Part B: engineering ELSEVIER Composites: Part B 30(1999)631-646 www.elsevier,com/locate/compositesb Characterization techniques for composites and other advanced materials A E. Pasto, D n. Braski, T.R. Watkins, W.D. Porter, E. Lara-Curzio,SB McSpadden High Temperature Materials Laboratory, Oak Ridge National Laboratory, Oak Ridge, TN 37831-6062, USA Abstract One of the key requirements in developing composites and other advanced materials is generation of a good understanding of the relationships between composition and structure on the one hand, and properties and behavior on the other. Another key requirement is application of this understanding to develop a material with the desired properties. a third key requirement is to understand the new material's failure mechanisms. All of these are wrapped in the term characterization, which is the subject of this paper. Application of umerous materials characterization techniques to the study of ceramic composites is described. o 1999 Elsevier Science Ltd. All rights Keywords: Materials characterization techniques 1. Introduction microstructural and microchemical analysis, equipment for One of the key requirements in developing composites ties of materials to elevated temperatures, X-ray amdp roper- measurement of the thermophysical and mechanical utron nd other advanced materials is generation of a good under- diffraction for structure and residual stress analysis, high- standing of the relationships between composition and speed grinding machines, and measurement of component structure on the one hand, and properties and behavior on shape, tolerances, surface finish and friction and wear prop- the other. Another key requirement is application of this erties. Users willing to publish the results of their work can understanding to develop a material with the desired proper perform no-cost materials characterization here, under the ties. A third key requirement is to understand the new sponsorship of OTT. materials failure mechanisms. All of these are wrapped in Over half of the work performed in the hTMl is the term "characterization, which is the subject of this sponsored by other R&d programs, from DOE and other paper. agencies. This research is also primarily characterization, The high temperature materials laboratory(HTML) is and it often involves composites. Composites work has art of the metals and ceramics division of Oak Ridge been sponsored by the continuous fiber ceramic composite National Laboratory (ORNL), where it serves the primary (CFCC) program, the ceramic technology for advanced heat purpose of providing equipment and staff to perform mate- engine program, now called the propulsion system materials rials characterization. It is a US Department of Energy program, and others (DOE)-designated National User Facility designed to assist The following sections will highlight some of the capabil- American industries, universities and governmental agen- ities resident at HTML, and illustrations of their application ies develop advanced materials, by providing a skilled to composite materials staff and numerous sophisticated, often one-of-a-kind pieces of materials characterization equipment. HTML is ponsored by DOE's office of transportation technologies 2. Microstructure/microchemistry characterization it is a 64, 500 sq. ft. building on the ORNL site, in which 2.1. Auger analysis of the interfaces in a composite eside six user centers", which are clusters of specialized equipments revolving around a specific type of properties The toughness offiber-reinforced ceramic-matrix compo- measurements. Available are electron microscopy for only controlled by the properties of the fiber and matrix materials, but the bonding forces between them Paradoxically, the strongest bonds do not increase tough Corresponding author. Tel: +1-423-574-5123; fax:+1-423-574-4913 E-imail address: pastoae @ornl. gov(AE Pasto) sliding of the fibers during fracture are more effective. A 1359-8368/99/- see front matter 1999 Elsevier Science Ltd. All rights reserved Pl:S13598368(99)00040-2
Characterization techniques for composites and other advanced materials A.E. Pasto*, D.N. Braski, T.R. Watkins, W.D. Porter, E. Lara-Curzio, S.B. McSpadden High Temperature Materials Laboratory, Oak Ridge National Laboratory, Oak Ridge, TN 37831-6062, USA Abstract One of the key requirements in developing composites and other advanced materials is generation of a good understanding of the relationships between composition and structure on the one hand, and properties and behavior on the other. Another key requirement is application of this understanding to develop a material with the desired properties. A third key requirement is to understand the new material’s failure mechanisms. All of these are wrapped in the term “characterization”, which is the subject of this paper. Application of numerous materials’ characterization techniques to the study of ceramic composites is described. q 1999 Elsevier Science Ltd. All rights reserved. Keywords: Materials characterization techniques 1. Introduction One of the key requirements in developing composites and other advanced materials is generation of a good understanding of the relationships between composition and structure on the one hand, and properties and behavior on the other. Another key requirement is application of this understanding to develop a material with the desired properties. A third key requirement is to understand the new material’s failure mechanisms. All of these are wrapped in the term “characterization”, which is the subject of this paper. The high temperature materials laboratory (HTML) is part of the metals and ceramics division of Oak Ridge National Laboratory (ORNL), where it serves the primary purpose of providing equipment and staff to perform materials characterization. It is a US Department of Energy (DOE)-designated National User Facility designed to assist American industries, universities and governmental agencies develop advanced materials, by providing a skilled staff and numerous sophisticated, often one-of-a-kind, pieces of materials characterization equipment. HTML is sponsored by DOE’s office of transportation technologies (OTT), energy efficiency and renewable energy. Physically, it is a 64,500 sq. ft. building on the ORNL site, in which reside six “user centers”, which are clusters of specialized equipments revolving around a specific type of properties measurements. Available are electron microscopy for microstructural and microchemical analysis, equipment for measurement of the thermophysical and mechanical properties of materials to elevated temperatures, X-ray and neutron diffraction for structure and residual stress analysis, highspeed grinding machines, and measurement of component shape, tolerances, surface finish and friction and wear properties. Users willing to publish the results of their work can perform no-cost materials characterization here, under the sponsorship of OTT. Over half of the work performed in the HTML is sponsored by other R&D programs, from DOE and other agencies. This research is also primarily characterization, and it often involves composites. Composites work has been sponsored by the continuous fiber ceramic composite (CFCC) program, the ceramic technology for advanced heat engine program, now called the propulsion system materials program, and others. The following sections will highlight some of the capabilities resident at HTML, and illustrations of their application to composite materials. 2. Microstructure/microchemistry characterization 2.1. Auger analysis of the interfaces in a fiber composite The toughness of fiber-reinforced ceramic-matrix composites is not only controlled by the properties of the fiber and matrix materials, but the bonding forces between them. Paradoxically, the strongest bonds do not increase toughness but, in fact, somewhat weaker bonds that promote some sliding of the fibers during fracture are more effective. A Composites: Part B 30 (1999) 631–646 1359-8368/99/$ - see front matter q 1999 Elsevier Science Ltd. All rights reserved. PII: S1359-8368(99)00040-2 * Corresponding author. Tel.: 11-423-574-5123; fax: 11-423-574-4913. E-mail address: pastoae@ornl.gov (A.E. Pasto) www.elsevier.com/locate/compositesb
L E Pasto et al./Composites: Part B 30(1999)631-646 using a profiling technique where material was intermit- tently ion-sputtered away and the fresh surface analyzed This produced composition depth profiles that extended from the fracture surface into the fibers or the matrix. It was found that the depth profiles of troughs (T in Fig. 1) djacent to particular fibers(f)were essentially the san those that mated directly above fibers. The profiles for the Fracture fibers were also similar to each other. Therefore. it was possible to combine the two types of depth profiles into Matrix one as shown in Fig. 2. Composition in at. is shown as a function of depth for a BN-coated SiC fiber/SiC matrix composite. Zero depth represents the fracture plane, which Fibe is also the weakest portion of the composite structure. The development engineer was able to pick out the location of the fracture plane and whether the coating had achieved the Fig. 1 Schematic diagram of fracture in a fiber-reinforced ceramic matrix desired effect in controlling fiber/matrix bond strength. It was also of interest to track the levels of oxygen from the fiber through the Bn coating and finally disappearing after convenient technique that has been used to control the fiber/ approximately 140 nm into the SiC matrix. The BN coating matrix bond strength involves coating the fibers such as Sic thickness was seen to be about 50 nm and appeared to have with an element such as C or BN before consolidation into a diffused to some extent into the SiC matrix. Such composi composite. Coating as well as consolidation are ofter tional details of the fiber/coating/matrix interface would accomplished using chemical vapor deposition methods. a have been extremely difficult, if not impossible, to analyze successful strategy that may be used to develop tougher by any other technique and were very helpful to the composites is to relate toughness to the fracture interface development and optimization of fiber-reinforced ceramic location and composition as an experimental measure of composites interface bonding. This type of experiment is ideally suited to a modern scanning Auger microprobe that has the ability to detect virtually all elements in areas as small as 0. 1 um. 3. Mechanical characterization Small, notched samples of the composite were loaded in special holders and placed in the HTML's PHI Model 660 The mechanical, physical and structural characterization Auger chamber and the sample was fractured in-situ in a of composites at the HTML has addressed different scales in special fracture stage under a vacuum of 10 Pa. Then these materials. In most cases, this characterization work areas on the fracture surface were observed at a high magni- has been driven by the need to understand property fication, where fiber surfaces(F) and other areas where performance relationships that may lead to the synthesi fibers had been pulled out(troughs or"T) were exposed, of better materials as illustrated in the schematic diagram in Fig. 1. Several The characterization of composites at the microscale has areas of each type were analyzed with the Auger microprobe been focused on the composite constituents, namely fibers 100 Coating 150 -100 Fig. 2. Compositional depth profiles of the interface region of a BN-coated SiC fiber/SiC matrix composite
convenient technique that has been used to control the fiber/ matrix bond strength involves coating the fibers such as SiC with an element such as C or BN before consolidation into a composite. Coating as well as consolidation are often accomplished using chemical vapor deposition methods. A successful strategy that may be used to develop tougher composites is to relate toughness to the fracture interface location and composition as an experimental measure of interface bonding. This type of experiment is ideally suited to a modern scanning Auger microprobe that has the ability to detect virtually all elements in areas as small as 0.1 mm. Small, notched samples of the composite were loaded in special holders and placed in the HTML’s PHI Model 660 Auger chamber and the sample was fractured in-situ in a special fracture stage under a vacuum of 1028 Pa. Then areas on the fracture surface were observed at a high magni- fication, where fiber surfaces (F) and other areas where fibers had been pulled out (troughs or “T”) were exposed, as illustrated in the schematic diagram in Fig. 1. Several areas of each type were analyzed with the Auger microprobe using a profiling technique where material was intermittently ion-sputtered away and the fresh surface analyzed. This produced composition depth profiles that extended from the fracture surface into the fibers or the matrix. It was found that the depth profiles of troughs (T in Fig. 1) adjacent to particular fibers (F) were essentially the same as those that mated directly above fibers. The profiles for the fibers were also similar to each other. Therefore, it was possible to combine the two types of depth profiles into one as shown in Fig. 2. Composition in at.% is shown as a function of depth for a BN-coated SiC fiber/SiC matrix composite. Zero depth represents the fracture plane, which is also the weakest portion of the composite structure. The development engineer was able to pick out the location of the fracture plane and whether the coating had achieved the desired effect in controlling fiber/matrix bond strength. It was also of interest to track the levels of oxygen from the fiber through the BN coating and finally disappearing after approximately 140 nm into the SiC matrix. The BN coating thickness was seen to be about 50 nm and appeared to have diffused to some extent into the SiC matrix. Such compositional details of the fiber/coating/matrix interface would have been extremely difficult, if not impossible, to analyze by any other technique and were very helpful to the development and optimization of fiber-reinforced ceramic composites. 3. Mechanical characterization The mechanical, physical and structural characterization of composites at the HTML has addressed different scales in these materials. In most cases, this characterization work has been driven by the need to understand property– performance relationships that may lead to the synthesis of better materials. The characterization of composites at the microscale has been focused on the composite constituents, namely fibers, 632 A.E. Pasto et al. / Composites: Part B 30 (1999) 631–646 Fig. 1. Schematic diagram of fracture in a fiber-reinforced ceramic matrix composite. Fig. 2. Compositional depth profiles of the interface region of a BN-coated SiC fiber/SiC matrix composite
A E. Pasto et al./Composites: Part B 30(1999)631-646 . oad(grams) 88导8R Hi-NicalonTM 0.8 0.6 04 ceived hours 0.2 700°C ambient air 2.533.544.5 Ln(Load) Fig. 3. Effect of static exposure to ambient air at 700C for 500 h on the distribution of strengths of Hi-Nicalon fiber coatings, matrix and their interfaces. Characterization performance and the effects of temperature stress and envir- at the macroscale, through the evaluation of test coupons onment on service life. and components fabricated with these materials, has Although the original charter of the HTML addresses addressed issues such as the effect of specimen geometry "high temperature materials"(i.e ceramics, ceramic matrix and sample preparation on mechanical properties and composites), the infrastructure, expertise and tools available at the html have also been used for the characterization of composite systems with polymeric and metallic matrices 3.l. 3.1.. Fibers The major advances in the development of continuous fiber-reinforced ceramic composites(CFCCs)over the last 20 years have been possible, in part, thanks to the avail- ability of strong, small diameter ceramic fibers(ca 8- 20 um)[1, 2]. Most of the characterization of these ceramic aligning sphere flat glass slide fibers has been focused on their tensile properties, namely fiber fine dimensions, the direct determination of other physical and mechanical properties of these fibers has remained a challenge. However, current efforts at the HTml have been focused on the determination of their poisson's ratio and transverse coefficient of thermal expansion by means of laser diffraction techniques 3] ne ultimate in-plane tensile strength of 2D CFCI Fig 4 Schematic of the experimental setup used for lateral compression is controlled primarily by the strength of the reinforcing fibers. knowing how the strength of the fibers evolves
fiber coatings, matrix and their interfaces. Characterization at the macroscale, through the evaluation of test coupons and components fabricated with these materials, has addressed issues such as the effect of specimen geometry and sample preparation on mechanical properties and performance and the effects of temperature stress and environment on service life. Although the original charter of the HTML addresses “high temperature materials” (i.e. ceramics, ceramic matrix composites), the infrastructure, expertise and tools available at the HTML have also been used for the characterization of composite systems with polymeric and metallic matrices. 3.1. Constituents 3.1.1. Fibers The major advances in the development of continuous fiber-reinforced ceramic composites (CFCCs) over the last 20 years have been possible, in part, thanks to the availability of strong, small diameter ceramic fibers (ca 8– 20 mm) [1,2]. Most of the characterization of these ceramic fibers has been focused on their tensile properties, namely their elastic modulus and tensile strength. Because of their fine dimensions, the direct determination of other physical and mechanical properties of these fibers has remained a challenge. However, current efforts at the HTML have been focused on the determination of their Poisson’s ratio and transverse coefficient of thermal expansion by means of laser diffraction techniques [3]. Since the ultimate in-plane tensile strength of 2D CFCCs is controlled primarily by the strength of the reinforcing fibers, knowing how the strength of the fibers evolves A.E. Pasto et al. / Composites: Part B 30 (1999) 631–646 633 Fig. 3. Effect of static exposure to ambient air at 7008C for 500 h on the distribution of strengths of Hi-Nicalone. Fig. 4. Schematic of the experimental setup used for lateral compression experiments
TV CAMERA ⊙ -LOAD CELL INDENTOR 匚 Z-STAGE 匚 Y-STAGE DATA A X- STAGE当 OPTICAL BENCH Fig. 5. Schematic of the ITS which consists of a set of micropositioned XYZ stages, an optical microscope, a TV camera and a computer for user-friendly data acquisition and control
A.E. Pasto et al. / Composites: Part B 30 (1999) 631–646 634 Fig. 5. Schematic of the ITS which consists of a set of micropositioned XYZ stages, an optical microscope, a TV camera and a computer for user-friendly data acquisition and control
A.E. Pasto et al. /Composites: Part B 30(1999)631-646 Fig. 6. Fracture surface of a 2D CFCC with time under service conditions has been the focus of recent research work. For example, Fig 3 shows the effect of static exposure to ambient air at 700C for 500 h on the distribution of strengths of Hi-Nicalon"[3]. The avail- ability of these data has aided the formulation of models to predict the mechanical behavior and service life of CFCCs at elevated temperatures in oxidizing environments for example [4-7 A major milestone in the development of fiber and polymer technology occurred in the late 60s and early 70s Fig 8. Fracture surface of pulled-out fibers when new high-modulus organic fibers having strengths and moduli five times larger than then existent nylon fibers were compressive properties consists in subjecting these fibers to first synthesized [8]. The new fibers were composed of stiff, lateral compression. Through the of the hTml's highly aligned aromatic molecules, and although these interfacial test system(ITS), it was possible to characterize fibers exhibit outstanding tensile properties, their compres- the lateral compressive behavior of thermally cross-linkable sive properties are poor as a result from their large structural poly(p-1, 2-dihydrocyclobutaphenylene terephthalamide) anisotropy. Whereas, the axial tensile properties are domi-(PPXTA)fibers to assess the role of intermolecular cross- nated by covalent bonds within the polymer backbone, the links on the elastic and plastic transverse properties of these axial compressive properties depend more on the weaker fibers, and hence on their axial compressive behavior secondary intermolecular bonds. One way of characterizing Fig 4 shows a schematic of the experimental setup used for the secondary intermolecular bonding, and hence their axial the lateral compression experiments which was integrated to the html 's its. The html's its is a universal micro- mechanical testing machine that has been used for a large variety of tests including lateral compression tests on single fibers and single-fiber indentation tests as described below Fig. 5 is a schematic of the Its which consists of a set of micropositioned XYZ stages, an optical microscope, a TV camera and a computer for user-friendly data acquisition 200 Fig. 9. The tensile load versus cross-head displacement response of a Fig. 7. Scanning electron micrograph showing the fracture surface of CC Nicalon/CVI SiC minicomposite with a 1.0 um thick carbon fiber minicomposite after tensile testing
with time under service conditions has been the focus of recent research work. For example, Fig. 3 shows the effect of static exposure to ambient air at 7008C for 500 h on the distribution of strengths of Hi-Nicalone [3]. The availability of these data has aided the formulation of models to predict the mechanical behavior and service life of CFCCs at elevated temperatures in oxidizing environments for example [4–7]. A major milestone in the development of fiber and polymer technology occurred in the late 60s and early 70s when new high-modulus organic fibers having strengths and moduli five times larger than then existent nylon fibers were first synthesized [8]. The new fibers were composed of stiff, highly aligned aromatic molecules, and although these fibers exhibit outstanding tensile properties, their compressive properties are poor as a result from their large structural anisotropy. Whereas, the axial tensile properties are dominated by covalent bonds within the polymer backbone, the axial compressive properties depend more on the weaker secondary intermolecular bonds. One way of characterizing the secondary intermolecular bonding, and hence their axial compressive properties consists in subjecting these fibers to lateral compression. Through the use of the HTML’s interfacial test system (ITS), it was possible to characterize the lateral compressive behavior of thermally cross-linkable poly(p-l,2-dihydrocyclobutaphenylene terephthalamide) (PPXTA) fibers to assess the role of intermolecular crosslinks on the elastic and plastic transverse properties of these fibers, and hence on their axial compressive behavior [9]. Fig. 4 shows a schematic of the experimental setup used for the lateral compression experiments which was integrated to the HTML’s ITS. The HTML’s ITS is a universal micromechanical testing machine that has been used for a large variety of tests including lateral compression tests on single fibers and single-fiber indentation tests as described below. Fig. 5 is a schematic of the ITS which consists of a set of micropositioned XYZ stages, an optical microscope, a TV camera and a computer for user-friendly data acquisition A.E. Pasto et al. / Composites: Part B 30 (1999) 631–646 635 Fig. 6. Fracture surface of a 2D CFCC. Fig. 8. Fracture surface of pulled-out fibers. Fig. 7. Scanning electron micrograph showing the fracture surface of a minicomposite after tensile testing. Fig. 9. The tensile load versus cross-head displacement response of a CC Nicalone/CVI SiC minicomposite with a 1.0 mm thick carbon fiber coating
L E Pasto et al./Composites: Part B 30(1999)631-646 of a 2D CFCC can be simplified to the collective failure of several minicomposites in parallel. Characterization of minicomposites has also allowed for the determination of in-situ fiber strengths by the analysis of the fracture surface of pulled out fibers(Fig. 8). This is important because the strength of the reinforcing fibers is often affected as a result of high temperatures and aggressive environments often required for the synthesis of the matrix Analysis of load displacement curves obtained during the tensile evaluation of minicomposites have provided vast 0 0.25 0.5 0.75 information about the various micromechanical mechan- isms that are responsible for the tough behavior of CFCCs [11]. Fig. 9 shows the tensile load versus cross-head displa- Fig 10. Tensile load versus displacement curve for Hi-Nicalon"/CVI SiC cement response of a CG-NicalonCVI SiC minicompo- with a 0 I um thick carbon interphas site with a 1.0 um thick carbon fiber coating. The curve non-linear behavior up to the peak load for loads larg and control. It also has a collection of small capacity load than the load required to initiate matrix cracking. The cells and fixturing for the conduction of a wide variety of small jumps in the load in the non-linear region are asso- mechanical tests that include the bending of single ceramic ciated with the occurrence of additional matrix cracks. At fiber [10]. In this case, it was demonstrated that the propor- the peak load, a critical number of fibers in the bundle have tional stress limit, or yield strength was substantially higher failed triggering the failure of all fibers and leading to a for cross-linked PPXTA fibers than for uncross-linked continuous decrease in load bearing capacity with increas- fibers. In addition, it was found that the cross-linked ing displacement. The tail in the curve after the peak load is PPXTA exhibited a large recoverable compressive strain, the result of frictional sliding of the fiber bundle that bridges reminiscent of elastomers, in contrast to the large unrecov- the critical matrix crack as it is being pulled out from the erable strain exhibited by uncross-linked PPXTA fibers [9]. matrix. With the use of micromechanical models, it is The analysis of the fracture surface of a 2D CFCC (Fig. 6) possible, for example, to determine the magnitude of the suggests that the failure process and ultimately the strength interfacial shear stress(an important parameter to predict of these materials is controlled to a large extent by the fiber the mechanical behavior of these materials)from the tail of bundles aligned along the loading direction. This realization the curve or from the distribution of matrix cracks. The tail has prompted work in the characterization of fiber bundles in the curve can only be observed when the magnitude of the and minicomposites to tailor composite properties and to fiber bond strength and the interfacial shear stress are both further develop an understanding of the micromechanical low, and if the stiffness of the load train is large. Otherwise mechanisms responsible for the macroscopic behavior of a sudden load drop follows the peak load as illustrated in the CFCCs. Minicomposites consists of a single fiber tow, tensile load versus displacement curve for Hi-Nicalon containing any where between 500 and 800 filaments coated CVI SiC with a 0. 1 um thick carbon interphase in Fig 10 with the specific interphase and infiltrated with the matrix of In this case, the thinner carbon coating and the larger surface interest. Fig. 7 is a scanning electron micrograph showing roughness of Hi-Nicalon fibers result in a higher inter- the fracture surface of a minicomposite after tensile testing. facial shear stress, in shorter matrix crack spacing and The micrograph illustrates the distribution of fiber pullout fiber pull-out lengths. However, note that the improved ther lengths(which are a direct reflection of the distribution of mal stability and mechanical properties of Hi-Nicalon fiber strengths)and the fracture surface of individual fibers. over CG-Nicalon fibers result in a significantly larger In relation to the micrograph in Fig. 6, the tensile failure tensile strength. In addition to their usefulness in understanding and quan- fy ing the micromechanical mechanisms that are responsi ble for the tough behavior of CFCCs, minicomposites are also used to probe novel fiber coatings and interfacial concepts, an area that continues to be the focus of intense research [12] 3. 1.2. Fiber/matrix interface After the fibers, perhaps the most critical element in CFCCs is the fiber/matrix interface. hereafter referred to as the interface. The interface in CCCs is what makes possible to combine brittle fibers in a brittle matrix to obtain Fig. Il. Schematic of the single fiber push-in and push-out tes a tough composite
and control. It also has a collection of small capacity load cells and fixturing for the conduction of a wide variety of mechanical tests that include the bending of single ceramic fiber [10]. In this case, it was demonstrated that the proportional stress limit, or yield strength was substantially higher for cross-linked PPXTA fibers than for uncross-linked fibers. In addition, it was found that the cross-linked PPXTA exhibited a large recoverable compressive strain, reminiscent of elastomers, in contrast to the large unrecoverable strain exhibited by uncross-linked PPXTA fibers [9]. The analysis of the fracture surface of a 2D CFCC (Fig. 6) suggests that the failure process and ultimately the strength of these materials is controlled to a large extent by the fiber bundles aligned along the loading direction. This realization has prompted work in the characterization of fiber bundles and minicomposites to tailor composite properties and to further develop an understanding of the micromechanical mechanisms responsible for the macroscopic behavior of CFCCs. Minicomposites consists of a single fiber tow, containing anywhere between 500 and 800 filaments coated with the specific interphase and infiltrated with the matrix of interest. Fig. 7 is a scanning electron micrograph showing the fracture surface of a minicomposite after tensile testing. The micrograph illustrates the distribution of fiber pullout lengths (which are a direct reflection of the distribution of fiber strengths) and the fracture surface of individual fibers. In relation to the micrograph in Fig. 6, the tensile failure of a 2D CFCC can be simplified to the collective failure of several minicomposites in parallel. Characterization of minicomposites has also allowed for the determination of in-situ fiber strengths by the analysis of the fracture surface of pulled out fibers (Fig. 8). This is important because the strength of the reinforcing fibers is often affected as a result of high temperatures and aggressive environments often required for the synthesis of the matrix. Analysis of load displacement curves obtained during the tensile evaluation of minicomposites have provided vast information about the various micromechanical mechanisms that are responsible for the tough behavior of CFCCs [11]. Fig. 9 shows the tensile load versus cross-head displacement response of a CG-Nicalone CVI SiC minicomposite with a 1.0 mm thick carbon fiber coating. The curve non-linear behavior up to the peak load for loads larger than the load required to initiate matrix cracking. The small jumps in the load in the non-linear region are associated with the occurrence of additional matrix cracks. At the peak load, a critical number of fibers in the bundle have failed triggering the failure of all fibers and leading to a continuous decrease in load bearing capacity with increasing displacement. The tail in the curve after the peak load is the result of frictional sliding of the fiber bundle that bridges the critical matrix crack as it is being pulled out from the matrix. With the use of micromechanical models, it is possible, for example, to determine the magnitude of the interfacial shear stress (an important parameter to predict the mechanical behavior of these materials) from the tail of the curve or from the distribution of matrix cracks. The tail in the curve can only be observed when the magnitude of the fiber bond strength and the interfacial shear stress are both low, and if the stiffness of the load train is large. Otherwise, a sudden load drop follows the peak load as illustrated in the tensile load versus displacement curve for Hi-Nicalone CVI SiC with a 0.1 mm thick carbon interphase in Fig. 10. In this case, the thinner carbon coating and the larger surface roughness of Hi-Nicalone fibers result in a higher interfacial shear stress, in shorter matrix crack spacing and fiber pull-out lengths. However, note that the improved thermal stability and mechanical properties of Hi-Nicalone over CG-Nicalone fibers result in a significantly larger tensile strength. In addition to their usefulness in understanding and quantifying the micromechanical mechanisms that are responsible for the tough behavior of CFCCs, minicomposites are also used to probe novel fiber coatings and interfacial concepts, an area that continues to be the focus of intense research [12]. 3.1.2. Fiber/matrix interface After the fibers, perhaps the most critical element in CFCCs is the fiber/matrix interface, hereafter referred to as the interface. The interface in CFCCs is what makes possible to combine brittle fibers in a brittle matrix to obtain a tough composite. 636 A.E. Pasto et al. / Composites: Part B 30 (1999) 631–646 Fig. 10. Tensile load versus displacement curve for Hi-Nicalone/CVI SiC with a 0.1 mm thick carbon interphase. Fig. 11. Schematic of the single fiber push-in and push-out test
A.E. Pasto et al. /Composites: Part B 30(1999)631-646 637 Fig. 12. The single-fiber push-out test has a single fiber pushed into the matrix using a small flat-bottomed diamond indenter attached to the load cell in the ITs until the fiber protrudes from the bottom of the sample(Fig. 13) Although in the case of polymer and metallic matrix have been determined in various ways [ 13, 14]. Among the composites, a strong bond is sought between the matrix various indentation techniques, the single fiber push-in and and the reinforcement, the opposite is desired with push-out tests are without a doubt the most popular because CFCCs. Although a certain degree of bonding and frictional of the relative simplicity in their conduction. These tests are iding is required in CFCCs to allow for load transfer schematically described in Fig. 11 between fibers and matrix, the most important characteristic In the case of the single-fiber push-out tests, a single fiber of the interface has to be its ability to deflect cracks that is pushed into the matrix using a small flat-bottomed propagate in the matrix which can only be obtained with an diamond indenter(Fig. 12)attached to the load cell in the interface possessing low toughness ITS(Fig 3)until the fiber protrudes from the bottom of the The two parameters that best quantify the micromech sample(Fig. 13) nical efficiency of interfaces in CFCCs, namely the fiber During the test, both the load and the displacement of the bond strength and the interfacial shear stress or sliding stress fiber surface with respect to the surface of the sample ar monitored as illustrated in Fig. 14 for the case of a CG- Nicalon fiber in a CVI SiC matrix. By the application of micromechanical models it is possible to obtain several parameters from the experimental stress versus fiber-end displacement curves as illustrated in Fig. 15[131 A variation of the push-out tests that has been used to study the wear characteristics of the interface is the single- fiber push-back test. Fig. 16 is a record of the stress versus fiber-end displacement obtained during a push-back test 2000 3 Fig. 14. Plot of the load and displacement of the fiber surface with respect to Fig 13 Fiber protruding from the bottom of a specimen after being pushed the surface of the sample for the case of a CG-Nicalon fiber in a CVI SiC
Although in the case of polymer and metallic matrix composites, a strong bond is sought between the matrix and the reinforcement, the opposite is desired with CFCCs. Although a certain degree of bonding and frictional sliding is required in CFCCs to allow for load transfer between fibers and matrix, the most important characteristic of the interface has to be its ability to deflect cracks that propagate in the matrix which can only be obtained with an interface possessing low toughness. The two parameters that best quantify the micromechanical efficiency of interfaces in CFCCs, namely the fiber bond strength and the interfacial shear stress or sliding stress have been determined in various ways [13,14]. Among the various indentation techniques, the single fiber push-in and push-out tests are without a doubt the most popular because of the relative simplicity in their conduction. These tests are schematically described in Fig. 11. In the case of the single-fiber push-out tests, a single fiber is pushed into the matrix using a small flat-bottomed diamond indenter (Fig. 12) attached to the load cell in the ITS (Fig. 3) until the fiber protrudes from the bottom of the sample (Fig. 13). During the test, both the load and the displacement of the fiber surface with respect to the surface of the sample are monitored as illustrated in Fig. 14 for the case of a CGNicalone fiber in a CVI SiC matrix. By the application of micromechanical models it is possible to obtain several parameters from the experimental stress versus fiber-end displacement curves as illustrated in Fig. 15 [13]. A variation of the push-out tests that has been used to study the wear characteristics of the interface is the single- fiber push-back test. Fig. 16 is a record of the stress versus fiber-end displacement obtained during a push-back test A.E. Pasto et al. / Composites: Part B 30 (1999) 631–646 637 Fig. 12. The single-fiber push-out test has a single fiber pushed into the matrix using a small flat-bottomed diamond indenter attached to the load cell in the ITS until the fiber protrudes from the bottom of the sample (Fig. 13). Fig. 13. Fiber protruding from the bottom of a specimen after being pushed out. Fig. 14. Plot of the load and displacement of the fiber surface with respect to the surface of the sample for the case of a CG-Nicalone fiber in a CVI SiC matrix
L E Pasto et al./Composites: Part B 30(1999)631-646 2000 interfacial shear stress as close as possible to its critical value at which a brittle to ductile transition occurs Fig. 19 illustrates the effect of a proprietary fiber surface treatment on the topography of carbon-coated CG-Nica- ond Lengt e push-out: 0.26mm lon fibers in a CVI SiC matrix after push out [16-20] in which it the tensile strength of materials with treated fibers that was 1.5 times higher than that of CFCCs with untreated Frictional shear Stress: 12.3=0.7 Other techniques that have been used to characterize various fibers and fiber coatings include the use of atomic 0 2 force microscopy. Fig 20 is the imaged surface of a BN Fiber-end Displacement (um) coated carbon fiber using the AFM [21,22]. Data directly obtained from topographic profiles can be used directly in Fig. 15. Several parameters are obtained from the experimental stress micromechanical models to assess the effects of roughness on the clamping stress and hence in the interfacial shear tress which is very similar to that obtained during a push-out test Although most of the emphasis in the mechanical char- However, in addition to the interfacial shear stress and state acterization of CFCCs has been focused on their in-plane of residual stresses at the interface, it is possible to deter- tensile strength, often it is found that improvements on the mine both the coefficients of static and dynamic friction and inplane tensile strength of these materials are obtained at the the characteristic amplitude of the fiber surface roughness expense of the mechanical properties in other directions. A from push-back tests. Fig. 17 shows how the interfacial good example is the case the interlaminar shear and trans shear stress evolves with repeated push-out push-back laminar tensile strength of 2D CFCCs. It has been found that tests and provides quantitative information of how the there is a transition in the mode of interlaminar shear failure interfacial characteristics would evolve during mechanical as a function of fiber coating thickness and therefore. inter fatigue in CFCCs facial shear stress. For CFCCs with thin fiber coatings it has Furthermore, it has been possible to determine by been found that interlaminar shear and tensile properties are single-fiber push-out tests how various parameter dominated by the matrix, whereas for thick fiber coatings the interfacial properties, such as the thickness of the properties are dominated by the fiber coating and hence coating as indicated in Fig. 18 the interface [23. Fig the dependence of the By means of single-fiber push-out tests it has been possi- apparent interlaminar shear strength of CFCCs obtained ble to determine both the residual stresses and coefficient of by the compression of double-notched specimens for two friction with carbon coatings of various thicknesses [151 different fiber coating thicknesses when the fiber coating A significant amount of work has been dedicated to opti- thickness is 0.3 um interlaminar failure was found to occur mizing the nature of the fiber matrix interface in order to along porosity-rich matrix regions, whereas for thicker fiber maximize the potential of CFCCs. Contrary to common coatings the interlaminar shear strength was found to domi understanding of the micromechanics of these materials nated by the fiber coating and failure occurred between the suggests that in order to achieve the maximum potential fiber and the fiber coating [23]. Similar behavior of inter- of CFCCs it is necessary to increase the magnitude of the face-dominated properties has been observed in the case of carbon/carbon composites unidirectional evaluated at 3000 ambient temperature. In this case Fig. 22 shows the stress-strain behavior of ID carbon/carbon composites in a direction normal to the lamination direction Examination of the fracture surfaces by scanning electron microscopy a200 (Fig. 23)revealed that failure had occurred between the fibers and the fiber coating [24] 3.1.3. Matrix In the absence of environmentally stable fibers and fiber oatings it appears that the matrix cracking stress will be considered the largest allowable design stress for many 2 ingress of the environment to the interior of the composite Displacement(um) leading to fiber oxidation and ultimately to composite Fig. 16. Record of the stress versus fiber-end displacement obtained during failure a push-back test. Matrix cracks lead o the onset of non-linear behavior of
which is very similar to that obtained during a push-out test. However, in addition to the interfacial shear stress and state of residual stresses at the interface, it is possible to determine both the coefficients of static and dynamic friction and the characteristic amplitude of the fiber surface roughness from push-back tests. Fig. 17 shows how the interfacial shear stress evolves with repeated push-out push-back tests and provides quantitative information of how the interfacial characteristics would evolve during mechanical fatigue in CFCCs. Furthermore, it has been possible to determine by means of single-fiber push-out tests how various parameters affect the interfacial properties, such as the thickness of the fiber coating as indicated in Fig. 18. By means of single-fiber push-out tests it has been possible to determine both the residual stresses and coefficient of friction with carbon coatings of various thicknesses [15]. A significant amount of work has been dedicated to optimizing the nature of the fiber matrix interface in order to maximize the potential of CFCCs. Contrary to common understanding of the micromechanics of these materials suggests that in order to achieve the maximum potential of CFCCs it is necessary to increase the magnitude of the interfacial shear stress as close as possible to its critical value at which a brittle to ductile transition occurs. Fig. 19 illustrates the effect of a proprietary fiber surface treatment on the topography of carbon-coated CG-Nicalone fibers in a CVI SiC matrix after push out [16–20] in which it the tensile strength of materials with treated fibers that was 1.5 times higher than that of CFCCs with untreated fibers. Other techniques that have been used to characterize various fibers and fiber coatings include the use of atomic force microscopy. Fig. 20 is the imaged surface of a BNcoated carbon fiber using the AFM [21,22]. Data directly obtained from topographic profiles can be used directly in micromechanical models to assess the effects of roughness on the clamping stress and hence in the interfacial shear stress. Although most of the emphasis in the mechanical characterization of CFCCs has been focused on their in-plane tensile strength, often it is found that improvements on the inplane tensile strength of these materials are obtained at the expense of the mechanical properties in other directions. A good example is the case the interlaminar shear and translaminar tensile strength of 2D CFCCs. It has been found that there is a transition in the mode of interlaminar shear failure as a function of fiber coating thickness, and therefore, interfacial shear stress. For CFCCs with thin fiber coatings it has been found that interlaminar shear and tensile properties are dominated by the matrix, whereas for thick fiber coatings the properties are dominated by the fiber coating and hence the interface [23]. Fig. 21 shows the dependence of the apparent interlaminar shear strength of CFCCs obtained by the compression of double-notched specimens for two different fiber coating thicknesses. When the fiber coating thickness is 0.3 mm interlaminar failure was found to occur along porosity-rich matrix regions, whereas for thicker fiber coatings the interlaminar shear strength was found to dominated by the fiber coating and failure occurred between the fiber and the fiber coating [23]. Similar behavior of interface-dominated properties has been observed in the case of carbon/carbon composites unidirectional evaluated at ambient temperature. In this case Fig. 22 shows the stress–strain behavior of 1D carbon/carbon composites in a direction normal to the lamination direction. Examination of the fracture surfaces by scanning electron microscopy (Fig. 23) revealed that failure had occurred between the fibers and the fiber coating [24]. 3.1.3. Matrix In the absence of environmentally stable fibers and fiber coatings it appears that the matrix cracking stress will be considered the largest allowable design stress for many components and applications. Matrix cracks allow the ingress of the environment to the interior of the composite leading to fiber oxidation and ultimately to composite failure. Matrix cracks lead o the onset of non-linear behavior of 638 A.E. Pasto et al. / Composites: Part B 30 (1999) 631–646 Fig. 15. Several parameters are obtained from the experimental stress versus fiber-end displacement curves. Fig. 16. Record of the stress versus fiber-end displacement obtained during a push-back test
A.E. Pasto et al. /Composites: Part B 30(1999)631-646 苏9 it Push-Back 0a604B 5.0 kv x2.S Fig. 17. Evolution of the interfacial shear stress with repeated push-out these materials and are often associated with the onset of acoustic emissions. Fig. 24 shows a sequence of loading/ unloading stress-strain curves at increasingly larger stress levels for a CG-Nicalon"/PIP SINCO CFCC [25]. During mechanical loading and unloading, acoustic emissions were recorded and are plotted in the same graph. Note that the Keiser effect is characteristic of these materials and that th onset of non -linear behavior 75 MPa is well related to the formation of the first hysteresis loop and the first acoustic emissions During mechanical fatigue at elevated temperatures, it has been found that the matrix cracking stress or so-called 0850445. kv x4/90k750 proportional stress limit corresponds to the stress below which the CFCC exhibits endurance as illustrated in Fig. 19. The effect of a proprietary fiber surface treatment on the topo- Fig. 25[25]. It is also found that at larger stresses, the graphy of carbon coated CG-NicalonT" fibers in a CVI SiC matri decrease in composite strength exhibits a steep slope as a result of the environmental degradation of the fiber/matrix glasses may be unable to flow and heal cracks. Another interface and fibers [251 lternative consists in taking advantage of the redistribution Several different approaches have been used to minimize of internal stresses in composites when these are subjected the deleterious effects that result from matrix cracking. to creep at elevated temperatures In the case when the fibers Among these, the use of viscous glassy-phases to heal are much more creep resistant than the matrix, it is possible matrix cracks has been used extensively with relative to transfer the loading to the fibers and after an appropriate success. However, there may be some limitations in this thermal treatment, depicted in Fig. 26, it may be possible to approach at interm emperatures at which these subject the matrix to compression [26, 27]. Since the matrix cracking stress is strongly dependent on the state of residual stresses, it is possible to suppress matrix cracking. This concept was verified for CG-Nicalon"/BMAS CFCCS 28, 29]. Fig. 27 shows the stress-strain curves for two specimens before and after the application of the thermo- 20 mechanical treatment depicted in Fig. 26. As indicated, it was possible to increase the matrix cracking stress of the material at room temperature from 75 to 160 MPa. This was verified both by determining the onset of non-linear beha vior based on hysteresis loops and the onset of acoustic emissions Interlayer Thickness(um) In recent years, a tremendous amount of work has been Fig. 18. Thickness of the fiber coating as determined by means of single dedicated towards the development of standardized test fiber push-out tests methods for the mechanical characterization of cCCS
these materials and are often associated with the onset of acoustic emissions. Fig. 24 shows a sequence of loading/ unloading stress–strain curves at increasingly larger stress levels for a CG-Nicalone/PIP SiNCO CFCC [25]. During mechanical loading and unloading, acoustic emissions were recorded and are plotted in the same graph. Note that the Keiser effect is characteristic of these materials and that the onset of non-linear behavior 75 MPa is well related to the formation of the first hysteresis loop and the first acoustic emissions. During mechanical fatigue at elevated temperatures, it has been found that the matrix cracking stress or so-called proportional stress limit corresponds to the stress below which the CFCC exhibits endurance as illustrated in Fig. 25 [25]. It is also found that at larger stresses, the decrease in composite strength exhibits a steep slope as a result of the environmental degradation of the fiber/matrix interface and fibers [25]. Several different approaches have been used to minimize the deleterious effects that result from matrix cracking. Among these, the use of viscous glassy-phases to heal matrix cracks has been used extensively with relative success. However, there may be some limitations in this approach at intermediate temperatures at which these glasses may be unable to flow and heal cracks. Another alternative consists in taking advantage of the redistribution of internal stresses in composites when these are subjected to creep at elevated temperatures. In the case when the fibers are much more creep resistant than the matrix, it is possible to transfer the loading to the fibers and after an appropriate thermal treatment, depicted in Fig. 26, it may be possible to subject the matrix to compression [26,27]. Since the matrix cracking stress is strongly dependent on the state of residual stresses, it is possible to suppress matrix cracking. This concept was verified for CG-Nicalone/BMAS CFCCs [28,29]. Fig. 27 shows the stress–strain curves for two specimens before and after the application of the thermomechanical treatment depicted in Fig. 26. As indicated, it was possible to increase the matrix cracking stress of the material at room temperature from 75 to 160 MPa. This was verified both by determining the onset of non-linear behavior based on hysteresis loops and the onset of acoustic emissions. In recent years, a tremendous amount of work has been dedicated towards the development of standardized test methods for the mechanical characterization of CFCCs. A.E. Pasto et al. / Composites: Part B 30 (1999) 631–646 639 Fig. 17. Evolution of the interfacial shear stress with repeated push-out push-back tests. Fig. 18. Thickness of the fiber coating as determined by means of single- fiber push-out tests. Fig. 19. The effect of a proprietary fiber surface treatment on the topography of carbon coated CG-Nicalone fibers in a CVI SiC matrix
1. E. Pasto et al. /Composites: Part B 30(1999)631-646 whole Image 5000mm 8.5868mm Area RMS: 113764 nm Avg. Height: 40.2851 nm Max Range: 99.9852 nm 2500nm O nm o! Fig. 20. The imaged surface of a BN-coated carbon fiber using the aFm Personnel at the HTml have played a key role in drafting components, heat exchangers and filtration equipment as and validating these documents. These include the para- part of the energy industries. Many of these components metric evaluation of geometric factors in composites such possess cylindrical geometry and hence test methods have as the effect of notch separation in the determination of been devised for the evaluation of their hoop strength. For interlaminar shear strength by the compression of double- example, Fig. 28 shows a schematic of the test for the notched specimens, the effects of machining on the tensile internal pressurization of composite cylindrical specimens strength of CFCCs, and the effects of geometry and loading ing an incompressible rubber insert parameters on these materials. These have led to the drafting a distinction must be made between the intrinsic tensile of standard test methods as part of the committee C28.07 of strength of the material and the hoop strength of a tubular the American Society for Testing and Materials [30,31] structure made with the same material. In the latter case. the strength will depend on both the intrinsic strength of the 3. 1.4. CPCC components material and some features of the structure such as seams Another activity is the evaluation of components. Many for creating a closed structure when the starting material is of these materials have found use in gas-turbine engine flat structures in the case when the starting material is a fabric. Fig. 29 shows a comparison of the tensile stress- calon TM/siC cFcc strain behavior of cg-Nicalon"/PIP sinco and that of the hoop stress versus hoop strain for a tubular structure made of the same material when internally pressurized using an incompressible rubber insert. In this case, failure of the tube could be traced to the seams in the structure although note that the slope and onset of non-linear behavior are identical in both cases [32] 0. 3 um carbon interface 1.1 um carbon interface 4. Thermophysical characterization 3 5678 Notch Separation(mm) 4.1. Thermal diffusivity of a metal matrix composite Fig. 21. The dependence of the apparent interlaminar shear strength of For several years, aluminum metal matrix composite(A CFCCs obtained by the compression of double-notched specimens for MMC) materials have been considered for automotive brake two different fiber coating thicknesses. applications, driven by the possibility of vehicle weight
Personnel at the HTML have played a key role in drafting and validating these documents. These include the parametric evaluation of geometric factors in composites such as the effect of notch separation in the determination of interlaminar shear strength by the compression of doublenotched specimens, the effects of machining on the tensile strength of CFCCs, and the effects of geometry and loading parameters on these materials. These have led to the drafting of standard test methods as part of the committee C28.07 of the American Society for Testing and Materials [30,31]. 3.1.4. CPCC components Another activity is the evaluation of components. Many of these materials have found use in gas-turbine engine components, heat exchangers and filtration equipment as part of the energy industries. Many of these components possess cylindrical geometry and hence test methods have been devised for the evaluation of their hoop strength. For example, Fig. 28 shows a schematic of the test for the internal pressurization of composite cylindrical specimens using an incompressible rubber insert. A distinction must be made between the intrinsic tensile strength of the material and the hoop strength of a tubular structure made with the same material. In the latter case, the strength will depend on both the intrinsic strength of the material and some features of the structure such as seams for creating a closed structure when the starting material is flat structures in the case when the starting material is a fabric. Fig. 29 shows a comparison of the tensile stress– strain behavior of CG-Nicalone/PIP SiNCO and that of the hoop stress versus hoop strain for a tubular structure made of the same material when internally pressurized using an incompressible rubber insert. In this case, failure of the tube could be traced to the seams in the structure, although note that the slope and onset of non-linear behavior are identical in both cases [32]. 4. Thermophysical characterization 4.1. Thermal diffusivity of a metal matrix composite For several years, aluminum metal matrix composite (Al MMC) materials have been considered for automotive brake applications, driven by the possibility of vehicle weight 640 A.E. Pasto et al. / Composites: Part B 30 (1999) 631–646 Fig. 20. The imaged surface of a BN-coated carbon fiber using the AFM. Fig. 21. The dependence of the apparent interlaminar shear strength of CFCCs obtained by the compression of double-notched specimens for two different fiber coating thicknesses