《复合材料 Composites》课程教学资源（学习资料）第五章 陶瓷基复合材料_SiC-SiC-47

journal of materials ELSEVIER Journal of Nuclear Materials 307-311(2002)1057-1072 www.elsevier.com/locate/jnucma Section 11. Structural ceramics and graphite Promise and challenges of Sic /Sic composites for fusion energy app. lications R.H. Jones a,, L. Giancarli A. Hasegawa, Y. Katoh d ohyama B. Riccardi. LL. Snead w.J. Weber a Pacific Northwest National Laboratory, MS P8-15, P.O. Box 999, Richland, WA 99352, US.A b CEA, Centre d Etudes de saclay, F-9119 Gif sur Yvette cedex, france Tohoku Unirersity, Aoba-ku, Sendai 980-8579, Japan d Institute of Adranced Energy, Kyoto University, Gokasho, Uji, Kyoto 611-00l1, Japan r Oak Ridge National Laboratory, Oak Ridge, TN37831,USA abstract Silicon carbide fiber/silicon carbide matrix composites have been specified in several recent fusion power plant design studies because of their high operating temperature (1000-1100oC)and hence high energy conversion efficiencies Radiation resistance of the B-phase of Sic, excellent high-temperature fracture, creep, corrosion and thermal shock resistance and safety advantages arising from low induced radioactivity and afterheat are all positive attributes favoring the selection of Sicr/SiC composites. With the promise of these materials comes a number of challenges such as their thermal conductivity, radiation stability, gaseous transmutation rates, hermetic behavior and joining technology Re- cent advances have been made in understanding radiation damage in Sic at the fundamental level through MD sim- ulations of displacement cascades. Radiation stability of composites made with the advanced fibers of Nicalon Type s and the UBE Tyranno SA, where no change in strength was observed up to 10 dpa at 800C, in the development of materials with improved thermal conductivity, modeling of thermal conductivity, joining techniques and models for ife-prediction. High transmutation rates of C and Si to form H, He, Mg, and Al continue to be a concern. c 2002 Elsevier science bv all rights reserved 1. Introduction led to their being considered in the TAURO, arIes and dream power plant designs. Sicr/Sic composites offer the promise of a high Challenges for these materials include their thermal temperature fusion reactor design because of the radia- conductivity, radiation stability, gaseous transmuta- tion resistance of the cubic, B-phase SiC matrix, their tion rates, hermetic behavior and joining technology excellent high-temperature fracture, creep, corrosion Their radiation stability is dominated by the differential and thermal shock resistance and safety advantages swelling between the Sic fibers, that are not fully dense arising from their low induced radioactivity and after- or crystalline, carbon interphases and B Sic matrices. heat. Also, developments for other applications, such as Within limits, these materials can be engineered to have aerospace, have driven improvements in material per- select properties; therefore, much of the research in formance that are beneficial to fusion applications. understanding their behavior and improving their per These positive attributes of Sic/Sic composites have formance has been focused on this aspect of their character. An overview of new understanding of the radiation behavior of Sic and SiCr/Sic composites Corresponding author. Tel +1-509 376 4276: fax:+1-509 will be given. This includes fundamentals of radiation 3760418 damage in SiC, advances in new composite materials E-mail address: rh. jones@pnl. gov (R H. Jones). experiments and modeling of thermal conductivity 0022-3115/02/. see front matter e 2002 Elsevier Science B v. All rights reserved PI:S0022-3115(02)00976-5

Section 11. Structural ceramics and graphite Promise and challenges of SiCf/SiC composites for fusion energy applications R.H. Jones a,*, L. Giancarli b , A. Hasegawa c , Y. Katoh d , A. Kohyama d , B. Riccardi e , L.L. Snead f , W.J. Weber a a Pacific Northwest National Laboratory, MS P8-15, P.O. Box 999, Richland, WA 99352, USA b CEA, Centre d
Etudes de Saclay, F-9119, Gif sur Yvette cedex, France c Tohoku University, Aoba-ku, Sendai 980-8579, Japan d Institute of Advanced Energy, Kyoto University, Gokasho, Uji, Kyoto 611-0011, Japan e ENEA-CR Frascati, via E. Fermi, 27, I00044 Frascati (Roma), Italy f Oak Ridge National Laboratory, Oak Ridge, TN 37831, USA Abstract Silicon carbide fiber/silicon carbide matrix composites have been specified in several recent fusion power plant design studies because of their high operating temperature (1000–1100
C) and hence high energy conversion efficiencies. Radiation resistance of the b-phase of SiC, excellent high-temperature fracture, creep, corrosion and thermal shock resistance and safety advantages arising from low induced radioactivity and afterheat are all positive attributes favoring the selection of SiCf /SiC composites. With the promise of these materials comes a number of challenges such as their thermal conductivity, radiation stability, gaseous transmutation rates, hermetic behavior and joining technology. Re￾cent advances have been made in understanding radiation damage in SiC at the fundamental level through MD sim￾ulations of displacement cascades. Radiation stability of composites made with the advanced fibers of Nicalon Type S and the UBE Tyranno SA, where no change in strength was observed up to 10 dpa at 800
C, in the development of materials with improved thermal conductivity, modeling of thermal conductivity, joining techniques and models for life-prediction. High transmutation rates of C and Si to form H, He, Mg, and Al continue to be a concern. 2002 Elsevier Science B.V. All rights reserved. 1. Introduction SiCf/SiC composites offer the promise of a high￾temperature fusion reactor design because of the radia￾tion resistance of the cubic, b-phase SiC matrix, their excellent high-temperature fracture, creep, corrosion and thermal shock resistance and safety advantages arising from their low induced radioactivity and after￾heat. Also, developments for other applications, such as aerospace, have driven improvements in material per￾formance that are beneficial to fusion applications. These positive attributes of SiCf/SiC composites have led to their being considered in the TAURO, ARIES and DREAM power plant designs. Challenges for these materials include their thermal conductivity, radiation stability, gaseous transmuta￾tion rates, hermetic behavior and joining technology. Their radiation stability is dominated by the differential swelling between the SiC fibers, that are not fully dense or crystalline, carbon interphases and b SiC matrices. Within limits, these materials can be engineered to have select properties; therefore, much of the research in understanding their behavior and improving their per￾formance has been focused on this aspect of their character. An overview of new understanding of the radiation behavior of SiC and SiCf /SiC composites will be given. This includes fundamentals of radiation damage in SiC, advances in new composite materials, experiments and modeling of thermal conductivity, Journal of Nuclear Materials 307–311 (2002) 1057–1072 www.elsevier.com/locate/jnucmat * Corresponding author. Tel.: +1-509 376 4276; fax: +1-509 376 0418. E-mail address: rh.jones@pnl.gov (R.H. Jones). 0022-3115/02/$ - see front matter 2002 Elsevier Science B.V. All rights reserved. PII: S 0 0 2 2 - 3 1 1 5 ( 0 2 ) 0 0 9 7 6 - 5

RH. Jones et al. I Journal of Nuclear Materials 307-311(2002)1057-1072 ransmutation rates, chemical compatibility a which are an anticipation of successful future R&D. The sion, irradiation creep and crack growth, join most significant assumptions are the following ology and thermal and mechanical transient Recent experimental results and mechanistic and mod- (i The Sic/Sic thermal conductivity at 1000C and materials. predictions show promise of new, improved at end-of-life(EOL)conditions is 20 W/m K. This value is considerably higher than that shown by present-day SiC/SiC. In fact, ent available data on existing industrial 3D Sicr/SiC indicate 2. Current design possibilities and needs at 1000C a value of 15 W/m K in the plane and of 9 W/m K through the thickness [1], without tak- Silicon carbide composites (SiCr/SiC), are being ing into account the effect of irradiation which are considered in future fusion power reactors because their expected to decrease by about a factor 3 the out high temperature properties (=1000C), offer the po of-pile value tential of very high energy conversion efficiency (50% or (in The maximum and minimum acceptable tempera- more). SiCr/SiC composite has been proposed as struc- tures are respectively 1100C and 600C; these tural material for the first wall and blanket in several values need to be confirmed under irradiation for conceptual design studies EOL conditions (iii)Compatibility between Pb-17Li and Sicr/Sic is ac- ceptable at 800C; this statement should be valid 2.1. Proposed blanket concepts after irradiation and at Pb-17Li velocity of few m/s and should also be valid for any brazing mate- The most recent proposals are TAURO in the Eu- ropean Union, ARIES-AT in the United States, and rials in contact with Pb-17Li: available dat firm a good compatibility for static Pb-17Li at DREAM in Japan. The first two concepts are Pb-17Li 800oC for 3000 h self-cooled blankets, while Dream is cooled by 10 MPa (iv) Use of preliminary SicrSiC models and design Helium [1]. Both TAURO and Aries-AT blankets are criteria are not yet validated by experiments: mod- essentially formed by a SiC/Sic box with indirectly els and criteria currently used for metals and de- cooled Fw that acts as a container for the pb-17Li fined in industrial design codes (.g, ASME, which has the simultaneous functions of coolant tritium RCC-MR) are not applicable for Sic/sic struc- breeder, neutron multiplier and, finally, tritium carrier tures Because of the relatively low SiC/sic electrical con (v) The electrical conductivity of Sic /Sic is about 500 ductivity, high Pb-17Li velocity is allowed without Q2m this value would allow sufficiently low needing large coolant pressures (<1.5 MPa). TAURO MHD effects for self-cooled Pb-17Li blanket and blanket is characterized by 2m-high single modules correspond to the presently measured out-of-pile which are reinforced by SiCr-SiC stiffeners ARIES-AT data. This result could, however, be jeopardized is characterized by a coaxial Pb-17Li flow, which occurs by Pb-17Li infiltration in the top layer of SiC: in two 8 m-high boxes inserted one into the other. The SiC: this infiltration could dramatically increase DREAM blanket is characterized by smaller modul the wall electrical conductivity which could quickly (0.5 m of height), each divided in three zones: FW reeding zone and shield; neutron multiplier material become unacceptably high. A SiC coating on SiCr/ Sic is probably sufficient to avoid this kind of effect (Be), tritium breeding material (LiO or other lithium (vi) Acceptably low coolant leakage in case of He-cool- ceramics)and shielding material(Sic)are packed in the g(10 MPa of pressure); very low quantity of He module as small size pebbles of I mm-diameter for Be tolerated in the plasma so SiCr/Sic hermeticity and Li2O, and 10 mm for SiC. The He coolant path need to be ensured by a reliable coating which includes a flow through the pebble beds and a porous should have the same irradiation resistance of the partition wall. These blankets allow very high coolant main structures; no experimental results are ye outlet temperatures and therefore a high energy con version efficiency. The maximum coolant outlet tem- (vii) Possibility of manufacturing relevant shapes with perature is 1100C obtained in the Pb-17Li of the appropriate thickness ranging between I and 6 ARIES-AT blanket which lead to a thermal efficiency of mm. Present requirements appear achievable in 58.5% present day industrial composites; however, mate- rial properties in these conditions need to be exper 2.2. Main assumptions for blanket de imentally verified. vili)Existing methods of joining finite components The TAURO. ARIEs-AT and dream designs have with characteristics similar to the base material been performed g optimistic SiCr/Sic properties good results are already available

transmutation rates, chemical compatibility and corro￾sion, irradiation creep and crack growth, joining tech￾nology and thermal and mechanical transient behavior. Recent experimental results and mechanistic and mod￾eling based predictions show promise of new, improved materials. 2. Current design possibilities and needs Silicon carbide composites (SiCf /SiC), are being considered in future fusion power reactors because their high temperature properties (’1000
C), offer the po￾tential of very high energy conversion efficiency (50% or more). SiCf /SiC composite has been proposed as struc￾tural material for the first wall and blanket in several conceptual design studies. 2.1. Proposedblanket concepts The most recent proposals are TAURO in the Eu￾ropean Union, ARIES-AT in the United States, and DREAM in Japan. The first two concepts are Pb–17Li self-cooled blankets, while DREAM is cooled by 10 MPa Helium [1]. Both TAURO and ARIES-AT blankets are essentially formed by a SiC/SiC box with indirectly￾cooled FW that acts as a container for the Pb–17Li which has the simultaneous functions of coolant, tritium breeder, neutron multiplier and, finally, tritium carrier. Because of the relatively low SiCf /SiC electrical con￾ductivity, high Pb–17Li velocity is allowed without needing large coolant pressures (<1.5 MPa). TAURO blanket is characterized by 2m-high single modules which are reinforced by SiCf–SiC stiffeners. ARIES-AT is characterized by a coaxial Pb–17Li flow, which occurs in two 8 m-high boxes inserted one into the other. The DREAM blanket is characterized by smaller modules (0.5 m of height), each divided in three zones: FW, breeding zone and shield; neutron multiplier material (Be), tritium breeding material (Li2O or other lithium ceramics) and shielding material (SiC) are packed in the module as small size pebbles of 1 mm-diameter for Be and Li2O, and 10 mm for SiC. The He coolant path includes a flow through the pebble beds and a porous partition wall. These blankets allow very high coolant outlet temperatures and therefore a high energy con￾version efficiency. The maximum coolant outlet tem￾perature is 1100
C obtained in the Pb–17Li of the ARIES-AT blanket which lead to a thermal efficiency of 58.5%. 2.2. Main assumptions for blanket designs The TAURO, ARIES-AT and DREAM designs have been performed assuming optimistic SiCf /SiC properties which are an anticipation of successful future R&D. The most significant assumptions are the following: iiii(i) The SiCf /SiC thermal conductivity at 1000
C and at end-of-life (EOL) conditions is 20 W/m K. This value is considerably higher than that shown by present-day SiCf/SiC. In fact, present available data on existing industrial 3D SiCf/SiC indicate at 1000
C a value of 15 W/m K in the plane and of 9 W/m K through the thickness [1], without tak￾ing into account the effect of irradiation which are expected to decrease by about a factor 3 the out￾of-pile value. iii(ii) The maximum and minimum acceptable tempera￾tures are respectively 1100
C and 600
C; these values need to be confirmed under irradiation for EOL conditions. ii(iii) Compatibility between Pb–17Li and SiCf/SiC is ac￾ceptable at 800
C; this statement should be valid after irradiation and at Pb–17Li velocity of few m/s and should also be valid for any brazing mate￾rials in contact with Pb–17Li; available data con- firm a good compatibility for static Pb–17Li at 800
C for 3000 h. ii(iv) Use of preliminary SiCf /SiC models and design criteria are not yet validated by experiments; mod￾els and criteria currently used for metals and de- fined in industrial design codes (e.g., ASME, RCC-MR) are not applicable for SiCf /SiC struc￾tures. iii(v) The electrical conductivity of SiCf /SiC is about 500 X1 m1; this value would allow sufficiently low MHD effects for self-cooled Pb–17Li blanket and correspond to the presently measured out-of-pile data. This result could, however, be jeopardized by Pb–17Li infiltration in the top layer of SiCf/ SiC; this infiltration could dramatically increase the wall electrical conductivity which could quickly become unacceptably high. A SiC coating on SiCf / SiC is probably sufficient to avoid this kind of effect. ii(vi) Acceptably low coolant leakage in case of He-cool￾ing (10 MPa of pressure); very low quantity of He is tolerated in the plasma so SiCf /SiC hermeticity need to be ensured by a reliable coating which should have the same irradiation resistance of the main structures; no experimental results are yet available to give indication about this requirement. i(vii) Possibility of manufacturing relevant shapes with appropriate thickness ranging between 1 and 6 mm. Present requirements appear achievable in present day industrial composites; however, mate￾rial properties in these conditions need to be exper￾imentally verified. (viii) Existing methods of joining finite components with characteristics similar to the base material; good results are already available. 1058 R.H. Jones et al. / Journal of Nuclear Materials 307–311 (2002) 1057–1072

RH. Jones et al. I Journal of Nuclear Materials 307-311(2002)1057-1072 59 x) Acceptable structure lifetime in terms of sumed as the maximum allowable compressive stress fluence and plasma-first wall interaction;a since no damage is observed under compression. The cant experimental campaign is required to number of fibers through the thickness of the composite the sice/Sic limits is usually lower and their arrangement different. So, if one can accept the uncoupling of stresses in plane and 2.3. SiC /SiC thermo-mechanical model and design crite stresses through the thickness, the former can be eva rId uated using the von mises criteria, while the latter corresponds to the measured rupture value. Taking into The proposal for using SiC/SiC as structural mate- account the above remarks, for compressive stresses the rial for a nuclear component with a reasonably long limit is the rupture limits while for tensile stresses the lifetime is fairly recent, therefore no adequate modeling limit is the elastic limit. a margin of about 20% may be and design criteria are available yet. Some preliminary allowable on the elastic limit. work has recently been performed [2] aiming both to For example, in the case of the TaURo blanket identify appropriate models available in the aerospace based on the CERaSEP@ composites produced by research field and to theoretically define sound design SNECMA, the following limits have been assumed criteria to improve the design thermo-mechanical ana for normal stresses through the thickness, 110 MPa for tensile stresses(roughly corresponding to the ma- 2.3.1. Modeling trix tensile resistance limit outside the composite)and SiCr/Sic composites exhibit a complex nonlinear 420 MPa for compressive stresses(rupture limit for behavior combining brittle damage, residual strains CERASEP N2-1) and opening-closing of microcracks. These composites for shear stresses through the thickness, 44 MPa(as- present different properties, and therefore different sumed rupture limit, to be confirmed ) strengths, for different loading directions; moreover, for stresses in plane, 145 MPa for tensile stresses(be- tensile and compression strengths are very different. ginning of fiber/matrix debonding) and 580 MPa for Under loading, the interaction of fibers and matrix lead compressive stresses(rupture limit measured on the at first to matrix microcracking. then to matrix/fiber CERASEP N2-1). It appears clear that this design decohesion, followed by opening of the microcrack and criteria proposal is very preliminary; it needs to finally to fiber failure. This sequence corresponds to an further evaluated both theoretically and, more im- initial isotropic behavior in plane and then to a crack portant, experimentally through a systematic specific growth perpendicular to the fibers depending on the load direction. A relatively simple model, able to take into account such a behavior and based on continuum damage mechanics which consider the composites as a 3. Promise of Sicr/Sic composites continuous media, has been implemented in the FEM code CASTEM. Significant improvements on the results Composite materials made from continuous fibers of have been obtained when compared with models used SiC, can be woven into several variant fabric architec- for metals [1]. On the other end, damage description has tures and the matrix formed with a variety of infiltration been limited to scalar variables that is appropriate when methods. The se of Sic has been shown by damage is oriented in the fiber direction but not satis- merous studies [l] to have a saturation swelling value of factory when damage is loading oriented. A substantial about 0.1-0.2% at 800-1000oC. This suggests that effort is still required to develop and implement this for a composites of Sicr/Sic have the potential for excellent radiation stability. The continuous fiber architecture coupled with engineered interfaces between the fiber and 2.3.2. Design criteria matrix, provide excellent fracture properties and frac- To avoid degradation of the composite physical ture toughness values on the order of 25 MPam /.The properties, the elastic limit must be used as the maxi- strength and fracture toughness are independent of mum allowable stress. On the other hand, one of the temperature up to the limit of the fiber stability. With most attractive characteristics of SiCr/Sic composites is aprovements in fiber stability these materials exhibit that they are damage-tolerant, that is, they are capable. excellent mechanical properties to at least 1200C. Als of accommodating a high degree of deformation because these fiber/matrix microstructures impart excellent of crack arrest phenomena driven by the interface be- thermal shock and thermal fatigue resistance to these ween fibers and matrix. In principle, it can be assumed materials so plasma discharge and start-up and shut the limit for matrix microcracking saturation(beginning lown cycles should not induce significant structural of fiber/matrix debonding) as the maximum allowable damage In oxygen bearing environments, Sic will form tensile stress. The actual failure limit can instead be as- a protective layer of Sio, that greatly retards further

ii(ix) Acceptable structure lifetime in terms of neutron fluence and plasma-first wall interaction; a signifi- cant experimental campaign is required to define the SiCf/SiC limits. 2.3. SiCf /SiC thermo-mechanical model and design crite￾ria The proposal for using SiCf /SiC as structural mate￾rial for a nuclear component with a reasonably long lifetime is fairly recent, therefore no adequate modeling and design criteria are available yet. Some preliminary work has recently been performed [2] aiming both to identify appropriate models available in the aerospace research field and to theoretically define sound design criteria to improve the design thermo-mechanical ana￾lyses. 2.3.1. Modeling SiCf/SiC composites exhibit a complex nonlinear behavior combining brittle damage, residual strains and opening-closing of microcracks. These composites present different properties, and therefore different strengths, for different loading directions; moreover, tensile and compression strengths are very different. Under loading, the interaction of fibers and matrix lead at first to matrix microcracking, then to matrix/fiber decohesion, followed by opening of the microcrack and finally to fiber failure. This sequence corresponds to an initial isotropic behavior in plane and then to a crack growth perpendicular to the fibers depending on the load direction. A relatively simple model, able to take into account such a behavior and based on continuum damage mechanics which consider the composites as a continuous media, has been implemented in the FEM code CASTEM. Significant improvements on the results have been obtained when compared with models used for metals [1]. On the other end, damage description has been limited to scalar variables that is appropriate when damage is oriented in the fiber direction but not satis￾factory when damage is loading oriented. A substantial effort is still required to develop and implement this for a complete design model. 2.3.2. Design criteria To avoid degradation of the composite physical properties, the elastic limit must be used as the maxi￾mum allowable stress. On the other hand, one of the most attractive characteristics of SiCf /SiC composites is that they are damage-tolerant, that is, they are capable of accommodating a high degree of deformation because of crack arrest phenomena driven by the interface be￾tween fibers and matrix. In principle, it can be assumed the limit for matrix microcracking saturation (beginning of fiber/matrix debonding) as the maximum allowable tensile stress. The actual failure limit can instead be as￾sumed as the maximum allowable compressive stress since no damage is observed under compression. The number of fibers through the thickness of the composite is usually lower and their arrangement different. So, if one can accept the uncoupling of stresses in plane and stresses through the thickness, the former can be eval￾uated using the Von Mises criteria, while the latter corresponds to the measured rupture value. Taking into account the above remarks, for compressive stresses the limit is the rupture limits while for tensile stresses the limit is the elastic limit. A margin of about 20% may be allowable on the elastic limit. For example, in the case of the TAURO blanket, based on the CERASEP composites produced by SNECMA, the following limits have been assumed: • for normal stresses through the thickness, 110 MPa for tensile stresses (roughly corresponding to the ma￾trix tensile resistance limit outside the composite) and 420 MPa for compressive stresses (rupture limit for CERASEP N2-1); • for shear stresses through the thickness, 44 MPa (as￾sumed rupture limit, to be confirmed); • for stresses in plane, 145 MPa for tensile stresses (be￾ginning of fiber/matrix debonding) and 580 MPa for compressive stresses (rupture limit measured on the CERASEP N2-1). It appears clear that this design criteria proposal is very preliminary; it needs to be further evaluated both theoretically and, more im￾portant, experimentally through a systematic specific experimental campaign. 3. Promise of SiCf/SiC composites Composite materials made from continuous fibers of SiC, can be woven into several variant fabric architec￾tures and the matrix formed with a variety of infiltration methods. The b-phase of SiC has been shown by nu￾merous studies [1] to have a saturation swelling value of about 0.1–0.2% at 800–1000
C. This suggests that composites of SiCf /SiC have the potential for excellent radiation stability. The continuous fiber architecture, coupled with engineered interfaces between the fiber and matrix, provide excellent fracture properties and frac￾ture toughness values on the order of 25 MPa m1=2. The strength and fracture toughness are independent of temperature up to the limit of the fiber stability. With improvements in fiber stability these materials exhibit excellent mechanical properties to at least 1200
C. Also, these fiber/matrix microstructures impart excellent thermal shock and thermal fatigue resistance to these materials so plasma discharge and start-up and shut￾down cycles should not induce significant structural damage. In oxygen bearing environments, SiC will form a protective layer of SiO2 that greatly retards further R.H. Jones et al. / Journal of Nuclear Materials 307–311 (2002) 1057–1072 1059

RH. Jones et al. I Journal of Nuclear Materials 307-311(2002)1057-1072 oxidation. Therefore, SiCrSic composites have the po Table 2 tential for excellent oxidation resistance in He +O, en- Defect formation energies in 3C-SiC [4] Defect type Formation energy (el 4. Challenges of engineering the properties of Sic / Sic Given the positive attributes of Sic /SiC composites, they would be the obvious first choice for structural applications in fusion energy systems if there were no issues in their use. However there are some unresolved 71 issues associated with their use as outlined in table l Considerable progress has been made in understanding these issues and in some cases improvements have beer C-Si(100) made. Since the properties of these composites are 9.32 engineerable, there is the potential, to some extent,to (110 engineer around these issues. The purpose of this paper is to summarize some of the understanding and im provements made in these materials suggest that the c interstitials should migrate from sublattice to sublattice. The most stable configuration for Si interstitials is in a tetrahedrally-coordinated in- 5. Fundamentals of radiation damage in SiC terstitial site surrounded by c atoms on the C sublattice Ongoing theoretical and computational 5.1. Defect formation energi migration of these stable defect configurations will yield the necessary parameters to model radiation damage Density functional theory dFT), based on the processes at higher temperatures and over longer time pseudopotential plane-wave method within the frame- scales in SiC using rate-theory approaches or kinetic york of the local density approximation (LDA), has Monte Carlo methods been used to study the formation and properties of na tive defects in 3C-SiC (cubic SiC), as described in detail 5.2. Damage production and accumulation elsewhere [3-5]. The formation energies for vacancie antisite defects and interstitials in 3C-SiC are summa Molecular dynamics(MD) simulations of displace rized in Table 2. Two types of vacancies form, namely C nent cascades and cascade overlap events have been and si vacancies. In addition, two types of antisite de performed using a modified version of the code mOl fects are formed by atoms located on the wrong sub- DY [6]. with either constant volume or constant pressure lattice. For interstitial defects, there are ten possible and periodic boundary conditions. Details of the md configurations, four tetrahedral and six dumbbell(split) simulations and interatomic potentials employed are configurations. It is found that the most stable config urations for C interstitials are C-C and C-Si split in The short cascade lifetime in SiC is illustrated in Fig. terstitials along the(100)and(110) directions, which 1, where the numbers of interstitials and antisite defects produced in a 10 keV Si cascade are shown as a function Table I of time. The number of interstitials and vacancies(not Critical issues associated with the use of SiC SiC composites in shown) reaches a peak at about 0. 1 ps and then de- nuclear environments creases due to defect recombination [7, 9, 10]. The defect Priman Secondary concentrations attain steady state values after about 0.4 Thermal conductivity Chemical compatibility(He) ps. The cascade lifetime has been found to be slightly Radiation stability Carbon interfaces longer(about 0.7 ps)for a 50 keV Si PKa in SiC [7, 91 Fibers-polymer derived Thermal fatigue and shock These lifetimes are about an order of magnitude smaller Lack of a database than the values reported for metals using similar PKA + Matrices-CVI and polymer Long-term thermal stability energies [6, 11]. impregnated Design codes The results from MD simulations [10 for the ne Transmutations displacements and antisite defects produced by a 10 keV Hermetic behavior Si primary knock-on atom(PKA) are shown in Fig. 2 as Joining technology a function of PKA energy. The number of net dis- Chemical compatibility placements is defined as the sum of the total number of interstitials (or vacancies) and antisite defects. The

oxidation. Therefore, SiCf/SiC composites have the po￾tential for excellent oxidation resistance in He þ O2 en￾vironments. 4. Challenges of engineering the properties of SiCf/SiC Given the positive attributes of SiCf/SiC composites, they would be the obvious first choice for structural applications in fusion energy systems if there were no issues in their use. However, there are some unresolved issues associated with their use as outlined in Table 1. Considerable progress has been made in understanding these issues and in some cases improvements have been made. Since the properties of these composites are engineerable, there is the potential, to some extent, to engineer around these issues. The purpose of this paper is to summarize some of the understanding and im￾provements made in these materials. 5. Fundamentals of radiation damage in SiC 5.1. Defect formation energies Density functional theory (DFT), based on the pseudopotential plane-wave method within the frame￾work of the local density approximation (LDA), has been used to study the formation and properties of na￾tive defects in 3C–SiC (cubic SiC), as described in detail elsewhere [3–5]. The formation energies for vacancies, antisite defects and interstitials in 3C–SiC are summa￾rized in Table 2. Two types of vacancies form, namely C and Si vacancies. In addition, two types of antisite de￾fects are formed by atoms located on the wrong sub￾lattice. For interstitial defects, there are ten possible configurations, four tetrahedral and six dumbbell (split) configurations. It is found that the most stable config￾urations for C interstitials are C–C and C–Si split in￾terstitials along the h100i and h110i directions, which suggest that the C interstitials should migrate from sublattice to sublattice. The most stable configuration for Si interstitials is in a tetrahedrally-coordinated in￾terstitial site surrounded by C atoms on the C sublattice. Ongoing theoretical and computational studies of the migration of these stable defect configurations will yield the necessary parameters to model radiation damage processes at higher temperatures and over longer time scales in SiC using rate-theory approaches or kinetic Monte Carlo methods. 5.2. Damage production and accumulation Molecular dynamics (MD) simulations of displace￾ment cascades and cascade overlap events have been performed using a modified version of the code MOL￾DY [6], with either constant volume or constant pressure and periodic boundary conditions. Details of the MD simulations and interatomic potentials employed are described elsewhere [7–10]. The short cascade lifetime in SiC is illustrated in Fig. 1, where the numbers of interstitials and antisite defects produced in a 10 keV Si cascade are shown as a function of time. The number of interstitials and vacancies (not shown) reaches a peak at about 0.1 ps and then de￾creases due to defect recombination [7,9,10]. The defect concentrations attain steady state values after about 0.4 ps. The cascade lifetime has been found to be slightly longer (about 0.7 ps) for a 50 keV Si PKA in SiC [7,9]. These lifetimes are about an order of magnitude smaller than the values reported for metals using similar PKA energies [6,11]. The results from MD simulations [10] for the net displacements and antisite defects produced by a 10 keV Si primary knock-on atom (PKA) are shown in Fig. 2 as a function of PKA energy. The number of net dis￾placements is defined as the sum of the total number of interstitials (or vacancies) and antisite defects. The Table 1 Critical issues associated with the use of SiCf /SiC composites in nuclear environments Primany issues Secondary issues Thermal conductivity Chemical compatibility (He) Radiation stability þ Carbon interfaces þ Fibers-polymer derived Thermal fatigue and shock Interphases-C, porous Lack of a database þMatrices-CVI and polymer impregnated Long-term thermal stability Design codes Transmutations Hermetic behavior Joining technology Chemical compatibility (Pb–Li) Table 2 Defect formation energies in 3C–SiC [4] Defect type Formation energy (eV) Vc 5.48 VSi 6.64 CSi 1.32 SiC 7.20 CTC 6.41 CTS 5.84 SiTC 6.17 SiTS 8.71 Cþ–Sih100i 3.59 Cþ–Ch100i 3.16 C–Siþh100i 10.05 Siþ–Sih100i 9.32 Cþ–Ch110i 3.32 Cþ–Sih110i 3.28 1060 R.H. Jones et al. / Journal of Nuclear Materials 307–311 (2002) 1057–1072

RH. Jones et al. /Jounal of Nuclear Materials 307-311(2002)1057-1072 10 keV Si PKA Ct9 60 40 o Antisite Defects 0.001 0.01 Time( ps) Fig. 1. Number of interstitials and antisite defects produced in a 10 keV Si cascade as a function of time [10] 98 3C-SiC Net Displacements Fig. 3. MD simulation of primary damage state in SiC at 300K Si PKA Damage Energy(kev) due to a 10 keV Si PKA 8. The Si and c defects are dark and light gray, respectively, and the interstitials, antisite defects and Fig. 2. Net displacements and antisite defects produced as a vacancies are given by large, medium, and small spheres, unction of Si PKA damage energy [10 spectively number of C displacements is much larger than the (dpa)for irradiation with 550 keV Si* ions at 190 K[12 number of Si displacements, which is consistent with 14. The solid curve (Fig. 4) is based on the direct recent experimental observations [5]. Similar behavior is impact/defect-stimulated model for amorphization [15]. observed for C PKAs. Antisite defects are produced by where point defects, such as interstitials and antisite nearest-neighbor replacements during the collisional defects, stimulate the growth of amorphous nuclei (or phase and some random interstitial-vacancy recombi- defect clusters) produced directly in a displacement nation during the subsequent relaxation phase cascade. As the dose increases, cascade superposition MD simulations, as illustrated in Fig. 3, have also and defect-stimulated growth at crystalline-amorphous shown that Si PKAs generate only small interstitial interfaces become more probable. The relative ratio of clusters, with most defects being isolated single inter- direct- impact and defect-stimulated cross sections from stitials and vacancies distributed over a large region the model fit to the data for Si are consistent with those [8, 12, 13]. These predictions are in agreement with the derived from the MD simulations based on relative interpretation of the experimental results on disordering cluster distributions [121 behavior in SiC, as shown in Fig. 4, where the relative MD methods with 10 key si pkas have been em- order on the Si sublattice in Sic at the damage peak ployed to simulate cascade overlap, damage accumula shown as a function of dose in displacements per atom tion and amorphization processes in 3C-SiC. In this

number of C displacements is much larger than the number of Si displacements, which is consistent with recent experimental observations [5]. Similar behavior is observed for C PKAs. Antisite defects are produced by nearest-neighbor replacements during the collisional phase and some random interstitial-vacancy recombi￾nation during the subsequent relaxation phase. MD simulations, as illustrated in Fig. 3, have also shown that Si PKAs generate only small interstitial clusters, with most defects being isolated single inter￾stitials and vacancies distributed over a large region [8,12,13]. These predictions are in agreement with the interpretation of the experimental results on disordering behavior in SiC, as shown in Fig. 4, where the relative disorder on the Si sublattice in SiC at the damage peak is shown as a function of dose in displacements per atom (dpa) for irradiation with 550 keV Siþ ions at 190 K [12– 14]. The solid curve (Fig. 4) is based on the direct￾impact/defect-stimulated model for amorphization [15], where point defects, such as interstitials and antisite defects, stimulate the growth of amorphous nuclei (or defect clusters) produced directly in a displacement cascade. As the dose increases, cascade superposition and defect-stimulated growth at crystalline-amorphous interfaces become more probable. The relative ratio of direct-impact and defect-stimulated cross sections from the model fit to the data for Si are consistent with those derived from the MD simulations based on relative cluster distributions [12]. MD methods with 10 keV Si PKAs have been em￾ployed to simulate cascade overlap, damage accumula￾tion and amorphization processes in 3C–SiC. In this Fig. 3. MD simulation of primary damage state in SiC at 300 K due to a 10 keV Si PKA [8]. The Si and C defects are dark and light gray, respectively, and the interstitials, antisite defects and vacancies are given by large, medium, and small spheres, re￾spectively. Fig. 1. Number of interstitials and antisite defects produced in a 10 keV Si cascade as a function of time [10]. Fig. 2. Net displacements and antisite defects produced as a function of Si PKA damage energy [10]. R.H. Jones et al. / Journal of Nuclear Materials 307–311 (2002) 1057–1072 1061

RH. Jones et al. I Journal of Nuclear Materials 307-311(2002)1057-1072 550 keV Si 190K o0g 200K 0.6 ▲ Si Disorder a Si Disorder 0.2 Model Prediction 0.3 Dose(dpa) Fig. 5. Relative disorder in Sic based on MD simulations [16- Fig. 4. Relative disorder on the si sublattice at the damage ak as a function H-SiC irradiated with 550 keV Sit 190K[2-14 agreement with experimental measurements [18]. Th the good agreement between MD simulations and ex- simulation study, 10 keV Si cascades, similar to those in perimental results provides atomic-level insights into Fig 3, were randomly overlapped at 200 K in an Md the interpretation of radiation damage processes in SiC simulation cell containing 40000 atoms until a fully Ongoing MD simulations on cascade annealing and disordered state was achieved after 140 cascades [16-18]. defect migration in Sic will yield new atomic-level un- t low doses, damage is dominated by single interstitials derstanding of the temperature dependence of radiation and small clusters consisting of interstitials and antisite damage processes in SiC. defects and their concentration increases with increas- ing dose. The coalescence of small and large clusters at rtant mechanism leading to 6. Recent advances in SiC /sic amorphization in SiC, and the homogeneous nucleation of small clusters at low doses is consistent with the 6. New materials homogeneous amorphization process that is observed experimentally by high-resolution TEM [19]. Under 6.1.I. Composites with advanced fiber these conditions, the primary driving force for irradia Composites produced with the advanced fibers, Hi- tion-induced amorphization is the accumulation of both Nicalon Type S and Tyranno-SA have been irradiated interstitials and antisite defects. The relative disorder and the test results are presented and compared to other from the MD simulations exhibits a sigmoidal depen- data in Fig. 6 [20-29]. Comparison is made to both dence on dose, as shown in Fig. 5, that is in good monolithic SiC and composites made with Ceramic agreement with the experimental measurements for 550 Grade-Nicalon and Hi-Nicalon fibers. The results cover keV Sit irradiation(Fig. 4). The interpretation of the a range of temperatures but the trend of the irradiated to MD results is consistent with the direct-impact/defect unirradiated ultimate strength, Surra / Sunirr, clearly stimulated model for amorphization, where the pro- shows that composites with the advanced fibers Hi- and antisite defects stimulates Nicalon Type S and Tyranno Sa showed no loss in morphous growth at crystalline-amorphous interfaces. strength up to a dose of 10 dpa. The results of Price The model fit shown in Fig. 5 is based on the average [28, 29 and Jones et al. [22] give some support to the relative cross sections determined previously for single possibility that the strength of irradiated advanced fiber 10 keV Si cascades [12]. High-resolution TEM image material could remain unchanged up to at least 10 dpa simulations of specific damage states in the MD simu- and perhaps higher. Further advances will likely require lation cell have been performed to reveal the change in ailoring the interface swelling characteristics to com- microstructural features with increasing dose from cas- pensate for differential swelling between the fiber and cade overlap [17]. The microstructural evolution in the matrix. Advanced interface developments that could MD simulations is very similar to that observed previ- provide this tailoring have been reported by Snead [27] ously in experimental HRTEM images obtained from and they include multilayer Sic/C interfaces and porou 19. Likewise, the swelling and or pseudo-porous Sic interfaces. An example of a stored energy determined as a function of dose from multilayer interface is shown in Fig. 7 and the resulting cascade overlap in the MD simulations are in good room temperature bend strengths are given in Fig. 8

simulation study, 10 keV Si cascades, similar to those in Fig. 3, were randomly overlapped at 200 K in an MD simulation cell containing 40 000 atoms until a fully disordered state was achieved after 140 cascades [16–18]. At low doses, damage is dominated by single interstitials and small clusters consisting of interstitials and antisite defects, and their concentration increases with increas￾ing dose. The coalescence of small and large clusters at higher doses is an important mechanism leading to amorphization in SiC, and the homogeneous nucleation of small clusters at low doses is consistent with the homogeneous amorphization process that is observed experimentally by high-resolution TEM [19]. Under these conditions, the primary driving force for irradia￾tion-induced amorphization is the accumulation of both interstitials and antisite defects. The relative disorder from the MD simulations exhibits a sigmoidal depen￾dence on dose, as shown in Fig. 5, that is in good agreement with the experimental measurements for 550 keV Siþ irradiation (Fig. 4). The interpretation of the MD results is consistent with the direct-impact/defect stimulated model for amorphization, where the pro￾duction of interstitials and antisite defects stimulates amorphous growth at crystalline-amorphous interfaces. The model fit shown in Fig. 5 is based on the average relative cross sections determined previously for single 10 keV Si cascades [12]. High-resolution TEM image simulations of specific damage states in the MD simu￾lation cell have been performed to reveal the change in microstructural features with increasing dose from cas￾cade overlap [17]. The microstructural evolution in the MD simulations is very similar to that observed previ￾ously in experimental HRTEM images obtained from ion-irradiated 3C–SiC [19]. Likewise, the swelling and stored energy determined as a function of dose from cascade overlap in the MD simulations are in good agreement with experimental measurements [18]. Thus, the good agreement between MD simulations and ex￾perimental results provides atomic-level insights into the interpretation of radiation damage processes in SiC. Ongoing MD simulations on cascade annealing and defect migration in SiC will yield new atomic-level un￾derstanding of the temperature dependence of radiation damage processes in SiC. 6. Recent advances in SiCf/SiC performance 6.1. New materials 6.1.1. Composites with advanced fibers Composites produced with the advanced fibers, Hi￾Nicalon Type S and Tyranno-SA have been irradiated and the test results are presented and compared to other data in Fig. 6 [20–29]. Comparison is made to both monolithic SiC and composites made with Ceramic Grade-Nicalon and Hi-Nicalon fibers. The results cover a range of temperatures but the trend of the irradiated to unirradiataed ultimate strength, Sirrad u =Sunirr u , clearly shows that composites with the advanced fibers Hi￾Nicalon Type S and Tyranno SA showed no loss in strength up to a dose of 10 dpa. The results of Price [28,29] and Jones et al. [22] give some support to the possibility that the strength of irradiated advanced fiber material could remain unchanged up to at least 10 dpa and perhaps higher. Further advances will likely require tailoring the interface swelling characteristics to com￾pensate for differential swelling between the fiber and matrix. Advanced interface developments that could provide this tailoring have been reported by Snead [27] and they include multilayer SiC/C interfaces and porous or pseudo-porous SiC interfaces. An example of a multilayer interface is shown in Fig. 7 and the resulting room temperature bend strengths are given in Fig. 8. Fig. 5. Relative disorder in SiC based on MD simulations [16– Fig. 4. Relative disorder on the Si sublattice at the damage 18]. peak as a function H–SiC irradiated with 550 keV Siþ ions at 190 K [12–14]. 1062 R.H. Jones et al. / Journal of Nuclear Materials 307–311 (2002) 1057–1072

RH Jones et al. Journal of Nuclear Materials 307-311(2002)1057-1072 Hi-Nicalon Type-S/PyC/FCVI-SiC icalon/PyC/FCV-Sie O-SAP\C/FCVI-SiC o Monolithic CVD-sic 5 1:500C·HFIR 6: 300C: HFIR 740C HFIR 5:430-5 Fig. 6. Relative strength, irradiated/unirradiated, for SiC/SiC composites These results show that it is possible to achieve signifi cant bend ultimate strengths for unirradiated SiCr/SiC CVD SI composites with either multilayer and pseudo-porous interfaces but that the radiation stability for a composite matrix with Hi-Nicalon is better with the multilayer than the pseudo-porous interface 6.1.2. New SiC /SiC composites by transient liquid phase sIntering process iC-interlayer The LPS process, which has been a common tech nique in producing monolithic silicon carbide and other C-interlayer at relatively low costs, was successfully applied to matrix densification for SiCr/SiC composites for the first time. The lab-grade materials with uni-directiona reinforcement exhibits typically 700 MPa in three-point 0.5 flexural strength and pseudo-ductile fracture mode with a good fiber pull-out. Tensile strength tests, both toughness evaluation are in progress. The composites are almost fully-dense, with very minor porosity within Hi-Nicalon fib intra-fiber-bundles. Thermal conductivity and herme ticity data will soon be available. Larger scale pro- ig. 7. Multilayer SiC interphase with thin pyrolitic layer ap duction and complex shaping are presently being plied followed by SiC interlayers. n other non-radiation environments where components 6.2. Thermal conductivity: modeling based guidance to or structures are subjected to a high heat flux, is the performance improvements expected in-service behavior of its effective transverse hermal conductivity, Ke. Knowledge about the ex- jor issue to be considered when using SiCr/Sic pected range of Kef is necessary to optimize Sic/Sic temperature neutron radiation environment, or

These results show that it is possible to achieve signifi- cant bend ultimate strengths for unirradiated SiCf /SiC composites with either multilayer and pseudo-porous interfaces but that the radiation stability for a composite with Hi-Nicalon is better with the multilayer than the pseudo-porous interface. 6.1.2. New SiCf /SiC composites by transient liquidphase sintering process The LPS process, which has been a common tech￾nique in producing monolithic silicon carbide and other ceramics at relatively low costs, was successfully applied to matrix densification for SiCf /SiC composites for the first time. The lab-grade materials with uni-directional reinforcement exhibits typically 700 MPa in three-point flexural strength and pseudo-ductile fracture mode with a good fiber pull-out. Tensile strength tests, both at room and at elevated temperature, and fracture toughness evaluation are in progress. The composites are almost fully-dense, with very minor porosity within intra-fiber-bundles. Thermal conductivity and herme￾ticity data will soon be available. Larger scale pro￾duction and complex shaping are presently being attempted. 6.2. Thermal conductivity: modeling based guidance to performance improvements A major issue to be considered when using SiCf /SiC in a high-temperature neutron radiation environment, or in other non-radiation environments where components or structures are subjected to a high heat flux, is the expected in-service behavior of its effective transverse thermal conductivity, Keff . Knowledge about the ex￾pected range of Keff is necessary to optimize SiCf /SiC configurations for their intended uses. Several modeling Fig. 7. Multilayer SiC interphase with thin pyrolitic layer ap￾plied followed by SiC interlayers. Fig. 6. Relative strength, irradiated/unirradiated, for SiCf /SiC composites versus radiation dose. R.H. Jones et al. / Journal of Nuclear Materials 307–311 (2002) 1057–1072 1063

RH. Jones et al. I Journal of Nuclear Materials 307-311(2002)1057-1072 600 500 500 Unirradiated (5073 75 MPa) E苏 00 Unirradiated (267+ 70 MPa) 200 1.1dpa【216"C) 1.1dpa(385°c) 200 10D 00010020030.04005006 00010D20.B0.040.050.06 Cross Head Displacement(cm) Cross Head Displacement (cm) ated(514土224Pa) 1.1dpa(260°c) 0010D20B3004t.D5.D6 Fig. 8. Bend-displacement curves for SiC/SiC with(a) pyrolytic C,(b)multiplayer, and (c) porous interfaces. studies have shown how Ker depends upon constituent For fiber-to-matrix conductivity ratios(r)less than fiber and matrix thermal conductivity values, and their 10(r< 10) and for fiber volume fractions, ()<0.5, the volume fractions and distributions [30-33. However, Hasselman-Johnson(H-J) model predictions given by many experimental measurements have indicated that Eq (1) deviate from numerical FEM results by less than interfaces between fibers and matrices in a composite 5%. By Eq. (I), for dispersed fibers in a matrix the ef- introduce a thermal barrier that may reduce Keff [34-37. fective transverse thermal conductivity(Ker)is primarily Furthermore, Kefr may be altered by physical changes of controlled by the thermal conductivity of the continuous the interface and even the surrounding atmosphere. As matrix phase(Km)and the interfacial conductance(h).A with mechanical behavior, to attain desired thermal simple thermal barrier model was introduced to describ behavior of SicSiC proper attention needs to be given h and the gaseous and direct contact components fgh to the design of the interphase and the control of and fahd, respectively ) Values of h determined by ec interfacial thermal effects. Classical composite models (2)for a uniaxial Hi-Nicalon fiber/amorphous SiC recently have been updated to include the effect of in- matrix composite in vacuum, argon and helium com- terfacial thermal barriers [38]. Interfacial thermal bar- pared favorably with values estimated by the simple riers are quantitatively characterized by a value called thermal barrier model. reasonable agreement between the interfacial conductance, which includes the effect of numerical FEM and experimental results with H-J imperfect matching of surfaces at an interface as well as model predictions suggest that Kefr for a Sicr/Sic com- the effect of interfacial gaps brought about by debonding posite with fiber volume fractions f<0.4 and with of the fiber from the matrix or microcracking within the simple unidirectional or cross-ply fiber architecture are fiber coating [38] well described by eq (1) below

studies have shown how Keff depends upon constituent fiber and matrix thermal conductivity values, and their volume fractions and distributions [30–33]. However, many experimental measurements have indicated that interfaces between fibers and matrices in a composite introduce a thermal barrier that may reduce Keff [34–37]. Furthermore, Keff may be altered by physical changes of the interface and even the surrounding atmosphere. As with mechanical behavior, to attain desired thermal behavior of SiCf /SiC proper attention needs to be given to the design of the interphase and the control of interfacial thermal effects. Classical composite models recently have been updated to include the effect of in￾terfacial thermal barriers [38]. Interfacial thermal bar￾riers are quantitatively characterized by a value called the interfacial conductance, which includes the effect of imperfect matching of surfaces at an interface as well as the effect of interfacial gaps brought about by debonding of the fiber from the matrix or microcracking within the fiber coating [38]. For fiber-to-matrix conductivity ratios (r) less than 10 (r < 10) and for fiber volume fractions, ðf Þ 6 0:5, the Hasselman–Johnson (H–J) model predictions given by Eq. (1) deviate from numerical FEM results by less than 5%. By Eq. (1), for dispersed fibers in a matrix the ef￾fective transverse thermal conductivity (Keff ) is primarily controlled by the thermal conductivity of the continuous matrix phase (Km) and the interfacial conductance (h). A simple thermal barrier model was introduced to describe h
and the gaseous and direct contact components (fghg and fdhd, respectively). Values of h determined by Eq. (2) for a uniaxial Hi-Nicalone fiber/amorphous SiC matrix composite in vacuum, argon and helium com￾pared favorably with values estimated by the simple thermal barrier model. Reasonable agreement between numerical FEM and experimental results with H–J model predictions suggest that Keff for a SiCf /SiC com￾posite with fiber volume fractions f 6 0:4 and with simple unidirectional or cross-ply fiber architecture are well described by Eq. (1) below: Fig. 8. Bend-displacement curves for SiCf /SiC with (a) pyrolytic C, (b) multiplayer, and (c) porous interfaces. 1064 R.H. Jones et al. / Journal of Nuclear Materials 307–311 (2002) 1057–1072

RH. Jones et al. /Jounal of Nuclear Materials 307-311(2002)1057-1072 Kefr=Km(Kr/Km-1-Kr/ah)Vr 63. Transmutation rates (1+Kr/Km+Kr/ah) x[(1-Kr/Km+Kr/ah)Ve Transmutation calculations were performed 1 +(1+Kr/Km+ Kr/ah) () REAC-3 code for pure SiC irradiated in the spectrum of the first wall of the ARIES-IV co usion energy device. The ARIES-IV first wall has a total neutron flux of 3. 6x 10 s n/cm s and a fast flux K are the thermal conductivity values of the matrix and (E>0.1 Mev) of 1.9 x 10 n/cms. Calculations were fiber constituents; and Vr and a are the fiber volume performed for a continuous irradiation of up to 12 ef- fraction and radius, respectively fective full power years(efpy) to a total neutron dose of For a woven 2D-SiC /SiC composite, the localized 1. 37 x 1024 n/cm. The elemental composition of the effects of dense fiber packing within individual tows material is not significantly affected by the post-irradi- (<0.6) and the occurrence of many direct fiber-fiber ation decay of radioactive isotopes, since they are either contacts at the numerous fiber bundle crossover points too short-lived or too long-lived to affect the composi will introduce positive deviations from Eq. (1). How- tion over a reactor lifetim ever, the analytic solution expressed by eq. () should be very appropriate to examine thermal conductivity 6.3.1. Sic burn-out degradation induced in these composites by neutron In this neutron spectrum, Si burns out somewhat radiation or by other mechanical or environmental more rapidly than C, Fig 9. Si burns out at a rate of treatments. If a 2D-SiCrSiC composite with initially about 0.0047/efpy. a 3% burnout of Sic has been sug high Kr- and h-values were irradiated, Ker could easily gested as a design limit for SiC composites In ARIES- be reduced by a factor of five or six due to the degra T, a 3% burnout occurs in about 6.5 efpy(total neutron dation of the interface conductance and the matrix fluence 7. 4 x 102n/cm2). The differing burnout rates of conductivity. Si and c result in an excess concentration of carbon To further examine this issue, the effects of temper- totaling about 3500 appm after 6.5 efpy. The Si and C ature and irradiation on Kefr were predicted for a hy- burnout rates are constant, and the excess C increases at pothetical 2D-SiCr/SiC composite made with high the rate of 540 appm excess Clefpy conductivity Tyranno SA fiber, a thin(0. 2-um) PyC fiber coating and a CVI-SiC matrix. For example, it was 6.3.2. Impurity burn-in predicted for this composite that Kefr would decrease The most abundant transmutation products, which from 34 W/mK before irradiation to <6 w/m k(at 200 ourn-in at constant rates, Fig. 10, are listed in table 3. C) after irradiation at 200 C. Similarly, Ker would After 6.5 efpy, although 3% of the Sic burns out, the decrease from 26 W/m K before irradiation to <10 W/ concentration of transmutant atoms totals almost 8% mk(at 1000C)after irradiation at 1000C This is because many transmutation reactions create 100 099 098 3% SiC burnout at 6.5 efp 095 Irradiation Time, efpy Fig. 9. The burnout of Si and C in Sic irradiated in ARIES-IV first wall as a function of dose in efp

Keff ¼ Km½ðKf=Km 1 Kf=ahÞVf þ ð1 þ Kf=Km þ Kf=ahÞ  ½ð1 Kf=Km þ Kf=ahÞVf þ ð1 þ Kf=Km þ Kf=ahÞ 1 ; ð1Þ where h is the effective interfacial conductance; Km and Kf are the thermal conductivity values of the matrix and fiber constituents; and Vf and a are the fiber volume fraction and radius, respectively. For a woven 2D-SiCf/SiC composite, the localized effects of dense fiber packing within individual tows (f 6 0:6) and the occurrence of many direct fiber–fiber contacts at the numerous fiber bundle crossover points will introduce positive deviations from Eq. (1). How￾ever, the analytic solution expressed by Eq. (1) should be very appropriate to examine thermal conductivity degradation induced in these composites by neutron radiation or by other mechanical or environmental treatments. If a 2D-SiCf /SiC composite with initially high Kf- and h-values were irradiated, Keff could easily be reduced by a factor of five or six due to the degra￾dation of the interface conductance and the matrix conductivity. To further examine this issue, the effects of temper￾ature and irradiation on Keff were predicted for a hy￾pothetical 2D-SiCf/SiC composite made with high conductivity Tyranno SAe fiber, a thin (0.2-lm) PyC fiber coating and a CVI-SiC matrix. For example, it was predicted for this composite that Keff would decrease from 34 W/m K before irradiation to 0:1 MeV) of 1:9  1015 n/cm2 s. Calculations were performed for a continuous irradiation of up to 12 ef￾fective full power years (efpy) to a total neutron dose of 1:37  1024 n/cm2. The elemental composition of the material is not significantly affected by the post-irradi￾ation decay of radioactive isotopes, since they are either too short-lived or too long-lived to affect the composi￾tion over a reactor lifetime. 6.3.1. SiC burn-out In this neutron spectrum, Si burns out somewhat more rapidly than C, Fig. 9. Si burns out at a rate of about 0.0047/efpy. A 3% burnout of SiC has been sug￾gested as a design limit for SiC composites. In ARIES￾IV, a 3% burnout occurs in about 6.5 efpy (total neutron fluence 7:4  1023 n/cm2). The differing burnout rates of Si and C result in an excess concentration of carbon totaling about 3500 appm after 6.5 efpy. The Si and C burnout rates are constant, and the excess C increases at the rate of 540 appm excess C/efpy. 6.3.2. Impurity burn-in The most abundant transmutation products, which burn-in at constant rates, Fig. 10, are listed in Table 3. After 6.5 efpy, although 3% of the SiC burns out, the concentration of transmutant atoms totals almost 8%. This is because many transmutation reactions create Fig. 9. The burnout of Si and C in SiC irradiated in ARIES-IV first wall as a function of dose in efpy. R.H. Jones et al. / Journal of Nuclear Materials 307–311 (2002) 1057–1072 1065

RH. Jones et al. I Journal of Nuclear Materials 307-311(2002)1057-1072 30000 15000 10000 Irradiation Time, efpy Fig. 10. Concentrations of transmutation products in Sic irradiated in ARIES-IV first wall as a function of irradiation time in effective Table 3 additional h He atoms. which number of atoms in the system to increase. Element Concentration Several other elements burn-in through two-step (appm/epr 6.5 efpy(at. % transmutations, so they accumulate nonlinearly as the uare of the irradiation time. Their concentrations are H relatively small compared to the transmutants discussed 1630 1.1 above, Fig. Il. The concentrations of elements produced 32 by two-step transmutations after 6.5 efpy are listed in Of the elements identified as transmutation products this report, only Al and Na are also considered 10 Irradiation Time, efpy Fig. Il Number density of transmutation products in Sic that result from multiple interactions of a nucleus with neutrons in the power years

additional H or He atoms, which causes the total number of atoms in the system to increase. Several other elements burn-in through two-step transmutations, so they accumulate nonlinearly as the square of the irradiation time. Their concentrations are relatively small compared to the transmutants discussed above, Fig. 11. The concentrations of elements produced by two-step transmutations after 6.5 efpy are listed in Table 4. Of the elements identified as transmutation products in this report, only Al and Na are also considered as Fig. 11. Number density of transmutation products in SiC that result from multiple interactions of a nucleus with neutrons in the ARIES-IV first wall as a function of irradiation time in effective full power years. Fig. 10. Concentrations of transmutation products in SiC irradiated in ARIES-IV first wall as a function of irradiation time in effective full power years. Table 3 Transmutation products – SiC Element Burn-in rate (appm/efpy) Concentration at 6.5 efpy (at.%) He 6384 4.2 H 2307 1.5 Mg 1630 1.1 Be 632 0.4 Al 469 0.3 P 146 0.1 1066 R.H. Jones et al. / Journal of Nuclear Materials 307–311 (2002) 1057–1072