c00.52es ELSEVIER Composites: Part A 32(2001)1095-1103 Effect of interphase characteristics on long-term durability of oxide-based fibre-reinforced composites B. Saruhana, *M. Schmucker M. Bartsch. Schneider K. Nubian.G. Wahl German Aerospace Centre, Institute for Materials Research, D-51147 Koln, Germany Technical University of Braunschweig, Institnte of Surface Technology, D-38108 Braunschweig, Germany Received 26 May 2000; revised 26 November 2000; accepted 30 November 2000 Abstract Mullite based fibre-reinforced composites having double layer fibre-coatings were produced and characterised. The multi-layer inter phases were produced by a CVD-process(carbon(fugitive)ZrO2 or Al2O3) on aluminosilicate Nextel 720 fibres. Composites were fabricated by infiltration of coated fibres with a pre-mullite slurry and hot-pressing in argon at 1300C. Short term heat-treatment of composites in air yielded a gap between the fibre and the oxide layer by oxidation of the carbon layer(so-called fugitive layer). The mposites were statically and cyclically heat-treated at 1300.C for 1000 h in order to identify the mechanical and microstructural changes Mechanical characterisation of the heat-treated composites was carried out by three-point bending. The effectiveness of the fugit determined by the oxide layer and its high-temperature stability in interaction with the matrix but it also depends on the loading Under cyclic heat-treatment conditions the composites are found to be more stable and damage tolerant than under constant high e. c 2001 Elsevier Keywords: B Interface/interphase: B. Mechanical properties; E. Chemical vapour deposition(CVD): B. Damage tolerance; Mullite 1. Introduction at very high temperatures, owing to the high diffusion rates of the oxides [21 Relying on the outstanding physical, chemical and ther In this study, two interphase combinations, based on the mal properties, mullite based fibre-reinforced composites multi-layer principle were used. The double layer interphase became favourite candidates for the high-temperature appli- systems, having carbon ZrO2- and carbon Al2O3- cation in the combustion chambers of gas turbine engines. double layers were produced by Chemical vapour deposi- These applications require damage tolerant, temperature- tion(CVD)-coating of Nextel 720 fibres successively and oxidation-resistant ceramic components. In order to The mechanical and microstructural effects of the double realise damage tolerance in ceramics, it is necessary to layer interphase systems were compared with that of a refer employ continuous fibre-reinforcement and possibly an ence composite which had only a carbon(fugitive)-layer at nert interface material (interphase). Thus, a suitable weak the fibre/matrix-interface bonding debonding and sliding at the fibre/matrix- interface can be achieved, which leads to fibre pull-out and conse- quently to the damage tolerance. These properties are to be 2. Materials and methods realised and maintained not only at room temperature, but also at elevated temperatures and over long terms. Non Coating of Nextel 720 fibre fabrics(8 harness Atlas) oxide interphases(e.g BN and C) are structurally suitable, were produced by a CVD-process, with special evaporation however,do not survive under long-term exposure at high and deposition equipment. The starting precursors were le to the lack of oxidation resistance [1]. zirconium and aluminium tetrametylheptandionate(tmhd)4 Among many suggested concepts, the fugitive coating, for oxide coating and propane for carbon coating.Carbon resulting in a gap at the fibre/matrix-interface delivered coating of the woven fibre mats was carried out with pure promising results. One drawback is the closure of the gap propane at 950C under 12.5 mbar pressure. Subsequently the ZrO2 and Al2O3-oxide coatings were vapour-deposited Corresponding author at about 510 and 610.C in a mixture of oxygen and argon E-mail address: bilge saruhan@dlr.de(B Saruhan). under 5 mbar, respectively. After completion of the fibre front matter o 2001 Elsevier Science Ltd. All rights reserved. PI:S1359-835X(01)00016-1
! " # $ !" # $% & % '()( # $% %& '( ')))* & '( "& ')))* +) "& '))) & , - , ./0 1 1&234 5' 6'5+2 " 7') . , , 8+)). , 1 & 2 - , 8+)). 8))) , - & & , 9 '))8 & : 6 & $* ; 8? 6 & 3 & 5 , & , >'? < , , - & 4 5' 6'5+ , . & 1./02 " 7') & - , , , 1&2 3 . " 7') 1@ 6 2 , ./0 , & = - , A 1 2B . ,& , , CD). 8'D = 4 5' 6'5+ , & D8) (8). D & 6 . ; E 6 +' 1'))82 8)CDF88)+ 8+DC@+DG3)83H '))8 & : 6 & E<<; 8+DC@+DG1)82)))8(8 ,,, & 3 3 . + ** , I 1 2
1096 B Saruhan et al. /Composites: Part A 32(2001)1095-1103 Carbon Carbon Al2O3 51251E Fig. 1. Scanning electron micrographs of coated fibres:(a)carbon/ZrO -double coating; and( b) carbon/AlO] double coating. (a) (b Carbon Carbon Zro AlO Fig. 2. Scanning electron micrographs of the interphases after hot-pressing at 1300C for 15 min under 10 MPa pressure. (a) Reference sample b)carbon/ZrO,; and(c)carbon/Al_O3-
8)C( #- & - . , / '0 10))2 )34)' J 8 ; 1 2 34 5' * 12 36'5+ J ' 8+)). 8D 8) E 1 2 % * 12 34 5'* 12 36'5+
B Sarhan ef al./ Composites: Part A 32(2001)1095-110 coating, the thickness of carbon coating was measured to microscope (Leitz LEO 982, Germany)and tested mechani- vary between 100 and 200 nm. The thickness of oxide coat- cally at room temperature by a three-point-bending test wit ings was in turn approximately 800 nm(Fig. 1(a)and (b)). a span of 20 mm, using a UTS-10 testing equipment with a The composites were prepared by aqueous slurry 200N load cell. The displacement in the middle of the infiltration of a submicrometer pre-mullite powder(Siral, 20 mm span was measured with one inductive strain Condea, Germany) into the unidirectional laid double gauge, neglecting the system compliance of the testing coated fibre tows and hot-pressing the composites in argon machine which was estimated to be very small compared at 1300C, for 15 min, under 10 MPa uniaxial pressure. The with that of the test samples. pyrolytic carbon layer was intact after hot-pressing. The omposites were heat-treated at 1200C for 2 h in air in order to obtain a gap between the fibre and the oxide layer 3. Results (so-called fugitive layer). Considering the application temperatures for the composites in combustion chambers, After hot-pressing in argon, the composites contained the composites were heat-treated at 1300C for 1000 h under intact carbon, carbon/monoclinic ZrO2 and carbon/AlO3- continuous-and cycling-heating conditions Thermal cycling interphases (Fig. 2(a)-(c)). The reference composite was carried out by heating up at a rate of 10 K/min to 1300C contained a 200 nm thick carbon interphase, after hot and holding at this temperature for I h before cooling down pressing in argon(Fig. 2a). The thickness of the double to room temperature. This cycle was repeated 1000 times. layer-coating varied after hot-pressing such that the carbon The mechanical and microstructural changes in the layer in carbon/monoclinic Zro2-composite was reduced to opposites were determined by microstructural investiga- 100 nm(initially 180 nm)and that of monoclinic ZrO2-layer tions and mechanical testing. The composites to 400 nm(initially 800 nm). The morphology of characterised microstructurally with a scanning elect clinic ZrO2-layer became somewhat porous(Fig 2b) Fig. 3. Scanning electron micrographs of the interphases after heat-treatment at 1200C for 2 h in air:(a) reference sample:(b)carbon/ZrO2; and (c) carbon/AlO3
, & , 8)) ')) - , @)) 1J 81 2 122 - , = , 1 . # 2 , 8+)). 8D 8) E - , - , 8')). ' , 1 & 2 . , 8+)). 8))) - , 8) !3 8+)). 8 , - , 8))) - , & - , , 1:A :5 C@' # 2 , ') 9-8) = , ')) " - ') , , & , , & , 6 3 4 5' 36'5+ 1J '1 2F122 - ')) 1J ' 2 - & 3 4 5' , 8)) 1 8@) 2 4 5' B)) 1 @)) 2 - 4 5' , 1J '2 #- & - . , / '0 10))2 )34)' 8)C7 J + 8')). ' ; 1 2 * 12 34 5'* 12 36'5+
1098 B Saruhan et al. /Composites: Part A 32(2001)1095-1103 Table 1 Mechanical data of the composites and the thickness of the layers before and after heat-treatment Fiber coating As-hot-pressed Heat-treated at Continuous heat- Cyclic heat- 200°Cfor2 h treated at1300° treated at1300° for 1000 h for1000×1h Fracture strength 160 MPa 42.4 GPa 84.6 GPa Carbon(fugitive)ZrO2 Thickness of C-laye Thickness of oxide-layer 800 400nm 150nm Fracture strength 230 MPa 200 MPa 140 MPa 106 GPa 81 GPa Carbon(fugitive )AlO Thickness of C-layer 180 50 Thickness of oxide-lay 79 MPa 170 MPa 170 MPa 200 MPa Young's modulus 75 GPa 74 5 GPa 91 GPa 92 GPa The carbon layer in carbon/Al2O3-composites after no change compared with that of the as-coated one hot-pressing was about 150 nm, being only slightly thin- (Fig. 2(c)) ner than the thickness of the as-coated layer. The thick Heat-treatment of composites at 1200"C for 2 h in air ness of the AlO3-layer was reduced to 150 nm (initially yielded a gap between fiber and oxide layer by oxidation 800 nm). The morphology of the Al2Oj-layer showed of the carbon layer (fugitive layer)(Fig 3(a)-(c)). This 1300.1000K Fig. 4. Scanning electron micrographs of the interphases after continuous exposure at 1300.C in air for 1000 h: (a)reference sample; (b) fugitive/ZrO2: and (c)fugitive/Al_O3
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
B Sarhan ef al./ Composites: Part A 32(2001)1095-110 process did not cause any warping or cracking of the fibre and causes formation of notch-like damage on fibres composites. The thickness of the fugitive layer in the refer-(Fig. 4(a)) ence composite was significantly smaller than the carbon The gap closure at the interface of the fugitive/ZrO2- yer after hot-pressing (about 80 nm double layer composites was very advanced already after 200 nm), immediately after the oxidation of carbon layer hot-pressing but led to a complete disappearance of the gap at 1200.C/2 h. The thickness of the fugitive and ZrO at 1300C, after 1000 h of continuous heating. The Zro- layers in the fugitive/ZrO2-composites were reduced to layer showed sintering-related morphological changes, result 0 nm and 200 nm, respectively. The thickness of the ing in a rough surface development and thickness reduction fugitive layer in the fugitive/Al2O3-double layer com-(Fig 4(b)). In the fugitive/Al2O3-double layer-composites posites showed no change. The thickness variations of also sintering necks formed after 1000 h continuous heating the interphases as well as numerical mechanical data to at 1300C. However, in this case, the interaction was not the composites before and after heat-treatment are listed only between the fibre and the Al2O3-layer, but also between in Table 1 the matrix and the Al2O3-layer. So that the thickness of the Generally the fugitive layer is thermally unstable and gap was heterogeneous; larger where sintering-related ses progressively as the heat-treatment temperature and contacts between the AlO3-layer and the matrix took duration increases. This is mainly because the matrix as well place, thinner or closed where the fibre and the Al2O as the oxide layer maintains a substantial sintering activity layer were sintered together(Fig 4(c)) after hot-pressing. At the interface of the reference compo- The interfacial relations after cyclic treatment at 1300.C site between the matrix and the fibre substantial sintering differ from those under continuous heat-treatment at the necks form at 1300C after 1000 h continuous exposure in same temperature and same exposure time,(Fig. 5(a)-(c)) air. This interaction between the fibre and the matrix The fugitive layer appeared to be maintained in all compo- educes the amount of alumina on the surface zone of the sites, although to different extents in each composite. The bR m Fig. 5. Scanning electron micrographs of the interphases after 1000 thermal cycles at 1300.C in air:(a) control sample:(b) fugitive/ZrO2; and(c)fugitive/ A12O3
, - & , 1 @) ')) 2 8')).3' - & 4 5' &34 5' , @) ')) & - & &36'5+ , - & , - 8 # & & - , & 6 , 8+)). 8))) - , A 1J B1 22 - &34 5' , & & 8+)). 8))) - 4 5' , & 1J B122 < &36'5+ 8))) 8+)). ,& , , 6'5+ , 6'5+ , * , , 6'5+ , 6'5+ , 1J B122 - 8+)). 1J D1 2F122 - & - #- & - . , / '0 10))2 )34)' 8)CC J D 8))) 8+)). ; 1 2 * 12 &34 5'* 12 &3 6'5+
B Saruhan et al. /Composites: Part A 32(2001)1095-1103 c as hot-pressed, 42. 4 GPa D1200c,2h,27GPa F 1300C, 1000h, cycl. 84 GPa E 1300C, 1000h, const. 84.6 GPa 0,2 04 6 strain [% c as hot-pressed,(106 GPa) stress [GPa] D1200c,2h,(87GPa F1300c,1000h,cyc.(117GPa) H 1300C, 1000h, const(81 GPa)-Different Batch strain [% Fig. 6. (a) Three-point bending curves of the composites with carbon(fugitive)-interphase after(C)hot-pressing, (D)after heat-treat heat-treatment at 1300C for 1000 1 ment at1300°cfor composites with carbon(fugitive)ZrOr-interphase(C)after hot-pressing, (D)after ment at 1200.C for 2 h(H)after 1000 h, cons, and(F) after heat-treatment at 1300C for 1000x I h, cycl. (c) Three-point bending the composites with c iterphase after(C)hot-pressing, (D)after heat-treatment at 1200.C for 2 h(H) after heat-treatment at 1300"C for 1000 h, cons and( F)after heat-treatment at 1300.C for 1000X I h, cycl. ZrOr-layer displayed a rather unique behaviour under cyclic 1200.C after 2 h of exposure. This led to the formation of conditions, yielding a less tighter appearance compared with dense matrix islands and the locally cumulated large pores, the continuously heated specimen(Fig. 5(b)) due to the pore diffusion and growth Under cyclic heat- The mullite matrix in all as-hot-pressed composites was treatment at 1300C, the extension of sintering was less homogeneously packed, but contained many finely distrib- obvious. In these composites, the pore volume distribution uted pores(approx. 40% porosity). Heat-treatment of the remained homogeneous and the size of dense islands wa composites led to localised sintering and an extensive relatively small pore diffusion within the matrix, starting already at The stress/strain behaviour of the composites after
4 5' = & , 1J D122 - , 1 B)M 2 & , 8')). ' - , 9 8+)). , & < & A , & - 3 & 88)) #- & - . , / '0 10))2 )34)' J ( 1 2 - & , 1&2 1.2 102 8')). ' 12 8+)). 8))) 1J2 8+)). 8))) 8 12 - & , 1&234 5' 1.2 102 8')). ' 1 2 8+)). 8))) 1J2 8+)). 8))) 8 12 - & , 1&236'5+ 1.2 102 8')). ' 1 2 8+)). 8))) 1J2 8+)). 8))) 8
Santhan et al. /Composites: Part A 32(2001)1095-110 1101 200 Stress c as hot-pressed, (75 GPa) D1200°c,2h(74.5GPa) 100 1300°c,1000h,cyc.(92GPa) 0 H 1300.C, 1000h, const. (91 GPa) -0.20.00 040.6 1,012141,6 hot-pressing as well as after heat-treatments was determined composite. The Youngs modulus was reduced from 106 by three-point-bending tests of the composites at room to 87 GPa(after 1200C/2 h) and strength from 230 to temperature. For the calculation of the Youngs modulus for 200 MPa. Whereas cyclic prolonged heat-treatment at each composite, the linear elastic part of the stress/strain 1300C resulted in slight increase in the Youngs modulus curves was taken. The stress/strain curves and calculated (117 GPa) but decrease in the strength (170 MPa) Youngs modulus are shown in Fig. 6(a)-(c) Continuous heat-treatment at 1300.C resulted in a severe For the reference composite with initial carbon mono layer decrease of fracture strength (140 MPa). Some load and the carbon/Al -double layer composite, the 1200'C capability after the load maximum was maintained in all heat-treatment resulted in slight changes in ultimate fracture carbon(fugitive)/ZrO2-double layer composites except for strength and Youngs modulus compared with their the prolonged continuously heat-treated composite counterparts as-hot-pressed composite(Fig. 6(a) and(c)) Both hot-pressed and 1200C/2 h heat-treated reference opposites exhibited some load capability during further 4. Discussion straining after reaching the load maximum. In contrast, the prolonged (1000 h) and continuous heat-exposure at For the interpretation of the fibre/matrix-interaction in a 1300.C resulted in the reference composite an increased composite, the Youngs modulus of the composite may be Youngs modulus(84 GPa)and strength(120 MPa)but considered. Assuming that the interphase fulfils the require fully brittle fracture. Cyclic heat-treatment of the composite ment for a good load transfer between the fibre and the at 1300.C for 1000 x I h gave the same elastic modulus and matrix, the Youngs modulus of the composite can be calcu- almost the same fracture strength (160 MPa)as con- lated using the rule of mixtures For our composites, having inuously heat-treated composite, however, it should be 35 vol. %6 Nextel 720-fibers with a mo 260 GPa and 65 vol %o porous matrix with a modulus, Em, apability remains during further straining of the cyclic tested of 46 GPa, Youngs modulus of the composite, EcMc, is sample. Both composites with double layer interphase gave remarkably higher ultimate fracture strength and Youngs modulus, compared to the reference composite(only carbon ECMC=EV+ EmV (1) mono layer)before and after prolonged heat-treatment. ECMc=260GPa×0.35+46GPa×065=12lGPa Somewhat different stress/strain behaviour was observed for the carbon(fugitive )/ZrOr-c omposite Youngs modulus, Er, for the Nextel 720-fibers were (Fig. 6b), compared with the 1200C/2 h heat-treated determined experimentally in a study after heat-treating
, , J KL 3 & , - 3 & KL , J (1 2F12 J , 36'5+ 8')). KL , 1J (1 2 122 8')).3' < 18))) 2 8+)). KL 1@B #E 2 18') E 2 . 8+)). 8))) 8 & 18() E 2 ,& , & KL 1 2 , 3 & , & 1&234 5' 1J (2 , 8')).3' - KL , 8)( @7 #E 1 8')).3' 2 '+) ')) E $ 8+)). KL 1887 #E 2 187) E 2 . 8+)). & 18B) E 2 , 1&234 5' J 3 KL 6 = , KL J & +D &M " 7') , + '() #E (D &M , + B( #E KL +. . ; +. . +5 +5 +. . '() #E )+D B( #E )(D 8'8 #E 8 KL + " 7') , #- & - . , / '0 10))2 )34)' 88)8 J ( 1 *2
B Saruhan et al. /Composites: Part A 32(2001)1095-1103 the fibres at 1300.C for 2 h [3]. The calculated composite extended heat-treatment (170 MPa against 135 MPa, Youngs modulus of 121 GPa should be considered as the respectively) indicates that the load transfer between the maximum value which can be achieved if there is a perfect fibre and the matrix is not due to the substantial sintering load transfer between the fibre and the matrix. If the modu- contact points, which impair the fibre strength by notch-like lus is fully dominated by either fibre or the matrix, then the defects, but more due to dissipative effects in the fugitive composite modulus would have upper bounds, respectively layer(Fig. 5(a) and(b). The maintenance of the fracture where Vm =0 and Ve=0 in 1). This applies to the strength in the composite at all temperatures indicates that composites tested under tensile loading. Since we used the oxide layer has a positive effect on fibre strength by redu- three-point-bending tests to characterise our composites, cing the damages caused by formation of the sintering necks the upper bound, where Vm=0, cannot be applied. This On the other hand, the stress/strain curve of the continuously describes the theoretical case that no matrix contribution heat-treated composite exhibited brittle fracture, indicatin exists for the load transfer. In such a case. the fibres cannot that there was some change in the interfacial relations. This transfer the bending load alone. We can, however, assume that is also supported by the SEM-observation of the composite, there exists no load transfer from the matrix to the fibre, in displaying partial closure of the fugitive layer(Fig. 4(c). other words the fibre does not contribute to the stiffness of the Occurrence of a fibre-contribution into the load transfer composite in bending, then with Ve=0 the composites should become noticeable as the Youngs modulus of the should have a lowest Youngs modulus of around 30 GPa. composite approaches to that of the calculated modulus The experimentally determined Youngs modulus of the (120 GPa). This case is more obvious with carbon(fugi reference composite showed that in the as-hot-pressed refer- tive)/LrO2-composites, where the experimentally deter- ence composite and in the composite where the carbon layer mined composite modulus lies around 106 GPa with the was removed at 1200C(ECMC 42 and 27 GPa, respec- as-hot-pressed and reduces to 87 GPa with the 1200C tively ), there existed almost no interaction between the fibre heat-treated sample. Removal of carbon layer retards the and the matrix. Extended heat-treatment of the reference effectiveness of load transfer in these composites. On the composite at 1300.C yielded a modulus of around 84 GPa, other hand, prolonged heat-treatment at 1300C gives a indicating that a certain load transfer through the sintering- reduced strength compared with the as hot-pressed sample related contact points at the fibre/matrix-interface occurred. as a result of severe fibre damage. The increase of the The higher load transfer between the fibre and the matrix Youngs Modulus to 117 GPa(cyclic heat-treated compo- caused also a simultaneous increase of the strength after the site)indicates strong interfacial bonding. The Youngs 1300.C heat-treatment. Thus, the load transfer is likely due modulus and strength of the continuously heat-treated to the formation of substantial sintering necks(Figs. 4(a) composite could not be compared, since the sample was and 6(a)). The fracture behaviour of the 1300C prolonged from a different batch but brittle fracture behaviour was, heat-treated composite is brittle in the case of continuously in this case too, obvious. This result corresponds with the heat-treated composite. Some remaining load capability in microstructural observation of a totally closed gap between the case of the cyclic heat-treated composite is believed due the fibre and the oxide layer. In the cyclic heat-treated to changes in matrix porosity distribution and formation of sample, presence of a limited degree of damage tolerance less effective sintering points between the fibre and the can be observed(Fig. 6(b))which is believed due to a less matrix which may break easily providing some damage tight bonding between fibre and matrix(Fig. 5(b)) tolerant behaviour By three-point-bending tests, it is diffi- observations can be explained by the repeated ve cult to distinguish clearly, if maintenance of some load (length) change in the ZrOz-layer during cooling and capability is due to dissipation at the fibre/matrix-interface, ing periods of the cyclic heat-treatment so called damage tolerance, or merely due to the delamina In a previous study, Keller et al. [4] have used 20 and tion fracture within the matrix of the composite. Since, 40 nm thick carbon coated Nextel 720-fibers in composites Youngs modulus and strength of the as-hot-pressed and having either a dense CAs(calcium aluminium silicate) 1200.C/2 h heat-treated reference composites are almost glass-matrix or a porous alumina/mullite-matrix. They the same as those of the matrix material we infer that the have demonstrated that the strength retention in porous apparent damage tolerant stress/strain behaviour was merely matrix composites was not dependent on the presence and thickness of a fugitive layer. In other words, presence of a The Youngs modulus of the carbon(fugitive)/Al2O3- fugitive layer in porous matrix composites makes little composites in the as-hot-pressed and 1200.C heat-treated difference to those without an interphase. In the dense samples was reasonably higher (75 GPa)than that of the matrix composites, in turn, the thicker carbon coatin reference composite showing that some fibre/matrix -(40 nm) gave higher strengths and the strength was main- interaction was present. The modulus increased slightly to tained after the removal of carbon. In general they observed 90 GPa for the samples heat-treated at 1300C. Considering that the decrease of strength in dense composites was a the microstructural observations, the measured higher frac consequence of chemical interaction and formation of ture strength of the carbon(fugitive)Al2O3-composites necks between the fibre and the matrix. compared with that of the reference composite after the The reference composite in our study contains a matrix
8+)). ' >+? - KL 8'8 #E & , & , B? & ') B) " 7') & .6 1 2 3 - & , & < , & , < 1B) 2 & , & < & , = , - 88)' #- & - . , / '0 10))2 )34)'
B Sarhan et al. Composites: Part A 32(2001)1095-1103 with relatively high porosity and can be regarded as a from the matrix to the fibre and increases the strength of the porous-matrix composite. The thickness of the carbon(fugi composites significantly (for instance in carbon/zro tive)-layer applied in our study is relatively large(200 nm), composites 230 MPa after hot-pressing and in carbon/ omparing with that of Keller et al. and therefore a lower Al O3-composites 180 MPa compared to 135 MPa of the fibre/matrix-interaction is to be expected(see Fig. 6(a)). ref Despite the presence of a large fugitive layer, the continuous Sintering between fibre and matrix in the reference heat-treatment at 1300.C leads to an almost total closure of sample decreases the damage tolerance as the strength the gap between the fibre and matrix increases in expense of damage tolerance. In the case of In contrast, our carbon(fugitive)ZrOz-interphase com- double layers, the behaviour of the oxide layer and its inter posite yields strength degradation after removal of the carbon ction with the matrix and fugitive layer is decisive for the as well as after heat-treatment at 1300C. This result is characteristics of the interphase. In those composites where comparable with that of the dense matrix composites of Keller a reasonable fugitive layer was maintained, crack deflection et al. where chemical interactions at the fibre/matrix-interface and fibre sliding at the fibre/matrix- interface were probable were held responsible for the strength reduction. We believe The efficiency of the fugitive/oxide-double layer system however, in the case of fugitive/LrO2-interphase composite, depends on the materials and the thermal loading condition the ZrO2-layer sinters and contracts in length heavily, clos-(e. g. cyclic or continuous thermal exposure). This may be ing the gap and causing notch-like defects on the fibre which due to the occurrence of different interactive mechanisms as a result reduce the damage tolerance as well as the within the matrix and at the interface. Under cyclic strength. This behaviour of monoclinic ZrO, was attributed heat-treatment conditions, the composites are found to be to the high thermal expansion coefficient(13 x 10 K ), more stable despite high temperature and relatively long compared to that of mullite(4.5×10-°K-) Cyclic heat-treatment in turn seems to break the sintering-related contact points within the matrix and at the interface during successive cooling down periods, yield References ing better maintenance of the gap and the damage tolerance In the case of ZrOr-interphase, repeated volume changes by reversible monoclinic-tetragonal phase-transformation [1] Chawla KK, Xu ZR, Ha JS, Schmuicker M, Schneider H Appl Compos Mater1997;4:263. under cyclic conditions affected the interfacial relations [2] Hay RS, Petry MD, Keller KA, Cinibulk MK, Welch JR Mater Re resulting in better sliding ability of the fibre Dc Symp Proc1995:365:377-82. [3] Milz C. PhD-study, Characterisation e silicate fibre to be used at high-temperature in fibre- opposites. Shaker 5. Conclusions Verlag, 2000 [4] Keller KA, Mah T, Parthasarathy TA, Cooke CM. Fugitive interfacial Presence of a dense oxide layer behind a carbon (i.e carbon coatings for oxide/oxide composites. J Am Ceram Soc fugitive)-layer plays a supporting role in the load transfer 2000;83(2):329-36
, & - 1 &2 & 1')) 2 , ! , 3 1 J (1 22 0 & 8+)). , 8? . , !! G 4% O 6 . 8CC7*B;'(+ >'? % E 0 ! !6 . ! $ O% % E 8CCD*+(D;+77F@' >+? A . E 0 . / '))) >B? ! !6 - E -6 . . J& 3 O 6 . ')))*@+1'2;+'CF+( #- & - . , / '0 10))2 )34)' 88)+