering additives because of the relatively high isostatic pressure w hich is applied com pared w ith norm al hot pressing. O nly im purities like oxygen in the form of a thin silica film are present, a lthough som etim es sintering

aids are added in order to produce parts w hich are less costly. The final density of the H IPped part is alm ost equal to the theoretical density.All three grades of silicon carbide w hich are used in the experim ents, have the a -

m odification w hich m eans that they are com posed of rhom bohedral or hexagonal units (see C hapter 2).

65

The m etal used in the experim ents w as austenitic stainless steel (AISI 316).The relevant chemical and physical properties of the silicon carbide and the stainless steel (determ ined w ith the aid of X-ray fluorescence spectroscopy and standard w et analytical chemical techniques) are given in Table 3.1 and Table 3.2, respectively.

Table 3.1 Chem ical com position of silicon carbide and austenitic stainless steel (AISI316).

Si C A1 Mg Fe Y P

HIPSiC 69.0 28.9 0.2

HPSiC 67.2 29.4 0.01 < 0.02 0.09 0.13 0.06

RBSiC 71.2 25.9 0.05 < 0.02 0.11 0.05 0.03

Cr Ni Mo M n C P S Si Fe

AISI316 16.1 10.2 2.2 1.08 0.08 0.08 0.02 0.49 bal.

Table 3.2 Physical properties of silicon carbide and austenitic stainless steel (AISI316).

HIPSiC HPSiC RBSiC AISI316

Density (g /cm 3) 3.21 3.10 3.10 7.90

Coefficient of therm al expansion (10'6 °C 1) (20 - 1000)°C

5 5.8 5.8 18

E-m odulus (GPa) 445 450 410 193

M axim um allowable tem perature in shielding gas (°C)

2200 2000 2000 1385

Average bending strength (MPa)

610 590 320

C om pressive strength (MPa)

2800 2700 700

Yield stress (MPa) 205

66

Both the silicon carbide and the steel have the sam e shape and size: disks w ith a

diam eter of 10 m m and a height of 5 mm. These disks w ere g round and polished w ith diam ond paste containing grains w ith an average diam eter of 1 pm (yielding an RA- value of = 0.05 pm ) to secure intim ate contact betw een them over the entire surface during the bonding process. After polishing the disks w ere cleaned ultrasonically in alcohol during 15 m inutes. In addition the silicon carbide specim ens w ere heated in air to about 500°C during 30 m inutes in order to rem ove hydrocarbon and beeswax rem nants (beeswax w as used in order to attach the ceramic disks to a glass plate; w ith this p late some 15 to 20 disks could be polished in one run).

The m aterials applied as interlayer, w ere nickel, copper and binary copper-nickel alloys w ith different am ounts of nickel added (nam ely 1, 3, 5, 10 and 15 wt%). The relevant chemical and physical properties of these m etals are listed in Table 3.3. The nickel w hich w as applied (designated as Nickel 200) contained 99.95 wt% Ni; the copper w hich was used in the experim ents, contained 99.93 wt% Cu.

Table 3.3 Chem ical and physical properties of the different interlayer m aterials.

Cu Ni

D ensity (g /cm 3) 8.94 8.90

Coefficient of therm al expansion (10"6 °C 1) (20 - 1000)°C

17.4 - 20.7 13.0 - 18.1

E-m odulus (GPa) 131 207

M elting poin t (°C) 1084 1453

Yield stress (MPa) 71 150

The nickel interlayers w ere punched ou t of a foil having a thickness of 0.2 m m by m aking use of a centreless punch w hich rendered disks w ith a diam eter o f 10 mm. After

punching the disks w ere heat treated in order to soften them so they could easily be flattened. A bout 60 disks were flattened at the same tim e under a pressure of 250 tons; the change in thickness by this treatm ent w as negligible. N ext they were etched in a 5% hydrochloric acid solution in order to rem ove the oxidic layer adhering to the surface of the metal. After this they w ere degreased ultrasonically during 15 m inutes. Finally, they w ere annealed in high vacuum at 600°C.

67

The copper and copper-nickel interlayers were punched out of foils having a thickness of 0.1 m m and 0.2 m m , respectively. After this they w ere degreased using the same

degreasing procedure as in the case of nickel.

3.2 D iffusion b o n d in g equ ipm en t

The diffusion bonding experim ents w ere carried out in tw o different furnaces: a high vacuum furnace and a gas shielded furnace. The tw o furnaces are illustrated in Fig. 3.1 and Fig. 3.2, respectively. H eating of the specimens and the application of m echanical pressure on the specim ens occurred in a similar w ay for bo th furnaces (Fig. 3.3).

MECHANICALPRESSURE

TURBOMOLECULARP U M P

TO HF GENERATOR *

THERMOCOUPLE

AIS1316

Water cooled

A12 ° 3 Molybdenum

Interlayer

□ SiC

• Viton seal

O Cu seal

Fig. 3.1 Schematic view of high vacuum furnace.

68

I B Thermocouple

L \N A190 ^ tube

I :-i I Insulating cap

Graphite heating element

EUi Fumace body

■ffl AISI316

I I SiC

EH Ceramic support

1200 10

oCL

5 ^

Q.

0

t (m in )

p r o c e s s m e c h a n i c a lt e m p e r a t u r e p r e s s u r e

Fig. 3.3 Process tem perature and mechanical pressure as a function of process time for bo th vacuum and shielding gas conditions.

In the high vacuum furnace, the m aterial com bination to be w elded w as heated to the desired tem perature w ith the aid of a high-frequency generator (Flender H im m el, M odel HG 8000). H eating w as controlled by a therm ostat (Eurotherm , M odel 818 P). The specim ens to be bonded w ere heated via a susceptor m ade of tantalum (enclosing the

specim ens) w hich was surrounded by a tubular coil of copper (com prising 4 w indings) connected to the HF generator. Both the susceptor and the coil w ere su rrounded by a rad iation shield of m olybdenum .H eating occurred at a rate of 25°C /m in. Before the heating was started, a vacuum of about 4xlO"4Pa w as created; this w as achieved by a rotary vane vacuum p um p and a turbo-m olecular pum p. The total pressure in the system w as controlled by a Pirani m easuring device and a cold cathode device. The complete p u m p in g /m easu rin g group w as controlled by an electronic drive u n it (Balzers, M odel TCP 300).The tem perature in the process cham ber w as m easured w ith the aid of therm ocouples of the type tungsten 5% rh en iu m /tu n g sten 26% rhenium (Thermo Electric). After the

process tem perature had reached the desired value, mechanical pressure w as applied on

( J

1000

800

600

4 0 0

200

0

70

the specim ens by m eans of a pneum atic device; w ith this device m echanical pressures

could be exerted ranging from 0 to 10 MPa and from 10 to 30 M Pa (taking into account a specim en diam eter of 10 mm), depending on w hich pneum atic cylinder w as m ounted. The pressure on the specimens could be controlled by m eans of a m anual loading regulator (ITT). The m echanical pressure w as released as soon as the cooling started. Cooling occurred initially at a rate of 5 °C /m in and continued at this rate till a tem perature of about 150°C. After this, furnace cooling took place till 80°C after w hich both vacuum pum ps w ere sw itched off and dry nitrogen w as let in.D uring the w elding experim ents the tem perature and total gas pressure inside the vacuum chest w ere registered by a 1-6 channel recorder.

The gas shielded furnace (gas composition: 95 vol.% Ar, 5 vol.% H 2) w as heated by a graphite elem ent of helical shape. The m antle was filled w ith isolating m aterial in order to m inim ize the heat loss. W ithin the graphite elem ent, an alum ina tube (type: Alsint) w as placed through w hich the shielding gas was flowing. Inside this tube, another alum ina tube w as placed w hich was attached to a specim en-holder (also m ade of alum ina). The pressure on the specim ens was exerted by m eans of a lever construction outside the furnace. The tem perature m easurem ent/regu la tion w as carried ou t by m eans of p latina/platina-13% rhodium respectively ch rom el/alum el therm ocouples. These therm ocouples w ere connected to a proportional pow er regulator so that heating and

cooling could be perform ed in a controlled way. Cooling to room tem perature took place as follows: the specim ens were furnace cooled linearly from the process tem perature to approxim ately 400°C, then the device w as sw itched off, the specim ens w ere allow ed to cool to 150°C and, subsequently, were taken out of the holder and w ere cooled in air to room tem perature.

3.3 M echanical testing

The m echanical strength of the specim ens w hich w ere successfully joined, was determ ined by m eans of shear stress testing. This mechanical testing technique is rather unconventional w hen diffusion bonded m etal/ceram ic joints are concerned. The shear stress testing m ethod w as chosen for the following reasons:

the specimens need no special treatm ent, that is, there is no specific prescrip tion of specim en dim ensions;

the apparatus w ith w hich the tests have been executed, is quite easy to handle and, therefore, m any sam ples can be tested in a reasonable period of time.

The tests were perform ed in order to assess qualita tively the strength of the bonded

71

sam ples; since no internationally standardized mechanical testing m ethod for

m etal/ceram ic joints has been developed and accepted, it is justified to apply the shear stress test for the above m entioned reasons.Furtherm ore, statistical m eans have been developed (Chapter 2) to further evaluate the

strength data in a correct way.

F or ce

\ A l i g n i n g 7 * c l a m p s

Knife

S p ec im en

Force

Knife

S u p p o r t i n g r o d

Metal

C e ra m i c

Fig. 3.4 Schematic view of shear strength apparatus.a. Front viewb. Side view.

The shear stress apparatus used is schematically given in Fig. 3.4. The ceramic part of

the joint is clam ped in a holder w hich itself is m ounted in a rotary bearing. The metal p a rt of the joint is loaded by the knife which is shaped sim ilarly as the joint surface. In order to prevent the in troducton of bending com ponents, the joint is sustained

through a transpositionable rod w hich is placed against the back of the m etal part

during testing. To test the shear strength of the bonded specim ens, the testing device is p laced in an Instron-like m achine (lira Test, M odel 2300) in such a w ay tha t a pressing

force could directly be applied on the knife of the shear-test device.

72

3.4 Analytical methods

The specim ens were analyzed w ith the aid of optical microscopy, X-ray diffraction (XRD), electron probe micro analysis (EPMA) and A uger electron spectroscopy (AES). In all cases optical m icroscopy w as executed first in order to get inform ation about the

degree of reaction betw een the different m aterials and, in addition , to assess the specim ens qualitatively (and sem i-quantitatively) concerning their m icrostructure and com position. This starting point m akes it possible to select an appropria te and supplem entary analytical technique.

3.4.1 Optical microscopy

O ptical m icroscopy w as applied to reveal the effects in the metallic parts of the specim ens after the bonding experim ents. In order to perform this analysis, some preparations had to be m ade. First the metallic part w as saw n perpendicularly to the interface, em bedded in diallylphtalate therm osetting resin w ith short glass fibres (M ounting Resin-4, Struers), ground on SiC em ery paper and polished w ith d iam ond

paste till 1 pm diam eter and ultim ately etched w ith an suitable etchant (Table 3.4). The specim ens, either etched or unetched, w ere analyzed using a Leitz N eophot 2 m etallographic microscope, affording magnifications up to 1000 tim es (2000 tim es w hen using oil immersion).

Table 3.4 Etchants used for optical m icroscopy analysis.

Etching agent Com position M aterial O bservation

Railing 5 g CuCl2 in 100 ml ethyl alcohol and

100 ml HC1

y-stainlesssteel

grain boundaries carbides

A m m onium 25% N H 4O H Cu grain boundarieshydroxide (+ Perhydrol)

(+ 30% H 2Oz) Cu-xwt% Ni

silicides

3.4.2 X-ray diffraction (XRD)

The application of XRD w as needed in order to elucidate the identity of the reaction

products w hich w ere form ed during the w elding process. Two different techniques were

73

used in this respect.O ne m ethod w as applied in order to analyze hard and brittle reaction products form ed in the vicinity of the interface during the w elding process w hich broke out during cooling to room tem perature. These brittle reaction products w ere c ru sh ed /g ro u n d w ith the aid of a m ortar until a fine pow der rem ained. This pow der w as analyzed by means of a Debije-Scherrer camera (Table 3.5) equipped w ith a Co anode operating at 45 kV and 30 m A under an angle of 4° and exposure time of 1.5 hours or at 50 kV and 40 mA

under an angle of 6° and exposure tim e of 1 hour.

Table 3.5 A djustm ent data used in perform ing XRD analysis.

Cam era Anode(kV-mA)

20 range (°) (step size °)

Angle (°) Exposure time (m inutes)

Debije- Co (45-30) 4 90Scherrer 6 60

Guinier- Cu (45-30) 25-85 (0.1) 45-90De Wolff

Guinier Co (45-25) 30-110 (0.1) 60Cu (45-30) 22-60 (0.1) 43

The other technique concerned the analysis of fracture surfaces of m etal/ceram ic joints. In this respect, only bonds fractured along the interface w ere interesting to be examined.

For this purpose a G uinier - De Wolff camera was used (Table 3.5) equipped w ith a Cu anode operating at 45 kV and 30 mA; the 2 0 range from 25-85° was scanned in steps of

0.1° or a G uinier cam era (Table 3.5) equipped w ith a Co anode operating at 45 kV and 25 m A or equipped w ith a Cu anode operating at 45 kV and 30 mA; the 2.0 range from 22-66° w as scanned in steps of 0.1°. For the actual analysis the specim ens w ere placed in a ro tating holder on a Siemens D500 goniometer.

3.4.3 Electron probe micro analysis (EPMA)

In order to analyze the sam ples w ith the aid of an electron probe, both fractured and non-fractured sam ples w ere cut perpendicularly w ith regard to the interface betw een the ceramic and the steel (a cross section of the diffusion bonded couples w as obtained in this way). In some cases, only the steel w as exam ined. In order to be able to analyze the

edge of the steel specim en, a th in layer of nickel was deposited on the steel. Subsequently, it w as em bedded in a conductive acrylic therm oplastic resin w ith iron pow der (M ounting Resin-1, Struers). The diffusion bonded couple w as then lapped w ith

74

a diam ond slurry (9 urn) and, subsequently, polished to 1 pm . After this the specim ens w ere cleaned ultrasonically in alcohol during 15 m inutes.

The com position of the reaction layers w hich were form ed during diffusion bonding, w ere analyzed using a Jeol JXA 733 electron probe X-ray m icroanalyser, equ ipped w ith

four w avelength-dispersive spectrom eters, an energy-dispersive detector and a fully com puterized analysis system (Tracor N orthern TN 5500 and TN 5600). F irst of all, back scattered electron im ages w ere recorded to reveal the different phases. The com position of each individual phase w as quantitatively determ ined w ith a focused electron beam

(12 keV, 35 nA). The intensities of CKaz SiKa/ FeK,„ C rKa/ N iKa, M nKa, M o ^ X-rays

m easured on the specim ens w ere com pared w ith those m easured on standards of the

pure elem ents and silicon carbide. From the ratio of the X-ray intensities of the specim ens and the standards, the concentration of each elem ent w as calculated after having corrected for atomic num ber, absorption and fluorescence according to the so- called p(pz) approach [3.2].

3.4.4 Auger electron spectroscopy (AES)

A PHI 4300 Scanning A uger M icroprobe (Perkin Elmer), equipped w ith a cylindrical m irror analyser (CMA) and electrostatic gun optics, w as em ployed for the determ ination of:

(i) com position at selected locations (point analyses);(ii) com position of dep th profiles.

The analyses of selected locations w ere perform ed at cross sections p repared of the copper-nickel layer betw een the austenitic stainless steel and the silicon carbide. The

com position dep th profiles w ere determ ined by ion-sputtering perpendicu lar to both fracture surfaces of sam ples sectioned at the interface betw een the copper-nickel layer and the silicon carbide.

The m easurem ents at selected locations were perform ed w ith a 10 keV, 20 nA prim ary electron beam at 30° to the sam ple surface norm al, corresponding w ith a beam diam eter of about 200 nm . D uring these m easurem ents im age registration w as em ployed to correct for instrum ental drift [3.3]. Prior to these m easurem ents surface contam ination was rem oved by Ar+ ion-sputtering.The m easurem ents at selected locations comprise the acquisition of five A uger spectral regions: C (254-287 eV), O (492-520 eV), Si (80-100 eV), Cu (905-927 eV) and Ni (705-725 eV). All regions w ere recorded at 0.5 eV intervals and 0.6% AE/E analyser energy resolution. The com position w as determ ined from A uger peak to peak height of the

75

A uger spectra in the differential d istribution using elem ental sensitivity factors. A lternate sputtering and data acquisition w ere used to collect the dep th profiles. A 3 keV, 2 pA prim ary electron beam w as used at 30° to the sam ple surface norm al. A 3.5 keV A r+ beam was used for sputtering. The data acquisition during each cycle comprises the recording of the sam e A uger spectral regions as for the po in t analyses (see above). However, the A uger spectral region of carbon was extended to 220-290 eV and w as

recorded w ith 1.0 eV intervals in the analyses of the silicon carbide. The com position w as determ ined from A uger peak to peak height of the Auger spectra in the differential d istribu tion using elem ental sensitivity factors.

76

References

3.1 K.A. Schwetz, Radex-Rundschau, 1 (1989) 26-39.3.2 G.F. Bastin, H.J.M. Heijligers, F.J.J. van Loo, Scanning, 8 (1986) 45-67.

3.3 R.R. Olsen, Perkin Elmer: Technical Bulletin, 8903 (1989).

77

CHAPTER 4

Direct -bonding

4.1 In troduction

In this chapter the results are presented of attem pts w hich w ere m ade to join silicon carbide to austenitic stainless steel AISI316 in a direct way, that is, w ithout m aking use of an interlayer.The description will cover bo th high vacuum experim ents and shielding gas experim ents using tw o types of silicon carbide, hot pressed silicon carbide and reaction bonded silicon carbide. Em phasis will be given to bonding experim ents perform ed in a vacuum environm ent. Bonding experim ents carried ou t in a shielding gas yielded results w hich are quite sim ilar to those obtained in vacuum . The joining of AISI316 w ith either of the

tw o types of silicon carbide is treated separately because it appeared that the chemical behaviour of bo th types of silicon carbide in contact w ith the steel differs significantly. The AISI316/SiC couples were analyzed by m eans of three techniques: optical microscopy, electron probe microanalysis and X-ray diffraction.

In the next section the structure of the AISI316/SiC sam ples is discussed on the basis of the results obtained w ith these techniques.After a detailed discussion of the results of the h igh vacuum bonding experim ents, a relatively brief survey is presented concerning the results of bonding experim ents carried

ou t in a shielding gas. This is follow ed by conclusions in w hich the results are sum m arized and an outlook on C hapter 5 is given.

4.2 H igh vacuum b o n d in g experim ents

In order to establish the conditions under which silicon carbide can be joined to austenitic stainless steel AISI316, a param eter study was perform ed using the diffusion

bonding equipm ent described in Chapter 3 (Fig. 3.1). The m ost im portant param eters of the diffusion bonding process are process time, process tem perature, mechanical pressure and furnace atm osphere (assum ing standard geom etry and standard surface

preparation; see C hapter 3).

The values of the process tim e were chosen betw een 22.5 m inutes and 24 hours. The values of the process tim e were chosen to differ by a factor 4, so that the resulting thickness of a reaction /d iffusion zone w ould differ by a factor 2, p rovided that diffusion is the rate determ ining process.

78

The values of the process tem perature were changed in steps of 50°C betw een 850°C and

1250°C.A m echanical pressure on the specim ens during diffusion w eld ing is necessary in order to establish good contact over the com plete surface area and to obtain a joint w ith in a w orkable period of tim e (24 hours). However, the m agnitude of the m axim um mechanical pressure on the specim ens during diffusion w elding of m eta l/ceram ic

couples is restricted by the plastic deform ation of the metallic com ponent. The values of the mechanical pressure used in the bonding experim ents w ere betw een 0.1 M Pa and 7.5 MPa, respectively.The furnace atm osphere w as a high vacuum of about 10'3 Pa.

4.2.1 HPSiC/AISI316

Two series of w elding experim ents were carried out: one series at a relatively low level of mechanical pressure (0.1 MPa) and one series at a relatively h igh level of m echanical pressure (7.5 MPa). The results of both series d id not differ significantly, indicating that a mechanical pressure of 0.1 MPa is sufficient to establish contact over the com plete surface area.It appears that hot pressed silicon carbide can not be bonded perm anently to stainless steel AISI316. In m ost cases, a bond is accom plished during w elding b u t this bond fractures w hen cooling to room tem perature.Fracture occurred either along the interface or in the ceramic m aterial in the vicinity of the interface. In som e cases (T > 1150°C) a liquid phase w as form ed during the w elding process.

Exam ination of the diffusion w elded sam ples was carried ou t first by m eans of optical m icroscopy following the procedure described in C hapter 3. This exam ination w as difficult because of the vigorous reaction w hich often had taken place. A typical m icrograph of the steel part of a fractured diffusion bonded HPSiC/AISI316 couple is given in Fig. 4.1. In this m icrograph only a small piece of the silicon carbide is visible. O n the basis of the optical m icroscopy results and the results of the analyses, w hich will be dealt w ith further on, a general picture of the joint em erges w hich is show n schematically in Fig. 4.2. It appears that reactions both on the m etal side and the ceramic side occur. The corresponding reaction zones are denoted a, b, c, d, respectively. These zones are treated separately in the next parts of this section.

79

Fig. 4.1 Optical m icrograph of the steel part of a fractured H PSiC /A ISI316 couple diffusion bonded at 1100°C during 6 hours under a mechanical pressure of 5 MPa (etchant: Kalling).

HPSiC

zone a

f / s _ W s s ^ \ s s s r s f S f S / f 7 S S S f S S i 0

AISI316

zone b

zone c

zone d

Fig. 4.2 Schematic illustration of the reaction zones in diffusion bonded HPSiC/ AISI316 couples.

80

Zone aIn zone a, a layered structure is observed. This structure w as exam ined w ith the aid of

EPMA and w as found to consist of alternating layers of a silicon rich phase, presum ably (Fe,Ni)2Si (w ith a small am ount of chrom ium ) + C; the total w id th of this zone is dependen t bo th on process time and process tem perature, a typical value of the w id th being 65 pm (1050°C, 24 hours). The occurrence of banded reaction zones consisting of periodic layers has been reported earlier by K irkaldy and Young [4.1], M ehan [4.2], Osinski [4.3] and Schiepers [4.4]. A plausible explanation for this phenom enon has been given by Van Loo et al. [4.5]. In this m odel it is argued tha t the sequence of reaction

layers in a S iC /m etal couple is: M / MxSiy / C / SiC (M = m etal), tha t is, a carbon free layer is form ed next to the metal. The periodicity of these reaction layers is no t fully

understood at p resent b u t is believed to be related to mechanical stresses built up during the various reactions.

Zone bZone b is situated along the original interface AISI31 6 /HPSiC. Fracture of the joint often occurs th rough this zone. This m akes analysis of this zone by m eans of EPMA difficult. In order to cope w ith this difficulty and to facilitate EPMA analysis, the tw o fractured parts w ere clam ped together w ith an alum inium foil in between. From the EPMA results obtained it appears that also in this zone a layered structure is present. These layers w hich ru n m ore or less parallel to the original interface, consist of m any voids a n d /o r carbon islands em bedded in a metallic m atrix (Fig. 4.3). The com position of the m atrix of this reaction zone could not exactly be determ ined w ith the aid of EPMA because of the presence of pores. However, the m atrix seems to consist of m ainly iron and silicon

and sm all am ounts of chrom ium and nickel.The periodicity of the layers in this zone m ay be explained in the sam e w ay as w as done

for zone a.W hen the carbon precipitates are looked at in m ore detail, it appears that they consist

of sm all lamellae w hich are m ore (in the presence of nickel) or less (in the presence of

iron) concentrated in layers. The carbon precipitates them selves show a fibrous structure

w hich also has been observed by Backhaus-Ricoult [4.6].

81

Fig. 4.3 Back-scattered electron im age of zone b (upper part: HPSiC, low er part: A1 foil).

It appears that fracture occurs either through this zone or through the adjacent ceramic

m aterial (zone a) because of the presence of hard and brittle reaction products and the developm ent of therm al residual stresses w hich cannot be accom m odated by the m aterials. W hen fracture occurs th rough zone b, an irregular fracture surface is p roduced (Fig. 4.4), w hile a sm ooth fracture surface results w hen fracture occurs th rough the silicon carbide (Fig. 4.5).

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I 12 m m

Fig. 4.4 Typical exam ple of fracture th rough zone b (HPSiC/AISI316, 1100°C, 24 hours, 0.1 M Pa, high vacuum ; dark field).

I 12 m m

Fig. 4.5 Typical exam ple of fracture th rough zone a (HPSiC/AISI316, 1050°C, 24 hours, 0.1 MPa, h igh vacuum ; dark field).

83

Zone cZone c is situated in the steel and consists of two phases w hich are show n in Fig. 4.6. The m atrix (the first phase) is dark in the back-scattered electron im age w hile the grain boundaries (the second phase) are light in the back-scattered electron image.

g j p i

Fig. 4.6 Back-scattered electron image of zone c (upper part) and part of zone d (lower part).

The com position of the tw o phases w as m easured by m eans of EPMA and is given in Table 4.1. From this table it is clear that the tw o phases differ significantly in composition. The m atrix is characterized by a rather h igh silicon content (6-7 wt%) and

a rather low nickel content (about 7 wt%) com pared to the standard silicon and nickel contents of AISI316 (0.4 wt% and 10.9 wt%, respectively). The phase along the grain boundaries in zone c has a silicon content (about 10 wt%) which is even higher than that in the m atrix, w hile the iron content (about 56 wt%) is m uch lower and the chrom ium

content (about 20 wt%) is m uch higher than the standard iron and chrom ium contents of AISI316 (68.8 wt% and 18.0 wt%, respectively). The higher silicon and low er nickel contents of the m atrix are accom panied by a low level of the m olybdenum content.

I 150 pm

84

Table 4.1 Chemical composition in wt% of zone c (T=1100°C, t=90 minutes).

Region Fe Cr Ni Mo Si C

bulk HPSiC 0.1 - - - 69.4 30.3

m atrix of zone c (dark) 70.2 17.3 6.8 0.1 6.5 ’)69.6 16.8 7.1 0.1 6.4 ')

grain boundaries in 55.5 20.7 12.4 0.2 9.9 ’)zone c (white) 56.4 20.2 12.1 0.1 9.9 ’)

56.6 20.1 12.3 0.1 10.0 ’)56.6 20.7 12.4 0.1 9.8 ’)

bu lk AISI316 68.8 18.0 10.9 1.9 0.4 0.1

’) carbon concentration (wt%) not changed significantly w ith regard to the carbon concentration in original AISI316

From these results it m ay be concluded that the m atrix of zone c is AISI316 w hich is depleted of nickel and m olybdenum and enriched w ith silicon.The identity of the grain boundary phase could not be revealed by m eans of XRD, bu t

from the ratio of the atomic percentages of the m ain elem ents in this phase (iron,

chrom ium , silicon and nickel), it seems that (Fe5Cr2Ni)Si2 silicide [4.7] has been form ed.

Zone dZone d lies next to zone c and is light in the back-scattered electron im age (see Fig. 4.6). The com position of this zone is com parable w ith that of unaffected AISI316; only the iron content is higher and the chrom ium content is low er than that of standard AISI316. H ow ever, the com position near the grain boundaries in this zone is distinctly deviating from standard AISI316: they are strongly enriched w ith both chrom ium and carbon. The

com position of bo th the m atrix and the grain boundaries is given in Table 4.2.

85

Table 4.2 Chemical composition in wt% of zone d (T=1100°C, t=90 minutes).

Region Fe Cr Ni Mo Si C

m atrix of zone d 74.4 14.5 9.5 - 0.5 ’)(white) 74.7 14.2 9.7 - 0.5 ’)

grain boundaries in 32.9 59.4 2.0 0.3 - 5.0zone d (black) 33.2 59.4 2.0 0.3 - 5.0

33.7 58.0 1.9 0.2 0.1 4.435.1 57.7 1.8 0.3 - 4.2

bulk AISI316 68.8 18.0 10.9 1.9 0.4 0.1

') carbon concentration (wt%) not changed significantly w ith regard to the carbon concentration of original AISI316

The com position near the grain boundaries in this zone clearly indicates tha t w eld decay has occurred. This w as m ade visible by etching w ith Railing agent. In Fig. 4.7 a

m icrograph is presented in w hich w eld decay is shown.

1 m m

Fig. 4.7 Example of w eld decay in the stainless steel part of directly bonded HPSiC/AISI316 couples.

86

Weld decay m anifests itself in the form ation of carbide precipitates at the grain boundaries w hich m ay result in in tergranular corrosion. It frequently arises w hen an austenitic stainless steel w ith a carbon content of m ore than 0.03 wt%, is heated (or slowly cooled) w ith in the tem perature region of 425°C and 800°C. Several theories concerning the m echanism of w eld decay exist. The generally accepted theory is that

precipitation of the carbide M23C6 (M=Cr,Fe,Mo) takes place preferably at grain boundaries. This leads to chrom ium and, to a lesser degree, iron and m olybdenum depleted zones along the grain boundaries, w hich then becom e anodic w ith respect to the bulk of the steel and thus becom e susceptible to corrosive m edia [4.8].In the present case it is obvious that w eld decay d id indeed occur du ring the (slow) cooling from the process tem perature to room tem perature. In the first place, there is

about 0.08 wt% free carbon present in the steel and in the second place, a large source of carbon in the form of silicon carbide is present in the vicinity of the steel. U nder

certain conditions (long process tim e and high process tem perature) carbide precipitation w as even observed to occur inside grains.In stainless steel containing m olybdenum , also M6C m ay be form ed. In type AISI316 stainless steel, M6C appears only after very long (1500 h) ageing and form s in a lim ited tem perature region around 650°C [4.9]. Presumably, M6C form s from M23C6, the suggested reaction sequence being:

M23C6 ( + M oat650°C ) ^ (Fe,Cr)21M o2C6 (+ M o) -> M6C (4.1)

In alloys containing large am ounts of Mo (> 3 wt%) the precipitation of M6C occurs at higher tem peratures and is generally not observed in AISI316 steel.

In a com prehensive article by Weiss and Stickler [4.11] about phase instabilities during h igh tem perature exposure of austenitic stainless steel AISI316, it is stated tha t carbon plays a crucial role in nucleation and grow th of carbides and interm etallic com pounds (a-phase, %-phase and r|-phase). In their investigation the au thors exam ined AISI316 steel w hich w as solution treated at 1260°C during 90 m inutes, w ater quenched and then aged at different tem peratures during various times. They reached several conclusions w hich will be sum m arized below.

First of all, it w as clearly dem onstrated that the austenitic m atrix of AISI316 stainless steel is unstable during heating to higher tem peratures and various carbides and interm etallic phases m ay be formed.

The sequence of form ation of the various phases was found to be as follows. Initially,

M23C6 precipitates below 900°C as a result of the supersaturation of carbon in the austenitic m atrix. This carbide form ation is rapid due to the presence of carbide form ing

87

elem ents (Cr, Mo, Fe) along initial nucleation sites of carbides (especially grain boundaries) and the fast diffusion of interstitial carbon atoms.In addition , it w as found that form ation of interm etallic com pounds is considerably

re tarded or even prevented. This phenom enon was attributed to tw o causes.The m ain cause is tha t the stability range for the interm etallic phases decreases rapidly w ith increasing am ount of carbon in solid solution. It is argued tha t a certain critical m inim um am ount of carbon in the m atrix has to be reached by carbide precipitation before interm etallic com pounds can form. Furtherm ore, the slower diffusion of the substitutional elem ents required for nucleation and grow th of interm etallic phases com pared to the fast carbon diffusion to sites w hich are already enriched in carbide form ing m etal atom s, is another (although less prom inent) cause of the delayed form ation of interm etallic phases.The results obtained by Weiss and Stickler may account for the form ation of M23C6

during the diffusion bonding experim ents. Apparently, only carbide form ation takes

place and form ation of interm etallic com pounds is suppressed by the relatively high carbon concentration in the steel due to diffusion of carbon, generated by the decom position of SiC, into the steel matrix. In addition, it m ust be m entioned that a- phase form ation only takes place below 830°C [4.13] and, m ore im portantly, after long annealing or ageing times. This m eans that in the present diffusion bonding experim ents G-phase could only have been form ed during cooling to room tem perature. However,

cooling has taken place at a too h igh rate, because no evidence of the presence of a a- phase w as found. The sam e reasoning applies for both the q - and y-phases It w as found tha t after the solution treatm ent at 1260°C, w ater quenching and subsequent ageing at 816°C, M23C6 is form ed after about 0.1 hour, while the y-phase is the first interm etallic com pound to be form ed after more than 100 hours ageing time.

In order to obtain additional inform ation about zones c and d, X-ray diffraction

m easurem ents were carried ou t on the steel part of the joints. The results of these X-ray diffraction m easurem ents show tha t besides y-Fe (austenite) also a-Fe (ferrite) is present

in zone c. The presence of a-Fe can be explained as follows.The constituents of stainless steels can be classified into tw o groups: ferrite form ing elem ents (Cr, Si, Mo, V, Al, W, Ti, Nb) and austenite form ing elem ents (C, Mn, N i, Cu, Co, N). These elem ents do not all contribute in the same degree to the ferrite or austenite form ation. For that purpose the so-called chrom ium equivalent and nickel equivalent w ere introduced. These equivalents are actually only valid for arc w elding processes, bu t can som etim es also be applied under non-arc w elding conditions.The chrom ium equivalent for the AISI300-series stainless steels is usually expressed as:

Cr = w t% C r + w t% M o + 1.5wt%Si + 0.5 w t% N be q

(4-2)

w hereas the nickel equivalent for the AISI300-series stainless steels is denoted as:

N i = w t% N i + 30w t% C + 30w t% N + 0.5 w t% M n (4.3)c q

These equivalents can be applied in the so-called Schaeffler-DeLong d iagram in w hich the structure of the alloy (ferrite, austenite, m artensite) is p lo tted as a function of the chrom ium equivalent and of the nickel equivalent. In fact, this diagram should only be used w hen determ ining the constitution of stainless steel arc w eld deposits [4.9].

For diffusion w elding, a sim ilar approach m ay be followed provided that the Schaeffler- DeLong diagram is modified. Pryce and A ndrew s [4.10] m odified the Schaeffler-DeLong diagram for stainless steels at the hot-rolling tem perature (1150°C). It m ay be expected tha t this m odified diagram can also be used in the case of diffusion welding. A ccording to Pryce and A ndrew s the m odified chrom ium equivalent can be w ritten as:

Cr . = w t% C r + 3wt% Si + w t% M o (4.4)e q -m o d

whereas the m odified nickel equivalent is represented by:

Ni = w t% N i + 0.5w t% M n + 21w t% C + 11.5 w t% N (4.5)e q - m o d

In their article Pryce and A ndrew s state that the factor 1 for M o is valid only for a Mo content of about 0.5-1 wt%, w hereas for higher Mo contents, up to 3 wt%, this w eight factor m ay be as high as 4.

The m odified Schaeffler-DeLong diagram is given in Fig. 4.8. In this figure the relevant data from Table 4.1 and Table 4.2 are given in the form of experim ental points. It can be seen that the unaffected AISI316 (denoted as a solid triangle) and the m atrix of zone d (open triangles) bo th fall in the austenite region, w hereas the m atrix of zone c (open squares) is partly ferritic, partly austenitic. It m ay therefore be concluded that in the diffusion affected parts of the stainless steel (that is near the original interface w ith the silicon carbide), the austenitic m atrix has (partly) transform ed into ferrite. The abrupt

transition from ferrite /austen ite (zone c) to austenite (zone d) is reflected in the different appearance of both zones in the back-scattered electron images.The results p resented in the foregoing m ay be explained by considering the diffusion process in m ore detail. Initially, only austenitic stainless steel is present. W hen the

89

diffusion bonding process starts, silicon carbide decom poses into silicon and carbon. Both elem ents diffuse into the stainless steel matrix, w here carbon is caught by the carbide form ing elem ents like chrom ium , iron and m olybdenum , w hich leads to the form ation of M23C6, especially along the grain boundaries. A t the sam e tim e the matrix is enriched w ith silicon. Silicon is know n for its ferrite form ing tendency and, therefore, the austenitic m atrix w ill tend to transform into ferrite. It is for this reason that ferrite form ation occurs m ainly in the proxim ity of the original interface and is m aintained

there because it is stabilized by solution of ferrite form ing elements. These elem ents are

better soluble in ferrite than in austenite.Pryce and A ndrew s further argue that form ation of cr-phase is accellerated by the presence of a ferrite phase com pared to the nucleation of this phase from the pure austenite m atrix [4.10]. In the present case, however, no o-phase w as form ed for reasons indicated above.

25

20

1 15 E

10

5

010 20 3 0 4 0 50

0 ra A IS 1316 n m a t r i x v m a t r i x eq-mod

z o n e c z o n e d

Fig. 4.8 M odified Schaeffler-DeLong diagram for stainless steels at 1150°C; A=austenite, F=ferrite, M =m artensite [4.10].

A + F + M

90

KineticsIn order to obtain inform ation about the kinetics of the diffusion process and, more specifically, to find ou t w hether the dep th of penetration is proportional to the square root of time, the thickness of the reaction layer was m easured of a num ber of joints obtained under different process conditions. It was decided to m easure the thickness of the reaction layer in the steel p a rt of the specim en because it appeared tha t this thickness is easier to determ ine than the thickness of the reaction layer in the ceramic part. Furtherm ore, fracture alw ays occurs in the ceramic (either in zone a or in zone b) and this severely ham pers the assessm ent of the penetration d ep th in the ceram ic part.The steel specim ens w ere cut perpendicularly to the original interface, g round and polished and then etched w ith Railing agent so that the reaction layers could be m easured w ith the help of optical microscopy.It appears that the boundaries of the reaction layers are not flat b u t quite irregular. In order to quantify the reaction layer thickness, m ean values w ere calculated from four m easurem ents.A lthough, as expected, the thickness of the reaction zone w as found to increase w ith process time, the results of the optical m icroscopy exam ination do not yield a clear relationship betw een penetration depth and process time, due to large scatter in the m easured thickness.

91

4.2.2 R B SiC /A ISB 16

W elding experim ents w ere also carried out w ith the com bination reaction bonded silicon

carbide and austenitic stainless steel. The results of these experim ents show that also in the case of this com bination it is not possible to produce a perm anent bond that survives cooling to room tem perature. A distinct difference in behaviour betw een HPSiC and RBSiC is that the latter reacts m uch more vigorously w ith steel than HPSiC. This can be explained by the fact that RBSiC contains free silicon w hich is very reactive w ith respect to m ost m etals at h igh tem peratures. This is the reason w hy experim ents w ere carried

o u t at tem peratures of not higher than 1100°C.A sim ilarity in behaviour betw een RBSiC/AISI316 couples and HPSiC/AISI316 couples is the fracture mode: in both cases fracture occurs along the interface th rough the reaction zone near the original interface a n d /o r in the ceramic m aterial itself.

As in the case of HPSiC/AISI316 couples, the diffusion bonded RBSiC/AISI316 couples were exam ined by m eans of optical microscopy. A typical optical m icrograph of the steel part of a fractured RBSiC/AISI316 couple is presented in Fig. 4.9.

Fig. 4.9 Optical m icrograph of the steel part of a fractured RBSiC/AISI316 couple diffusion bonded at 1000°C during 24 hours under a m echanical pressure of 5 MPa (etchant: Kalling).

I— I400 p m

9 2

A schematic picture of a diffusion bonded RBS1C/AISI316 couple (fractured during cooling) is given in Fig. 4.10. This picture is based on the results of analyses of several

RBSiC/AISI316 couples.

RBSiC

zone a

AISI316

zone b

zone c

zone d

Fig. 4.10 Schematic representation of the reaction zones in diffusion bonded RBSiC/AISI316 couples.

A gain different reaction zones can be distinguished. These zones are treated separately in the following parts of this section.

Zone aThe reaction products w hich are form ed in the silicon carbide are very heterogeneous in com position. It is clear tha t the free silicon present in the ceramic, dom inates the course of events. In addition , the silicon carbide decom poses resulting in free silicon and free carbon. Silicon and carbon diffuse through the interface into the steel while chrom ium , iron and nickel diffuse aw ay from the steel m atrix into the ceramic.Zone a consists of a layered structure of alternating bands of a silicon rich phase, (Fe,Ni,Cr)xSiy and C. The chemical com position of this zone as obtained by EPMA is g iven in Table 4.3. The results of the EPMA m easurem ents strongly suggest tha t x=2 and y = l, yielding the gross form ula (Fe,Ni,Cr)2Si w hich is in agreem ent w ith the results of

93

other investigators [4.4, 4.12]. The silicon carbide contains voids because the free silicon norm ally present in reaction bonded silicon carbide has diffused away. A n explanation for the presence of the periodic layers in zone a has already been given in section 4.1.

Zone b

This zone, in w hich fracture occurred in m ost cases, is extrem ely irregular bo th in shape and in com position. Fracture resulted because of the presence of reaction products w hich are hard and brittle. Moreover, the level of therm al residual stresses w hich arise during cooling, is too high to be coped w ith by the materials. W ith the aid of EPMA it was found that, beside m any voids, carbon, iron, nickel, chrom ium and silicon are present.

Due to large scatter in the m easured contents of the elements, no quantitative statem ents can be m ade about the identity of the various phase(s) in this zone. It is assum ed,

however, that the matrix is a m ixed silicide containing Fe, N i and Cr [4.6].

Zone cAs in the case of HPSiC/AISI316 couples, zone c consists of tw o phases: the matrix w hich is dark in the back-scattered electron im age of Fig. 4.11 and precipitates present preferably along the grain boundaries w hich are light in the back-scattered electron image.C om pared to the com position of standard stainless steel AISI316, these precipitates appear to be enriched w ith silicon, chrom ium and m olybdenum and reduced in iron

(Table 4.3). In this case it is very likely that the precipitates at the grain boundaries

consist of the x-phase (see section 4.2.1) w hich can be denoted as (Fe,Ni)60C r2o(Mo,Si)2(). O n the other hand, a silicide of the type (Fe5Cr2Ni)Si2, m entioned in the foregoing section, m ay have been formed.

As in the case of HPSiC/AISI316 couples, X-ray diffraction analysis was also perform ed

on RBSiC/AISI316 couples in order to obtain additional inform ation about the grain boundary precipitates.Unfortunately, it w as not possible to reveal the identity of these precipitates w ith certainty using this technique. It appears, however, that the precipitates have a body centered cubic structure, w hich is consistent w ith the assum ption that the precipitates are of the x-type.

94

Table 4.3 Chemical composition in wt% of zones a, c and d (T=1050°C, 24 hours).

Region Fe Cr Ni Mo Si C

bulk RBSiC 0.1 - - - 73.1 26.6

zone a (silicide) 65.0 4.1 7.9 - 23.1 ’)58.9 2.3 19.4 - 19.1 ’)

m atrix zone c (dark) 69.7 14.7 7.5 0.7 5.3 0.469.6 15.0 8.0 0.8 5.3 0.469.1 15.1 8.1 0.9 5.3 0.4

grain boundaries zone c 54.1 19.0 9.6 9.3 7.9 ’)(white) 53.7 18.9 9.3 9.7 7.6 ’)

53.9 19.0 9.2 9.9 7.7 ’)

m atrix zone d 68.9 16.9 10.6 1.9 0.5 0.569.1 16.6 10.6 1.8 0.7 0.468.1 17.2 10.6 1.9 0.5 0.5

bu lk AISI316 68.8 18.0 10.9 1.9 0.4 0.1

') carbon concentration (wt%) not changed significantly w ith regard to the carbon concentration in original AISI316

Fig. 4.11 Back-scattered electron im age of zone c (upper part) and part of zone d (lower part) of a RBSiC/AISI316 couple.

95

This uncertainty in the in terpretation of the X-ray diffraction results, can be attributed to the fact that the 29-values of m any interm etallic com pounds are alm ost identical. A dditionally, the analysis w as ham pered by the low density of the precipitates in the sample.

As in the case of H PSiC /A IS1316 couples, the chrom ium and nickel equivalents were calculated from the values of Table 4.3 using the equations (4.4) and (4.5). The obtained values are p lotted in the m odified Schaeffler-DeLong diagram w hich is presented in

Fig. 4.12. It appears that the m atrix of zone c is partly ferritic and partly austenitic as is indicated by the open squares in the figure. The partly ferritic structure of zone c can be explained as follows. In the beginning only austenite is present. D uring the diffusion bonding process, free silicon and silicon and carbon originating from the decom position

of silicon carbide diffuse into the steel. The carbon form s carbides w ith iron, chrom ium and m olybdenum w hereas, due to the presence of silicon, the m atrix is partly

transform ed into ferrite, silicon being a strong ferrite form ing element.

7

10 2 0 30 4 0 50

C ra AISI316 □ m a t r i x v m a t r i x eq-mod

z o n e c z o n e d

Fig. 4.12 M odified Schaeffler-DeLong diagram for stainless steels at 1150°C; A =austenite, F=ferrite, M =m artensite [4.10],

A + F + M

96

Zone d

In all couples exam ined there are slight indications of weld decay in the steel part, b u t it appears th a t w eld decay plays a less prom inent role than in the case of hot-pressed silicon carbide/A ISI316 combination. This m ay be caused by the presence of free silicon w hich suppresses the dissociation of silicon carbide resulting in a sm aller am ount of free

carbon. The m atrix of zone d appears to be com pletely austenitic as is indicated by the open triangles in Fig. 4.12.

KineticsAs in the case of HPSiC/AISI316 couples, it w as attem pted to m easure the thickness of the reaction layer to find out w hether there is a linear relationship betw een the square

root of process tim e and the reaction layer thickness.However, determ ination of the penetration dep th w as found to be extrem ely difficult due to the irregular geom etry of the reaction layers. This m ade a reliable assessm ent of the relationship betw een the process tim e and the penetration dep th impossible.

4.3 B onding experim ents in sh ie ld in g gas

For reasons of com parison, experim ents were also carried o u t in a shielding gas flow (composition: 95 vol.% Ar, 5 vol.% H 2). The results obtained are sim ilar to the results of the experim ents carried ou t in a vacuum environm ent [4.14] and will be sum m arized below.

4.3.7 HPSiC/AISl316

H ot pressed silicon carbide cannot be joined to austenitic stainless steel in a direct w ay w ithou t failing during cooling to room tem perature. Failure of the diffusion bonded couples occurs in the sam e w ay as in the case of HPSiC/AISI316 couples diffusion bonded in a vacuum environm ent. Below about 950°C the interaction betw een HPSiC and AISI316 is insufficient, w hereas at tem peratures of 1150°C and higher m elting of an iron silicide phase occurs.The joint can be subdivided in reaction zones in the sam e w ay as in the case of diffusion bonding in h igh vacuum . A significant difference w ith the h igh vacuum results seems

tha t in zone c an interm etallic com pound of the %-type is form ed. The com position of this phase w as found to vary appreciably w ith a h igh tolerance for m etal interchange

97

[4.11]. A typical com position has been determ ined by K autz et al. [4.15] as (Fe<Ni)36Cr18Mo4. In the present case the com position is about (Fe,Ni)62Cr2(l(Mo,Si)20.

The presence of this phase can be explained by the different w ay of cooling to room tem perature w hich w as applied in the shielding gas experim ents. In the diffusion bonding experim ents perform ed in high vacuum , cooling occurred at a rate of 5°C /m in, w hereas the specimens in the shielding gas w ere cooled in 5 hours time from process

tem perature till 300°C and then furnace cooled to room tem perature. This low er cooling rate m ight be responsible for the form ation of interm etallic com pounds like the %-phase (see section 4.2.1). Further evidence for the presence of a %-phase is that it has a bcc

structure and nucleates preferably at grain boundaries [4.11].

It is also conceivable that a m ixed silicide of the type (Fe5Cr2Ni)Si2 is form ed as m entioned in section 4.2.1.

4.3.2 RBSiC/AIS1316

Also in the case of reaction bonded silicon carbide perm anent bonds can not be produced w ith austenitic stainless steel in a shielding gas environm ent. A t a tem perature of 1150°C and higher, m elting occurs of possibly an iron silicide phase. In all cases a bond is form ed betw een the tw o m aterials b u t this bond fails during cooling to room tem perature. The fracture of the diffusion bonded couples takes place at the same

locations as in the case of diffusion bonded couples in high vacuum . It appears that the sam e reaction zones are present in the joint as in the case of sim ilar experim ents perform ed in h igh vacuum (Fig. 4.10).In zone c, a com pound w hich is probably of the %-type having the gross form ula (Fe,Ni)6flCr20(Mo,Si)20 is found. This phase precipitates preferably along the grain boundaries of the phase and has the bcc structure. As is indicated above in section 4.3.1, it is no t im aginary that a m ixed silicide is form ed of the type (Fe5C r2Ni)Si2.

98

4.4 Conclusions

The results described in the foregoing show that no perm anent bonds can be accom plished w hen diffusion bonding silicon carbide to stainless steel w ithout m aking

use of an interm ediate (metallic) layer.

The causes for this are twofold:during diffusion bonding of silicon carbide to austenitic stainless steel reaction products are form ed w hich are hard and brittle and have a therm al contraction behaviour w hich differs from that of both AISI316 and SiC; during cooling from process tem perature to am bient tem perature, residual stresses are generated due to the differences in therm al expansion coefficients betw een AISI316, SiC and reaction products; it appears that these stresses can not be accom m odated by any of the m aterials [4.16].

The reaction layers form ed during the diffusion w elding process can be subdivided in different zones.In the silicon carbide a zone is form ed (zone a) w hich consists of alternating bands of (Fe,Ni,Cr)2Si and C w hich ru n parallel to the original interface.A long the original interface SiC/AISI316 a zone is form ed (zone b) in w hich also layers are present. These layers w hich also run parallel to the original interface, consist of carbon precipitates and voids. The m atrix consists m ainly of iron and silicon and small am ounts of nickel and chrom ium . Fracture occurs either through this zone or th rough zone a in the ceramic.The next zone (zone c) is situated in the stainless steel and consists of tw o phases. The

m atrix appears to be AISI316 w hich is depleted of Ni and Mo, and enriched w ith Si. The grain boundaries in zone c appear to be either a silicide or an interm etallic com pound of the X'type.The last zone (zone d) has a m atrix of AISI316 w hich is depleted in chrom ium . The phenom enon of w eld decay is reflected in the chemical com position of the grain boundaries w hich contain relatively high levels of chrom ium and carbon indicating that a M23C6 phase is form ed.

The results obtained w ith HPSiC are sim ilar to those obtained w ith RBSiC, an im portan t difference being that w ith the latter m aterial a more vigorous reaction is observed due to the presence of free silicon.It appears tha t diffusion bonding in a vacuum or in a shielding gas yields com parable results, im plying tha t the environm ent has little or no influence on the bonding process. The low er cooling rate after the diffusion bonding process in shielding gas m ight be

99

responsible for the form ation of the interm etallic com pound %.In the case of HPSiC /AISI316 couples, fracture is observed to occur either th rough zone b or th rough the ceramic material. In the latter case fracture is initiated in zone b and the resulting crack propagates further through the ceramic m aterial. In the case of RBSiC/AISI316 couples, fracture often took place in zone b; in a single case fracture occurred th rough the ceramic m aterial b u t was initiated in zone b.

A pplication of a metallic interlayer m ight be beneficial bo th in reducing residual therm al stresses through yielding (relaxation of therm al stresses) and in suppressing form ation of brittle reaction products w hich are form ed in the vicinity of the steel/S iC interface (diffusion barrier).

100

References

4.1 J.S. K irkaldy and D.J. Young, Diffusion in the condensed state, The Institute of Metals, London 1987.

4.2 R. M ehan, Mat.Sci.Res., 14 (1980) 513-532.4.3 K. Osinski, The influence of alum inium and silicon on the reaction betw een iron

and zinc, Ph.D. Thesis, E indhoven U niversity of Technology, 1983.4.4 R.C.J. Schiepers, The interaction of SiC w ith Fe, N i and their alloys, Ph.D. Thesis,

E indhoven U niversity of Technology, 1991.4.5 F.J.J. van Loo, J.A. van Beek, G.E. Bastin, R. Metselaar, D iffusion in Solids -

Recent D evelopm ents, 1985, p .231-259.4.6 M. Backhaus-Ricoult, M etal-Ceramic Interfaces, Proc.Int. W orkshop, M. Ruble et

al. (eds.), Pergam on Press, Oxford 1990, p .79-92.4.7 P. Villars and L.D. Calvert, Pearson 's H andbook of C rystallographic Data for

Interm etallic Phases (Vol.3), Am erican Society for Metals, Ohio 1985.4.8 J.F. Lancaster, M etallurgy of W elding (4th ed.), Allen & U nw in (Publishers) Ltd.,

London 1987.4.9 C.J. N ovak, H andbook of Stainless Steels (Chapter 4), D. Peckner and I.M.

Bernstein (eds.), M cGraw-Hill Book Company, N ew York 1977, p .1-78.4.10 L. Pryce and K.W. A ndrew s, J.Iron Steel Inst., 195 (1960) 415-417.4.11 B. Weiss and R. Stickler, Metall.Trans., 3 (1972) 851-866.4.12 G.V. Raynor and V.G. Rivlin, Int.Met.Rev., 30 (1985) 181-208.4.13 O. Kubaschewski, Iron - Binary Phase Diagrams, Springer Verlag, Berlin 1982,

p .31-34.

4.14 G.J. de Jonge, H et diffusielassen van silicium carbide aan austenitisch roestvast

staal, A fstudeerverslag, Technische Universiteit Delft, 1989.4.15 H.R. K autz and H. Gerlach, Arch.Eisenhiittenw., 2 (1968) 151.4.16 B.T.J. Stoop, G. den O uden, M aterialen, 5 (1989) 39-42.

101

CHAPTER 5

Interlayers

5.1 In troduction

It w as show n in the previous chapter tha t direct bonding of silicon carbide ceramics and austenitic stainless steel by m eans of diffusion w elding is im possible. A lthough bonds m ay be form ed at the process tem perature, these bonds fracture during cooling to room tem perature. The im practicability of joining the two m aterials by m eans of diffusion w elding is due to:

the form ation of brittle reaction products;the inherent generation of residual therm al stresses during cooling from process

tem perature to am bient tem perature.For that reason it w as deciced to m ake use of metallic interlayers w hich have to satisfy the following criteria:

the coefficient of therm al expansion (a) of the interlayer has to m atch the coefficient of therm al expansion of both the metal and the ceramic; in practice this m eans that the value of a of the interlayer should lie som ew here betw een tha t of the m etal and that of the ceramic;the yield strength of the interlayer m aterial should be low, so tha t the interlayer

can accom m odate the therm al stresses generated during cooling to room

tem perature by plastic deform ation;the Young's m odulus of the interlayer m aterial should be low in order to keep the elastic stress at a relatively low level;the interlayer m aterial should have lim ited reactivity w ith respect to bo th the

silicon carbide and the steel.

It is obvious tha t no interlayer m aterial m eets all of the requirem ents stated above and tha t a com prom ise has to be found. O n the basis of the above m entioned criteria nickel,

copper and copper-nickel alloys w ere selected.This chapter deals w ith the results of diffusion bonding experim ents w ith the com bination silicon carbide and austenitic stainless steel applying Ni interlayers, Cu interlayers and several binary Cu-x%Ni interlayers. The m ain purpose o f these diffusion bonding experim ents is to assess which type of interlayer is best su ited for joining silicon

carbide to austenitic stainless steel.

102

5.2 Experiments with interlayers

The experim ents w ere carried ou t using the equipm ent described in C hapter 3, applying a vacuum environm ent of about 10’3-10'4 Pa. The process tem perature w as chosen

betw een 811°C and 1061°C and the experim ents were perform ed at intervals of 25°C. The process time, as in the case of the direct bonding experim ents, w as varied betw een22.5 m inutes and 24 hours. The m echanical pressure w as chosen betw een 0.1 M Pa and 15 MPa.Diffusion bonding experim ents w ere perform ed w ith Ni, C u and Cu-x% w t N i (x = 1, 3, 5, 10, 15) as interlayer m aterial. The nickel disks w ere annealed at 600°C d u ring 30 m inutes just before they w ere inserted betw een the ceram ic and the metal. In the case of copper and copper-nickel alloys, this was considered unnecessary, as it appears from literature data [5.1] that the yield strength (Rp 02) is reduced to very low values at tem peratures of 850°C and higher. This is show n in Fig. 5.1.

3 0

° 20 Q_2

“ ■ 10 cc lu

o8 0 0 9 0 0 1000 1100

T PC)

° Ni * CuFig. 5.1 Yield strength (Rp „ 2) as a function of tem perature for p redeform ed nickel and

copper.

Furtherm ore, it is know n that at these tem peratures the creep rate of copper and copper- nickel alloys is so h igh that plastic deform ation occurs easily and intim ate contact betw een the surface areas is reached in a very short time.

103

Three types of silicon carbide w ere used: ho t pressed silicon carbide, reaction bonded

silicon carbide and hot isostatic p ressed silicon carbide, respectively.The m etal p a rt of the diffusion bonded specim ens w as austenitic stainless steel AISI316. The chemical com position and some relevant properties of the various m aterials are given in C hapter 3.

5.2.1 N i as interlayer material

As in the case of direct bonding of silicon carbide to austenitic stainless steel, experim ents were perform ed at relatively low mechanical pressure (0.1 MPa) and at relatively h igh mechanical pressure (7.5 or 15 MPa). It appeared that the difference in

m echanical pressure does not lead to significantly different results.From the results obtained it can be concluded that it is very difficult, if not im possible, to bond silicon carbide perm anently to AISI316 using nickel as interlayer material. It appears that in m ost cases a bond is established at the process tem perature b u t that during cooling the bond fails.In all cases, fracture occurs either in the reaction layer betw een the nickel and the silicon carbide or in the silicon carbide itself. If fracture occurs in the silicon carbide, it appears that initiation of the crack takes place at the surface in or very near the reaction layer betw een the nickel and the silicon carbide. The occurrence of fracture in the silicon

carbide im plies tha t the bond betw een nickel and silicon carbide is more able to stand

up against the effect of therm al contraction during cooling than silicon carbide itself. U nder specific conditions bonds can be accom plished betw een silicon carbide and austenitic stainless steel w hen applying a nickel interlayer. However, these bonds are very w eak and after the joining process cracks are visible in the ceramic m aterial, indicating tha t therm al stresses have increased to a level above the ru p tu re strength of the ceramic.Above 1015°C the reaction betw een nickel and silicon carbide was found to be very

severe and total em brittlem ent of the interlayer m aterial resulted.Above 1100°C (1050°C w hen free silicon is present) a liquid phase is form ed which, m ost probably, is a nickel silicide.Thus it can be concluded that failure of the bonds is due to the excessive reaction betw een Ni and SiC. The reaction products cause serious deterioration of the bond and during cooling to room tem perature fracture occurs due to the developm ent of therm al residual stresses.

The results presented above clearly indicate that Ni is not a suitable interlayer m aterial

to be used to join AISI316 and silicon carbide.

104

5.2.2 Cu as interlayer material

Diffusion bonding experim ents were also carried o u t w ith copper as interlayer m aterial betw een ho t pressed silicon carbide and austenitic stainless steel. It appears that under specific conditions perm anent bonds can be m ade w ith copper as the interm ediate

m aterial. How ever, the m axim um shear strength of these bonds w as found to be extrem ely low (approxim ately 1 MPa).Visual inspection show ed that there had been very little chemical interaction betw een

the copper and the silicon carbide. In m ost cases fracture occurred along the C u/S iC interface. O nly in tw o cases a tiny piece of the ceramic m aterial rem ained stuck at the copper layer. This is an indication that physical a n d /o r mechanical adherence betw een the copper and the silicon carbide is responsible for the (weak) bonding.

It is interesting to note that under specific conditions, notably at higher process tem peratures a n d /o r longer process times, fracture occurs in the ceramic m aterial. This m ight be an indication that at higher process tem peratures a n d /o r longer process times, in addition to physical and mechanical phenom ena, chemical interaction starts to play a role.

Before being used, the copper interlayer m aterial is rolled, punched out of a foil and flattened (Chapter 3) w hich results in a certain degree of plastic deform ation. D uring heating to the process tem perature and during w elding and cooling the interlayer m aterial recrystallizes and grain grow th occurs. W hen the actual w elding takes place,

a mechanical pressure is exerted on the specim ens and the m aterial deform s plastically. D uring the cooling stage, especially in the tem perature region from 250°C till room tem perature, the m aterial again experiences plastic deform ation due to the residual stresses that occur w hen m aterials w ith different therm al contraction are joined at high

tem perature and cooled to room tem perature. Optical microscopy shows that indeed recrystallization and grain grow th has occurred in the interlayer m aterial. This is

illustrated in Fig. 5.2. The original copper sheet is cold rolled to a thickness of 0.2 m m w hich results in a structure consisting of long (a few hundreds of m icrom eters) and

small (about 5 micrometer) fibrous grains.A sim ilar effect w as found in the case of copper-nickel alloys w hich are treated in the next section.

The diffusion bonding experim ents w ith copper interlayers show that too little

interaction occurs betw een the copper and the ceramic to provide a strong bond. In several instances, perm anent bonds were accom plished b u t these show ed very weak strengths w hich indicates that only physical a n d /o r mechanical adherence plays a role.

It can be concluded that copper is not a suitable interlayer m aterial to be used for the

105

Fig. 5.2 Example of recrystallization and grain grow th in the copper interlayer.

joining of silicon carbide to austenitic stainless steel. Apparently, the interaction betw een copper and silicon carbide is too weak to w ithstand the therm al stresses w hich develop during cooling to am bient tem perature.

106

5.2.3 C u-x% N i as interlayer material

IntroductionIt w as suggested by Schiepers [5.2] that a possible solution of the problem s described in the foregoing, could be the in troduction of a metallic interlayer w hich consists of a

ductile, non-reactive m atrix and a small addition of a reactive species. In the present case copper and nickel (in the form of a Cu-x%Ni alloy) seems to be a suitable combination. The idea is tha t the copper m atrix is able to cope w ith the residual therm al stresses that arise during cooling, w hile nickel is actually the bonding agent. Thus it m ay be expected tha t the copper reduces the therm al stresses du ring cooling from process tem perature to room tem perature th rough yielding, w hereas the nickel is the active species that renders the bonding betw een the interlayer and the ceramic m aterial (this idea has already been applied in active metal brazing; see C hapter 2).

The application of the copper-nickel alloy in diffusion bonding experim ents in w hich silicon carbide is to be joined to austenitic stainless steel, w as preceded by the question w hat the m inim um am ount of nickel in the binary alloy should be in order to have a

sufficient extent of chemical reaction. It has been calculated that the activity of N i has to be at least 0.039 in order to react w ith silicon carbide at 968°C, assum ing regular

solutions and neglecting the different crystal structures of nickel and silicon. These assum ptions are justified as long as the solubility of silicon in nickel (13 at.% Si) is not exceeded. Therefore, the aNi in copper-nickel alloys to be used as interlayer m aterial in the case of SiC/AISI316 joining m ust be > 0.039 so that chemical reaction can take place. The relation betw een activity and concentration of nickel in solid copper-nickel alloys has been determ ined by R app and M aak at 700°C and 1000°C [5.3]. It appears that the

nickel content in the solid copper-nickel alloys has to be > 0.01 at.% to bring about sufficient chemical reaction.

In order to find out w hich Cu-N i com bination w ould m atch best betw een silicon carbide and stainless steel, diffusion w elding experim ents were carried ou t w ith five binary copper-nickel alloys (containing respectively 1, 3, 5, 10 and 15 wt% Ni).

C u-l% N i

The diffusion bonding experim ents carried out w ith a C u-lw t% N i interlayer dem onstrate tha t this interlayer is not suited to join SiC to AISI316. In m ost cases a very w eak bond w as accom plished b u t during cooling to room tem perature fracture occurred along the interface SiC / C u-lw t% N i, indicating that the bond is the w eakest link.

107

Cu-3%NiIn the case of diffusion bonding using Cu-3wt%Ni as interlayer m aterial only one perm anent bond w as produced, having a shear strength of 19.5 MPa. In all o ther cases either too little interaction had occurred betw een the interm ediate and the ceramic or

fracture along the interface had taken place (during cooling or w hen positioning the

bonded couples in the shear test apparatus). As in the case of SiC /C u-1 wt%

Ni/AISI316, it is evident tha t the joint betw een SiC and Cu-3wt% Ni is the w eakest part of the specimen.

Cu-5%NiW elding experim ents carried ou t w ith a Cu-5wt% Ni insert p roved to be m ore successful than the w elding experim ents w ith copper-nickel inserts w hich contain less nickel.The results of these experim ents are presented in Fig. 5.3 in the form of a process tem pera tu re/p rocess time diagram . This figure shows that perm anent bonds can be produced in a specific region of the tem pera tu re/tim e diagram . O utside this region couples fracture during cooling to room tem perature or are not joined at all. H igh process tem perature (> 1038°C) in com bination w ith long process tim e results in melting phenom ena. In all cases fracture occurs along the SiC/Cu-5% N i interface.

1075y

V V 0 9o 1025

V V V V V<D .

V + + V VZ3

- t— 975O V + + + Vi_0)CL V + + + VE0) 925 V V V + V

-t-

Q7S

V

_1_1 ■ ■ I 1__

V V +

_l_l_l_______10 100 1000

t i m e ( m i n )

+ b o n d e d v n o t o m o l t e nb o n d e d

Fig. 5.3 Results of diffusion w elding experim ents w ith HPSiC/Cu-5% Ni/AISI316.

108

Cu-10%NiThe results of diffusion w elding experim ents w ith the com bination of SiC and AISI316 applying a Cu-10wt%Ni interm ediate are com parable w ith those applying a Cu-5wt% Ni interm ediate.

The results of the w elding experim ents show that also in this case perm anent bonds can

be produced w ith in a specific process tem pera tu re/p rocess tim e region. However, this region is sm aller than the region found in the case of Cu-5wt% Ni as interlayer material.

Cu-15%NiW elding experim ents w ith an interlayer containing the largest am ount of nickel, Cu- 15wt%Ni, p roved to be less successful than those perform ed w ith interlayers consisting of copper w ith 5wt% Ni and 10wt%Ni, respectively. It appeared that m elting starts to occur already at 1015°C and that at low er tem peratures no perm anent bonds can be m ade.

109

5.3 Conclusions

The experim ental results described in the previous sections show that both nickel and copper in the pure form are not suited as interlayer m aterial to be used for the joining of stainless steel AISI316 and silicon carbide.Nickel is too reactive w ith respect to silicon carbide, w hereas in the case of copper the

reactivity is too small.It appears that Cu-x%Ni alloys are better suited to serve as interlayer m aterial, due to their lim ited reactivity (Ni) in com bination w ith their good ductility (Cu). The best results are obtained w ith Cu-5wt%Ni and Cu-10wt%Ni. U sing these interlayers, relatively h igh shear strengths (up to 60 MPa) can be obtained w ith in a specific process tem pera tu re/p rocess tim e window.The experim ents w ere carried out w ith three types of silicon carbide: ho t pressed (HP) silicon carbide, reaction bonded (RB) silicon carbide and hot isostatically pressed (HIP)

silicon carbide, respectively. The m ost reliable results w ere obtained w ith H IP silicon

carbide.

O n the basis of these conclusions it w as decided to continue the experim ents w ith HIP silicon carbide and w ith Cu-5wt%Ni and Cu-10wt%Ni as interlayer material.

110

References

5.1 P. D ahlm ann, Zug-, Kriech- und Relaxationsverhalten von K upfer bei Tem peraturen dicht unterhalb des Schm elzpunktes, Ph.D. Thesis, Technische

Hochschule Aachen, 1988.5.2 R.C.J. Schiepers, The interaction of SiC w ith Fe, N i and their alloys, Ph.D. Thesis,

E indhoven U niversity of Technology, 1991.5.3 R.A. Rapp and F. M aak, Acta Met., 10 (1962) 63-69.

I l l

CHAPTER 6

D iffusion bonding experiments with copper-nickel interlayers

6.1 Introduction

In the previous chapter the results w ere presented of bonding experim ents w hich were carried out w ith different interlayers (Ni, Cu and several Cu-x%Ni alloys).It w as show n that perm anent bonds can be obtained betw een silicon carbide and austenitic stainless steel using binary Cu-Ni alloys as interlayer material. The best results w ere obtained w ith a Cu-5%Ni interlayer and a Cu-10%Ni interlayer, respectively. In addition, it appeared that HIPSiC is the m ost suited type of silicon carbide to join w ith AISI316 w hen applying metallic interlayers.This chapter deals w ith additional experim ents w hich were perform ed w ith HIPSiC and

Cu-5%Ni and Cu-10%Ni as interlayer materials. The results of these experim ents are to be considered as an extension of the results described in C hapter 5.First of all, the com position and structure of the joint is dealt w ith by m aking use of analytical techniques like optical microscopy, X-ray diffraction, scanning electron microscopy, electron probe micro analysis and A uger electron spectroscopy.After this, attention is focused on the kinetic aspects of diffusion w elding SiC to AISI316 w hen applying a Cu-5%Ni interlayer or a Cu-10%Ni interlayer.In addition, the mechanical behaviour of the diffusion bonded couples is considered,

thereby taking into account the influence of the process tem perature, the process time, the m echanical pressure and the thickness of the interlayer material.Finally, a general m odel of the SiC/Cu-x% Ni/AISI316 joint is p resented in w hich the

results of the foregoing sections are briefly sum m arized.

This chapter is divided in tw o parts. In the first part the above m entioned aspects w ith respect to the Cu-5%Ni interlayer m aterial are treated extensively, w hereas in the second part the Cu-10%Ni interlayer m aterial will be dealt w ith in a m ore global way, the

behaviour of the latter interlayer m aterial being in m any (but not all) respects sim ilar to tha t of the form er metallic insert.

112

6.2 Copper - 5 wt% nickel

6.2.1 Welding experiments

W elding experim ents were carried ou t in vacuum in the tem perature-tim e w indow

presented in the previous chapter. The process tem perature w as varied betw een 945°C and 992°C while the process tim e w as varied betw een 90 m inutes and 360 m inutes. As w as show n in C hapter 5, in this tem perature-tim e region perm anent bonds can be m ade. In addition, tw o other process param eters, mechanical pressure and the thickness of the interlayer m aterial, w ere varied in order to exam ine their influence on the joint formation.The mechanical pressure w as varied betw een 3 and 30 MPa by choosing the following

set of pressures: 3, 5, 7.5,15 and 30 MPa.The thickness of the interlayer w as varied betw een 0.2 and 1 m m by varying the num ber of interlayer disks (each 0.2 m m thick) from 1 to 5.

6.2.2 Structure of the joint

The joints obtained w ere exam ined in detail using various techniques. O n the basis of this exam ination it appears that the joint can be d ivided in four zones:

a reaction zone in the silicon carbide near the SiC /Cu-5% N i interface (zone a); a diffusion zone in the copper-nickel alloy near the S iC /C u-5% N i interface (zone

b);a diffusion zone in the copper-nickel alloy near the AISI316/Cu-5% Ni interface

(zone c);a diffusion zone in the austenitic stainless steel near the AISI316 / Cu-5%Ni interface (zone d).

A schematic illustration of the four zones is given in Fig. 6.1, w hereas Fig. 6.2 gives an

overall view of the joint, exhibiting the four different zones. Fracture occurs either betw een zone a and zone b (sometimes partly in zone a) or starts betw een zone a and zone b (along the circumference) and proceeds further through the ceramic material. A pparently, the carbon (graphite) w hich is present in zone a (see next section), is the w eakest part and a crack w ill start at that particular place, continuing its w ay along the

m ost fragile path.The zones w ere exam ined w ith the aid of optical m icroscopy (using an etchant com posed of a solution of am m onium hydroxide and, w hen necessary, hydrogen peroxide), electron probe m icroanalysis, scanning electron microscopy, X-ray diffraction

113

and A uger electron spectroscopy (for details, see C hapter 3). In the following each of the

four zones is described separately.

HIPSiC

original interfaces ^ interlayer

AISI316

zone a zone b

zone c

zone d

Fig. 6.1 Schematic illustration of the four different zones in a HIPSiC/Cu-5% Ni/AISI316 joint.

I 140 pm

Fig. 6.2 M icrograph of HIPSiC/ Cu-5%Ni / AISI316 joint, diffusion bonded at 992°C during 180 m inutes under a pressure of 7.5 MPa in h igh vacuum (etchant:n h 4o h + h 2o 2).

114

Zone aA m icrograph of zone a is given in Fig. 6.3. The thickness of this zone depends on the w elding param eters and lies in the range of a few m icrom eters. The interface betw een this zone and the unaffected silicon carbide is ra ther irregular in shape, w hereas the interface w ith the interlayer (original interface) is straight.

I 120 p m

Fig. 6.3 Reaction zone a in SiC near the copper-nickel.

Both spot m easurem ents and line scans w ere perform ed in this zone w ith an electron

probe. It appears that the zone is very inhom ogeneous, both w ith respect to its geom etry and w ith respect to its chemical composition. Roughly speaking, one can d istinguish two phases:

a phase w hich m ainly consists of N i, Si and C;

a phase in w hich C is the m ost p rom inent elem ent besides Si, C u and Ni.It should be stated, however, that there is a large scatter in the values obtained w ith EPMA, once again indicating that the reaction zone is very diverse as far as its com position is concerned.

The SEM m icrograph depicted in Fig. 6.4 shows the decom posed silicon carbide w ith the tw o different phases and a crack which separates zone a and zone b.

115

Fig. 6.4 SEM m icrograph of zone a (lower part) and zone b (upper part) of a diffusion bonded H IPSiC/Cu-5% Ni/AISI316 couple (992°C, 180 m in, 7.5 MPa, high vacuum).

In order to collect m ore inform ation about the structure of zone a, A uger electron spectroscopy was perform ed on a H IPSiC/Cu-5% Ni/AISI316 joint obtained w ith process param eters T=992°C, t=90 m in and p=7.5 MPa and fractured by m eans of the shear test equipm ent.

A uger d ep th profiles w ere acquired w ith different sputter rates of zone a, starting from the fracture surface of the silicon carb ide/ copper-nickel interface. The sputter rates applied were: 5, 35 and 60 n m /m in (during 100, 118 and 200 m inutes resulting in

analysis depths of 0.5, 4.1 and 12 pm , respectively). The elem ents which were

determ ined, are carbon, copper, nickel, silicon and oxygen.As an illustration Fig. 6.5 gives the dep th profile obtained at 60 n m /m in .From the three dep th profiles the following conclusions can be draw n:

both copper and nickel have diffused into the silicon carbide to a dep th of about 10 pm ; it is rem arkable that the total am ount of copper w hich has diffused into the silicon carbide is only about twice the am ount of nickel which has diffused into the silicon carbide while the com position ratio of both m etals in the original interlayer

is 95 : 5 (Cu:Ni, in wt%);

oxygen is p resent only in m inor quantities in the vicinity of the fracture surface,

w hich strongly suggests that surface contam ination is concerned; the silicon concentration is very low at the fracture surface and increases gradually

w ith increasing dep th into zone a;the carbon concentration is relatively high at the fracture surface and decreases to

a low er level at 0.5 pm , virtually rem aining at this level u p to a d ep th of 12 pm;

at a d ep th of about 10 pm the original SiC is present.

100

90

80

60

H 50 <

180 20016080 100 S P U T T E R T I M E (M IN )

1404020

Fig. 6.5 C om position dep th profile at spu tter rate of 60 n m /m in during 200 m inutes, starting from the fracture surface, th rough zone a into the silicon carbide.

It should be m entioned that in A uger analysis the shape of the Si and C peaks changes depending on the chemical state (bonding).In order to obtain inform ation about the chemical state of both silicon and carbon in

zone a, Target Factor Analysis (TFA) [6.1-6.3] w as perform ed using the set of Auger spectra acquired in the dep th profiling. It appeared that both elem ents occur in two

different chemical states at different depths in the silicon carbide specim en. This is illustrated in Fig. 6.6 and Fig. 6.7. In these figures differentiated spectra are presented of the Si L23W and C K W A uger transitions. It can be seen that there are clear dissim ilarities betw een cycle 3 (representing a layer very near the surface of the SiC) and

cycle 100 (representing a layer w hich corresponds to original SiC) indicating that the chemical state of both carbon and silicon changes going from the surface into the bulk.

117

\ cycle 100

6 "d N(E) cycle 35 -d E

4 -

3 -

2 -

10085

KINETIC ENERGY, eV

Fig, 6.6 AES spectra of silicon near the surface (cycle 3) and in the bu lk (cycle 100).

cycle 1009 "

cycle 3

6 "

a n (e ) B ;;d E

2 ■'

2 8 0 290260 270250

KINETIC ENERGY, eV

230220

Fig. 6.7 AES spectra of carbon near the surface (cycle 3) an d in the bulk (cycle 100).

118

This is illustrated once again in Fig. 6.8 and Fig. 6.9 in w hich the fractions of silicon and carbon in the tw o different chemical states are p lo tted versus the spu tter tim e (w hich in fact results in a dep th profile).

Si e lem en ta l

Si in SiC

zoPucpitl.

0.6

0.4

0.2

- 0.2100

S PU T T E R T IM E (M IN )

Fig. 6.8 P lot of the silicon fraction w ith respect to the total am ount of elem ental Si and of Si bonded to carbon as a function of the sputter time.

C e lem en ta l

C in SiC0.8

O 0.6h-u<u- 0.4

0.2

- 0.2100

S P U T T E R T IM E (M IN )

Fig. 6.9 Plot of the carbon fraction w ith respect to the total am ount of elem ental C and of C bonded to silicon as a function of the sputter time.

119

Finally, in Fig. 6.10 the fraction versus sputter tim e of bo th elem ents in tw o different chemical states is depicted, very clearly show ing that the chemical state of bo th elements are tightly related. This figure clearly indicates that silicon carbide decom poses into free silicon and carbon (graphite) near the surface and that the degree of decom position decreases w ith increasing depth. In addition, it has been verified tha t the form of the spectra corresponds to the form of the spectra of pure graphite, silicon and silicon

carbide, respectively. This evidence is supported by the EPMA results presented earlier

and by the XRD results w hich show ed that carbon is present in the form of graphite in

the vicinity of the in terlayer/S iC interface.

1.2

1

0.8

0.6

0.4

0.2

0

- 0.2 180 200140 160120100

SPUTTER TIME (MIN)

Fig. 6.10 Plot of silicon and carbon fraction in two chemical states versus sputter time (com bination of Fig. 6.8 and Fig. 6.9).

120

Zone bZone b is situated in the copper-nickel alloy next to the ceramic. The thickness of zone b is in the order of tens of m icrom eters and the thickness ratio zone b /z o n e a is approxim ately ten. The zone is enriched w ith silicon (produced by the decom position reaction of SiC) and depleted in copper and nickel due to the diffusion of these elem ents

into the silicon carbide. The silicon, diffusing into the copper-nickel m atrix, forms precipitates.These precipitates are clearly visible in the m icrograph depicted in Fig. 6.11. It can be seen that the precipitates become sm aller in size w hen going from the S iC /C uN i

interface to the CuN i/A ISI316 interface. The m ean length and thickness of the precipitates near the SiC /Cu-5% N i interface are 5 and 1 pm , respectively, w hile the m ean length and thickness near the Cu-5% Ni/AISI316 interface are 1 and 0.25 pm ,

respectively. M oreover, it is apparent that the density of the precipitates increases w ith increasing distance from the S iC /C uN i interface.

Fig. 6.11 O verall picture of the copper-nickel interlayer after diffusion bonding at 992°C during 180 m inutes under a pressure of 7.5 MPa in high vacuum (bright field, etchant: ammonia).

121

A nother interesting observation is the reddish colour of zone b, especially in the vicinity of the S iC /C u-5% N i interface. This observation confirms the finding (see above) that the am oun t of nickel diffused from the interlayer into the silicon carbide com pared to that of copper is m uch larger than w ould be expected on the basis of the ratio of both elem ents in the original interlayer (5Ni:95Cu, wt%).In Fig. 6.12 part of Fig. 6.11 is show n in more detail.

I 110 p m

Fig. 6.12 Detailed view of zone b (bright field, oil im m ersion, etchant: ammonia).

In this figure a num ber of features catch the eye. In the first place, it can be seen that the dim ensions of the precipitates w hich are situated near the ceram ic/in terlayer interface (upper part of the figure) are larger than those situated farther aw ay (lower part of the

figure). Secondly, it is clear that the precipitates w hich are located along the grain boundaries are larger than those located in the grains and that they are elongated in the direction of the grain boundary along w hich they are nucleated. Furtherm ore, it can be seen tha t the precipitates on the grain boundaries are su rrounded by a precipitation free zone. This can be explained by the fact that nucleation on grain boundaries is favoured com pared w ith nucleation in the grains themselves. It also appears that precipitates on grain boundaries penetrate further into the copper-nickel alloy than the precipitates in

122

the grains, indicating that silicon diffuses preferably along grain boundaries.

By m aking use of polarized optical microscopy it is possible to collect m ore inform ation about the precipitates than in the case of bright-field m icroscopy (Fig. 6.13 and Fig. 6.14). In Fig. 6.13 partly polarized light was applied and it can be seen tha t small w hite spots

are presen t w ith in and betw een the larger precipitates. W ith com pletely polarized light (Fig. 6.14) m any m ore w hite spots appear. These spots are sm aller than the precipitates and are for the greater part located w ith in the grains. These sm all w hite spots could not be identified; however, they m ight be copper silicide precipitates.

Fig. 6.13 D etailed view of zone b (partly polarized light, oil im m ersion, etchant: am m onia).

123

I 110 p m

Fig. 6.14 Detailed view of zone b (2000x completely polarized light, oil im m ersion, etchant: ammonia).

Some specim ens w hich were fractured along the copper-nickel / silicon carbide interface, were selected for fu rther analysis by m eans of X-ray diffraction. A specific advantage of specim ens w hich are fractured along the in terlayer/ceram ic interface is that they are quite flat so that no preparation is necessary before an analysis can be perform ed. This is considered beneficial because of the likely presence of carbon (graphite) at the fracture surface w hich disappears w ith grinding and polishing of the surface of the specimens.

The results of the XRD m easurem ents are presented in Fig. 6.15 and lead to the

following observations:the line intensities of Cu-N i (111) and Cu-Ni (200) planes are pronounced, bo th on

the ceramic side and on the m etal side;on the m etal side three "shoulder peaks” are present w hich cannot be identified; these peaks are real in the sense that they represent an existing phase (all specim ens show ed the same phenom enon);on the ceramic side oc-SiC lines are present as is to be expected; furtherm ore, a small b u t reproducible peak is present at 20=26°; this peak is no t detected on the

m etal side and is presum ably due to the presence of carbon (graphite); the o ther peaks on the SiC side w hich cannot be identified w ith certainty are presum ably reflections from one or more silicides or a ternary solid solution.

124

cn a . u

h—cnzLUI—Z

ETo 4b.oTWO - THETA (DEGREES)

80.044.024.0

Fig. 6.15 X-ray diffractogram of the m etal (upper curve) and the ceramic (lower curve) at either side of the C u-N i/S iC interface of a diffusion bonded couple (968°C, 90 m inutes, 7.5 MPa).

It m ust be rem arked in this respect that the penetration dep th of the X-ray beam in both the m etal and the ceramic is about 20-30 pm, w hich m eans that inform ation is only

obtained from the surface layer of the sample.

In addition to XRD m easurem ents, a num ber of m icroanalyses w ere perform ed on the diffusion bonded couples, w ith em phasis on the precipitates w hich w ere form ed in the

copper-nickel interlayer. The results are given in Table 6.1. Also SEM pictures were taken, an exam ple of w hich is given in Fig. 6.16. It appears that the precipitates are preferably present along grain boundaries. It can also be seen that denuded zones are situated next to these grain boundaries w hich confirms the results obtained w ith optical

microscopy.A rem ark needs to be m ade about Table 6.1. The values in the table should be regarded

w ith som e reservation because of the small dim ensions of the precipitates, w hich implies that part of the m atrix m ight have been included in the m easurem ents.

125

Fig. 6.16 SEM picture of the Cu-5%Ni interlayer show ing precipitates in the grains and at the grain boundaries. N ote the denuded zones along the grain boundaries.

Table 6.1 EPMA results of the Cu-5%Ni interlayer (zone b) after diffusion bonding (at.% are norm alized values).

Region Ni (at.%) Si (at.%) C (at.%) Cu (at.%)

precipitate (1) 40.6 23.9 14.4 21.0

precipitate (2) 43.1 24.3 12.3 20.1

precipitate (3) 42.0 22.5 11.1 24.3

precipitate (4) 41.1 22.6 11.9 24.4

m atrix zone b 9.6 5.5 2.6 82.2

m atrix Cu-5%Ni 5.9 - 1.8 92.1

The num bers (1-4) refer to different precipitates.

126

N evertheless, it is evident that significant diffusion of silicon into the copper-nickel

m atrix has taken place, especially along the grain boundaries. The silicon probably precipitated together w ith nickel resulting in the presence of nickel silicide(s). A dditional inform ation can be obtained w hen plotting the data from Table 6.1 in the ternary Cu-Ni- Si phase diagram given in Fig. 2.20, taking into account the obscuring influence of the copper matrix. There is a strong indication that Ni2Si is formed.From linescans w hich w ere m ade, it appears that peaks in both the nickel and silicon concentration correspond w ith dips in the copper concentration. Especially the m irror im ages of the copper and nickel concentrations are rem arkable. This confirms the assum ption that nickel silicides are present (which m ay be enriched in copper), possibly

together w ith a ternary Cu-Ni-Si solid solution.W hen penetrating further into the copper-nickel interlayer (tow ards the AISI316 part), the precipitates become sm aller and, finally, d isappear as is also the case w ith the oscillations in copper and nickel concentration. In this region the original copper-nickel

alloy is unaltered and its com position is 5wt%Ni-95wt%Cu.It is also evident from Table 6.1 that the carbon concentrations cannot be ignored. The relatively high values of the carbon content m ight be due to contam ination of the copper-nickel interlayer during preparation. It is m ore likely, however, tha t the high

carbon content is due to diffusion of carbon atom s along grain boundaries and subsequent segregation in the form of graphite (copper and nickel are know n as

graphitizing elements).

In order to obtain m ore detailed inform ation about both the surface com position of the copper-nickel interlayer w hich w as fractured from the silicon carbide and the

precipitates in this (originally) binary alloy, EPMA was perform ed on several specimens. The results are presented in Fig. 6.17 and lead to the following observations:

the silicon concentration is about 50 at.% in SiC, increases and decreases in zone a and is about 6 at.% in zone b; w hen a precipitate is m easured, the atomic concentration is doubled;the carbon concentration is about 50 at.% in SiC, decreases and increases in zone

a and is about 12 at.% in zone b; w hen a precipitate is m easured, the atomic concentration increases to about 16 at.%;nickel is not p resent in SiC; the nickel concentration shows an irregular course in both zone a and zone b; w hen a precipitate is m easured, the atom ic concentration of nickel is about twice the atomic concentration of silicon;copper is no t p resent in SiC; the copper concentration steeply increases going from zone a to zone b and show s a rather irregular course in zone b.

127

100z o n e a

0 2 4 6 8 10 12 14 16 18 2 0 2 2 2 4 26

d i s t a n c e ( 10~6m )

Si C Ni Cu

Fig. 6.17 EPMA linescans of copper, carbon, nickel and silicon in zone b.

O n the basis of the EPMA results it w as suggested that nickel silicide precipitation takes place in the C u-N i matrix. In order to collect more inform ation on this aspect, AES w as also applied to examine the phenom enon of precipitation in the copper-nickel matrix. M ultiplex AES w as perform ed on both the precipitates and the matrix. It w as found that besides silicon and nickel also copper m ay be present in the precipitates. Like in the case of EPMA, this m ay be due to the fact that part of the m atrix has been included in the m easurem ents.Furtherm ore, it was revealed that silicon is not only present in the precipitates b u t also in the m atrix of zone b, as w as found w ith EPMA. In addition, as in the case of EPMA,

relatively h igh concentrations of carbon w ere m easured (about 10 at.%).W hen com paring the results obtained by EPMA and by AES, it appears that reasonable agreem ent exists. This im plies that presum ably Ni2Si has precipitated and tha t the

precipitates m ay be enriched w ith copper or m ay be em bedded in a ternary solid solution.

128

Zone c

Zone c is located in the copper-nickel alloy near the copper-nickel/AJSI316 interface. The structure of zone c is depicted in Fig. 6.18. The m icrograph shows the presence of precipitates w hich are present both at the grain boundaries and in the matrix. The grain

boundary precipitates penetrate deeper into the copper-nickel than the m atrix precipitates, indicating that grain boundary diffusion is m ore effective than lattice diffusion.

Fig. 6.18 Detailed view of the copper-nickel/ AISI316 interface and zone c containing precipitates along grain boundaries and in the m atrix (bright field, oil im m ersion, etchant: ammonia).

In Table 6.2 the results are given of EPMA carried o u t on precipitates w hich are present in the copper-nickel interlayer near the stainless steel/ copper-nickel interface. The values listed in this table m ust again be considered w ith care because the dim ensions of the

precipitates are very small, typical values of the diam eter being 0.4 pm.

From the table it appears that iron and small am ounts of nickel and chrom ium have diffused into the copper-nickel m atrix, w hich apparently resulted in the form ation of precipitates at the grain boundaries and in the matrix.

Linescans w ere m ade across the interface betw een the copper-nickel and the AISI316 for three specim ens (m ade under different process conditions). The results are presented in Fig. 6.19. From this figure it becomes apparen t that iron and, to a lesser degree, chrom ium and nickel have diffused into the copper-nickel insert while copper has

129

diffused into the AISI316. The m axim um iron solubility in the copper-rich solid solution is about 3.5 at.% Fe at the peritectic tem perature of 1096°C while below the eutectoid tem perature of 850°C/ iron forms large clusters in the copper m atrix [6.4]. The m axim um chrom ium solubility in the copper-rich solid solution is approxim ately 0.89 at.% Cr at 1077°C [6.4] and decreases w ith decreasing tem perature. From these data and the data

m entioned in Table 6.2 it is reasonable to expect that Fe-Cr precipitates have form ed

during cooling from process tem perature to room tem perature.

Table 6.2 Results of EPMA spot m easurem ent of the Cu-5%Ni interlayer (zone c) after diffusion bonding (at.% are norm alized values).

Region Fe (at.%) C r (at.%) C u (at.%) Ni (at.%)

precipitate (1) 7.6 1.8 66.2 6.8

precipitate (2) 3.5 0.7 72.7 5.1

precipitate (3) 2.8 0.6 71.3 5.0

precipitate (4) 3.2 0.9 75.1 5.6

The num bers (1-4) refer to different precipitates in zone c.

100

E 80co

=! 60 oCLEo° 400 uE1 20

o r — S B H f f l s — m ^ — 1— + — + — + — + — + — +- 7 - 5 - 3 - 1 1 3 5 7

d i s t a n c e ( 10~6m )

v F e + C r o Ni u C u

Fig. 6.19 EPMA linescans of four elements diffusing across the Cu-5%Ni/AISI316 interface (from zone c to zone d and vice versa).

□ □ □ □ □ [

V v V V V□

AISI316 C u - 5 % N i

+ + + + + +o O O O o O

—FR- f l- -R --W -—B— 171o + _I_ i-SJ

130

Furtherm ore, an increased concentration of nickel is detected at the interface Cu- 5%Ni/AISI316. This is an indication of up hill diffusion of nickel w hich is presum ably less m obile at the CuN i/A ISI316 interface than iron and chrom ium .Unfortunately, the com position of the precipitates in the copper-nickel alloy (presum ably iron w ith som e chrom ium ) could not be determ ined by m eans of AES because of the

very sm all dim ensions of the particles.

Zone dZone d is located in the austenitic stainless steel near the Cu-5% Ni/ AISI316 interface. It consists of austenitic stainless steel enriched w ith copper originating from the copper- nickel alloy as has already been m entioned in the foregoing section. In addition, this zone show s a slightly reduced iron and chrom ium content, w hich is due to diffusion of these elem ents into the copper-nickel interlayer. The copper is dissolved in the steel matrix for about 4 pm in the case of a couple diffusion bonded at 992°C during 7.5 MPa (the atomic concentration being about 4 at.% near the copper-nickel/steel interface).

131

6.2.3 Kinetics

In C hapter 2 (section 2.5) it w as show n that in m any cases diffusion processes can be described by the equation

d = ^ 2 k pt C6-1)

in w hich d represents the penetration depth, kp the parabolic g row th constant and t the process time, respectively.In order to exam ine the kinetics of the diffusion process occurring during joining HIPSiC and AISI316 using Cu-5%Ni as interlayer m aterial, the grow th of the reaction /d iffusion zone w as m easured as a function of process tem perature and process time.In principle it is possible to use either the thickness dc of the reaction zone a in silicon carbide or the thickness d M of the diffusion zone b in the Cu-5%Ni interlayer m aterial. In other w ords, either the diffusion of nickel and copper into the silicon carbide or the diffusion of silicon into the Cu-5%Ni m ay be used to study the kinetic processes.It m ay be expected that dc and dM are related. In Fig. 6.20 m easured values of d c are

p lo tted versus m easured values of d M.

15

10

5

02001501005 00

Fig. 6.20 Reaction zone thickness in the ceramic (dc) versus the penetration depth of silicon into copper-nickel (dM) for HIPSiC/Cu-5% Ni / AISI316 couples.

132

It appears from this figure that an approxim ately linear relationship exists betw een d c and d M. It w as decided to use dM for tw o reasons. Firstly, as appears clearly from the figure, dM>>dc w hich im plies that the relative error in m easuring d M is sm aller than the relative error in m easuring d c. Secondly, it appeared that d M is easier to m easure than dc due to the less irregular form of the reaction front of zone b com pared w ith the form

of the reaction front of zone a.In order to m easure d M, the specim ens w ere etched w ith am m onium hydroxide and the actual m easurem ent of d M w as carried ou t w ith the aid of optical microscopy. M easurem ents w ere perform ed on specim ens diffusion bonded at process tem peratures of 922°C, 945°C, 968°C, 992°C and 1015°C, respectively and for process tim es up to 24 hours. The results obtained are given in Fig. 6.21 in the form of a d versus Vt plot.

250

200

j= 150io

100~D

50

200 3001000

t ,/2 ( s ,/2)

+ 922-C O 9 4 5 ‘C A 9 68-C V 9 9 2 ‘C ~ 1015'C

Fig. 6.21 Penetration dep th of Si in Cu-5%Ni versus the square root of process time at 922°C, 945°C, 968°C, 992°C and 1015°C.

It appears that for each process tem perature the experim ental points can be approxim ated by a straight line (least squares approxim ation). From this linear relationship, using equation (6.1), it is possible to com pute values of kp. The results of these calculations are presented in Table 6.3.

133

Table 6.3 C alculated parabolic grow th constant kp for HIPSiC/ Cu-5% Ni/ AISI316.

T (°C) kp (10'13m 2/s )

922 2.0

945 2.4

968 5.4

992 5.9

1015 6.6

From Fig. 6.21 it can be clearly seen that the straight lines do not pass th rough the origin. This can have various causes. First of all there is significant scatter in the

experim ental points (indicated by error bars). In addition, it is possible that in the beginning the diffusion process is delayed by the presence of diffusion barriers in the form of an oxide film or a reaction zone w hich results in retarded diffusion (incubation time). Furtherm ore, the reaction betw een elements m ay be rate lim iting w hen a new

com pound is formed. It is also possible that in the beginning of the process grain boundary diffusion plays a m ore im portant role than in a later stage, so that the reaction layer grow th is faster in the beginning than later on (see C hapter 2 for more details).

It appears from Fig. 6.21 that the straight lines intersect the abcissa both at negative and positive values. This strongly suggests that scatter plays the dom inant role and that incubation effects can be neglected. On this basis it is assum ed that at tim e t=0 the

penetration d ep th d=0.From the data presented in Fig. 6.21 the activation energy of the overall process can be

obtained. For this purpose the A rrhenius equation can be used (see C hapter 2):

In

f \d 22 t InD _Q_

RT(6 .2 )

in w hich D0 represents the frequency factor, Q the activation energy of the overall process, R the gas constant and T the absolute tem perature, respectively. The activation

energy w as calculated from the slope of the straight line obtained by plotting the natural logarithm of (d2/2 t) as a function of the reciprocal of the absolute tem perature (Fig. 6.22). It w as decided to p lo t ln(d2/2 t) instead of ln(kp) for the reason m entioned

above. The value of Qact w as calculated to be 281 ± 61 k j/m o le .

134

-2 7

-2 8IN

>x>

c— -2 9

-3 07 8 9

1 0 4/ T ( K~1)

Fig. 6.22 Plot of In (d2/2 t) versus 104/T for HIPSiC/ Cu-5%Ni used to determ ine the activation energy of the diffusion process.

Jackson et al. reported on the reaction of NiCrAl w ith HPSiC in air a t process tem peratures betw een 973 K and 1423 K [6.5]. They calculated the activation energy of

the overall reaction to be 184 k j/m o le . Yamada et al. s tud ied the system N i/S iC and found that the reaction w as diffusion controlled w ith an activation energy of 180 k j/m o le [6.6]. Schiepers stud ied the N i/S iC system in vacuum at process tem peratures betw een 973 K and 1308 K and determ ined an activation energy of 199 k j/m o le [6.7]. The values from literature of the activation energy of the reaction betw een Ni and SiC and the m ean value of the activation energy of the reaction betw een Cu-5%Ni and SiC differ by about 80-100 k j/m o le . This indicates that copper has to be taken into account w hen studying the kinetic behaviour of the Cu-5% N i/SiC system.In fact, it m ight be assum ed tha t there are tw o possible rate determ ining steps in the diffusion process:

the diffusion of silicon through the interlayer; the diffusion of nickel th rough the interlayer.

The first diffusion process has an activation energy of about 225 k j/m o le , w hereas the

second diffusion process has an activation energy of about 236 k j/m o le (see Table 2.1; assum ing that the interlayer consists of copper). O n the basis of these data it cannot be decided w hich of the tw o is the rate determ ining step of the process. In fact, only m arker experim ents can give a decisive answ er about this question.

135

6.2.4 Mechanical behaviour

In order to obtain inform ation about the bond strength, all specim ens w hich were bonded perm anently, w ere subjected to a shear test according to the m ethod described

in C hapter 3. The results of the mechanical strength tests show that the shear strength of the diffusion bonded com bination HIPSiC/ Cu-5% Ni/AISI316 is dependen t on the

process tem perature, the process tim e and the mechanical pressure.The influence of the process tem perature on the shear strength is illustrated in Fig. 6.23.

70

60

oQ_ 50

-C40o>c0)30

20

1000 1025975900 950925

t e m p e r a t u r e (°C)

Fig. 6.23 Shear strength of HIPSiC /Cu-5% N i /AISI316 joints versus process tem perature (p = 7.5 MPa, t = 90 min, high vacuum).

This figure show s tha t there is large scatter in the shear strength values a t various

process tem peratures w hich obscures a clear relationship betw een shear strength and process tem perature. Nevertheless, it appears that there is an optim um in the shear strength w hen p lotted against the process tem perature; this optim um lies in a process

tem perature region around 968°C.The influence of the process tim e on the shear strength of diffusion w elded H IPSiC /C u 5% Ni/AISI316 couples is dem onstrated in Fig. 6.24. Also in this case the experim ental results are characterized by large scatter. It appears tha t there is an optim um in the shear

strength as a function of process time. The figure shows the tendency of the shear strength to rise, up to a process time of 90 m inutes and to decrease again w hen longer process tim es are applied.

136

7 0

60oCL

50

szO ) 4 0c<D!l_ 30(/)k_o 20Q)

_ cCZ) 10

010

o

T01

o

100

t i m e ( m i n )

-e-1000

Fig. 6.24 Shear strength of H IPSiC/Cu-5% Ni/AISI316 joints versus process tim e (p7.5 MPa, T = 968°C, h igh vacuum).

The general trend observed is tha t for both process tem perature and process tim e there is a m axim um of the value of the shear strength. For this reason it w as decided to continue the experim ents w ith optim um process param eters: a process tem perature of 968°C and a process tim e of 90 m inutes, respectively.U nder these conditions the influence of the m agnitude of the m echanical p ressure on the shear strength w as studied. It appears that the shear strength increases w hen the

m echanical pressure is increased from 0 to 7.5 M Pa and that a constant level is reached above a pressure of 7.5 MPa. Also in this case the scatter in shear s trength values is large.Finally, the influence of the thickness of the metallic insert on the shear strength was studied. For this purpose diffusion bonding experim ents were perform ed w ith Cu-5%Ni interlayers having a thickness ranging from 0.2 m m to 1.0 mm. The results of these experim ents clearly indicate that a variation in the thickness of the interlayer in this

range does not have a significant effect on the shear strength of the bonded couples.

All in all, it can be stated that w hen diffusion w elding HIPSiC and AISI316 w ith a Cu- 5%Ni interlayer, best results are obtained w hen the process tem perature T = 968°C, the process time t = 90 m inutes, the m echanical pressure p > 7.5 M Pa and the interlayer thickness d > 0.2 mm.

In o rder to quantify the scatter in the m easured shear strength values of bonded couples

137

produced w ith in the w indow of optim al conditions, the shear strength values can be presented in a W eibull plo t (see C hapter 2). This is done in Fig. 6.25 for joints obtained at a mechanical pressure of 7.5 M Pa and 30 MPa and otherw ise identical conditions (T=968°C, t=90 m inutes, high vacuum). From Fig. 6.25 it can be concluded that:

the average shear strength is similar for bo th mechanical pressures: about 17 MPa in the case of a mechanical pressure of 7.5 MPa, and about 19 M Pa in the case of

a mechanical pressure of 30 MPa;the W eibull m odulus is similar in both cases: 0.8 in the case of p=7.5 M Pa and 0.7 in the case of p=30 MPa; these values illustrate the large scatter in m easured shear strength of diffusion bonded HIPSiC/ Cu-5%Ni/ AISI316 couples; the Weibull m odulus of m onolithic HIPSiC is about 12.

0

- 2

- 32 3 50 4

n t

□ 7 . 5 MPa A 3 0 MPa

Fig. 6.25 W eibull p lo t of the m easured shear strength values of H lPS iC /C u- 5%Ni/AISI316 joints, obtained at a mechanical pressure of 7.5 MPa and 30 MPa, respectively (T=968°C, t=90 minutes).

The mechanical strength of the diffusion bonded couples is due to both chemical and mechanical interaction. Chemical interaction occurs by m eans of various chemical

reactions betw een the copper-nickel alloy and the silicon carbide. In fact, the diffusion of nickel, silicon, copper and carbon results in the form ation of silicides and graphite w hich contributes to the mechanical strength of the joint. The mechanical interaction

138

becomes effective as soon as mechanical pressure is applied on the specim ens to be

bonded. The silicon carbide is pressed against the copper-nickel alloy w hich deform s plastically. Because of the irregularities alw ays present, even on polished surfaces, inden tation of asperities on the ceramic surface into the copper-nickel takes place, w hich leads to a certain degree of mechanical bonding.

Fracture of the diffusion bonded couples occurs either th rough zone a or th rough the ceramic m aterial in the vicinity of zone a. In the latter case nearly alw ays fracture is initiated in zone a and the crack propagates along the w eakest path , that is th rough the silicon carbide. It appears that joints w hich fracture (partly) in the silicon carbide have a higher average mechanical strength than joints w hich fracture m ainly along the interface.

The reason for this behaviour is that in zone a graphite is accum ulated w hich considerably reduces the mechanical strength of the joint. It appears, however, that under certain conditions fracture occurs in the silicon carbide. This im plies that u nder these conditions the mechanical strength of zone a is higher than that of the ceramic m aterial. On the basis of these observations, it m ay be expected that the mechanical

strength of the joints is directly related to the thickness of reaction zone a and, therefore, also to the thickness of diffusion zone b (see Fig. 6.20). In Fig. 6.26 the shear strength is

p lo tted as a function of the thickness of diffusion zone b (i.e. the penetration dep th of silicon in to Cu-5%Ni) in the form of a histogram . The joints are subdiv ided into three categories w ith respect to the shear strength. From this figure it appears that m axim um strength is obtained at values of d M betw een 75 and 100 pm. A t low er values of d M the shear strength values are low er due to insufficient interaction betw een silicon carbide and copper-nickel. A t higher values of d M the shear strength values are low er because of the form ation of too m uch graphite.Thus it can be stated that the mechanical strength of diffusion bonded H lP S iC /C u-

5% N i/ AISI316 couples is lim ited by the form ation of brittle reaction products, especially of graphite in zone a.

In addition to the form ation of brittle reaction products, the m echanical strength of diffusion bonded SiC/Cu-5% N i/A ISI316 couples is also affected by the developm ent of residual stresses.

139

100

V)

c’o

<DXIE3C

80

60

40

20

0 - 5 0 5 0 - 7 5

T > 1 0M P a

7 5 - 1 0 0 100-150d M (1 0 ‘* m )

cm 5<t< 1 o mM P a

>150

t < 5M P a

Fig. 6.26 The shear strength of H IPSiC/Cu-5% Ni/AISI316 couples as a function of the penetration dep th of silicon into the Cu-5%Ni interlayer.

These residual stresses are generated during cooling from process tem perature to am bient tem perature and are prim arily due to large differences in the therm al expansion coefficient of silicon carbide, copper-nickel, austenitic stainless steel and the reaction products w hich are form ed. In the next chapter, m ore attention will be paid to the

occurrance of residual stresses.

140

6.3 Copper -1 0 wt% nickel

6.3.1 Introduction

In addition to the diffusion bonding experim ents carried out w ith Cu-5% Ni interlayer m aterial, a series of diffusion bonding experim ents was also perform ed w ith Cu-10%Ni interlayer m aterial, follow ing the same procedure as described in the relevant section of C hapter 5.

The m echanical strength of the bonded specim ens was tested by m eans of the shear strength test and the structure of the joint w as exam ined by m eans of optical

microscopy, EPMA and XRD. A ttention w as also given to the kinetics of the diffusion process.

6.3.2 Structure of the joint

Analyses of the joint w ere perform ed w ith the help of optical microscopy, X-ray diffraction and electron probe microanalysis. As is to be expected, sim ilar results were

found as in the case of joints w ith a Cu-5%Ni interlayer. The m ain differences are that reaction zone a is larger in size and diffusion b is smaller in size at identical process conditions, w hereas the am ount of carbon in reaction zone a is larger and the precipitates in zone b are also larger in the vicinity of the S iC /Cu-10% N i interface. The latter aspect is clearly visible in the optical m icrograph presented in Fig. 6.27.These differences can be explained by the fact that a double am ount of nickel is present in the Cu-10%Ni m aterial com pared w ith the Cu-5%Ni m aterial. Consequently, the silicon carbide is decom posed faster, resulting in a thicker reaction zone a containing m ore carbon in the form of graphite. The free silicon that is form ed by the decom position of SiC diffuses into the Cu-10%Ni interlayer bu t penetrates less deep into the Cu-10%Ni than in the Cu-5%Ni m aterial due to the enhanced in teraction w ith nickel

(form ation of precipitates). This implies that reaction zone b is sm aller in size, w hereas the density of the precipitates in zone b is larger.The diffusion zones c and d are sim ilar to those found in the case of the Cu-5'X.Ni interlayer as is to be expected.

141

I 140 pm

Fig. 6.27 M icrograph of zone b in diffusion w elded HIPSiC/ C u-10% Ni/ AISI316 show ing large precipitates (T=936°C, t=6 hours, p=5 MPa).

142

6.3.3 Kinetics

A sim ilar procedure as reported for Cu-5%Ni in section 6.2.4 w as follow ed in order to achieve values for kp. The results are given in Fig. 6.28 and Table 6.4.

200

150

100

50

03 0 02001000

t 1/2 ( s 1/2)

+ 9 4 5 ”C O 968'C A 9 9 2 ”C V 1015*C

Fig. 6.28 Penetration dep th of Si in Cu-10%Ni versus the square root of process tim e at 945°C, 968°C, 992°C and 1015°C.

Table 6.4 C alculated parabolic grow th constant kp for H IPSiC/Cu-10% Ni/AISI316.

T (°C) kp (1013m 2/s )

945 1.1

968 3.8

992 5.2

1015 6.2

143

As in the case of Cu-5%Ni interlayers, the penetration dep th of silicon increases linearly w ith the square root of the process time w hereas the kp-values increase w ith increasing process tem perature.Also in this case the average activation energy w as determ ined from the slope of the straight line obtained by plotting the natural logarithm of (d2/2 t) as a function of the reciprocal of the absolute tem perature (Fig. 6.29).

-2 7

^ -2 8 CM

>

C— -2 9

-3 07 8 9

104/ T (K-1)

Fig. 6.29 Plot of In (d2/2 t) versus 104/T for HIPSiC/Cu-10% N i/A ISI316 used to determ ine the activation energy of the diffusion process.

This results in a value of 276 ± 49 k j/m o le . This value is very sim ilar to the value of the activation energy found in the case of the Cu-5%Ni interlayer. This indicates that the

activation energy of the overall process is not significantly influenced by the am ount of nickel in the binary Cu-Ni alloys, at least not in the concentration range up to 10% nickel. A t higher nickel concentrations the situation is very different as has been show n

in section 6.2.3 of this chapter. Furtherm ore, it seems that also in this case either diffusion of nickel th rough copper or diffusion of silicon through copper is the rate determ ining step in the diffusion process (see section 6.2.3).

144

6.3.4 Mechanical behaviour

The perm anen t bonds w hich w ere obtained, w ere tested m echanically according to the procedure described in C hapter 3.It appears that, roughly speaking, m etal/ceram ic combinations diffusion w elded betw een

800°C an d 900°C, have a shear strength w hich is relatively low (m axim um of 5 MPa), w hereas w ith in the range from 900°C to 1030°C the m easured shear strength increases to values of m axim al 57 MPa. It m ust be rem arked, however, tha t above a tem perature

of about 1030°C m elting of a silicide was observed and it is evident tha t this som ew hat obscures the results.As w as also noticed in the case of diffusion bonded specim ens w ith a Cu-5%Ni interlayer, fracture occurred either th rough the reaction zone betw een the ceramic and the m etallic insert (zone a) or through the ceramic itself. Occasionally, fracture w as found to take place in both modes.To check the reproducibility of the strength of diffusion bonded couples u nder optim al w elding conditions, a num ber of bonded specim ens produced at a process tem perature of 1030°C during a process tim e of 90 m inutes w as subjected to a shear test. The results

of these tests are presented in Fig. 6.30 in the form of a W eibull plot.

2

1

WCL

3 o_c_c

- 1

- 20 1 2 3 4

In t

Fig. 6.30 W eibull plot of the m easured shear strength values of H lP S iC /C u- 10%Ni/AISI316 joints, obtained at a mechanical pressure of 7.5 M Pa (T=1028°C, t=90 m inutes, d=0.2 mm).

145

From this figure the Weibull m odulus w as calculated to be about 0.7. A sim ilar value w as found in the case of diffusion bonded couples using a Cu-5%Ni interlayer. H owever, the average shear strength in the case of Cu-10%Ni interlayer (5 M Pa) is m uch low er than in the case of Cu-5%Ni interlayer (17 MPa). The relatively low average of the

shear strength is due to m any diffusion bonded couples w hich have a shear strength

w hich is negligibly small.It appears that also in the case of joints w ith a Cu-10%Ni interlayer a relation exists betw een the shear strength of the diffusion bonded H IPSiC/Cu-10% Ni/AISI316 couples and the fracture m ode. It is clear from the shear test results that, on the average, diffusion bonded couples w hich fracture along the interface (through zone a) have lower

shear strength values than diffusion bonded couples w hich fracture (partly) in the

ceramic material.

W hen the m echanical behaviour of both types of diffusion bonded couples is considered, it is clear that in view of the m echanical shear strength preference should be given to the

Cu-5%Ni m aterial as interlayer.A nother aspect tha t needs to be m entioned is that in the case of Cu-5%Ni interlayers the process tem perature (968°C) a t w hich the best shear strength results can be obtained, is substantially low er than in the case of Cu-10%Ni interlayers (1028°C). In the latter case som etim es m elting w as observed w hich deteriorates the bond.

146

6.4 General model of the SiC/Cu-x%Ni/AISI316 joint

The results p resented in the previous sections lead to the following general m odel of the SiC / Cu-x% Ni/ AISI316 joint.D uring diffusion w elding of the SiC/ Cu-x% Ni/AISI316 com bination, diffusion takes

place across tw o interfaces:the S iC /C u-x% N i interface; the Cu-x% Ni/ AISI316 interface.

This leads to a joint consisting of four zones each w ith its ow n structure and properties.

Silicon carbide decom poses into silicon and carbon through the interaction w ith nickel at elevated tem peratures. Silicon diffuses away into the copper-nickel in term ediate while

carbon rem ains behind for the greater part. As a result of this, two zones (zone a and

zone b, respectively) can be discerned at the ceram ic/in terlayer interface.Zone a is situated in the silicon carbide and is irregular in shape and heterogeneous w ith respect to its chemical composition. It consists m ainly of a diffuse netw ork of carbon in the form of graphite w ith varying concentrations of copper, nickel and silicon in it. The o ther zone (zone b) is located in the copper-nickel alloy and is characterized by the

presence of precipitates w hich are predom inantly situated along the grain boundaries and, to a lesser degree, in the grains themselves; along the grain boundaries denuded zones are present. It appears that these precipitates are nickel silicides, m ost probably N i2Si, possibly enriched w ith copper or em bedded in a ternary Cu-Ni-Si solid solution.

D uring the bonding process, both copper and nickel m igrate into the stainless steel while at the sam e tim e iron, chrom ium and, to a lesser degree, nickel m igrate into the copper- nickel alloy. This results in tw o diffusion zones (zone c and zone d, respectively) w hich are situated at either side of the in terlayer/steel interface.Zone c is situated in the copper-nickel alloy and contains small precipitates w hich m ainly consist of iron and chrom ium , located both at the grain boundaries and in the

grains.

Zone d is located in the AISI316 and appears to consist of a m atrix of austenitic stainless steel w hich is enriched in copper and is degraded in iron and chrom ium .It appears that the grow th of the reaction zones is diffusion controlled and that the activation energy of the diffusion w elding process is 270 ± 50 k j/m o le . The rate determ ining step of the process seems to be either the diffusion of silicon through the copper m atrix a n d /o r the diffusion of nickel th rough the copper matrix.

The strength of the joint is determ ined by the w eakest link, i.e. by zone a in w hich

147

carbon is p resen t in the form of graphite. Fracture of the joint is observed to occur either th rough this zone or th rough the ceramic m aterial in the vicinity of this zone. W hen this latter fracture m ode occurs, in m ost cases the crack is initiated in zone a and then propagates th rough the ceramic material.

It appears th a t m axim al shear strength is obtained for a specific thickness of zone a. This is especially apparent in the case of a Cu-5%Ni interlayer: w hen the thickness of zone

a is betw een 7.5 and 10 pm , shear strength values up to 60 MPa are found. U nfortunately, the scatter in the m easured shear strength values is relatively large

(Weibull m odulus of about 0.8).

148

References

6.1 E.R. M alinow ski and D.G. Howery, Factor Analysis in Chemistry, Wiley, N ew York, 1980.

6.2 D.G. W atson, Surf.Interface Anal., 15 (1990) 516-524.

6.3 T. Piitz, A. Fuchs, H. Ehrhardt, Surf.Sci., 265 (1992) 219-228.6.4 T.B. M assalski (ed.), Binary Alloy Phase D iagram s (2nd ed.), ASM International,

Ohio 1990.6.5 M.R. Jackson, R.L. M ehan, A.M. Davis, E.L. Hall, Met.Trans., 14A (1983) 355-364.6.6 T. Yamada, H. Sekiguchi, H. Okam oto, S. A zum a, A. K itam ura, Proc. 2nd

Int.Symp. on Ceramic M aterials and Com ponents for Engines, W. Bunk and H. H ausner (eds.), Verlag Deutsche Keramische Gesellschaft, Liibeck 1986, p.441-448.

6.7 R.C.J. Schiepers, The interaction of SiC w ith Fe, Ni and their alloys, D issertation,

E indhoven 1991.

149

CHAPTER 7

The role of residual stresses

7.1 Introduction

The reliability of a diffusion bonded ceram ic/m etal couple is dependen t on both the presence of reaction products that are form ed at the ceram ic/m etal interface during the solid state bonding process, and the presence of therm al residual stresses induced during cooling of the joint from process tem perature to room tem perature.As w as dem onstrated in Chapter 4 and Chapter 6, brittle reaction products are form ed during diffusion bonding of SiC/AISI316 and SiC / inter layer /A ISI316, respectively, and fracture occurs th rough these reaction products. In this chapter attention is given to the

residual stresses w hich are expected to be present in the joints.The residual stress state of a (diffusion) bonded couple is influenced by various m aterial properties such as therm al expansion behaviour, yield strength, Young's m odulus and the Poisson constant and also by process conditions like tem perature, time, cooling rate

and, related to the latter, tem perature gradients.The residual stress level of a ceram ic/m etal joint is not easy to m easure and, therefore, m any researchers resort to calculations utilizing m aterial properties and heating procedure. One calculation technique that is often applied, is the so-called finite elem ent m ethod (FEM) w hich can be used to predict the m axim um therm al stress state in (diffusion) bonded ceram ic/m etal couples [7.1-7.8]. However, it appears that, generally speaking, only qualitative predictions can be m ade concerning critical regions in

ceram ic/m etal joints and that reliable solutions are not available except for special cases.

This is caused by the occurrence of so-called stress singularity arising at edges and m aterial transitions, im plying that infinite stresses result from the calculations. N evertheless, calculations were perform ed on the system SiC/AISI316 in order to obtain qualitative inform ation about the location of the critical regions and about the stress distribution in the diffusion bonded couple. The relevant m aterial properties were assum ed to be tem perature independent (thermal expansion coefficients w ere averaged over the tem perature region 20-900°C). Furtherm ore, linear elastic behaviour was presum ed, w hich m eans that the stress levels are actually determ ined by therm al contraction th rough cooling from 1000°C to 20°C.The results obtained in this w ay show close resemblance w ith those obtained by Stoop [7.9], as w as to be expected, and will be briefly sum m arized below.The critical regions (both high tensile stress and high shear stress) w ere found at the

ceram ic/m etal interface, especially near the free surface of the joint. The stresses occur

150

because of the significant difference in therm al expansion coefficient of the m aterials. The tensile stresses are highest at the free surface of the ceramic, im plying that com pressive stresses are present in the m etal part. In practice this m eans that fracture w ill be in itiated at the free surface near the m etal/ceram ic interface . The crack propagates either th rough the ceramic or along the interface w here brittle reaction products are present.

As m entioned earlier, the pursued procedure has the major disadvantage of not properly

dealing w ith stress singularity and this was also encountered in the presen t case. Refining the m esh or introducing m etal plasticity d id not rem edy this singular behaviour.

However, the problem of stress singularity can be ru led out m athem atically by in troducing an other strategy, the so-called linear elastic fracture m echanics (LEFM)

concept, to evaluate the residual stress zones.

In this chapter the calculation of stress intensity factors of diffusion w elded joints using the LEFM approach is presented and discussed both in the case of directly bonded couples and in the case of sam ples bonded w ith the aid of an interlayer. In addition, the influence of crack length, type of interlayer m aterial and interlayer thickness on the stress intensity factors are considered.

Finally, tw o exam ples will be given dem onstrating the influence of geom etry a n d /o r m aterial properties on the residual stress situation. In the first exam ple it is show n that decreasing the diam eter of the interlayer (from 10 to 7 mm) has a beneficial effect on the m echanical strength of the diffusion bond, w hereas in the second exam ple it is show n that in the case of a sym m etrical joint, i.e. a ceram ic/in terlayer/ceram ic joint, the residual stress is significantly low er than in the case of a non-sym m etrical joint (ceramic / in te rlay er/ metal).

151

7.2 Stress state evaluation using linear elastic fracture mechanics

In linear elastic fracture mechanics (LEFM) the characterizing param eter for crack extension is the stress intensity factor. Actually, the stress intensity factor describes the m agnitude of the elastic crack tip stress field. In m ixed m ode LEFM (as can be applied in the case of m ost diffusion bonded ceram ic/m etal couples), the stress intensity factors

are related to loading m odes I, II and III, respectively (I - opening m ode, II = sliding m ode, IE = tearing mode). In addition, the stress intensity factor can correlate crack

grow th and fracture behaviour of m aterials, provided tha t elasticity of the crack tip

stress field is p redom inant [7.10].For plane stra in situations, the following equations describe the stress field near the crack tip:

K,a =

yy\J2 Ttr

i 6 COS—

2

(- . 9 . 381 + s m _ sin—

2 2

K„

\ j lnx

. 6 6 36sin— cos— cos— 2 2 2

(7-1)

K„a =xy

J2nxe

. COS— 2

V

- . 6 . 361 - sin— sin— 2 2

Ki . 6 6 36 sin— cos— cos—2 2 2

(7.2)\]2nx

a = -----— cos— (7.3)zy 2

The sym bols used in these equations are defined in Fig. 7.1.M any studies dealt w ith the issue of calculating stresses around a crack tip at the interface betw een tw o dissim ilar m aterials [7.11-7.19]. Problem s arose w hen a mechanical load w as applied to the bi-m aterial set w ith an interface crack present. This resulted, as

it happens, in oscillations of bo th the stress com ponents and the crack surfaces. It is obvious that, though the solution is correct in a m athem atical sense, in terpenetrating crack surfaces have no physical relevance. How ever, this phenom enon is often d isregarded because this k ind of behaviour only occurs over a region w hich is sm aller

than 10"4 tim es the crack length [7.18].In order to overcom e the oscillatory behaviour of the stress field near the interface crack, a so-called com plex stress intensity factor K* is introduced. This factor is used as a crack tip characterizing param eter [7.12] in order to deal p roperly w ith sm all scale nonlinear

m aterial behaviour (plasticity, for example) a n d /o r sm all scale contact zones at the crack tip (crack face interpenetration) and is defined as:

152

y y

xz

XX zz

LEADING EDGE OF THE CRACK

Fig. 7.1 C oordinate system used to describe the crack tip stress field, including sym bols used in equations (7.1) to (7.3).

K* = K* + iK,: (7.4)

in w hich i=V-l.

A head of the crack located at the interface, the in-plane stresses along the interface are:

0 + 1 0 =yy »y

K - r “ (7.5)

where r represents the distance to the crack tip and

e = — In I n

f \ f \M 1G,

M rG,

1g 7

(7.6)

in w hich G12 represents the shear m odulus. For plane stress:

153

( 3 - v )u. = L

1 (l+Vj)(7.7)

and for p lane strain and axisym m etric geometry:

p = 3 - 4 v , (7.8)“ i

w here y stands for the Poisson constant of m aterial j.It should be kep t in m ind that the factors K* and Kn* are no t identical w ith the com m on opening and sliding m odes b u t that they take into account the oscillatory character of equation (7.5). O nly w hen the tw o m aterials are equal (i.e. w hen e=0) the factors K* and Kt[* becom e identical to K, and K„, respectively [7.17],

The stress intensity factors for the interface crack can be determ ined through

extrapolation of the num erical results of stresses at points at a distance from the crack tip to avoid the oscillation singularity. The complex stress intensity factor K* is obtained as follows (neglecting the oscillatory behaviour of the stress field):

K* = J k *2+K,;2 = lim +CT2V I II r—* 0 • V yy xy(7.9)

Provided that the m aterial interface is a plane, the J-integral is pa th independent and the

relationship betw een this quantity and the stress intensity factors is:

i = h (k “ *k ") (710)

w here

1H

(7.11)

2 cosh2(rce)

and E/=E for plane stress and E/= E /l-v 2 for plane strain and axisym m etric geometry. The relevant data for calculation of the stress intensity factors are given in Table 7.1. U sing the calculated values of J (determ ined through applying the virtual crack extension m ethod [7.20]) and the data listed in Table 7.1, the values for K* can be com puted.

154

Table 7.1 Data of applied materials needed for calculation of stress intensity factors.

M aterial E v a G F F e e(GPa) (10'6K 1) (GPa) (plane

stress)(planestrain)

(planestress)

(planestrain)

SiC 445 0.16 4.95 192 2.448 2.360 0 0

AISI316 199 0.283 19.0 78 2.118 1.868 0.0377 0.0240

Cu 128 0.343 17.7 48 1.978 1.628 0.0518 0.0267

Ni 210 0.31 14.4 80 2.053 1.760 0.0319 0.0147

Cu-5%Ni 130 0.341 17.4 48 1.983 1.636 0.0514 0.0268

From Table 7.1 it clearly em erges tha t the values of e for the various m aterial com binations are quite small w hich indicates that the oscillations in the stress field will be small too. This im plies that the oscillatory part of K* will also be m odest. As a consequence, it is allow ed to apply equations (7.1) to (7.3) p rovided that e (that is, the oscillatory part of K*) approaches zero.

The calculation of K and J was carried ou t by m aking use of a grid containing eight node axisym m etric elem ents, as is show n in Fig. 7.2. The geom etry of the grid w as designed in such a w ay that it corresponded to the geom etry of the m aterial com binations used th roughout the diffusion w elding experim ents.

155

crackedge

A ISI 316

center

a.

Fig. 7.2 FEM grid used to calculate K and J.a. general viewb. detailed view of region around crack

crack edge {]

\ /\ / /

-\ /

-/ \/ \

/ \

I LlSI 3 1 6 <««£$<

iiii\

H U ’ SiC 1

j

b e a = 0 .1 mm

156

First of all, the influence of the m esh w id th on bo th the values of K* (calculated by m eans of extrapolation) and values of K, (derived from the com puted J values) is

evaluated. The results are show n in Fig. 7.3 and it appears tha t the values of the stress intensity factors K* and K; m ay be considered not to depend on the applied m esh w idth.

2 0 0 0 i— -----------------------------------------------------------------------------------------------------------------

+-B -

1000

+- a -

+-B-

+

- E F

10

r e l . m e s h w i d t h

n Kj + K*

Fig. 7.3 C om puted values of K* and Kj plotted versus the relative m esh w idth.

Furtherm ore, it appears from this figure that the stress intensity factor m ay be determ ined by either m ethod, tha t is, th rough extrapolation or by m aking use of the v irtual crack extension m ethod. The great benefit of perform ing stress extrapolation over v irtual crack extension is that the form er provides the possibility of separating K, and

Kn, w hereas the latter approach yields the com bined influence of tensile and shear load on the interface.

157

7.2.1 Computation o f KUI fo r joints w ithout an interlayer

The stress intensity factors K, and K„ w ere com puted for joints w ithout an interlayer by using equations (7.1) and (7.2) and extrapolation to r=0 (that is, to the crack tip). The

m axim um tem perature w as taken to be 1000°C, w hereas the length of the crack was varied betw een 0.1 and 4.1 mm.The results of the calculations are presented in Fig. 7.4 in the form of a K versus crack

length plot.

4 0 0 0

3 0 0 0

E

| 2000

1000

542 30

c r a c k l e n g t h ( m m )

□ K| * K„

Fig. 7.4 C om puted stress intensity factors K, and K„ as a function of crack length for a HIPSiC / AISI316 joint w ithout an interlayer.

From this figure it appears that at each value of the crack length the stress intensity

factor K„ is higher than the stress intensity factor K,. Furtherm ore, it is clear that K, does not vary m uch w ith the crack length, w hereas in the case of m ode II loading, Kn increases w ith the crack length, reaching a m axim um at a crack length of about 2 mm, and then decreases w hen the crack length is further increased.These results clearly indicate that the strength of diffusion bonded SiC/AISI316 couples is dom inated by a shear load rather than a tensile load.

158

7.2.2 Computation o f Kl n fo r joints w ith an interlayer

In order to be able to assess the beneficial role of metallic interlayers on the residual stress state of HIPSiC / AISI316 joints, additional calculations were perform ed w ith Cu, N i and Cu-5%Ni as interlayer material. For this purpose, the grid w hich w as used for

the calculation of the residual stress in direct bonds betw een HIPSiC and AISI316, w as adjusted to include the metallic interlayer. The m axim um tem perature w as taken to be 1000°C and the m aterials data listed in Table 7.1 w ere used. The relevant stress intensity factors w ere com puted by m aking use of stress extrapolation. Similar results were

obtained as in the case of joints w ithout an interlayer.

Calculations w ere also carried ou t to evaluate the influence of the thickness of the

interlayer on the stress level in the joint. In Fig. 7.5 and Fig. 7.6 the results are show n of the calculation of the J-integral and of Kn as a function of the interlayer thickness for a crack length of 0.6 m m in the case of diffusion bonded H IPSiC /A ISI316 couples w ith copper, nickel and copper-5%nickel as interlayer material.These figures show that, for all three interlayer m aterials, both J and Kn decrease w ith increasing interlayer thickness and reach a constant value above a thickness of about 0.5 mm.

159

20

EEz

2 . 5 02.001.00 1.500 .00 0 . 5 0

i n t e r l a y e r t h i c k n e s s ( m m )

+ Cu - “V-- Ni ~ ~ CuNiS

Fig. 7.5 Calculated J-integral as a function of interlayer thickness for H IPSiC /interlayer/A ISI316 joints (crack length: 0.6 mm).

2 000 V

1500 -

10001.50 2.000 . 0 0 0 . 5 0 1.00

i n t e r l a y e r t h i c k n e s s ( m m )

+ Cu v Ni o CuNiS

2 . 5 0

Fig. 7.6 Calculated stress intensity factor K„ as a function of interlayer thickness for H IPSiC /interlayer/A ISI316 joints (crack length: 0.6 mm).

160

7.3 Diffusion bonding using interlayers of reduced size

It is evident that the stress level in the joint is dependent on the geom etry of the m aterial com bination. For instance, it m ay be expected that a reduction of the diam eter of the metallic insert will result in a decrease of the residual therm al stresses (and, therefore,

in an increase of the mechanical strength). In order to find o u t w hether this is the case a series of diffusion w elding experim ents w as perform ed using an interlayer diam eter of 7 mm. A schematic illustration of the applied geom etry is presented in Fig. 7.7.

HIPSiC

AISI316

Fig. 7.7 Schematic illustration of the specim en configuration.

The applied process tem perature w as 1028°C, the diffusion bonding tim e 90 m inutes, the m echanical pressure on the specim ens 15 MPa and w elding w as carried o u t in high vacuum .

It appeared that in all cases perm anent bonds w ere produced. These w ere all subjected to a m echanical shear test yielding an average shear strength of about 20 M Pa w hich is higher than that obtained in the case of an interlayer diam eter of 10 m m (17 MPa; see section 6.2.4).

The shear strength results are presented in the form of a Weibull plot in Fig. 7.8.From this p lo t a W eibull m odulus of 2.4 w as obtained, a value w hich is substantially higher than the value (<1) obtained in the case of bonding experim ents w hich were perform ed under identical circumstances, bu t in w hich the interlayer d iam eter w as 10 m m (see C hapter 6, Fig. 6.25).

161

2

- 2

- 33 42

I n t

Fig. 7.8 W eibull p lo t of HIPSiC/Cu-10% Ni/AISI316 diffusion bonded at 1028°C during 90 m inutes in h igh vacuum under a pressure of 15 M Pa (thickness insert: 0.2 m m , diam eter: 7 mm).

Thus, it m ay be concluded that a reduction in insert diam eter leads to a higher average shear strength and to sm aller scatter in strength values, presum ably due to the reduction in residual stress.

In addition to the experim ents described above, calculations were perform ed on H IPSiC/Cu-10% Ni/AISI316 couples w ith an insert diam eter of 7 m m , using the LEFM approach. The results of these calculations show that apart from the in troduced crack at the in terlayer/ceram ic interface, tw o additional singular regions em erge under these conditions. These singular regions are situated at the sharp transition from ceramic to interlayer and at the sharp transition from interlayer to steel. In order to overcom e this problem , it w as decided to reduce the diam eter of bo th the ceramic and the steel also from 10 to 7 m m in the calculation model. A lthough this geom etry is not equal to the geom etry used in the experim ents, it is considered to be a reasonable approxim ation because during the diffusion w elding process the interlayer will deform plastically and the transitions from interlayer to either the ceramic or the steel w ill in fact be relatively smooth.

162

In Fig. 7.9 the results are presen ted of calculations of K„ as a function of crack length for bo th the case that the diam eter of ceramic, interlayer and m etal is 7 m m and the case that the respective diam eters are 10 mm.It appears that for each crack length the stress intensity factor K„ in the case of an overall diam eter of 7 m m is sm aller than in the case of an overall diam eter of 10 m m , especially w hen the crack length exceeds a value of 0.6 mm, im plying that indeed the stress level is reduced by a reduction in the diam eter of the m etallic insert.

3 0 0 0

2 0 0 0

1000

2 .001 . 5 00.00 0 . 5 0 1.00

c r o c k l e n g t h ( m m )

o d = 1 0 m m A d = 7 m r

Fig. 7.9 Plot of Kn as a function of crack length for HIPSiC/Cu-10% N i/A ISI316 (interlayer thickness 0.2 m m , and diam eter of the m aterials 7 and 10 mm, respectively).

163

7.4 Diffusion bonding silicon carbide to silicon carbide

The fact tha t perm anent bonds can be produced betw een silicon carbide and austenitic stainless steel AISI316 by m aking use of Cu-Ni interlayers, as show n in C hapter 6, suggests that it is also possible to join silicon carbide to itself w ith the help of this

interlayer material.In this section the results are presented of a series of prelim inary experim ents dealing w ith the diffusion w elding of HIPSiC to HIPSiC using a copper-10%-nickel interlayer having a thickness of 0.2 mm.The experim ents w ere perform ed at a process tem perature of 1000°C during 90 m inutes in a h igh vacuum environm ent and under a mechanical pressure of 7.5 MPa. Both the ceramic parts and the metallic inserts w ere prepared as described in C hapter 3. In all

cases a perm anent bond was accomplished. N o cracks w ere observed, neither in the ceramic m aterial nor along the interface.

All specim ens w ere tested m echanically using the equipm ent described in Chapter 3. The average shear strength w as found to be 58 MPa, a value w hich is m uch higher than the value obtained in the case of HIPSiC/ C u-10%Ni/ AISI316 bonds (5 MPa). Fracture always occurred along the interlayer/ceram ic interface and in a few cases along both interfaces, i.e. the interlayer w as detached from both ceramic disks. This is a strong indication that the graphite containing reaction zone is the w eakest zone of the joint. In order to evaluate the reliability of the diffusion bonds, Weibull statistics w as applied. The result of this analysis is presented in Fig. 7.10. The Weibull m odulus calculated from this plot, appears to be 3.2 w hich again is significantly higher than the value obtained in the case of m etal/ceram ic joining w ith the same interlayer m aterial (0.7).The larger strength and higher reliability of the sym m etric S iC /C u-N i/S iC joint

com pared to the asym m etric SiC / C u-N i/A ISI316 joint is due to the low er residual stress level in the sym m etric joint.This was confirm ed by FEM calculations, again m aking use of the LEFM approach. In

Fig. 7.11 K„ is p lo tted against the crack length for both a sym m etric and a asym m etric joint. From this figure it is clear tha t the value of K„ is m uch sm aller in the case of a sym m etric joint com pared to the value of Kn in the case of an asym m etric joint.

164

(N/m

m3/

2)1

0

- 2

- 354

In t

Fig. 7.10 W eibull plo t of diffusion bonded H IPSiC /C u-10% N i/H IPSiC couples at 1000°C du ring 90 m inutes at a pressure of 7.5 MPa in h igh vacuum .

3 0 0 0

2 0 0 0

1000

00 .0 0

V

-

V

V

» ...

.....T

'

+ <3

+

1+

1+

1

0 . 5 0 1.00 1.50

c r a c k l e n g t h ( m m )

s y m m V a s y m m

2 .0 0

Fig. 7.11 Plot of K„ as a function of crack length for H IPSiC /C u-10% N i/ AISI316 and H IPSiC /Cu-10% N i/H IPSiC , respectively (interlayer thickness = 0.2 m m, interlayer diam eter = 10 mm).

165

.5 Conclusions

It is possible to evaluate the stress level in diffusion bonded m etal/ceram ic couples w ith the help of FEM calculations.The calculation approach w hich can be used to predict m axim um shear stress levels in diffusion bonded m etal/ceram ic joints, leads to problem s due to the presence

of singular points at the interface. These problem s can be circum vented by m aking use of the linear elastic fracture mechanics concept.Useful inform ation about the residual stress level of diffusion bonded m etal/ceram ic couples can be obtained by calculation of stress intensity factors (K)

and the J-integral.The results of the calculations show that the residual stress is reduced w hen

m aking use of metallic interlayers betw een SiC and AISI316.The LEFM approach includes only linear elastic effects, w hereas in fact plastic

deform ation takes place.The geom etry of m etal/ceram ic couples plays an im portan t role in the residual stress level. For instance, reduction in the diam eter of m etal/ceram ic couples

results in a low er residual stress level.The residual stress level in sym m etric (S iC /in terlayer/S iC ) joints is low er than the residual stress level in asym m etric (S iC /interlayer/A ISI316) joints.

166

References

7.1 K. Suganum a, T. Okam oto, M. Koizumi, M. Shim ada, J.Am.Ceram.Soc., (1984) C256-C257.

7.2 M. N akam ura, S. Ito, T. Ohji, Y. Hirai, K. Kanayam a, H. Tabata, Yogyo-Kyokai- Shi, 94 (1986) 230-234.

7.3 Ya.V. Lyamin, R.A. M usin, V.N. Ivanov, Autom.W eld., 9 (1986) 13-17.7.4 K. Suganum a, T. Okam oto, K. Kamachi, J.Mater.Sci., 22 (1987) 2702-2706.

7.5 T. Yamada, H. Sekiguchi, H. Okam oto, S. A zum a, A. K itam ura, K. Fukaya, N ippon Kokan Technical Report, O verseas, 48 (1987) 67-74.

7.6 J.R. M cDermid, M.D. Pugh, R.A.L. Drew, A dvanced Structural M aterials, D.S. W ilkinson (ed.), Pergam on Press, N ew York 1989, p .169-177.

7.7 B.T.J. Stoop and G. den O uden, M aterialen, 5 (1989) 39-42.

7.8 T. Yamada, M. Satoh, A. Kohno, K. Yokoi, J.Mater.Sci., 26 (1991) 2887-2892.7.9 B.T.J. Stoop, Diffusion bonding of silicon nitride to austenitic stainless steel, Ph.D.

Thesis, Delft U niversity of Technology, 1991.

7.10 H.L. Ewalds and R.J.H. Wanhill, Fracture M echanics, Delftse U itgevers M aatschappij, Delft 1985, p .28-55.

7.11 K.Y. Lin and J.W. Mar, Int.J.Fract., 12 (1976) 521-531.7.12 J.R. Rice, Trans. ASME, J.Appl.Mech., 55 (1988) 98-103.7.13 K. M izuno, K. M iyazawa, T. Suga, J.Fac.Eng., The U niversity of Tokyo, 34 (1988)

401-412.

7.14 A.K. G autesen and J. D undurs, Trans. ASME, J.Appl.Mech., 55 (1988) 580-586.7.15 T. Suga, S. Schm auder, G. Elssner, Journal de Physique, 49 (1988) C5-539-C5-544.

7.16 R. Yuuki and S.-B. Cho, Eng.Fract.Mech., 34 (1989) 179-188.7.17 P.P.L. M atos, R.M. M cMeeking, P.G. Charalam bides, M.D. Drory, Int.J.Fract., 40

(1989) 235-254.

7.18 R.R. Reynolds, K. Kokini, G. Chen, Trans. ASME, J.Eng.Mat.Techn., 112 (1990) 38- 43.

7.19 C. A tkinson, A dvances in Fracture Research, K. Salaram a et al (eds.), Pergam on, Oxford 1990, p.3053-3061.

7.20 A. Bakker, Int.Journal of Pressure Vessels and Piping, 14 (1983) 153-179.

167

Summary

In this thesis the results are presented of a study dealing w ith joining silicon carbide to austenitic stainless steel AISI316 by m eans of diffusion w elding. W elding experim ents w ere carried out w ithou t and w ith the use of a metallic interm ediate, like copper, nickel and copper-nickel alloys a t various conditions of process tem perature, process time, mechanical pressure and interlayer thickness. M ost experim ents w ere carried ou t in high vacuum . For reasons of com parison, however, some experim ents w ere also carried out in a gas shielded environm ent of 95 vol.% Ar and 5 vol.% H 2.

It appears that w ithou t the use of a metallic interlayer no reliable bonds can be produced betw een silicon carbide and austenitic stainless steel AISI316, neither in a vacuum environm ent, nor in shielding gas. Joints are established a t the process tem perature, bu t these joints fracture in m ost cases during cooling to am bient tem perature. Fracture occurs as a result of the form ation of brittle reaction products and the presence of therm al stresses w hich develop due to the different therm al contraction behaviour of the

ceramic and the metal.Reaction products are form ed bo th in the silicon carbide and in the austenitic stainless steel AISI316. In the silicon carbide, mixed silicides are found of the type (Fe,Ni)2Si and (Fe,Ni,Cr)2Si. In the steel, w eld decay is observed in the form of precipitation of M23C6 along the grain boundaries. D epending on the process conditions either interm etallic com pounds of the %-type or m ixed silicides w ith varying com positions are formed.

It appears that due to diffusion of silicon into the steel a m ixed ferritic /austen itic zone is form ed in the steel. It w as found to be possible to characterize this zone w ith the aid of a m odified Schaeffler-DeLong diagram .A t the original interface of the ceramic and the m etal, periodic layers consisting of carbon and voids w ere detected. These layers are present in a m atrix containing m ainly iron and silicon and, to a lesser degree, chrom ium and nickel. In all cases, fracture occurred either th rough this layer or th rough the ceramic adjacent to this layer.

In o rder to control the diffusion process and to bridge the difference in therm al expansion (and thus to im prove the joint strength and the joint reliability), w elding experim ents w ere also carried ou t using metallic interlayers of different com position. Ni, C u and Cu-x%Ni were selected as interlayer m aterial. It appears that only w ith Cu- x%Ni interlayers, perm anent bonds betw een SiC and AISI316 can be obtained.W hen a nickel interlayer is used, too m uch interaction occurs betw een silicon carbide and nickel, resulting in hard and brittle reaction products like carbon in the form of

graphite and nickel silicides. D uring cooling from process tem perature to room

168

tem perature fracture occurs due to the developm ent of residual stresses. Fracture occurs th rough the layer in w hich the brittle reaction products are present.In the case of a copper interlayer, insufficient interaction betw een copper and silicon carbide takes place for a perm anent bond to be established.

W hen copper-nickel interlayers are applied, perm anent bonds are obtained under specific process conditions. Copper-nickel inserts containing 5 w t% N i yield the best results. The average shear strength of these joints is 17 MPa. The Weibull m odulus of these joints is 0.8, indicating the large scatter in the shear s trength values.

As in the case of direct bonding SiC to AISI316, reaction /d iffusion zones are form ed as a result of diffusion of several constituents of the three materials.Reaction zone a, w hich is form ed in the silicon carbide near the interface w ith the copper-nickel alloy, consists of carbon, copper, nickel and silicon, respectively. C arbon is p resen t in the form of graphite. The d istribution of the elem ents in this zone is very

heterogeneous and irregular as is the interface w ith the unaffected silicon carbide. Reaction zone b, w hich is form ed in the copper-nickel alloy near the interface w ith the silicon carbide, contains precipitates of nickel silicides, situated m ainly along grain boundaries bu t also random ly distributed in the grains. The m atrix of this zone is presum ably a ternary Cu-Ni-Si solid solution.Diffusion zones c and d are situated on either side of the Cu-N i/A ISI316 interface. Zone c contains very sm all precipitates w hich are possibly Fe-Cr clusters, w hereas zone d is a solid solution of copper in AISI316.

It appears that the g row th of the reaction zones is diffusion controlled and that the activation energy of the diffusion bonding process is 270 ± 50 k j/m o le . The rate determ ining step of the process seems to be either the diffusion of silicon th rough the copper m atrix a n d /o r the diffusion of nickel through the copper m atrix.

The mechanical behaviour of SiC /Cu-x% N i/A ISI316 is governed by bo th the residual stress state of the bond and the thickness of zone a (the am ount of carbon in the form

of graphite). The strongest bonds are obtained w hen the thickness of this zone lies betw een 7.5 and 10 pm.

The residual stress level in the joint can be evaluated by m eans of FEM (finite elem ent m ethod) calculations. The m ost reliable results are obtained w hen apply ing a LEFM (linear elastic fracture mechanics) approach. The results of these calculations show that the residual stress level in the joint is significantly reduced by the application of a metallic interlayer.

The m echanical strength of SiC /Cu-x% N i/A ISI316 diffusion bonds can be im proved by reducing the contact area betw een the SiC and the copper-nickel. Results of experim ents

169

using interlayers of reduced diam eter show that both the average shear strength and the

Weibull m odulus increase. This is consistent w ith the results of FEM calculations which dem onstrate that reducing the diam eter of the interlayer m aterial yields low er values of

stress intensity factors and energy release rates.Finally, it is show n experim entally that a sym m etric S iC /C u-10% N i/SiC joint is stronger and m ore reliable than an asym m etric SiC/Cu-10% Ni/AISI316 joint. The sym m etric joint has an average shear strength of 58 MPa and a W eibull m odulus of 3.2, w hereas the asym m etric joint has an average shear strength of 5 M Pa and a Weibull m odulus of 0.7. This difference in strength and reliability is prim arily due to the fact that the residual stress level in the sym m etric joint is considerably sm aller than that in the asym m etric joint, as w as confirm ed by the results of FEM calculations.

170

Samenvatting

In deze dissertatie zijn de resultaten gegeven van een onderzoek betreffende het verbinden van silicium carbide aan austenitisch roestvast staal AISI316 door m iddel van

diffusielassen. Lasexperim enten zijn u itgevoerd zowel m et als zonder gebruikm aking

van een m etallische tussenlaag, zoals koper, nikkel en koper-nikkel legeringen waarbij de procesparam eters tem peratuur, tijd, m echanische druk en tussenlaagdikte zijn gevarieerd. De m eeste experim enten zijn u itgevoerd onder hoogvacuiim . Ter vergelijking zijn echter ook enkele experim enten u itgevoerd onder een bescherm gas m et de

sam enstelling 95 vol.% Ar en 5 vol.% H 2.

H et is niet mogelijk om betrouw bare verbindingen te verkrijgen tussen silicium carbide en austenitisch roestvast staal indien geen gebruik w ordt gem aakt van een m etallische

tussenlaag, noch onder vacuum noch onder bescherm gas. Er w orden verbindingen gevorm d bij de procestem peratuur, m aar deze bezw ijken in de m eeste gevallen tijdens afkoelen naar kam ertem peratuur. Breuk treed t op als gevolg van de vorm ing van brosse reactieprodukten en de aanw ezigheid van therm ische restspanningen die he t gevolg zijn v an het verschil in krim pgedrag tussen de keram iek en het metaal.

Vorming van reactieprodukten v indt plaats zowel in het silicium carbide als in het austenitisch roestvast staal. In het silicium carbide w orden m engsilicides gevorm d van het type (Fe,Ni)2Si en van het type (Fe,Ni,Cr)2Si. In het staal v ind t lasbederf p laats onder vorm ing van MZ3C6 langs korrelgrenzen, en afhankelijk van de procesparam eters w orden of interm etallische verbindingen van het y-type of m engsilicides m et varierende sam enstelling gevorm d.

H et blijkt d a t als gevolg van diffusie van silicium in het staal een gem engde ferritische/ austenitische zone in het staal w ord t gevorm d. De identificatie van deze zone is mogelijk door gebruik te m aken van een aangepast Schaeffler-DeLong diagram .Langs het oorspronkelijk grensvlak van de keram iek en het m etaal w orden periodieke lagen gevonden bestaande u it koolstof en holtes. Deze lagen zijn aanw ezig in een m atrix

die voornam elijk uit ijzer en silicium bestaat en, zij het in m indere m ate, chroom en

nikkel. In alle gevallen v in d t b reuk plaats in deze lagen of in de keram iek grenzend aan deze lagen.

Teneinde de diffusie in te dam m en en het verschil in therm ische expansie te overbruggen (en daarm ee de verbindingssterkte en betrouw baarheid van de verbindingen te verhogen), zijn lasexperim enten gedaan waarbij m etallische tussenlagen van verschillende sam enstellingen zijn toegepast.

D iffusielasexperim enten m et tussenlagen zijn uitgevoerd m et nikkel, koper en koper -

171

nikkel legeringen als tussenlaagm ateriaal. H et blijkt dat alleen in het geval van koper- nikkel legeringen perm anente verbindingen tussen SiC en AISI316 kunnen w orden

verkregen.W anneer een tussenlaag van nikkel w ord t toegepast, blijkt er teveel interactie tussen silicium carbide en nikkel op te treden. Dit resulteert in harde en brosse reactieprodukten zoals koolstof, in de vorm van grafiet, en nikkelsilicides. Tijdens he t afkoelen van p rocestem peratuur naar kam ertem peratuur treedt breuk op als gevolg van de aanw ezigheid van therm ische restspanningen. Breuk v ind t plaats in de laag w aar zich de brosse reactieprodukten bevinden.In het geval van een tussenlaag van koper blijkt er zo w einig interactie plaats te vinden tussen koper en silicium carbide dat er geen perm anente verbinding w ord t gevorm d. Indien tussenlagen v an koper-nikkel legeringen w orden gebruikt, kurm en er wel perm anente verbindingen w orden verkregen. Koper-nikkel legeringen m et een gehalte v an 5 gew.% nikkel leveren de beste resultaten. De gem iddelde afschuifsterkte van deze verbindingen bedraag t 17 M Pa en er w ordt een Weibull m odulus van 0,8 gevonden, hetgeen d u id t op een grote spreiding in de w aarden voor de afschuifsterkte.

Evenals in het geval van het direct verbinden van SiC aan AISI316 w orden, als gevolg v an diffusie, verschillende reactie/d iffusie zones gevorm d.Reactiezone a, die is gevorm d in het siliciumcarbide, bestaat uit koolstof, koper, nikkel en silicium. Koolstof is aanw ezig in de vorm van grafiet. De verdeling van de elem enten in deze zone is bijzonder heterogeen en onregelm atig, hetgeen evenzo geldt voor het

grensvlak tussen deze zone en het onveranderde siliciumcarbide.Reactiezone b, die is gevorm d in de koper-nikkel legering nabij het grensvlak m et het silicium carbide, bevat nikkelsilicide precipitaten die voornam elijk zijn gesitueerd langs korrelgrenzen m aar ook w illekeurig in de korrels aanw ezig zijn. De m atrix van deze zone is verm oedelijk een vaste oplossing van een ternaire Cu-Ni-Si fase.De diffusiezones c en d zijn gelegen aan w eerszijden van het Cu-N i/A ISI316 grensvlak.

Zone c bevat u iterst kleine precipitaten, mogelijk Fe-Cr clusters, terwijl zone d bestaat

u it een vaste oplossing van koper in AISI316.H et blijkt dat de groei van de diverse zones w ord t bepaald door diffusie van verschillende elem enten en dat de activeringsenergie van het diffusielasproces 270 ± 50 k j/m o l bedraagt. Voor de snelheidsbepalende stap van het proces lijken tw ee opties mogelijk: diffusie van silicium door de koperm atrix e n /o f diffusie van nikkel door de koperm atrix.

De m echanische eigenschappen van SiC/Cu-x% N i/A ISI316 verbindingen w orden

beheerst door zowel de restspanningstoestand van de verbindingen als de dikte van

172

zone a (oftewel, de hoeveelheid koolstof in de vorm van grafiet). De sterkste verbindingen w orden verkregen als de dikte van reactiezone a ligt tussen 7,5 en 10 pm. H et niveau van de therm ische restspanningen in de verbindingen kan w orden bepaald m iddels FEM (eindige elem enten m ethode) berekeningen. De m eest betrouw bare resultaten w orden verkregen w anneer een LEFM (lineair elastische breukm echanica)

benadering w ordt toegepast. De resultaten van deze berekeningen tonen aan dat het n iveau van de therm ische restspanningen in de verbindingen significant w ord t verlaagd w anneer een m etallische tussenlaag w ordt gebruikt.De m echanische sterkte van SiC/ Cu-x% Ni/ AISI316 diffusielassen kan w orden verhoogd door het contactoppervlak tussen het SiC en de Cu-x%Ni te verkleinen. R esultaten van

experim enten waarbij gebruik is gem aakt van tussenlagen m et een kleinere diam eter tonen aan dat zowel de gem iddelde afschuifsterkte als de W eibull m odulus toenem en. D it kom t overeen m et de resultaten van FEM berekeningen die aantonen dat verkleining van de d iam eter van de tussenlaag lagere w aarden voor de spanningsintensiteitsfactoren

en "energy release rates" opleveren.Tenslotte is experim enteel aangetoond d a t een sym m etrische SiC /C u-10% N i/SiC verbinding sterker (gem iddelde afschuifsterkte 58 MPa) en betrouw baarder (Weibull m odulus 3,2) is dan een asym m etrische SiC/Cu-10% Ni/AISI316 verbinding (gem iddelde

afschuifsterkte 5 M Pa en Weibull m odulus 0,7).Dit verschil in sterkte en betrouw baarheid is voornam elijk toe te schrijven aan het aanzienlijk lagere niveau van therm ische restspanningen in de sym m etrische verbindingen in vergelijking m et het niveau in asym m etrische verbindingen. D it w ord t

bevestigd door de resultaten van FEM berekeningen.

173

Dankbetuiging

Geen enkel proefschrift kom t tot stand zonder de hu lp van vele m ensen die, elk op hun eigen wijze, een bijdrage hebben geleverd in de periode dat het prom otie onderzoek

w erd uitgevoerd.Allereerst w il ik mijn prom otor G ert den O uden bedanken. Hij heeft mij via zijn

enthousiasm erende en vrijheidsschenkende bonding, ongebreideld optim ism e en opbouw ende kritiek altijd w eten te stim uleren en m otiveren om door te gaan m et het onderzoek.Ik ben Ben T.J. Stoop zeer erkentelijk voor zijn onbaatzuchtige sam enw erking. Hij is een ongem een behulpzam e collega geweest die mij heeft ingew ijd in de boeiende w ereld van

de m ateriaalkunde. Hij is nooit te beroerd geweest vragen van m ijn kant, onder andere betreffende com puters, te beantw oorden en heeft voor mij een groot deel van de

"sommetjes" betreffende restspanningen uitgevoerd.M arcel H erm ans, Jaap Hooijmans, Xiao You H ong, O nno Griebling, Johan Zijp, Ton A endenroom er, W outer Bruins en Erik van Brug w aren college's die het prom ovendusleven op verschillende w ijzen veraangenaam den zowel b innen als buiten

de gangbare w erktijden.A nneke van Veen heeft al het, vaak ondankbare, adm inistratieve w erk secuur uitgevoerd en bijgehouden.De vakkundige en vaak hum orvolle bijdrage van de technici van de sectie Lastechnologie & NDO, Willem Brabander, Frans Bosman en Gijs Kerkhof, m ag niet onverm eld blijven. Evenzo zijn de technici uit de Instrum entm akerij, H arry van Baarle, A ad van der Voort, Jan H erm sen, Nol van der Velden en Jan Vissenberg, zeer behulpzaam geweest. Zij hebben een groot deel van de appara tuu r (op)gebouw d en vele "pillen gedraaid". Zonder hen zou het onderzoek nog veel langer hebben voortgeduurd. G arm t de Jonge heeft m iddels zeer zorgvuldig en m inutieus w erk, daarbij regelm atig tegenslagen trotserend, een substantiele bijdrage geleverd aan de to tstandkom ing van de hoofdstukken 4 en 5. Vooral de interpretatie betreffende de sam enstelling van

gevorm de fasen die gevorm d w aren na diffusielassen, heeft veel duidelijk gem aakt en

de w eg geeffend naar volgende experim enten.Marcel van Vliet heeft via veel experim ented w erk een duidelijk beeld w eten te scheppen betreffende de procesparam eters die nodig zijn om een perm anente m etaal/ keram iekverbinding te verkrijgen m et een acceptabele sterkte w at de voortgang v an het w erk in de eindfase van het prom otie onderzoek heeft bevorderd. Een weergave van zijn w erkzaam heden is gedeeltelijk in de hoofdstukken 5 en 6 terug te vinden.Bij het analyseren van preparaten m et behulp van optische microscopie is de oordeelkundige en im m er bereidwillige hulp van P.F. Colijn onm isbaar geweest.

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W.G. Sloof heeft als hoofd van de oppervlakte-analysegroep bijzonder veel en uiterm ate bru ikbare m etingen verricht, zowel via rontgenm icroanalyse als Auger elektronenspectroscopie. De resultaten h iervan zijn cruciaal gew eest voor een groot deel van het onderzoek en staan w eerspiegeld in de hoofdstukken 4 en 6. Bovendien ben ik hem zeer erkentelijk voor het kritische com m entaar op de hoofdstukken 4, 5 en 6. E.J.M. Fakkeldij heeft belangrijk w erk gedaan op het gebied van de Auger- elektronenspectroscopie, terwijl J. H elm ig SEM w erk heeft verricht en bovendien veel bruikbare opnam en heeft gem aakt w aarvan er een aantal staat afgebeeld in hoofdstuk 6 .

A.H.L.M. Klijnhout heeft de chemische analyse van de u itgangsm aterialen voor zijn

rekening genom en, zowel op nat-chem ische wijze als m et behu lp van rontgenfluorescentie.N.M. van der Pers en J.F. van Lent hebben vele rontgendiffractiem etingen verricht die

een schat aan bruikbare inform atie hebben opgeleverd.Jaap J. "van zuu rp ru im via veenkoloniaal en w.z.b.s. tot M onseigneur" Meijer ben ik dank verschuldigd voor zijn n iet aflatende serviceverlening, op m aat en w el direct. "De m oedige bevrijders van de w itte vlakte" wil ik, zowel in bestuursverband als individueel, bedanken voor het vertier da t zij hebben geboden.M. Lont heeft, m iddels het beschikbaar stellen van een hoogvacuum oven plus stuurkast, ons de m ogelijkheid geboden de eerste diffusielasexperim enten u it te voeren. Vervolgens wil ik dr. F.J.J. van Loo en prof. A. Bakker bedanken voor h u n zinvolle com m entaar op het concept proefschrift.

Tenslotte ben ik het M inisterie van Economische Zaken bijzonder erkentelijk. H et heeft, via het Innovatiegericht O nderzoeksprogram m a (IOP), het onderzoek financieel ondersteund en via sym posia ruchtbaarheid gegeven aan het m etaal/keram iek verbinden in N ederland.

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