Journal of Nuclear Materials 384 (2009)195-211 Contents lists available at ScienceDirect Journal of Nuclear materials ELSEVIER journalhomepagewww.elsevier.com/locate/jnucmat The effect of neutron irradiation on the fiber/matrix interphase of silicon carbide composites T Nozawa*Y Katoh, L L. Snead Materials Science and Technology Division, Oak Ridge National Laboratory, P 0. Box 2008, Oak Ridge, IN 37831-6138, US ARTICLE IN F O ABSTRACT Article his eceived 11 July 2008 Given the good stability of mechanical properties of silicon carbide (SiC) under neutron irradiation. ultimate irradiation tolerance of Sic composite materials may be limited by the fiber/matrix interphas Accepted 13 November 2008 which is critically important to the performance of these composites. This study investigates the irradi- ation stability of pyrolytic carbon(Pyc)monolayer and Pyc/sic multilayer interphases by tensile and sin- le fiber push-out test techniques. Neutron irradiation was performed to doses of 0.7-7.7 dpa at temperatures from 380 to 1080C. Both interfacial debond shear strength and interfacial friction stress parently decrease by irradiation, although this is not so dramatic when Tir 1000C). Noteworthy these irradiation effects depend on the type of interphase material, i.e for the pyrolytic carbon or mul layer SiC variants studied. In the range of irradiation temperature and dose, the degradation in interfa- al shear properties, while measurable, is not of a magnitude to degrade the mechanical performance of he composites. This was observed for both interphase types studied. In particular, the proportional limit tensile stress decreases slightly by irradiation while the tensile fracture strength undergoes very minor e 2008 Elsevier B V. All rights reserved. 1 Introduction adopted as an F/M interphase, which serves to intercept and tie- up propagating cracks. Since the enhanced fracture tough Silicon carbide(Sic) has been widely used for high-temperature and ultimate performance of composites is critically dependent engineering applications due to its inherently high thermo-chem- on this interphase, its irradiation stability, and the combined ef- ical stability, good oxidation resistance and strength retention at fects of environment and irradiation on the fiber, matrix, and inter- high-temperatures. Moreover, its resistance to neutron irradiation, phase is of critical importance. e.g., low-induced radioactivity and low irradiation-induced after- The effect of neutron irradiation on the F/M interface was first heat, gives scope for a potential application in fusion and advanced evaluated on Nicalon /CVI-SiC composites with Pyc as an F/M nuclear fission energy systems [1. Additionally, it has been proven interphase [5,6]. In this composite system. the crystalline CVI-SiC that Sic in a stoichiometric form offers exceptional retention of matrix swells by irradiation depending on irradiation temperature mechanical properties neutron irradiation. High-purity and neutron dose, typical of ceramic materials. However, the chemical-vapor-deposited (cvd) sic reportedly retains its strength tallized glassy Nicalon" fiber underwent densification by neutron irradiation at to 20 dpa [2-4 Due to this differential irradiation-induced dimensional change be- Due to inherent brittleness of Sic in its monolithic form, Sic is tween Nicalonand CVI-SiC, comparably large stresses generate a being developed for use in the composite form, combining a Sic fi- the F/M interface, resulting in shear failure of the composites, i.e., ber, a Sic matrix, and a fiber/matrix(F/M)interphase. Near-stoichi- shape instability probably induced by strong contribution of ometric and highly-crystalline Sic, e.g., Hi-Nicalon M Type-s or trans-interface tensile stress [5]. Meanwhile, this shape instability TyrannoM-SA SiC fibers, and a chemical-vapor-infiltrated(cvi) issue has been solved by adopting near-stoichiometric and highly SiC matrix, are generally adopted because of their better irradiation crystalline advanced third generation Sic fibers such as Hi-Nic tolerance. a poorly graphitized pyrolytic carbon(Py c)is typically alon m Type-s and Tyranno m-SA Since both advanced Sic fibers CVI-SiC matrix swell in similar manners. irradiation-induced shear stresses at the F/M interface are minimized. Indeed, no major macroscopic deformation by irradiation has been identified for Hi-Nicalon" Type-S/CVI-SiC composite system [7]. Moreover, no 1195. Japan.Tel:+81292826416:fax:+81292843589 significant deterioration of flexural and tensile properties has been a. takashi67@jaea. go jp(T Nozawa). reported[8-10]. 'S- see front matter e 2008 Elsevier B v. All rights reserved. nuchaL2008.11.015
The effect of neutron irradiation on the fiber/matrix interphase of silicon carbide composites T. Nozawa *, Y. Katoh, L.L. Snead Materials Science and Technology Division, Oak Ridge National Laboratory, P.O. Box 2008, Oak Ridge, TN 37831-6138, USA article info Article history: Received 11 July 2008 Accepted 13 November 2008 abstract Given the good stability of mechanical properties of silicon carbide (SiC) under neutron irradiation, the ultimate irradiation tolerance of SiC composite materials may be limited by the fiber/matrix interphase, which is critically important to the performance of these composites. This study investigates the irradiation stability of pyrolytic carbon (PyC) monolayer and PyC/SiC multilayer interphases by tensile and single fiber push-out test techniques. Neutron irradiation was performed to doses of 0.7–7.7 dpa at temperatures from 380 to 1080 C. Both interfacial debond shear strength and interfacial friction stress apparently decrease by irradiation, although this is not so dramatic when Tirr 1000 C). Noteworthy, these irradiation effects depend on the type of interphase material, i.e., for the pyrolytic carbon or multilayer SiC variants studied. In the range of irradiation temperature and dose, the degradation in interfacial shear properties, while measurable, is not of a magnitude to degrade the mechanical performance of the composites. This was observed for both interphase types studied. In particular, the proportional limit tensile stress decreases slightly by irradiation while the tensile fracture strength undergoes very minor change. 2008 Elsevier B.V. All rights reserved. 1. Introduction Silicon carbide (SiC) has been widely used for high-temperature engineering applications due to its inherently high thermo-chemical stability, good oxidation resistance, and strength retention at high-temperatures. Moreover, its resistance to neutron irradiation, e.g., low-induced radioactivity and low irradiation-induced afterheat, gives scope for a potential application in fusion and advanced nuclear fission energy systems [1]. Additionally, it has been proven that SiC in a stoichiometric form offers exceptional retention of mechanical properties under neutron irradiation. High-purity chemical-vapor-deposited (CVD) SiC reportedly retains its strength by neutron irradiation at least to 20 dpa [2–4]. Due to inherent brittleness of SiC in its monolithic form, SiC is being developed for use in the composite form, combining a SiC fi- ber, a SiC matrix, and a fiber/matrix (F/M) interphase. Near-stoichiometric and highly-crystalline SiC, e.g., Hi-NicalonTM Type-S or TyrannoTM-SA SiC fibers, and a chemical-vapor-infiltrated (CVI) SiC matrix, are generally adopted because of their better irradiation tolerance. A poorly graphitized pyrolytic carbon (PyC) is typically adopted as an F/M interphase, which serves to intercept and tieup propagating cracks. Since the enhanced fracture toughness and ultimate performance of composites is critically dependent on this interphase, its irradiation stability, and the combined effects of environment and irradiation on the fiber, matrix, and interphase is of critical importance. The effect of neutron irradiation on the F/M interface was first evaluated on NicalonTM/CVI-SiC composites with PyC as an F/M interphase [5,6]. In this composite system, the crystalline CVI-SiC matrix swells by irradiation depending on irradiation temperature and neutron dose, typical of ceramic materials. However, the poorly-crystallized glassy NicalonTM fiber underwent densification. Due to this differential irradiation-induced dimensional change between NicalonTM and CVI-SiC, comparably large stresses generate at the F/M interface, resulting in shear failure of the composites, i.e., shape instability probably induced by strong contribution of trans-interface tensile stress [5]. Meanwhile, this shape instability issue has been solved by adopting near-stoichiometric and highlycrystalline advanced third generation SiC fibers such as Hi-NicalonTM Type-S and TyrannoTM-SA. Since both advanced SiC fibers and CVI-SiC matrix swell in similar manners, irradiation-induced shear stresses at the F/M interface are minimized. Indeed, no major macroscopic deformation by irradiation has been identified for Hi-NicalonTM Type-S/CVI-SiC composite system [7]. Moreover, no significant deterioration of flexural and tensile properties has been reported [8–10]. 0022-3115/$ - see front matter 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.jnucmat.2008.11.015 * Corresponding author. Present address: Fusion Research and Development Directorate, Japan Atomic Energy Agency, 2-4 Shirakata Shirane, Tokai, Ibaraki 319- 1195, Japan. Tel.: +81 29 282 6416; fax: +81 29 284 3589. E-mail address: nozawa.takashi67@jaea.go.jp (T. Nozawa). Journal of Nuclear Materials 384 (2009) 195–211 Contents lists available at ScienceDirect Journal of Nuclear Materials journal homepage: www.elsevier.com/locate/jnucmat
T Nozawa et aL /Joumal of Nuclear Materials 384(2009)195-211 Hinoki et al. [11] closely investigated the irradiation effect on While there is a growing body of work evaluating the F/M inter- le interfacial shear properties by the push-in technique and com- phase fundamental knowledge about the irradiation effect on the pared the influence on varied interphases: a PyC monolayer, a PyC F/M interphase is still insufficient. Specifically, a fundamental Sic multilayer, and a pseudo 'porous'Sic interphase. By neutron question about the irradiation tolerance of the multilayer inter irradiation up to 0.5 dpa at 300-500C, slight deformation for phase is unanswered. The goal of the following study is to provide any interphase studies was observed. a decrease of interfacial data and analysis of the effect of neutron irradiation on the F/M shear properties was also identified by the double notch shear test interphase to help bridge this knowledge gap method for irradiation up to 1 dpa at 800-1000C for the same se ies of composites [12]. In both Sic/Sic composites with do-porous Sic interphase severely deteriorated their interfacial properties. In recent work by the authors [13], interfacial shear 21. materials properties for Hi-Nicalon Type-S/CVI-SiC composites with either yer or a PyC/SiC multilayer interphase were investi- Materials evaluated were unidirectional SiC/Sic composites gated for the limited irradiation condition (o-7.7 dpa, (Hyper-Therm High-Temperature Composites Inc, Huntington Tirr=800C). The results indicate(1)a very minor irradiation effect Beach, CA)with varied F/M interphases(Table 1). Reinforcing Sic on interfacial shear properties for the monolayer composites and fibers were Hi-Nicalon"M Type-S(3. 1 g/cm Si/C=1.05).A high- (2)a slight irradiation-induced decrease of shear stresses for the purity and high-crystallinity Sic matrix was chemical-vapor-infiI multilayer composites. This difference is somewhat surprising trated. A PyC monolayer or a Py/Sic multilayer interphase was since it has been believed that the thinly-layered carbon interpha- chemically deposited on the fiber surface in advance of matrix den ses would be more radiation-stable sification. Thickness of the Pyc monolayer was 520-720 nm. In contrast, the multilayer interphase d of a sequel Table 1 of w20 nm-thick PyC and 100 nm-thick Sic sub-layers( Fig. 1). Key characteristics of unidirectional Hi-Nicalon Type-S(CVI-SiC composites The fiber volume fraction and porosity were 30-40% and 15-20% Interphase Fiber volume fraction Porosity Density (glcm respectively. The high-porosity of these CVI-SiC/Sic composites ~0.29 ~257 was due primarily to the presence of columnar pores along the Multilayer ~038 ~0.14 2.66 fiber bundles. Key characteristics of these composites have been reported elsewhere [12-14 a b 》 5 1um d 4 um 500nm Fig 1. Typical microstructural images of (a).(b)Pyc monolayer and(c). (d) Pyc/Sic multilayer interphases of as-received SiC/SiC composites[13]
Hinoki et al. [11] closely investigated the irradiation effect on the interfacial shear properties by the push-in technique and compared the influence on varied interphases: a PyC monolayer, a PyC/ SiC multilayer, and a pseudo ‘porous’ SiC interphase. By neutron irradiation up to 0.5 dpa at 300–500 C, slight deformation for any interphase studies was observed. A decrease of interfacial shear properties was also identified by the double notch shear test method for irradiation up to 1 dpa at 800–1000 C for the same series of composites [12]. In both cases, SiC/SiC composites with pseudo-porous SiC interphase severely deteriorated their interfacial properties. In recent work by the authors [13], interfacial shear properties for Hi-NicalonTM Type-S/CVI-SiC composites with either a PyC monolayer or a PyC/SiC multilayer interphase were investigated for the limited irradiation condition (/ = 7.7 dpa, Tirr = 800 C). The results indicate (1) a very minor irradiation effect on interfacial shear properties for the monolayer composites and (2) a slight irradiation-induced decrease of shear stresses for the multilayer composites. This difference is somewhat surprising since it has been believed that the thinly-layered carbon interphases would be more radiation-stable. While there is a growing body of work evaluating the F/M interphase, fundamental knowledge about the irradiation effect on the F/M interphase is still insufficient. Specifically, a fundamental question about the irradiation tolerance of the multilayer interphase is unanswered. The goal of the following study is to provide data and analysis of the effect of neutron irradiation on the F/M interphase to help bridge this knowledge gap. 2. Experimental 2.1. Materials Materials evaluated were unidirectional SiC/SiC composites (Hyper-Therm High-Temperature Composites Inc., Huntington Beach, CA) with varied F/M interphases (Table 1). Reinforcing SiC fibers were Hi-NicalonTM Type-S (3.1 g/cm3 , Si/C = 1.05). A highpurity and high-crystallinity SiC matrix was chemical-vapor-infiltrated. A PyC monolayer or a PyC/SiC multilayer interphase was chemically deposited on the fiber surface in advance of matrix densification. Thickness of the PyC monolayer was 520–720 nm. In contrast, the multilayer interphase was composed of a sequence of 20 nm-thick PyC and 100 nm-thick SiC sub-layers (Fig. 1). The fiber volume fraction and porosity were 30–40% and 15–20%, respectively. The high-porosity of these CVI-SiC/SiC composites was due primarily to the presence of columnar pores along the fiber bundles. Key characteristics of these composites have been reported elsewhere [12–14]. Table 1 Key characteristics of unidirectional Hi-NicalonTM Type-S/CVI-SiC composites. Interphase Fiber volume fraction Porosity Density (g/cm3 ) PyC 0.29 0.15 2.57 Multilayer 0.38 0.14 2.66 Fig. 1. Typical microstructural images of (a), (b) PyC monolayer and (c), (d) PyC/SiC multilayer interphases of as-received SiC/SiC composites [13]. 196 T. Nozawa et al. / Journal of Nuclear Materials 384 (2009) 195–211
T Nozawa et aL/Journal of Nuclear Materials 384(2009)195-211 2. 2. Neutron irradiation on the specimen holder above a narrow groove. The fiber was then randomly selected and monotonically loaded up to the maximum Two types of neutron irradiation campaigns were performed in system allowable load capacity(650 mN). A Berkovich indenter the High Flux Isotope Reactor(HFIR)at Oak Ridge National Labora- tip was used in this experiment. An applied load rate was 0.05 N/ tory. The HFIR-14J fixed-core capsule irradiation was performed in N. S. Details of the fiber push-out test technique have been de- unshielded removable beryllium position of the HFIR. Irradia- scribed elsewhere [18. Microstructures of the F/M interphase tion dose was 7.7 dpa assuming 1.0 x 1025 n/ m2(E>0.1 Mev) and the pushed-out fiber surface were observed by scanning elec corresponds to one displacement per atom(1 dpa), while the irra- tron microscopy(SEM)for both as-received and neutron irradiated diation temperature was 800C. In contrast, small-capsule rabbit materials. irradiations were performed in the hydraulic tube of the hFiR un- Of many stress parameters, two interfacial shear properties: der the Fun and neri SiC/ SiC series of irradiations. The neutron an interfacial debond shear strength (ts)as a critical shear stress to doses were 0.7-4.2 dpa and irradiation temperatures were in the induce an interfacial crack along the bonded F/M interface, and (2) range of 380-1080C Irradiation temperature was measured and an interfacial friction stress(tr) at the debonded interface are de- ontrolled in-situ by thermocouples and capsule sweep-gas mix- ed. Both ts and tr are calculated using experimental push-out tures for the HFIR-14 experiment. The uncertainty of irradiation test results of (1)a debond initiation stress(od)and (2)a complete emperatures was reasonably low(+20C). In contrast, irradiation debonding and sliding stress(omax) as defined in Fig. 2. In the fol- mperatures in rabbit irradiation were estimated by the isochro- lowing discussion, a compressive stress is denoted with a negative lI annealing of CVD-SiC temperature monitors. In principle, the sign For evaluation of this type of data, various analytical models temperature monitor gives the temperature near the end of the have been developed [18-24. Of these models, a non-linear shear- irradiation period with a similar uncertainty to the instrument lag model proposed by Hsueh [ 21] becomes a basis to evaluate ts experiment(±20°C)[15,16 considering the precise stress interaction at the F/M interface: (1) the radial dependence of the axial stresses in both the fiber 23. Tensile test and the matrix, (2) the shear stress distribution in the matrix and(3)the exact equilibrium equation in relating the tangential Following a guideline of ASTM C1275-00, cyclic unloading/ stress to the radial stress at the interface Hsueh further developed reloading tensile tests were conducted at ambient-temperature a double shear-lag model by separately idering the stress using an electromechanical testing machine. Rectangular minia- interactions among the fiber, matrix and F/M interphase [22]. or 20 x 2Wx 15mm were prepared. Note that the total speci- cability to isotropic constituents. In the modified double shear-lag men length of 50 mm was fixed for all specimens. Specimen size model by the authors(see Appendix), anisotropy of the constitu- effect is a potential concern when testing these specimens. Consid- ents is considered. Besides, residual stresses induced by thermal ering the structural minimum unit width(or thickness)of the uni- expansion mismatch and by differential swelling among the fiber. directional composites is larger than the width(or thickness)of the F/M interphase, and matrix can be discussed together in the model, mono fiber bundle(<l mm), a very minor effect of specimen size although the effect of dynamic phenomena such as irradiation on tensile properties, which depend on the axial fiber volume frac- creep of Sic and Py c was ignored for simplicity. In contrast, Shetty tion, is expected in the size range of concern [17 Indeed, no signif- [24] originally discussed the interfacial friction issue for the deb icant size effect was found in this study. The tensile specimen was onded interface. In the previous work by the authors [25 ]. Shetty passively gripped via aluminum grip-end tabs using a pneumatic model was updated for anisotropic composites and this method wedge-type gripping device Tensile strain was measured by a pair was applied in this study. f strain gauges with a gauge length of 5.0 mm, which were adhe- The detailed analogy of the modified Hsueh model was dis- sively bonded on specimen surfaces of the middle gauge section. a cussed elsewhere [ 22, 25]. The resultant form can provide an inter- onstant crosshead displacement rate was 0.5 mm/min th(ts) -it)-z(od+Oth+Orr ) exp(it)+exp(-it)-2 exp (it)-exp(-it) Youngs modulus was defined as an initial tangent modulus in where the constants i and Z are determined by dimensions and the tensile stress-strain curve. Proportional limit tensile stress elastic constants of the constituents, and th and irr are stress (PLS)was determined as a stress of 5% deviation in stress from ini- parameters defined by the thermal expansion mismatch and the tial linearity following ASTM C1275-00. Ultimate tensile strength differential swelling among the fiber, matrix and F/M interphase. (UTS)and total elongation were defined as a stress and a strain respectively(see Appendix). Assuming no contribution from the thermal residual stress(or irradiation-induced stress), oth (or girr becomes 2. 4. Single fiber push-out test Assuming a coulomb friction, an interfacial friction stress (tr) an be determined as [25]: Single fiber push-out tests were conducted at room- ture using a nano-indentation test system. A thin-strip 可=-(++m) vith a thickness of 30-220 um was cut from the tensile n where, u is a coefficient of friction and oth. rough and ormad are radial with both surfaces polished by the standard metallographic tech- clamping stresses induced by thermal expansion mismatch, fiber niques to a surface finish of +0.5 um. The specimen was bonded surface roughness and differential swelling. respectively.These
2.2. Neutron irradiation Two types of neutron irradiation campaigns were performed in the High Flux Isotope Reactor (HFIR) at Oak Ridge National Laboratory. The HFIR-14 J fixed-core capsule irradiation was performed in an unshielded removable beryllium position of the HFIR. Irradiation dose was 7.7 dpa assuming 1.0 1025 n/m2 (E > 0.1 MeV) corresponds to one displacement per atom (1 dpa), while the irradiation temperature was 800 C. In contrast, small-capsule rabbit irradiations were performed in the hydraulic tube of the HFIR under the FUN and NERI SiC/SiC series of irradiations. The neutron doses were 0.7–4.2 dpa and irradiation temperatures were in the range of 380–1080 C. Irradiation temperature was measured and controlled in-situ by thermocouples and capsule sweep–gas mixtures for the HFIR-14 experiment. The uncertainty of irradiation temperatures was reasonably low (±20 C). In contrast, irradiation temperatures in rabbit irradiation were estimated by the isochronal annealing of CVD-SiC temperature monitors. In principle, the temperature monitor gives the temperature near the end of the irradiation period with a similar uncertainty to the instrument experiment (±20 C) [15,16]. 2.3. Tensile test Following a guideline of ASTM C1275-00, cyclic unloading/ reloading tensile tests were conducted at ambient-temperature using an electromechanical testing machine. Rectangular miniature tensile specimens with a gauge size of either 20L 4W 1.5T or 20L 2W 1.5T mm3 were prepared. Note that the total specimen length of 50 mm was fixed for all specimens. Specimen size effect is a potential concern when testing these specimens. Considering the structural minimum unit width (or thickness) of the unidirectional composites is larger than the width (or thickness) of the mono fiber bundle (<1 mm), a very minor effect of specimen size on tensile properties, which depend on the axial fiber volume fraction, is expected in the size range of concern [17]. Indeed, no significant size effect was found in this study. The tensile specimen was passively gripped via aluminum grip-end tabs using a pneumatic wedge-type gripping device. Tensile strain was measured by a pair of strain gauges with a gauge length of 5.0 mm, which were adhesively bonded on specimen surfaces of the middle gauge section. A constant crosshead displacement rate was 0.5 mm/min. Young’s modulus was defined as an initial tangent modulus in the tensile stress–strain curve. Proportional limit tensile stress (PLS) was determined as a stress of 5% deviation in stress from initial linearity following ASTM C1275-00. Ultimate tensile strength (UTS) and total elongation were defined as a stress and a strain at composite fracture, respectively. 2.4. Single fiber push-out test Single fiber push-out tests were conducted at room-temperature using a nano-indentation test system. A thin-strip specimen with a thickness of 30–220 lm was cut from the tensile specimen with both surfaces polished by the standard metallographic techniques to a surface finish of 0.5 lm. The specimen was bonded on the specimen holder above a narrow groove. The fiber was then randomly selected and monotonically loaded up to the maximum system allowable load capacity (650 mN). A Berkovich indenter tip was used in this experiment. An applied load rate was 0.05 N/ N s. Details of the fiber push-out test technique have been described elsewhere [18]. Microstructures of the F/M interphase and the pushed-out fiber surface were observed by scanning electron microscopy (SEM) for both as-received and neutron irradiated materials. Of many stress parameters, two interfacial shear properties: (1) an interfacial debond shear strength (ss) as a critical shear stress to induce an interfacial crack along the bonded F/M interface, and (2) an interfacial friction stress (sf) at the debonded interface are de- fined. Both ss and sf are calculated using experimental push-out test results of (1) a debond initiation stress (rd) and (2) a complete debonding and sliding stress (rmax) as defined in Fig. 2. In the following discussion, a compressive stress is denoted with a negative sign. For evaluation of this type of data, various analytical models have been developed [18–24]. Of these models, a non-linear shearlag model proposed by Hsueh [21] becomes a basis to evaluate ss considering the precise stress interaction at the F/M interface: (1) the radial dependence of the axial stresses in both the fiber and the matrix, (2) the shear stress distribution in the matrix and (3) the exact equilibrium equation in relating the tangential stress to the radial stress at the interface. Hsueh further developed a double shear-lag model by separately considering the stress interactions among the fiber, matrix and F/M interphase [22]. One major drawback of the original Hsueh models is limited applicability to isotropic constituents. In the modified double shear-lag model by the authors (see Appendix), anisotropy of the constituents is considered. Besides, residual stresses induced by thermal expansion mismatch and by differential swelling among the fiber, F/M interphase, and matrix can be discussed together in the model, although the effect of dynamic phenomena such as irradiation creep of SiC and PyC was ignored for simplicity. In contrast, Shetty [24] originally discussed the interfacial friction issue for the debonded interface. In the previous work by the authors [25], Shetty’s model was updated for anisotropic composites and this method was applied in this study. The detailed analogy of the modified Hsueh model was discussed elsewhere [22,25]. The resultant form can provide an interfacial debond shear strength (ss) as where the constants k and Z are determined by dimensions and elastic constants of the constituents, and rth and rirr are stress parameters defined by the thermal expansion mismatch and the differential swelling among the fiber, matrix and F/M interphase, respectively (see Appendix). Assuming no contribution from the thermal residual stress (or irradiation-induced stress), rth (or rirr) becomes zero. Assuming a coulomb friction, an interfacial friction stress (sf) can be determined as [25]: sf ¼ l rth r þ rrough r þ rirrad r ; ð2Þ where, l is a coefficient of friction and rth r , rrough r and rirrad r are radial clamping stresses induced by thermal expansion mismatch, fiber surface roughness and differential swelling, respectively. These ss ¼ rfk 2 rd½expðktÞ þ expðktÞ Zðrd þ rth þ rirrÞ½expðktÞ þ expðktÞ 2 expðktÞ expðktÞ ; ð1Þ T. Nozawa et al. / Journal of Nuclear Materials 384 (2009) 195–211 197
T Nozawa et aL / Joumal of Nuclear Materials 384(2009)195-211 Maximum applied load: 600 mN ato:0.05s-1 Indent Progressive debonding contact regime 400 Complete debonding initiation sliding 00 ⅸx|Fi 100 Indenter men holder 0 2000 3000 4000 Displacement curve of the fiber push-out test. Two experimental parameters: (1)a debond initiation load (ad)and (2)a complete debonding and liding load (omax)are defined. clamping stresses have a close relationship with a measured sliding undergoes shrinkage first and swells with increasing neutron flu- stress(σmax)as ence in the direction perpendicular to the deposition plane, while Er(+Ve(oth+rough +girad lexp 2ktcvyrt it shrinks monotonically in the direction parallel to the deposition plane [26]. An empirical fit of the swelling was provided in [27- here e and v are Youngs modulus and Poisson 's ratio Subscripts f List of material properties applied in the analytical model the modified Shetty's model assumes a fiber surrounded by a com- High-density isotropic carbon /2 irradiated and c denote the fiber and the composites, respectively. Note that valuables Irradiated verage,Le→ em, since there is Dens =d1+ lo clear solution for multi-phase system so far. When specimen Elastic modulus(GPa) Ea(1+023小) are sufficiently thin, tr consequently becomes proportional to Poisson's rat Assuming no change Omax/t regardless of material properties of the constituents, (10-°c) Linear swelling Fig. 3 in Ref. [261 (4) Low-density isotropic carbon (glassy carbon)/35] Density (g/cm) =d1+3e) Elastic modulus(GPa) Eo Fig. 4 in Ref. [35] 2.5. Materials properties for calculation o=0.2(assumed) suming no change Linear swelling Fig. 1 in Ref [ 35] The input values used in calculation are empirically obtained CVD-SiC (4) [2, 4, 26-38] and are summarized in Table 2. According to the study Density (g/cm2) by Yan et al. [39. Pyc as an F/M interphase is more graphitic Elastic modulus(GPa) =Eo(1-9e)when the fibers(6-10 nm), while the structure becomes more turbot atic along the radial direction, i.e., near-isotropic. It is noteworthy that the structure of Py c depends significantly on processing con- Poisson's ratio vo=0.2 Assuming no change ditions. Because of this structural uncertainty for PyC, this study CTE(10-/C) Assuming no change considers two types of turbostratic carbon: a high-density carbon Linear swelling 26 and a low-density carbon 35]. A major difference between Hi-Nicalon Type-S/38) the high-density PyC and low-density Py C is in thermal expansivity Density (g/cm") =d(1+3e) as well as density and this issue is discussed later. The actual struc- Elastic modulus(GPa) ure of PyC applied in this study should be classified between these Poissons ratio suming no change two materials suming no change a high-density(2 g/cm)turbostratic carbon with Bacon Linear swelling suming same with anisotropy factor of 1 med for Pyc CVD-SiC cient of thermal expansion(CTE)is assumed to be 5.5 x 10-6/ C. Linear swelling (c) From many irradiation studies, the high-density isotropic carbon Neutron fluence(o)in units of 10-n/m
clamping stresses have a close relationship with a measured sliding stress (rmax) as rmax ¼ Efð1 þ mcÞ Ecmf rth r þ rrough r þ rirrad r exp 2lEcmft rfEfð1 þ mcÞ 1 ; ð3Þ where E and m are Young’s modulus and Poisson’s ratio. Subscripts f and c denote the fiber and the composites, respectively. Note that the modified Shetty’s model assumes a fiber surrounded by a composite average, i.e., the two-phase cylindrical system, since there is no clear solution for multi-phase system so far. When specimens are sufficiently thin, sf consequently becomes proportional to rmax/t regardless of material properties of the constituents, sf ffi rf 2 rmax t : ð4Þ 2.5. Materials properties for calculation The input values used in calculation are empirically obtained [2,4,26–38] and are summarized in Table 2. According to the study by Yan et al. [39], PyC as an F/M interphase is more graphitic near the fibers (6–10 nm), while the structure becomes more turbostratic along the radial direction, i.e., near-isotropic. It is noteworthy that the structure of PyC depends significantly on processing conditions. Because of this structural uncertainty for PyC, this study considers two types of turbostratic carbon: a high-density carbon [26] and a low-density carbon [35]. A major difference between the high-density PyC and low-density PyC is in thermal expansivity as well as density and this issue is discussed later. The actual structure of PyC applied in this study should be classified between these two materials. A high-density (2 g/cm3 ) turbostratic carbon with Bacon anisotropy factor of 1 is assumed for PyC interphase. The coeffi- cient of thermal expansion (CTE) is assumed to be 5.5 106 /C. From many irradiation studies, the high-density isotropic carbon undergoes shrinkage first and swells with increasing neutron fluence in the direction perpendicular to the deposition plane, while it shrinks monotonically in the direction parallel to the deposition plane [26]. An empirical fit of the swelling was provided in [27]. 0 100 200 300 400 500 600 700 0 1000 2000 3000 4000 Displacement [nm] Applied Load [mN] Debond initiation Indenter penetration Progressive debonding regime Complete debonding & sliding Indenter contact Matrix Fiber Matrix Fiber Matrix Fiber Maximum applied load: 600 mN Load rate/load ratio: 0.05 s-1 Matrix Fiber Berkovich indenter Specimen holder with a groove Fig. 2. Schematic of the load-displacement curve of the fiber push-out test. Two experimental parameters: (1) a debond initiation load (rd) and (2) a complete debonding and sliding load (rmax) are defined. Table 2 List of material properties applied in the analytical model. Valuables Non-irradiated Irradiated High-density isotropic carbon [26] Density (g/cm3 ) d0 = 1.9 =d0(1 + 3e) Elastic modulus (GPa) E0 = 25 =E0(1 + 0.23/) Poisson’s ratio m0 = 0.2 Assuming no change CTE (106 /C) a0 = 5.5 Assuming no change Linear swelling – Fig. 3 in Ref. [26] Low-density isotropic carbon (glassy carbon) [35] Density (g/cm3 ) d0 = 1.5 =d0(1 + 3e) Elastic modulus (GPa) E0 = 25 Fig. 4 in Ref. [35] Poisson’s ratio m0 = 0.2 (assumed) Assuming no change CTE(106 /C) a0 = 2.8 Assuming no change Linear swelling – Fig. 1 in Ref. [35] CVD-SiC [4] Density (g/cm3 ) d0 = 3.2 =d0(1 + 3e) Elastic modulus (GPa) E0 = 460 =E0(1 .9e) when Tirr 1000 C Poisson’s ratio m0 = 0.2 Assuming no change CTE(106 /C) a0 = 4.4 Assuming no change Linear swelling – Fig. 22 in Ref. [4] Hi-NicalonType-S [38] Density (g/cm3 ) d0 = 3.1 =d0(1 + 3e) Elastic modulus (GPa) E0 = 420 =E0(1 20.9e) when Tirr < 1000 C Poisson’s ratio m0 = 0.2 Assuming no change CTE(106 /C) a0 = 5.1 Assuming no change Linear swelling – Assuming same with CVD-SiC Linear swelling (e). Neutron fluence (/) in units of 1025 n/m2 . 198 T. Nozawa et al. / Journal of Nuclear Materials 384 (2009) 195–211
T Nozawa et aL/Journal of Nuclear Materials 384(2009)195-211 Material properties of low-density turbostratic carbon were of swelling saturates with increasing neutron dose(1 dpa). For summarized by Virgil'ev and Lebedev [35]. The density is x1.5 g/ irradiation temperatures less than 1000C, swelling is a strong cm. The elastic modulus ranged from 25 to 30 GPa. The coefficient function of temperature, decreasing with increasing temperature of thermal expansion of the low-density carbon is about I. The irradiation-induced change of Youngs modulus of SiC 2.8x 10/C. Neutron irradiation causes a significant shrinkage can be well-described by the Tersoff potential and the resulting when Tirr =140-750C, resulting in densification. With formation irradiation effects are well documented [4, 40. Because of lack of of a new system of micropores, the thermal expansivity of low- irradiation data on advanced SiC fibers, the post-irradiation prop density carbon reportedly decreases by irradiation. Because of lim- erties of Hi-Nicalon" Type-S fiber are presently assumed to be ited irradiation data available this study assumes no change of equivalent to those of CVD-SiC. However it is worth noting that re thermal expansivity by irradiation. The elastic modulus of low- cent data [41 identified some discrepancies in post-irradiation density carbon first decreases by irradiation and gradually in- properties of Sic fibers with CVD-Sic. The Poisson s ratio and the creases with increasing neutron fluence up to 1 x 10 n/cm. Be- thermal expansivity of Sic are assumed to be unchanged by yond this fluence, the elastic modulus seems to be saturated irradiation. (15% increase of the non-irradiated value Material properties of composites can be estimated by the rule High-crystallinity and high-purity Sic swells by irradiation with of mixtures, which is modified by Hashin for the transversely iso- the accumulation of radiation-induced defects and the magnitude tropic system[ 42. In this calculation for multilayer composites, (a)UD Hi-Nicalon Type-S/ PyC/CVI-SiC 400fNon-irradiated 1.8da3809 7.7dpa800°c 1.0dpa1000° (b)UD Hi-Nicalon Type-S/ ML/ CVI-Si 37dpa640°C Non-irradiated 77da800°C 42dpa1080°c 200 04 g stress-strain curves of as-received and neutron-irradiated unidirectional Hi-Nicalon Type-s/CVI-SiC composites with either(a)a Pyc monolayer or b)a pyc/s
Material properties of low-density turbostratic carbon were summarized by Virgil’ev and Lebedev [35]. The density is 1.5 g/ cm3 . The elastic modulus ranged from 25 to 30 GPa. The coefficient of thermal expansion of the low-density carbon is about 2.8 106 /C. Neutron irradiation causes a significant shrinkage when Tirr = 140–750 C, resulting in densification. With formation of a new system of micropores, the thermal expansivity of lowdensity carbon reportedly decreases by irradiation. Because of limited irradiation data available, this study assumes no change of thermal expansivity by irradiation. The elastic modulus of lowdensity carbon first decreases by irradiation and gradually increases with increasing neutron fluence up to 1 1021 n/cm3 . Beyond this fluence, the elastic modulus seems to be saturated (15% increase of the non-irradiated value). High-crystallinity and high-purity SiC swells by irradiation with the accumulation of radiation-induced defects and the magnitude of swelling saturates with increasing neutron dose (>1 dpa). For irradiation temperatures less than 1000 C, swelling is a strong function of temperature, decreasing with increasing temperature. [4]. The irradiation-induced change of Young’s modulus of SiC can be well-described by the Tersoff potential and the resulting irradiation effects are well documented [4,40]. Because of lack of irradiation data on advanced SiC fibers, the post-irradiation properties of Hi-NicalonTM Type-S fiber are presently assumed to be equivalent to those of CVD-SiC. However it is worth noting that recent data [41] identified some discrepancies in post-irradiation properties of SiC fibers with CVD-SiC. The Poisson’s ratio and the thermal expansivity of SiC are assumed to be unchanged by irradiation. Material properties of composites can be estimated by the rule of mixtures, which is modified by Hashin for the transversely isotropic system [42]. In this calculation for multilayer composites, Fig. 3. Typical tensile stress–strain curves of as-received and neutron-irradiated unidirectional Hi-NicalonTM Type-S/CVI-SiC composites with either (a) a PyC monolayer or (b) a PyC/SiC multilayer interphase. T. Nozawa et al. / Journal of Nuclear Materials 384 (2009) 195–211 199
T Nozawa et aL /Joumal of Nuclear Materials 384(2009)195-211 Tensile properties of as-received and irradiated SiC/SiC composites. ampaign condition E(GPa) AE(%) Ours(MPa)dus(3) Ops (MP)ensile stress Failure strain 2366(46) 97(35) 0.400.14)+150 HFIR-14J 7.7dpa800°C 12381(26) 217(36) 0.38(002)+138 Multilayer None 375(19 271(51) 232(47) 009(003) 14 HFIR-NER °421(53)+12193(61) 93(61) 004(0.02)-56 96(57) HFIR-N 136(-) 0.14( MTROOM-95U 19(51) 223(36) MTROOM-95U °C338(41)-10231(51) 166(35) No dat No data HFIR-14J °356(39)-5302(-) 11 0.2(-)+3 Note: Numbers in parenthesis indicate one standard deviation. b E200 Neutron Dose [dpa-SIC Neutron Dose [dpa-SICI 06 0 Neutron Dose (dpa-Sic] Neutron Dose [dpa-SICl s/CVI-SiC composites with esther a byc monolayer or ycl SiC multilayer interphase
Table 3 Tensile properties of as-received and irradiated SiC/SiC composites. Interphase Irradiation campaign Irradiation condition Young’s modulus Ultimate tensile strength Proportional limit tensile stress Failure strain # of valid tests E (GPa) DE (%) rUTS (MPa) DrUTS (%) rPLS (MPa) DrPLS etotal (%) Detotal (%) Monolayer None – 362 (34) – 319 (64) – 228 (44) – 0.16 (0.06) – 25 JMTR00M-95U 1.0 dpa 800 C 320 (46) 12 366 (46) +15 197 (35) 14 0.41 (0.17) +156 8 JMTR00M-95U 1.0 dpa 1000 C 289 (16) 20 418 (24) +31 183 (45) 20 No data No data 6 HFIR-FUN 1.8 dpa 380 C 354 (28) 2 366 (82) +15 190(23) 17 0.40 (0.14) +150 3 HFIR-14 J 7.7 dpa 800 C 320 (11) 12 381 (26) +19 217 (36) 5 0.38 (0.02) +138 4 Multilayer None – 375 (19) – 271 (51) – 232 (47) – 0.09 (0.03) – 14 HFIR-NERI 0.7 dpa 700 C 406 (36) +8 234 (12) 14 203 (32) 12 0.06 (0.01) 33 2 HFIR-NERI 0.7 dpa 1030 C 421 (53) +12 193 (61) 29 193 (61) 17 0.04 (0.02) 56 3 HFIR-NERI 3.7 dpa 640 C 366 (52) 2 296 (57) +9 211 (53) 9 0.13 (0.07) +44 3 HFIR-NERI 4.2 dpa 1080 C 433 (54) +15 257 () 5 136 () 41 0.14 () +56 3 JMTR00M-95U 1.0 dpa 800 C 355 (65) 5 319 (51) +18 223 (36) 4 0.15 (0.07) +67 4 JMTR00M-95U 1.0 dpa 1000 C 338 (41) 10 231 (51) 15 166 (35) 28 No data No data 3 HFIR-14J 7.7 dpa 800 C 356 (39) 5 302 () +11 186 (42) 20 0.12 () +33 4 Note: Numbers in parenthesis indicate one standard deviation. Fig. 4. Effects of neutron dose on (a) Young’s modulus, (b) proportional limit stress, (c) ultimate tensile strength and (d) total elongation for unidirectional Hi-NicalonTM TypeS/CVI-SiC composites with either a PyC monolayer or a PyC/SiC multilayer interphase. 200 T. Nozawa et al. / Journal of Nuclear Materials 384 (2009) 195–211
T Nozawa et aL/Journal of Nuclear Materials 384(2009)195-211 400+ Non-irradiated 0200400600800 Irradiation Temperature[] Irradiation Temperature r'c d Non-irradiated Non-irradiated 0 200400 0001200 Irradiation Temperature rcl Irradiation Temperature r'C) Fig. 5. Effects( n temperature on(a) Youngs modulus, (b) proportional limit stress. (c)ultimate tensile strength and (d) total elongation fo licalonType-S/CVI-SiC composites with either a Py c monolayer or a Py c/ sic multilayer interphase. The data subsets of neutron dose range in 0.7-4.. or unidirectional Hi 00 f Pyc monolayer(720nm-thick) PyC/SiC multilayer and sliding Debond initiation ample thickness: 200um Displacement [nm ig. 6. Typical push-out fracture behaviors of Hi-Nicalon Type-s/CVI-SiC composites with either a Pyc monolayer or a Pyc sic multilayer interphase
Fig. 5. Effects of irradiation temperature on (a) Young’s modulus, (b) proportional limit stress, (c) ultimate tensile strength and (d) total elongation for unidirectional HiNicalonTM Type-S/CVI-SiC composites with either a PyC monolayer or a PyC/SiC multilayer interphase. The data subsets of neutron dose range in 0.7–4.2 dpa were plotted. Fig. 6. Typical push-out fracture behaviors of Hi-Nicalon Type-S/CVI-SiC composites with either a PyC monolayer or a PyC/SiC multilayer interphase. T. Nozawa et al. / Journal of Nuclear Materials 384 (2009) 195–211 201
T Nozawa et aL/Joumal of Nuclear Materials 384(2009)195-211 Push-out test results for as-received and irradiated Sic/siC composites. Interphase 2871(882) 895(892) Non-irradiated (batch #2) 如Gm 0. dpa380°C(HFR-FUN) 312 18dpa380°(HFR-FUN) 520(26 842(400 1055(245) 7.7dpa8oo°(HFR-14 1186(199) 1189(198 467(311) 3544(498) Multilayer 3085(835) 085(835) 3554(1 07dpa740°c(HFR-NER) 974(1027) 3078(1472) 07dpa1030°(HFR-NER) 2 4985(180) 3.7dpa640°C( HFIR-NER) 23(932 70(650) 2619(536) 4.2 dpa 1080"C(HFIR-NERI 2016(62) 7.7dpa800°(HHR-14 2223(472) 3040(83 2550(58 2982(536 Note: Numbers in parenthesis indicate one standard deviation. Table 5 terfacial shear properties of as-received and irradiated SiC/SiC composites. radiation campaig Irradiation condition Interfacial friction st one(batch #2) FIR-FUN 8dpa380°C +16 FIR-FUN 97(14) 32(8) HFIR- 7.7dpa800° 148( 185(1 24 56(6) 338(45) 4(46) 267(18) 279(44) 215(49) HFIR-NERI 270(27) 81(2) HFIR-14 277(30) 329(29) 111(24) ote: Numbers in parenthesis indicate one standard deviation
Table 4 Push-out test results for as-received and irradiated SiC/SiC composites. Interphase Irradiation condition Sample thickness (lm) Debond initiation stress (MPa) Sliding stress (MPa) # of valid tests Monolayer Non-irradiated (batch #1) 50 2871 (882) 2895 (892) 22 100 2611 (1033) 3351 (1071) 33 200 2875 (734) – 18 Non-irradiated (batch #2) 50 891 (423) 910 (440) 22 100 1576 (619) 2071 (911) 31 200 1870 (537) 4049 (917) 33 0.8dpa 380 C (HFIR-FUN) 53 1892 (642) 1965 (715) 23 120 2458 (1168) 3362 (1415) 45 153 3170 (886) 4238 (696) 24 189 1940 (732) 2837 (1186) 13 212 2645 (954) 4091 (891) 50 1.8dpa 380 C (HFIR-FUN) 58 863 (307) 1059 (398) 14 87 1520 (265) 1646 (361) 30 114 842 (400) 910 (457) 74 145 1055 (245) 1155 (278) 22 7.7 dpa 800 C (HFIR-14J) 40 1186 (199) 1189 (198) 16 80 1467 (311) 1712 (396) 23 118 1680 (426) 2821 (823) 24 210 1702 (284) 3544 (498) 10 Multilayer Non-irradiated 30 3085 (835) 3085 (835) 5 50 3554 (1090) 4205 (1107) 12 100 3518 (680) – 18 200 3355 (830) – 16 0.7 dpa 740 C (HFIR-NERI) 25 3006 (963) – 49 100 2974 (1027) – 22 137 3078 (1472) – 26 154 3323 (654) – 22 0.7 dpa 1030 C (HFIR-NERI) 100 2198 (1119) – 26 150 3056 (773) 4985 (180) 25 185 2155 (735) – 25 3.7 dpa 640 C (HFIR-NERI) 40 2457 (982) 2870 (650) 17 165 2836 (732) – 24 190 2619 (536) – 24 4.2 dpa 1080 C (HFIR-NERI) 75 1789 (402) 2016 (62) 28 150 2124 (692) – 30 210 2098 (375) – 28 7.7 dpa 800 C (HFIR-14J) 60 2188 (399) 2223 (472) 21 90 3040 (839) – 18 120 2550 (585) – 33 220 2982 (536) – 22 Note: Numbers in parenthesis indicate one standard deviation. Table 5 Interfacial shear properties of as-received and irradiated SiC/SiC composites. Interphase Irradiation campaign Irradiation condition Interfacial debond shear strength Interfacial friction stress High-density PyC case Low-density PyC case ss (MPa) Dss (%) ss (MPa) Dss (%) sf (MPa) Dsf (%) Monolayer None (batch #1) – 234 (45) – 311 (47) – 115 (21) – None (batch #2) – 125 (35) – 172 (27) – 61 (1) – HFIR-FUN 0.8 dpa 380 C 208 (36) +16 267 (35) +10 63 (8) 28 HFIR-FUN 1.8 dpa 380 C 97 (14) 46 165 (15) 32 32 (8) 64 HFIR-14 J 7.7 dpa 800 C 148 (15) 18 185 (14) 24 56 (6) 36 Multilayer None – 338 (45) – 424 (46) – 267 (18) – HFIR-NERI 0.7 dpa 740 C 330 (51) 2 402 (51) 5 >210 <21 HFIR-NERI 0.7 dpa 1030 C 240 (51) 29 320 (50) 25 100 (3) 63 HFIR-NERI 3.7 dpa 640 C 279 (44) 17 346 (43) 18 215 (49) 19 HFIR-NERI 4.2 dpa 1080 C 201 (28) 41 270 (27) 36 81 (2) 70 HFIR-14J 7.7 dpa 800 C 277 (30) 18 329 (29) 22 111 (24) 58 Note: Numbers in parenthesis indicate one standard deviation. 202 T. Nozawa et al. / Journal of Nuclear Materials 384 (2009) 195–211
T Nozawa et aL/Journal of Nuclear Materials 384(2009)195-211 total thickness of the Pyc layers was applied. As mentioned later. rized as follows:(1)very minor change in Young,'s modulus by residual stresses at the interface for the multilayer composites irradiation, (2)enhanced quasi-ductility beyond the proportional are nearly compatible by taking the case of monolayer composites limit, i. e, no embrittlement by irradiation, resulting in apparent with a total thickness of multilayered Pyc. increase of total elongation and 3)reduction of PLS. The tensile ngth was nearly unchanged for monolayer composites at any 3. Results irradiation temperatures. Similarly no significant decrease of UTs was observed for multilayer composites when Tirr 1000C was specified. Fig 4 plots tensile data as a function of neutron dose. In the fig Fig3 shows typical stress-strain curves of non-irradiated and ure, error bars indicate t1 standard deviation. The proportional irradiated Hi-Nicalon" M Type-s/CVI-Sic composites with either a limit tensile stress slightly decreases even at low neutron doses Pyc monolayer or a Pyc sic multilayer interphase and Table 3 lists of 0. 7 dpa and approaches a constant(-200 MPa)with increasing reduced tensile properties. Data from an irradiation experiment in neutron dose for both interphase types. Tensile strength and total the japan materials test reactor (MTR, Oarai, Japan) for the identi longation conversely tend to increase and approach a constant cal materials are included [12]. Key common findings are summa- with increasing neutron dose. The increment of UTS for monolayer s (30%)was greater than that of mult (20%) Similarly, the increment of total elongation for monolayer composites(160%)was almost double of multilayer composites (70%) Meanwhile, their Youngs moduli seem nearly unchanged 500 Monolayer by irradiation. Large data scatter is due primarily to the variance of porosity and fiber volume fraction. The data subsets in the fluence range of 0.7-4.2 dpa in Fig. 4 冒4005~+740°℃ were extracted and separately plotted as a function of irradiation temperature(Fig. 5). In Fig. 5, Young,s moduli for both interface 640°C 800° structure types appeared unchanged by irradiation accounting for 10G0c the wide data scatter(30%). In contrast, the pls slightly de- creased after irradiation at all temperatures, however, the temper 1080° ature-dependence seems very minor even at 1000C. This observed for both interphase types tested. Additionally, within 80° 800°C the limited data set provided, it appears that a further decrease of the pls has taken place when Tirr >1000C. Both UTS and total 380° elongation were nearly unchanged by irradiation for multilayer composites when Tirr 1000C, while they significantly increased by irradiation for monolayer composites. When Tirr>1000C,a clear difference between the monolayer and multilayer composite was obtained: decreases of UTS for multilayer composites vs no major deterioration for monolayer composites. Neutron Dose (dpa-SIc] 3. 2. Interfacial shear properties Fig 6 shows comparison of typical load-displacement during fiber push-out process between monolayer and multilayer composites As schematically shown in Fig. 2, the fiber push-out Monolayer process generally has the following four stages: (1)elastic defor mation with a tiny plastic deformation, (2)debond initiation at the interface from the loading surface, 3) progressive debond cou- 三c pled with interfacial friction on the debonded fiber surface and ( 4) 740°C complete fiber debonding and sliding. When using a sharp indenter tip. the initial segment of the load-displacement curve becomes ----=- non-linear due to the plastic deformation below the indenter. Be- yond debond initiation, the displacement rapidly increases and a load-displacement curve shows a secondary non-linear segment. In Fig. 6, a concave shape of the secondary non-linear segment 1030%C 800°C for multilayer composites is due to the inherently high interfacial 1080° friction resulting from the rough surface of the fiber. The limited load capacity of the test system does not allow the complete deb- °C 2800 onding and sliding of the composites with exceptionally high de- bond initiation stress. More than 80% of measurements of oma 2380° or multilayer composites were invalid when specimen thickness was >100 um. The slightly convex shape observed for monolayer composites is due to the comparably low friction stress. a plateau Neutron Dose [ dpa-SIC beyond complete debonding, i.e, sliding of the pushed-out fiber. was observed for the case of thick Py C monolayer composites. ress n adiatie tempera ue are nted ien ghe figure The further load increment after complete sliding indicates a phys- Irradiation-induced change of elastic properties of Sic and PyC was considered. No ical contact of the indenter sides with the Sic matrix when using a contribution from swelling was however considered Glassy carbon was assumed. conic indenter
total thickness of the PyC layers was applied. As mentioned later, residual stresses at the interface for the multilayer composites are nearly compatible by taking the case of monolayer composites with a total thickness of multilayered PyC. 3. Results 3.1. Tensile properties Fig. 3 shows typical stress–strain curves of non-irradiated and irradiated Hi-NicalonTM Type-S/CVI-SiC composites with either a PyC monolayer or a PyC/SiC multilayer interphase and Table 3 lists reduced tensile properties. Data from an irradiation experiment in the Japan materials test reactor (JMTR, Oarai, Japan) for the identical materials are included [12]. Key common findings are summarized as follows: (1) very minor change in Young’s modulus by irradiation, (2) enhanced quasi-ductility beyond the proportional limit, i.e., no embrittlement by irradiation, resulting in apparent increase of total elongation and (3) reduction of PLS. The tensile strength was nearly unchanged for monolayer composites at any irradiation temperatures. Similarly no significant decrease of UTS was observed for multilayer composites when Tirr 1000 C was specified. Fig. 4 plots tensile data as a function of neutron dose. In the figure, error bars indicate ±1 standard deviation. The proportional limit tensile stress slightly decreases even at low neutron doses of 0.7 dpa and approaches a constant (200 MPa) with increasing neutron dose for both interphase types. Tensile strength and total elongation conversely tend to increase and approach a constant with increasing neutron dose. The increment of UTS for monolayer composites (30%) was greater than that of multilayer composites (20%). Similarly, the increment of total elongation for monolayer composites (160%) was almost double of multilayer composites (70%). Meanwhile, their Young’s moduli seem nearly unchanged by irradiation. Large data scatter is due primarily to the variance of porosity and fiber volume fraction. The data subsets in the fluence range of 0.7–4.2 dpa in Fig. 4 were extracted and separately plotted as a function of irradiation temperature (Fig. 5). In Fig. 5, Young’s moduli for both interface structure types appeared unchanged by irradiation accounting for the wide data scatter (30%). In contrast, the PLS slightly decreased after irradiation at all temperatures, however, the temperature-dependence seems very minor even at 1000 C. This is observed for both interphase types tested. Additionally, within the limited data set provided, it appears that a further decrease of the PLS has taken place when Tirr > 1000 C. Both UTS and total elongation were nearly unchanged by irradiation for multilayer composites when Tirr 1000 C, a clear difference between the monolayer and multilayer composites was obtained: decreases of UTS for multilayer composites vs. no major deterioration for monolayer composites. 3.2. Interfacial shear properties Fig. 6 shows comparison of typical load-displacement curves during fiber push-out process between monolayer and multilayer composites. As schematically shown in Fig. 2, the fiber push-out process generally has the following four stages: (1) elastic deformation with a tiny plastic deformation, (2) debond initiation at the interface from the loading surface, (3) progressive debond coupled with interfacial friction on the debonded fiber surface and (4) complete fiber debonding and sliding. When using a sharp indenter tip, the initial segment of the load-displacement curve becomes non-linear due to the plastic deformation below the indenter. Beyond debond initiation, the displacement rapidly increases and a load-displacement curve shows a secondary non-linear segment. In Fig. 6, a concave shape of the secondary non-linear segment for multilayer composites is due to the inherently high interfacial friction resulting from the rough surface of the fiber. The limited load capacity of the test system does not allow the complete debonding and sliding of the composites with exceptionally high debond initiation stress. More than 80% of measurements of rmax for multilayer composites were invalid when specimen thickness was >100 lm. The slightly convex shape observed for monolayer composites is due to the comparably low friction stress. A plateau beyond complete debonding, i.e., sliding of the pushed-out fiber, was observed for the case of thick PyC monolayer composites. The further load increment after complete sliding indicates a physical contact of the indenter sides with the SiC matrix when using a conic indenter. Fig. 7. Effects on neutron dose on (a) the interfacial debond shear strength and (b) the interfacial friction stress. Irradiation temperatures are noted in the figure. Irradiation-induced change of elastic properties of SiC and PyC was considered. No contribution from swelling was however considered. Glassy carbon was assumed. T. Nozawa et al. / Journal of Nuclear Materials 384 (2009) 195–211 203
T Nozawa et aL /Joumal of Nuclear Materials 384(2009)195-211 Unirradiated Irrad 380C, 1.dpa Irrad 800C, 7.7dpa N/A 2 Fig 8. Typical micrographs of pushed-out fiber surfaces The least-square fits by Eqs. (1)and (4)can provide an interfa- primary crack occurred in the inner most Pyc layer adjacent the fi- cial debond shear strength(ts)and an interfacial friction stress(tr), ber for both cases. For monolayer composites, there was no signif- respectively, from the experimental push-out data in Table 4 Table icant change in cracking path by neutron irradiation. The pushed 5 lists ts and tr of monolayer and multilayer composites and Fig. 7 out fiber surfaces were smooth regardless of the irradiation, indi- plots interfacial shear property data as a function of neutron dose cating crack propagation along the fiber surface. In contrast, the (for the low-density Py C case). Note that the preliminary analysis rough fiber surface for the as-received multilayer composites im- in this study disregarded the contribution from swelling because plies crack propagation within PyC interphase adjacent the fiber. the actual swelling in composite forms depend significantly on The changing crack path: the crack path 'within interphase before irradiation assisted stress relaxation: The importance of irradiation irradiation Vs the crack path along, the fiber/carbon interface after creep was pointed out by the bend stress relaxation test by Katoh irradiation (7.7 dpa at 800C)was identified. et al.[43 The anisotropic swelling of Sic under stress has also een reported [44]. Considering precisely the effect of irradia- 4.Discussion tion-induced dimensional change with combined effects of dy namic irradiation effects is presently very difficult due to many 4.1. Clamping stresses at the fiber/matrix interface uncertain factors and scarcity of data. Meanwhile, simple estima tion by applying a minimum value of omax gives a lower limit (tr>210 MPa)as plotted in the figure with an arrow in Fig. 7, residual atre sig ind a tpe em nin ten although the interfacial friction following irradiation at 740C face. These clamping stresses are generally induced by the thermal to 0.7 dpa could not be quantified due to the load limitation of expansion mismatch among the fiber, F/M interphase and matrix, the test system and by roughness of the fiber surface [20 Additionally, the sec From Fig. 7, it is apparent that shear properties at the F/M inter- ondary stresses induced by swelling and stress relaxation by irra- face depend on neutron dose. They apparently changed after irra- diation creep may impose added complexity under neutron diation,however,the dose dependence is not so significant when irradiation For advanced Sic/SiC composites, the fiber and matrix 1080.C. Both ts and tr decrease by irradiation in low dose are highly-crystalline Sic and it has been believed that their irradi- region and remain nearly unchanged with further increasing neu- ation effects are essentially similar. However, recent study identi tron dose regardless of the interface structure types, e.g. monolayer fied some differences in defect accumulation between fibers and or multilayer. Of particular emphasis is that both ts and tr for mul- purer SiC matrix, implying potential difference in swelling [41]. tilayer composites were most affected when Tirr-1080C. How- which probably gives a residual stress at the F/M interface by irra- ever, it is worth noting that the interfacial shear properties for diation. In spite of the presence of CTE mismatch between the Sic irradiated multilayer composites are still above average for mono- fiber and matrix, dose-dependent dimensional change of PyC layer composites. Even if the interfacial shear stresses decrease by would inevitably impose substantial stresses at the F/M interface irradiation, the high shear strength is still sufficient for the load- [26]. In short, the radiation-induced stresses would undoubtedly sharing at the F/M interface. Presently, no irradiation data is avail- influence the results and ultimately need to be consideree able for monolayer composites when Tirr>1000C and therefore it A clamping stress working at the F/M interface can be qualita- should be noted that it is unclear whether this is specific to the tively predicted under the simple assumption by adopting the multilayer interface case. Another uncertainty to be pointed out four-phase( fiber core/PyC interphase/Sic matrix/composites )cyl- is a reduction of interfacial friction by irradiation at 800C up to inder model [45-47]. This model discusses(1)anisotropy of the 7.7 dpa for multilayer composites. constituents, (2)the effect of interlayer thickness, (3)contributions
The least-square fits by Eqs. (1) and (4) can provide an interfacial debond shear strength (ss) and an interfacial friction stress (sf), respectively, from the experimental push-out data in Table 4. Table 5 lists ss and sf of monolayer and multilayer composites and Fig. 7 plots interfacial shear property data as a function of neutron dose (for the low-density PyC case). Note that the preliminary analysis in this study disregarded the contribution from swelling because the actual swelling in composite forms depend significantly on irradiation assisted stress relaxation; The importance of irradiation creep was pointed out by the bend stress relaxation test by Katoh et al. [43]. The anisotropic swelling of SiC under stress has also been reported [44]. Considering precisely the effect of irradiation-induced dimensional change with combined effects of dynamic irradiation effects is presently very difficult due to many uncertain factors and scarcity of data. Meanwhile, simple estimation by applying a minimum value of rmax gives a lower limit (sf > 210 MPa) as plotted in the figure with an arrow in Fig. 7, although the interfacial friction following irradiation at 740 C up to 0.7 dpa could not be quantified due to the load limitation of the test system. From Fig. 7, it is apparent that shear properties at the F/M interface depend on neutron dose. They apparently changed after irradiation, however, the dose dependence is not so significant when Tirr = 1080 C. Both ss and sf decrease by irradiation in low dose region and remain nearly unchanged with further increasing neutron dose regardless of the interface structure types, e.g. monolayer or multilayer. Of particular emphasis is that both ss and sf for multilayer composites were most affected when Tirr = 1080 C. However, it is worth noting that the interfacial shear properties for irradiated multilayer composites are still above average for monolayer composites. Even if the interfacial shear stresses decrease by irradiation, the high shear strength is still sufficient for the loadsharing at the F/M interface. Presently, no irradiation data is available for monolayer composites when Tirr > 1000 C and therefore it should be noted that it is unclear whether this is specific to the multilayer interface case. Another uncertainty to be pointed out is a reduction of interfacial friction by irradiation at 800 C up to 7.7 dpa for multilayer composites. Fig. 8 exhibits typical micrographs of pushed-out fiber surfaces of as-received and neutron-irradiated materials. Importantly the primary crack occurred in the inner most PyC layer adjacent the fi- ber for both cases. For monolayer composites, there was no significant change in cracking path by neutron irradiation. The pushedout fiber surfaces were smooth regardless of the irradiation, indicating crack propagation along the fiber surface. In contrast, the rough fiber surface for the as-received multilayer composites implies crack propagation within PyC interphase adjacent the fiber. The changing crack path: the crack path ‘within’ interphase before irradiation vs. the crack path ‘along’ the fiber/carbon interface after irradiation (7.7 dpa at 800 C) was identified. 4. Discussion 4.1. Clamping stresses at the fiber/matrix interface Interfacial shear properties are significantly dependent on residual compressive (clamping) stresses induced at the F/M interface. These clamping stresses are generally induced by the thermal expansion mismatch among the fiber, F/M interphase and matrix, and by roughness of the fiber surface [20]. Additionally, the secondary stresses induced by swelling and stress relaxation by irradiation creep may impose added complexity under neutron irradiation. For advanced SiC/SiC composites, the fiber and matrix are highly-crystalline SiC and it has been believed that their irradiation effects are essentially similar. However, recent study identi- fied some differences in defect accumulation between fibers and purer SiC matrix, implying potential difference in swelling [41], which probably gives a residual stress at the F/M interface by irradiation. In spite of the presence of CTE mismatch between the SiC fiber and matrix, dose-dependent dimensional change of PyC would inevitably impose substantial stresses at the F/M interface [26]. In short, the radiation-induced stresses would undoubtedly influence the results and ultimately need to be considered. A clamping stress working at the F/M interface can be qualitatively predicted under the simple assumption by adopting the four-phase (fiber core/PyC interphase/SiC matrix/composites) cylinder model [45–47]. This model discusses (1) anisotropy of the constituents, (2) the effect of interlayer thickness, (3) contributions from the thermal expansion mismatch and (4) irradiation effects. Of many models, the analytical method proposed by Oel [47] is Fig. 8. Typical micrographs of pushed-out fiber surfaces under various irradiation conditions. 204 T. Nozawa et al. / Journal of Nuclear Materials 384 (2009) 195–211