FRACTURE OF CERAMIC-TO-METAL INTERFACES

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FRACTURE OF CERAMIC-TO-METAL INTERFACES

G. Elssner, T. Suga, M. Turwitt

To cite this version:

G. Elssner, T. Suga, M. Turwitt. FRACTURE OF CERAMIC-TO-METAL INTERFACES. Journal

de Physique Colloques, 1985, 46 (C4), pp.C4-597-C4-612. �10.1051/jphyscol:1985465�. �jpa-00224718�

JOURNAL DE PHYSIQUE

Colloque C4, suppliment au n04, T o m e 46, avril 1985 page C4-597

FRACTURE O F CERAMIC-TO-METAL INTERFACES

G. Elssner, T. Suga and M. Turwitt

Max-PZanck-Institut ffir MetaZZforschung, Institut far

Stuttgart

F.R.G.

Abstract - The methods used for studying fracture at metal-ceramic interfaces are based on fracture mechanical principles. Sandwich-like specimens con- sisting of two rectangular pieces of ceramic solid-state bonded by a inter- mediate metal layer and notched or precracked at one interface between the two materials are fractured in a three- or four-point bend test device.

The interface fracture energy Gc is obtained from the measured load- deflection curve, the fracture load, the dimensions of the specimen and by use of a correction function YG which pays regard to the elastic pro- perties of the materials involved and the geometry of the specimen.

The fracture behaviour of ceramic-to-metal joints with and without inter- mediate reaction layers is studied. Internal stresses due to differences in the thermal expansion of the constituents are reduced by microcrack formation. For polycrystalline material transitions a mixed failure mode is observed including both adhesive and cohesive fracture contributions.

The fracture energy is influenced by impurity phases and by the micro- structure of the interface region. Experiments with combinations of single crystalline niobium and sapphire indicate that the interface fracture energy depends on the crystallographic orientationship of the two crystal surfaces bonded together. It is shown that the work of adhesion is only a very small fraction of the interface fracture energy. The con- tributions of energy dissipation processes to the fracture energy are discussed in more detail.

I - INTRODUCTION

Solid-state bonded metal-to-ceramic joints consisting of two ceramic bodies and an intermediate metal layer were chosen to study the fracture behaviour of ceramic-to-megal interfaces. They showed two different types of intermediate regions between ceramic and metal. In the majority of cases reaction layers are formed during the high-vacuum diffusion welding at elevated temperatures. Examples of material combinations which show this type A of a intermediate region are SiCINb, SiC/Mo, SiC/Ni, Si3N4/Zr, Si3N4/Hf and Si3N4/Ta. Type B of an abrupt transition between metal and ceramic occurs in diffusion welded combinations of

Article published online by EDP Sciences and available at http://dx.doi.org/10.1051/jphyscol:1985465

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A1203 and Nb. As confirmed by conventional TEM pictures and lattice images of the interface region between sapphire and single crystalline niobium /I/ no inter- mediate reaction layers could be detected due to the fact that alumina and niobium undergo at the welding temperature a solid-solution reaction in which A1203 is dissolved in Nb /2/.

Fig. 1 serves to avoid misunderstandings in the description of fracture phenomena.

It shows the middle part of 'a sandwich-like bend test specimen of height con-

sisting of the materials 1 and 2 with a thin intermediate layer 3. The crack running both within this intermediate layer and in the adjacent regions of the materials 1 and 2 is called an interface crack provided that the transition region of thickness dm in which the crack is running is small in comparison to the speci- men height W. According to that definition the cracks sketched at both sides the transition region are no interface cracks but cracks running in the materials 1 and 2.

Fig. 1 - Specimen composed of the materials 1 and 2 with a thin intermediate layer 3. An interface crack runs within a transition region of thickness

d

W (W

height of specimen). Typical cracks running near the interface

iff material 1 or 2 are also shown. A = interface between 1 and 3. B = interface between 2 and 3.

Although the transition region of a ceramic-to-metal joint consists of a diffusion zone of finite thickness or of thin layers of reaction products its approximation by a sharp interface between two elastic and ideally bonded materials provides a useful tool for the description of bond strength. The interface fracture energy

Gc is a quantitative measure of the bond strength. It can be derived by the

use of a global energy balance for the creating of new surfaces (Fig. 2):

where UE is the internal energy, UK the kinetic energy, r the fracture energy needed for the creating of new surfaces, Q. the heat flux, WD the dissipated work, e.g. associated with plastic deformation processes, Ws and WB work due to J applied surface and body forces. The dots denote differentiation with respect to time.

Fig. 2 - Global energy balance for the crack tip region of an interface crack.

The thermodynamical and mechanical properties of the region ere described by a suitable crack tip model /3/. UE, UK internal and kinetic energy, r fracture energy needed for the creating of new surfaces, Q. heat-flux, W dissipated work, W and WB work due to applied surface and body fokes, Gc interlace fracture e;iergy.

On the condition of an autonomy of the fracture processes and neglecting UK, and WB the interface fracture energy cc can be expressed by

Q

j

where A is the crack surface area. Gc is also called the generalized critical energy release rate / 3 / . The reversible fraction of the interface fracture energy is,

equal to the work of adhesion WA. Therefore, fracture energy measurements can be used to determine the amount of irreversibly dissipated energy accompanied with an interface fracture if the work of adhesion is known.

The bond strength of a ceramic-to-metal interface can also be characterized

by a fracture resistance parameter Kc. If the stress and strain field in the

vicinity of an interface crack tip are known, the energy release rate G can

be calculated in terms of a stress intensity factor

It is important to note

that an unique relationship holds between the energy release rate G and the

absolute value of the complex stress intensity factor K of an interfacial crack

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which i s independent o f the chosen crack t i p model and given by:

According t o equation (4) a f r a c t u r e r e s i s t a n c e parameter Kc can be d e r i v e d from t h e f r a c t u r e energy Gc:

where E* i s an e f f e c t i v e modulus o f e l a s t i c i t y

w i t h p = shear modulus o f m a t e r i a l j ( j

1,2), k .

( 3 - v . ) / ( l + v.) f o r plane

j J J J

s t r e s s ,

= 3 - 4 v . f o r generalized plane s t r a i n and v j = Poisson r a t i o , and

j J

where P i s one o f t h e s o - c a l l e d Dundurs parameters given by

w i t h k = p2/p1. I f t h e crack t i p r e g i o n i s small i n comparison t o t h e t o t a l crack length, t h e i n t e r f a c e f r a c t u r e energy Gc w i l l n o t depend on t h e t y p e o f t h e chosen i n t e r f a c e crack t i p model. Therefore Kc can be taken as a measure o f t h e i n t e r - face f r a c t u r e r e s i s t a n c e w i t h o u t any s p e c i f i c a t i o n o f a crack t i p model.

Three- o r f o u r - p o i n t bend t e s t specimens notched a t one o f t h e metal-ceramic i n t e r f a c e s (Fig. 3 ) are used f o r t h e determination o f t h e i n t e r f a c e f r a c t u r e energy 6,. I f t h e l o a d - d e f l e c t i o n curve i s l i n e a r up t o t h e f r a c t u r e l o a d t h e i n t e r f a c e f r a c t u r e energy 6, i s equal t o I r w i n ' s c r i t i c a l energy r e l e a s e r a t e Gc and can be c a l c u l a t e d from t h e experimental data u s i n g t h e equation

where Fc i s t h e f r a c t u r e load, e h a l f t h e d i f f e r e n c e between t h e upp,er and lower span o f t h e bend t e s t device, B t h e w i d t h and W t h e h e i g h t o f t h e specimen.

The . c o r r e c t i o n YG depends on t h e r a t i o s a/W, d/W and h/W where a i s t h e crack

l e n g t h o r notch depth, d t h e t h i c k n e s s of t h e metal l a y e r and h t h e d i s t a n c e

between notch p o s i t i o n and metal-to-ceramic i n t e r f a c e (Fig. 3), and on t h e Dundurs

parameters

and 8 (equation ( 7 ) ) . YG i s c a l c u l a t e d by means of t h e f i n i t e element

method o r can be obtained by compliance measurements. I f a d e v i a t i o n from l i n e a r i t y

is observed in the load-deflection curve the interface fracture energy Ec must be determined by use of the equation

which relates the generalized critical energy release rate Ec to Irwin's critical energy release rate Gc. The nonlinearity coefficient is obtained from the load- controlled load-deflection curve according to a method introduced by Liebowitz and

Fig. 3 - Layer bonded bend test specimens and equation for the determination of the fracture energy G .

(a) Four-point bend &st, (b) Three-point bend test.

Eftis/3/. Further informations concerning the determination of the interface fracture energy Gc and the interface fracture resistance Kc, the calculation of the correction function YE, and the analysis of interface crack models are given el sewhere / 4 / .

I1 - CONTRIBUTIONS TO THE INTERFACE FRACTURE ENERGY

It seems to be useful to precede the description and discussion of experimental

results by a consideration of the contributions of energy dissipating processes

to the fracture energy. A.G. Evans /5/ proposed the following empirical equation

for the fracture energy GIC of polycrystalline magnesia:

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whereupon GIC i s t h e sum o f t h e surface energy y m u l t i p l i e d by a geometrical f a c t o r Y , t h e energy -yP consumed f o r t h e emission of d i s l o c a t i o n s a t t h e crack t i p region, t h e energy

f o r t h e formation o f secondary cracks and a remaining energy yo. By analogy w i t h t h i s i d e a t h e s i n g l e c o n t r i b u t i o n s t o t h e i n t e r f a c e energy gc can be s p e c i f i e d as given i n Fig. 4. The f i r s t group o f energy c o n t r i b u - t i o n s i s d i r e c t l y r e l a t e d t o t h e i n t e r f a c e crack. pA i s t h e r a t i o o f t h e adhering area t o t h e t o t a l i n t e r f a c e area, YAWA t h e work o f adhesion increased by t h e

i n t e r f a c e roughness and W P t h e energy expended w i t h i n t h e i n t e r f a c e . The con- t r i b u t i o n s of p l a s t i c deformation a r e changed by t h e f a c t o r s aI o r aII i n com- p a r i s o n t o t h e energies o f p l a s t i c deformation o f t h e s i n g l e m a t e r i a l s yI P o r

The second group i n c l u d e s t h e c o n t r i b u t i o n s f o r t h e f r a c t u r e o f par- t i c l e s o f t h e m a t e r i a l s I and I1 and t h e t h i r d group embraces t h e energy con- t r i b u t i o n s f o r t h e f o r m a t i o n o f secondary cracks. For a pure i n t e r f a c e f r a c t u r e o f a ceramic-to-metal t r a n s i t i o n t h e equation given i n Fig. 4 reduces t o t h e r e l a t i o n

i f t h e term W P i s neglected. aIylP and aIIyIIP a r e c o n t r i b u t i o n s o f t h e p l a s t i c deformation energy o f t h e ceramic and t h e metal, r e s p e c t i v e l y .

Fig. 4 - C o n t r i b u t i o n s t o t h e f r a c t u r e energy 6,

The magnitude o f t h e f a c t o r s aI and aII can be estimated by means o f a continuum

mechanical consideration, i f one assumes t h a t t h e r a t i o o f t h e amounts of p l a s t i c

deformation energy d i s s i p a t e d a t an i n t e r f a c e f r a c t u r e and a t a cohesive f r a c t u r e

is equal to the ratio of the sizes of the corresponding plastic zones at the crack tips. The radius of a plastic zone is proportional to the square of the critical stress intensity factor, Kc or KIC2, respectively. If one makes allowance 2 for the locally increased stress at unbonded areas of the interface by an effective stress intensity factor ~ ~ the coefficients aI and aII are given by the 1 6 ~

relations

I11 - MICROCRACKS AND FRACTURE ENERGY

During cooling down from the solid-state bonding temperature thermal stresses are built up in the interface region due to the differences in the thermal ex- pansion coefficients of the ceramic and the metal. They lead to the formation of microcracks which are arranged approximately perpendicular to the interface.

Their influence on the fracture energy was studied for the systems hot-pressed silicon nitride (HPSN) -Zr-HPSN 161 and HPSN-Hf-HPSN by varying the metal layer thickness.

The HPSN-Hf-HPSN joints were solid-state bonded at 1200 "Cllh under a pressure of 10 N/mnz in a high vacuum. Zr-foils of 0.125 to 2 mm thickness were chosen as intermediate metal layers. Some properties and the composition of the materials used for the solid-state bonded joints are given in Table I. Three-point bend test specimens of the dimensions 30 x 5 x 2 mm were cut from the welded pieces.

Notches of width 50 to 100 ym and of depth a were introduced into the metal- to-ceramic interface. The bend tests were carried out at a cross-head speed of 1 ymlmin. Further experimental details are given in 171.

Table I Materials used for solid-state bonded HPSN-Hf-HPSN joints.

Material Composition (wlo') E(GPa) v c ~ ~ ( I o - ~ / K )

-

Hf 97.0 Hf, 2;8 Zr, 141 0.30 6 .O

0.015 0, 0.01 Fe, Ta, C, N

HPSN 97.0 Si3N4, 0.98 Mg, 314 0 -28 2 - 7 0.14 Ca, Y203, Si20N2

E = Young's modulus, v = Poisson ratio, aT thermal expansion coefficient

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Fig. 5 - L i g h t o p t i c a l micrographs o f t h e i n t e r m e d i a t e r e g i o n o f HPSN- Hf-HPSN j o i n t s . A r e a c t i o n l a y e r i s formed between t h e ceramic (upper p a r t o f each micrograph) and t h e metal (bottom).

(a) Thickness o f t h e Hf l a y e r 0.125 mm

( b ) Thickness o f t h e H f l a y e r 0.645 mm. Microcracks induced by thermal stresses a r e v i s i b l e .

5

Fig. 6 - F r a c t u r e energy Gc

N

- and f r a c t u r e r e s i s t a n c e K f o r t h e

n -

i n t e r f a c e f a i l u r e o f s o l i 8 - s t a t e bonded HPSN-Hf-HPSN j o i n t s as a f u n c t i o n o f t h e metal l a y e r t h i c k -

'' ness d. K i s c a l c u l a t e d from t h e

2 measured Ecdata by use o f equation (5).

3

The s o l i d - s t a t e bonded HPSN-Hf j o i n t s a r e c h a r a c t e r i z e d by a t h i n r e a c t i o n l a y e r o f about 5 ym thickness (Figs. 5a, b ) which i s composed o f s i l i c i d e s , n i t r i d e s and t e r n a r y h a f n i u m - s i l i c o n - n i t r o g e n compounds /8/. Furthermore, very t h i n glassy l a y e r s develop d u r i n g t h e welding process a t t h e HPSN i n t e r - face /5/. The formation o f these glassy l a y e r s i s caused by t h e hot-pressing aids added t o t h e HPSN m a t e r i a l . Fig. 5b shows microcracks i n t h e metal p a r t o f a j o i n t o f metal l a y e r thickness d

0.645 mm which t e r m i n a t e a t t h e i n t e r f a c e between HPSN and t h e r e a c t i o n l a y e r . I n Fig. 6 t h e measured i n t e r f a c e f r a c t u r e energy Gc i s p l o t t e d as a f u n c t i o n o f t h e hafnium l a y e r thickness d. The minimum i n f r a c t u r e energy a t a metal l a y e r thickness o f about 0.6 mm i s due t o a f r a c t u r e a t t h e i n t e r f a c e between HPSN and t h e r e a c t i o n l a y e r . This i n t e r f a c e i s r a t h e r weak caused by t h e formation o f t h e a l r e a d y mentioned formation o f glassy l a y e r s , The l e n g t h s o f t h e microcracks increase w i t h i n c r e a s i n g metal l a y e r t h i c k n e s s from 1 ym a t d = 0.125 nnn t o about 20 um a t d

2 m and t h e t i p s o f t h e micro- cracks s h i f t w i t h i n c r e a s i n g d from t h e H f - r e a c t i o n l a y e r i n t e r f a c e over t h e r e a c t i o n layer-HPSN i n t e r f a c e i n t o t h e adjacent regions o f t h e ceramic. A t d = 0.125 mm f r a c t u r e occurs predominantly a t t h e i n t e r f a c e between t h e metal and the r e a c t i o n l a y e r . A t d = 0.39 mm f r a c t u r e paths are observed which a l t e r n a t e between t h e s t r o n g r e a c t i o n layer-metal i n t e r f a c e and t h e weak r e a c t i o n l a y e r - HPSN i n t e r f a c e due t o microcracks w i t h i n t h e r e a c t i o n l a y e r . The increase o f f r a c t u r e energy a t d

1 nnn i n comparison t o t h e f r a c t u r e energy measured f o r j o i n t s w i t h a metal l a y e r thickness o f d

0.64 mm i s caused by microcracks extending i n t o t h e ceramic p a r t o f t h e j o i n t . Thereby f r a c t u r e occurs b o t h i n t h e r e a c t i o n l a y e r and i n t h e adjacent r e g i o n s o f t h e s t r o n g ceramic m a t e r i a l .

The magnitude o f r e s i d u a l thermal stresses i n t h e i n t e r f a c e r e g i o n o f HPSN-Hf combinations can be estimated by t h e s t r a i n mismatch AaT-AT where AaT i s t h e d i f f e r e n c e between t h e expansion c o e f f i c i e n t s o f t h e metal and t h e ceramic and AT i s t h e d i f f e r e n c e between welding temperature and room temperature. On t h e o t h e r hand, t h e r e d u c t i o n o f i n t e r n a l s t r a i n i s given by t h e product N . ^{X where}

N i s t h e number o f microcracks p e r u n i t l e n g t h and X t h e mean opening o f t h e micro- cracks. X was determined by SEM measurements a t a m a g n i f i c a t i o n o f 30 000:l.

Table I1 shows a comparison between N X and AT . AaT. As can be seen agreement between t h e two values i s good. T h i s leads t o t h e conclusion t h a t thermal s t r e s - ses are n e a r l y completely e l i m i n a t e d by microcrack formation.

Corresponding observations were made w i t h HPSN-Zr-HPSN j o i n t s and w i t h HPSN

s o l i d - s t a t e bonded by a t r i p l e l a y e r Zr-Nb-Zr. I n these systems t h e magnitude o f

t h e i n t e r f a c e f r a c t u r e energy Gc i s a l s o s t r o n g l y i n f l u e n c e d by t h e p o s i t i o n o f

t h e t i p s o f t h e microcracks.

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Table 11. Microcrack densities N in sol id-state bonded HPSN-~f-HPSN joints.

Thickness d of N Mean opening x N-X AT AaT * the Hf-'layer of the microcracks

(mm) (l/mn) (vm)

* AaT = difference in the thermal expansion coefficients of Hf and HPSN;

AT = 1170 "C.

IV - EFFECTS OF GRAIN GROWTH AND CRYSTAL ORIENTATION

The studies described in section I11 refer to ceramic-to-metal transitions of type A with intermediate reaction zones and constituents of strongly different thermal expansion behaviour. To simplify matters, combinations of Nb and sapphire were chosen for the study of the influence of metal grain growth and crystal orientation on the interface fracture energy. This system of type B without any intermediate reaction layers is characterized by nearly the same thermal expansion behaviour of its constituents.

Sandwich-like configurations of the sequence polycrystalline alumina-Nb-sapphire- Nb-polycrystalline alumina were solid-state bonded in a high-vacuum at 1700 "C/2h under a pressure of 10 N/mm2. The triple layer Nb-sapphire-Nb consists of three platelets of 1 mm thickness. The dimensions of the pieces of polycrystalline

-

1700 "0 / 2 h

:

-

I--

/

Fig. 7 - Interface fracture energy G of joints between (1100) sapphire afid polycrystalline niobium as a fynction of the niobium grain growth AD.

AD. D,, -q 4.9

otter ond before solid-state bond~ng

0 1

I

0 100 200 300 LOO

Groin growth

AB

alumina attached t o b o t h sides o f t h e t r i p l e l a y e r a r e 14

20 x 15 mm. S i x speci- mens of l e n g t h 31 mm were c u t from each welded piece, notched a t one i n t e r f a c e between Nb and sapphire and f r a c t u r e d i n a f o u r - p o i n t bend t e s t device. F u r t h e r experimental d e t a i l s a r e g i v e n elsewhere /9/.

Polished p o l y c r y s t a l l i n e Nb sheets o f mean i n i t i a l g r a i n s i z e s Db between 13 and 600 ym s o l i d - s t a t e bonded t o (1700) planes o f sapphire p l a t e l e t s a r e used t o determine t h e i n t e r f a c e f r a c t u r e energy GC as a f u n c t i o n o f t h e growth o f t h e Nb g r a i n s d u r i n g welding f o r 2h a t 1700 "C (Fig. 7). The g r a i n growth i s described by t h e d i f f e r e n c e AO between t h e measured g r a i n s i z e Da a f t e r s o l i d - s t a t e bonding and t h e i n i t i a l g r a i n s i z e Eb. F r a c t u r e occurs e x a c t l y a t the Nb-sapphire i n t e r f a c e . The f r a c t u r e energy increases w i t h i n c r e a s i n g g r a i n growth d u r i n g

s o l i d - s t a t e bonding. I t may be assumed t h a t an extensive g r a i n growth i n t h e niobium p a r t o f t h e j o i n t s i s accompanied by an accommodation o f t h e o r i e n t a t i o n and

s t r u c t u r e of t h e surfaces o f t h e metal g r a i n s t o t h e sapphire i n t e r f a c e . Thereby l o w energy c o n f i g u r a t i o n s o f phase boundaries between metal and ceramic may be formed which a r e c h a r a c t e r i z e d by a h i g h i n t e r f a c e f r a c t u r e energy. As shown by f r a c t o g r a p h i c examinations t h e bonding o f two mating surfaces i s p e r f e c t a f t e r extensive g r a i n growth whereas unbonded areas a d j a c e n t t o t h e g r a i n boundaries are observed f o r specimeps w i t h a l a r g e i n i t i a l g r a i n s i z e and o n l y low d i f f e r e n c e s between t h e i n i t i a l and f i n a l g r a i n s i z e . The assumption o f an o r i e n t a t i o n dependence of t h e i n t e r f a c e f r a c t u r e energy i s based on experimental r e s u l t s ob- t a i n e d f o r sapphire-copper /6/ and s i n g l e c r y s t a l 1 i n e Nb-sapphire j o i n t s .

(001) planes and (1T00) planes o f sapphire were s o l i d - s t a t e bonded t o p o l y - c r y s t a l l i n e Nb o f f i n a l g r a i n s i z e 370 ym and t o (110) planes o f s i n g l e c r y s t a l 1 in e Nb. ( 1 i 0 0 ) sapphire- (110) Nb j o i n t s were manufactured both w i t h t h e [0001] and t h e Dl?!] d i r e c t i o n o f sapphire p a r a l l e l t o t h e DOII

d i r e c t i o n o f t h e Nb s i n g l e c r y s t a l . I n Table I11 t h e measured values o f Table I 1 1

Data o f t h e i n t e r f a c e f r a c t u r e energy G C of niobium-sapphire j o i n t s

Sapphire [OOOlI

Sapphire

I l T 0 0 )

Niobium. polycrystalline D = 370 ym

G c = 26 J I m 2

Gc = 5 5 J l m 2

Niobium. s~ngle crystal 1110 1 plone

6,

>

Gc = 7 L ~ l m ~

[0011,,11 [OOO1ls,h,,a Gc = I 1 0 -Ilrn2

[0011,,11 11170

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the i n t e r f a c e f r a c t u r e energy GC o f these Nb-sapphire j o i n t s a r e given. The r e s u l t s i n d i c a t e t h a t GC i s i n f l u e n c e d by t h e t y p e o f c r y s t a l l o g r a p h i c planes mating a t t h e i n t e r f a c e and, f o r a given p a i r o f mating planes, by t h e d i r e c t i o n s a l i g n e d p a r a l l e l t o each o t h e r i n t h e Nb and sapphire planes. The observed dependence o f t h e f r a c t u r e energy may be caused by orientation-dependent energy d i s s i p a t i o n processes d u r i n g f r a c t u r e , by d i f f e r e n c e s i n t h e atomic s t r u c t u r e o f t h e mating surfaces o r by a combination o f both. It seems t o be reasonable t h a t t h e emission o f d i s l o c a t i o n s from t h e i n t e r f a c e crack t i p r e g i o n i s i n - fluenced b y t h e o r i e n t a t i o n o f t h e metal c r y s t a l . Thereby t h e magnitude o f t h e c o n t r i b u t i o n o f p l a s t i c energy t o t h e f r a c t u r e energy may be changed. On t h e o t h e r hand t h e c o n t r i b u t i o n o f p l a s t i c work,

P , i s a f u n c t i o n o f t h e i d e a l work t o f r a c t u r e according t o r e c e n t t h e o r e t i c a l c o n s i d e r a t i o n s on b r i t t l e f r a c t u r e i n deformable s o l i d s / l o / . Therefore, i t can be assumed t h a t t h e main reason f o r t h e o r i e n t a t i o n dependence o f GC i s a change i n t h e work o f ad- hesion MA which i s t h e measure o f t h e i d e a l work t o f r a c t u r e a metal-to-ceramic i n t e r f a c e and i s i n f l u e n c e d by t h e atomic s t r u c t u r e o f t h e i n t e r f a c e . Fig. 8 shows p a t t e r n s o f (1700) s a p p h i r e / ( l l O ) Nb i n t e r f a c e s w i t h t h e d i r e c t i o n s

[ 1 1 2 0 sapphire // [1'11oJ Nb and [1120] sapphire / / [0011 Nb. The a s t e r i k s denote coincidence s i t e s o f oxygen atoms o f t h e (1700) sapphire plane w i t h

0

Nb atoms o f t h e corresponding (110) plane f o r a p e r m i s s i b l e m i s f i t o f 0.33 A.

The p a t t e r n on t h e r i g h t s i d e o f Fig. 8 r e f e r s t o t h e i n t e r f a c e w i t h t h e h i g h e r f r a c t u r e energy.GC

110 J/m2. I t resembles p a t t e r n s o f t h e coincidence s i t e s o f metal-metal i n t e r f a c e s and e p i t a x i a l metal-ceramic i n t e r f a c e s /11/.

Sapphire ( S ) I Nb ~ : : ~ ~ ~ 1 ~ : ~ ~ ~

F i . ⁸ P a t t e r n s of t h e coincidence s i t e s o f Nb an_d oxygen a t ~ m s f o r sapphire (S~IN~ j o i n t s w i t h ( l T O O ) S / ( l l ~ ) N b , 111241s / / [ l l O J N b and (llOO)s/(llO)Nb,

0 I ~ o ] ~ ~ / L ~ ~ ~ ~ ~ ~

V - PURE INTERFACE FRACTURES

A1203-Nb and glass-epoxy joints are two examples of material transitions for which critical crack extension was observed exactly at the interface bet- ween the two materials bonded together. Table IV shows data of the interface fracture energy GC of polycrystall ine A1 203-Nb-A1203 joints sol id-state bonded at 1600 OC/2h and of glass-epoxy-glass joints of 3 mm adhesive thickness to- gether with data of the fracture energy GIC of the corresponding single materials.

The notched specimens were fractured in a four-point bend test device. The fracture energy data are compared with values of the work of adhesion WA for the two material combinations and with the surface energies 2y of the single materials /12,13,14/. WA represents only a small fraction of the interface energy eC of the metal-to-ceramic transition and of the glass-epoxy joint. In both cases the fraction of irreversibly dissipated energy is higher than 98%.

A comparison between the fracture energy values of the single materials and the interface fracture energy values of the combinations leads to the conclusion that the more deformable components, metal or epoxy, respectively, supply the major contribution to the fracture energy of the joint.

Table IV -

Comparison of values of the surface energy and the work o f adhesion with measured values of the fracture energy and the fracture resistance for A1 0 /Nb and glass/

epoxy joints and its bulk materials. The Gc data are obtained byGubcritica1 crack growth experiments /4/.

Fracture resistonce Klc or Kc

(

M N l m 3 i 2 1

3 . 2 5 1 3 . L

2 . 8 8 0 . 6 2 0 . 5 2 0 . 7 1

I

Fraction fi ' of irreversibly dissipated energy

( % I

9 1 . 6 9 9 . 6 9 8 .

8 9 . 0 9 9 . 8 9 9 . 7

Material or materiol combination

AI2O3 (99.7%)

Nb A12 O3 I Nb Glass

Epox Y

Glass/ Epoxy

Surface energy

2 y or work of adhesion

(

Jlm2 1

2 . 1 8 5 . 1 6 0 . 8 2 0 . 5 7 4 0 .

0 . 1 7 8

Fracture energy G~~ or

(

Jlm2)

2 6 . 1

1 1 7 7 . 0

L 3 . L

5 . 2

6 1 . 0

5 8 . 3

C4-6 10 JOURNAL DE

An estimate of the contributions to the interface fracture energy for A1203- Nb joints by means of equations (11) and (12) with p - 0.55, YA = 2 and the data of Table IV leads to pAaIyP

9.1 J/m2, p A a I I y I : ' = 33.2 J/m2, and WA

0 . 8 J/m2 which is in a good agreement with the value of the work of adhesion given in Table IV. The estimate confirms that a relatively large contribution pAaIIyIIP of the interface fracture energy is due to plastic deformation pro- cesses in the metal part. A corresponding calculation for the glass-epoxy joints with

1 and YA = 3 yields pAaIIyIIP = 57.3 J/m2 and WA = 0.3 J/m2 if one neglects the formation of a plastic zone in the glass because of a difference of one order of magnitude between the low yield stress of the epoxy and the relatively high fracture stress of the glass. In the case of Nb-A1203 joints the yield stress of the oxygen-embrittled metal is of the same order of magnitude as the fracture stress of the ceramic.

for A1203/Nb and epoxy/glass

was determined by SEM observation of the fracture surface.

The simple continuum mechanical consideration leading to equation (12) can be used to express GC as a function of MA. According to the equations (5), (11) and (12) the fracture energy eC is given by the relation

(1-v12)yI P (1-vI12 )yII

with a* = +

1 1 G 1 c ~ l

Thus, the interface fracture energy EC depends directly on the work of ad- hesion although its apparent contribution to the fracture energy is negligibly small.

VI - SUMMARY

Small sandwich-like bend test specimens of the configuration ceramic-metal-

layer-ceramic notched at one interface situated in the middle of the speci-

mens are used to study the fracture of ceramic-to-metal transition regions

manufactured by solid-state bonding. Two types of interface fracture occur

in the transition region, pure interface fracture where the critical crack

runs exactly into the ceramic-metal interface and mixed interface fracture

with both adhesive and cohesive contributions. The bond strength of a ceramic-

to metal-transition can be characterized by the interface fracture energy GC.

The experimenta'l determination of EC is based on a global energy balance for the creation of new surfaces. The interface fracture energy is affected by microcracks due to thermal expansion mismatch of the constituents of the joint, by impurity or reaction layer phases in the transition region, by the grain size of the constituents and by unbonded areas in the transition region.

Experiments with sapphire single crystals solid-state bonded to niobium

single crystals show that the interface fracture energy depends on the crystallo- graphic orientation of the interface. This may be caused both by a orientation- dependent atomic structure of the interface accompanied by a change in the work of adhesion and by orientation-dependent energy dissipation processes during fracture.

Generally, the apparent contribution of the reversible work of adhesion, W A Y to the interface fracture energy is very small. For a pure interface fracture of a polycrystalline A1203-Nb joint about 76% of the fracture energy are due to plastic deformation processes in the metal and 21% are due to irreversible energy dissipa- tion processes in the ceramic. It can be deduced from a simple continuum mechanical consideration that the interface fracture energy GC depends directly on the work of adhesion.

VII - REFERENCES

1. M. Riihle, "On comparisons between observed and computed grain boundary structures and energies in metals", paper I 12 presented at this con- ference.

2. M.Florjancic, W. Mader, M. RUhle and M. Turwitt, "HREM and diffraction studies of A1203/Nb interfaces, paper A 13 presented at this conference 3 . H. Liebowitz and J. Eftis, Engng. Fract. Mech. 3(19711 267

4. T. Suga and G. Elssner, "Determination of the fracture energy and the fracture resistance of interfaces, paper D 5 presented at this conference 5. A. G. Evans, Phil. Mag. 22 (1970) 841

6. W. Diem, doctoral thesis, University of Stuttgart (1982)

7. T. Suga and G. Elssner, "Bond strength determination of metal-to- ceramic joints by a bend test method, to be published in Z. Werkstoff- techni k

8. H. Nowotny, B. L u x a n d H . Kudielka,Mh. Chemie87 (1956) 447 9. M. Turwitt. G. Elssner and G. Petzow, "Manufacturing and mechanical

properties of interfaces between sapphire and niobium, paper A12

presented at this conference

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JOURNAL DE PHYSIQUE

10. M. L. Jokl, V. Vitek, C.J. Mc Mahon, Acta Met. 28 (1980) 1479 11. E. Bauer and H. Poppa, Thin Solid Films 12 (1972) 167

12. A. J. Kinloch, R.A. Gledhill, W.A. Dukes in "Adhesion, Science and Technology", L.H. Lee, Edt., Plenum Press, New York (1975) 597 13. E.N. Hodkin, M.G. Nicholas, D.M. Poole, J. Less-Common Metals 20

(1970) 93

14. T. Suga, G. Elssner, in Proc. Symp. on Composite Materials,

DGM, Konstanz, May 1984, in press