1 .mr mu 'szuk'r g ,-m. ‘1 n! .rPr" 5 ,. my") .. . - m < ‘ mu m c m KMA ( ~‘k r - , v; my. . .A a ‘ .n...‘...,.;,.._, ..,H.,.,.. .. ,,.,,,. .. ....,. ,, x' ,., '4 . :;‘. fl.“ ':. ..: . .mr, f.‘!;’ ., mi. 5., u» - o 14-; m- w #f-z-‘nm— ”WA“! ..,3_ .N. .-. ,... r 4: ”my. WNW. ‘ _. 0,, . . . w , . » .. n n. 1 an .. 1 - w»... -qfiw. .1 2.3..1" .l- -. n -. "‘ ' :1»: . s "ol‘ur-oq-‘up “U”, n.— ,....,.1 3w Will!"lllfllfllllmwllmNHHHNWUHHHI l 1/ s 28* ’ 778 2042 J ‘ LIBRARY Michigan State University \ J This is to certify that the dissertation entitled Physical and Mechanical Properties of Hipped Ag Fibers Relnforced YBazCu3O7_X Superconductor Composite presented by Gun-Eik Jang has been accepted towards fulfillment of the requirements for Ph . D . degree in Materials Science Major professor Date 5: / 31/ ’92 . MSU is an Affirmative Action/Equal Opportunity Instituzion 042771 PLACE IN RETURN BOX to remove this checkout from your record. TO AVOID FINES return on or before date due. (mews menus meme F‘ l QM; __T__ “ i l MSU Is An Affirmative Action/Equal Opportunity Institution c:\clm\ddedue.pm3-p.1 -— PHYSICAL AND MECHANICAL PROPERTIES OF HIPPED Ag FIBERS REINFORCED YBagCuaO7-x SUPERCONDUCTOR COMPOSITE BY GUN-EIK JANG A DISSERTATION Submitted to Michigan State University in partial fulfillment of the requirements for the degree of ‘ DOCTOR OF PHILOSOPHY Department of Metallurgy, Mechanics and Material Science 1992 ABSTRACT PHYSICAL AND MECHANICAL PROPERTIES OF HIPPED Ag FlBER REINFORCED YBa2CU3O7-x SUPERCONDUCTOR COMPOSITE BY GUN-EIK JANG Ceramics exhibit a strong relationship between mechanical and physical properties and the processing and fabrication sequence. The final microstruc- ture also depends on processing. This interrelation between processing, mi- crostructure and properties is particularly true for the new high Tc superconduct- ing ceramic oxides of the type, YBagcu307-x. An attempt to improve the me- chanical stability of such superconducting ceramics, through composite tech— nology, requires a detailed understanding of the relationship between mi- crostructure and processing. To characterize some of these relations, Several properties have been measured for sintered oxide superconductor. Polycrystalline YBach307-x has proven to be difficult to density by conventional sintering to the full theoretical density while maintaining the superconducting orthorhomic phase. This precludes conventional sintering from being an effective process for achieving full density. In this research, the YBazCU307-x samples were prepared by two differ- ent processing techniques: one by conventional sintering and the other by hot isostatic pressing (HIP). It was determined that reinforcement with silver fibers and hot isostatic pressing produced the best results without degrading the su— GUN-EIK JANG perconducting properties. in this thesis, physical and mechanical properties of HlP processed and silver fiber reinforced 1-2-3 based superconductors are dis- cussed and compared with the properties of unreinforced and conventionally sintered samples. Finally, significantly improved mechanical properties such as fracture toughness and fracture energy, obtained by a multi-stage processing, are reported. ACKNOWLEDGEMENTS First of all, I would like to express my gratitude to Professor Mukherjee for his continual encouragement and guidance, which he provided throughout my Ph. D. program and research project. I would also like to thank my committee members for their timely review of this dissertation. A special thanks to my korean friends and colleagues in laser laboratory for their assistance and friendship. Finally, a special thank you is to be extended to my wife Hwa-Young and my beautiful daughter Jee-Hee for their patience and continual encouragement during the many months of this research effort. iii TABLE OF CONTENTS T OF FIGURES T OF TABLES INTRODUCTION LITERATURE SURVEY 2. 1 High Tc Superconductor: History, Structure and Properties 2.1.1 2.1.2 2.1.3 .1.4 1 2 [\J ..2 ..2 M .3 ..3 ..3 3. 2. 2. ..3 2. 2. 2. 2. 0300000001(130043-0)I\)--L N 90 History of high temperature superconductors Crystal structure of YBach3O7-x Mechanical and electrical properties in Y-Ba-Cu-O superconductor Weak-link behavior at the grain boundaries Superconductor Processing. Techniques and Mechanisms Sintering and Hot Isostatic Pressing HIP mechanisms as applied to YBach307- -x consolidation echanical Properties of Ceramic Matrix Composites The general theory of strength of fiber reinforced composites Theory of energy absorption mechanisms Stress-strain curves of fiber reinforced composites Testing methods .4. 1 Tensile strength .4. 2 Flexural strength Fracture toughness .5. 1 The double cantilever beam (DCB) test ..5 2 The single edge notched beam (SENB) test ..5 3 Chevron notch beam (CNB) test 5.4 Double torsion test .5. 5 Indentation test EXPERIMENTAL PROCEDURE 3. 1 Sample Preparation _ 3. 1. 1 Sample preparation (conventional sinterIng) 1.2 3. wwwwwww @ Sample preparation (hot isostatic pressing) ample Characterization Methods X- -ray diffraction studies Electrical resistance measurements Critical current density measurements Magnetic moment measurements Density and porosity measurements Flexural strength and toughness measurements iv Page vi xii 11 16 2O 20 22 25 25 32 35 38 38 39 41 45 45 48 48 50 52 52 53 58 58 58 62 64 66 67 3 Morphological examination and EDAX analysis 7. 1 Optical microscopy 7. 2 Scanning Electron Microscopy and EDAX 3. 2. 8 Scanning Auger Electron Microprobe . 2. 7 3. 2. 3. 2. iESULTS AND DlSCUSSlON .. 1 Crystal Structure versus Processing Conditions 2 Superconducting Properties 4. 2. 1 Critical temperature (Tc) measurement 4. 2. 2 Critical current density (Jc) measurement 4. 2. 3 DC magnetic moment measurement 3 Variation of Density and Porosity of Samples with Different Types of Processing 4 Mechanical Properties of Ag-Fiber Reinforced Composites 5 Microstructure and interfaces 4. 5. 1 Optical microscope analysis 4. 5. 2 Scanning Electron Microscopy (SEM) analysis 6 Microchemistry of Bulk and Interface Regions 4. 6. 1 Energy dispersive x-ray analysis _ 4. 6. 2 Auger Electron Spectroscopy (AES) anaIySIs :ONCLUSIONS ENDIX ' OF REFERENCES 97 97 106 106 111 120 120 120 137 139 147 LIST OF FIGURES Ire Page The effect of various doping elements (in weight percent) on the critical temperature of YBBQCU3O7-X as measured by ac. susceptibility (Ref. 24). 7 The perovskite structure. 8 The crystal structure of YBa20U3O7-x (Ref. 29). 9 Magnetization hysteresis loops at 4.5 K for a single crystal of 1-2-3 with the Cu-O planes oriented (a) perpendicular to the magnetic field and (b) parallel to the applied field (Ref. 55). 14 Critical current densities deduced from magnetization hysteresis as a function of magnetic field applied either parallel or perpendicular to the Cu-O planes (Ref. 55). 15 Transport Jc vs. H in YBazCU307-x. The shaded area represents the improved intergrain Jc values in bulk YBagCU307-x with added flux pinning defects (Ref. 24). 17 A schematic illustration of time, temperature and phase equilibria, associated with melt -texture processing (a) A qualitative sectional phase diagram (b) Three basic processing steps, shown as the temperature-time curve (Ref. 24). 19 The effects of temperature gradient on the melt-textured microstructure. (a) Long-range grain alignment, obtained with a temperature gradient. (b) Locally melt-textured structure, obtained without a temperature gradient (Ref. 24). 19 Diagram illustrating Vmin and ch and their relationship to 0f. 0m and 0mu- 27 Schematic illustration of the distribution of fiber tensile stress (0) and interface shear stress (1:) along a short fiber embedded in a matrix (Ref. 76). 29 The variation of 13m and om with position along the fiber when the matrix (and fiber) deforms elastically. 29 vi FAI- FIE fir He The effect, (shown schematically), of fiber volume fraction, Vi, fiber length, l, and fiber strength, O'f, on the strength, 0c, of the composite (Ref. 76). ldealized stress-strain behavior of a ceramic matrix composite (Ref. 80). Page 31 37 The schematic diagram of the 3-point and the 4-point bending mode. 42 The different types of specimen geometry for fracture toughness measurement (Ref. 90). Schematic of chevron notch flexural beam for the determination of fracture toughness. Schematic diagram for the HIP processing cycle. Time-temperature-pressure profile during HlPing of Ag fiber reinforced YBach3O7-x. Schematic diagram describing the experimental set-ups for resistance measurement. Schematic diagram of the four-probe technique. Schematic diagram of sample holder assembly for AC resistance measurement. Schematic diagram describing the experimental set-ups for critical current density measurement. Magnetic Property Measurement System (MPMS) components (Ref. 91). Schematic diagram for notch-beam test in 3-point bending. A schematic diagram illustrating the basic processes associated with Auger electron transition. A fracture surface examined by Auger spectroscopy apparatus (x 500). Typical x-ray diffraction patterns from HlPed specimens. X-ray diffraction pattern from a specimen reinforced with 15 volume % Ag. Reference diffraction pattern line positions of pure Ag and pure 1-2-3 compound are shown for comparison. Vii 46 49 55 56 59 6O 61 63 65 69 72 74 78 79 Figure 29. 30. 32. 33. 34. 35. 36. 37. 38. 39. 40. 41 . 42. fir - _ __ _-_.-- we Page X-ray diffraction patterns from specimens reinfOrced with 15 volume °/o Ag; for different processing conditions. 80 X-ray diffraction patterns of 1-2—3 specimens for different processing conditions. 83 The temperature dependance of resistance of 1-2-3 compound with different amounts of reinforcing Ag fibers. 86 The electrical resistance of specimens, with different volume % Ag fibers, at 150 K. 88 Resistivity vs. temperature curves for 1 -2-3 specimens processed by two different conditions; a presintered sample, and a sample HlPed and annealed. 90 The voltage-current (V-l) curves for specimens with different volume % Ag fibers. 92 The slopes of the voltage-current (V-l) curves in Fig. 34, are plotted as a function of volume % Ag fibers. 94 The temperature dependance of the DC magnetization for HlPed specimens. 96 Variation of porosity with volume °/o Ag fibers (a) conventionally sintered samples; (b) HlPed samples. 98 A plot of flexural strength vs. volume % Ag (a) conventionally sintered samples; (b) HlPed samples. 100 K10 versus volume % Ag (a) conventionally sintered samples; (b) HlPed samples. 103 Typical load-deflection curves of YBazCU307-x samples reinforced with various volume % Ag fibers. Also shown are corresponding typical fractographs. 104 Microstructure of as-received pure Ag fibers (x 120). 107 Optical microstructures taken from the polished surface of 1-2-3 specimens at different magnifications (x 200 and x 500) (a) sintered, (b) HlPed and annealed. 108 ). Optical microstructures of the sintered specimen reinforced by 15 volume % Ag fiber, taken at different magnifications (a) x 200, (b) x 500. 109 viii FlgI 43(l 451 46 47 re ). Optical microstructures of HlPed specimen reinforced by 15 volume % Ag fiber, taken at different magnifications (a) x 200, (b) x 500. SEM micrographs, taken from the fractured surface of HlPed 1-2-3 specimens, at different magnifications ( x 500 and x 1,000) (a) sintered, (b) HlPed. ). SEM micrographs, taken from the fractured surface of HlPed 1-2—3 specimens reinforced by different Ag volume % (a) 1-2—3 matrix, (D) 5 % Ag, (c) 8 °/o Ag. ). SEM micrographs, taken from the fractured surface of HlPed 1-2-3 specimens reinforced by different Ag volume % (d) 12% Ag, (e) 15 % Ag. SEM micrographs of fracture surfaces of HlPed and Ag reinforced specimen showing: (a) fiber pull-out, and (b) fiber debonding. SEM micrographs showing damaged Ag fiber in HlPed 1-2-3 matrix (a) x 400, (b) x 1,500. SEM micrographs taken from the fractured surfaces (x 500) (a) HlPed specimen without any post-annealing (b) HlPed with post-annealing, at 850°C for 20 hours. Fractured surfaces of HlPed specimen examined for EDAX elemental analysis as presented in Figs. 50(a), 50(b) and 50(0) (x 400 and x 700). ). Elemental analysis by EDAX of a 1-2—3 superconductor. l. Elemental analysis by EDAX on a fractured surface (with overall area scanning). . Elemental analysis by EDAX of a region of Ag fiber. AES spectra taken from the polished surface of a HlPed specimen reinforced with 5 volume % Ag. AES spectra taken from the polished surface of a HlPed specimen reinforced with 5 volume % Ag after ion sputtering. AES spectra taken from the fractured surface of a HlPed specimen reinforced with 5 volume °/e Ag. Page ‘110 112 113 114 115 116 118 121 122 123 124 125 126 128 Fig 55. 56 57 58 “e . Page AES spectra taken from the fractured surface of a HlPed specimen reinforced with 5 volume % Ag after ion sputtering. 129 AES spectra in the line scanning mode, taken from the fractured surface of a HlPed specimen reinforced with 5 volume % Ag. 130 AES spectra with 5 keV excitation energy, taken from the fractured surface of a HlPed specimen reinforced with 5 volume % Ag. 132 Fracture surface of a HlPed specimen reinforced with 5 volume % Ag which exhibits electron beam radiation damage at 15 keV excitation energy. Rounded projections indicate melting. 133 . AES image produced by compositional dot—mapping, taken from the fractured surface of a HlPed specimen reinforced with 5 volume % Ag; (a) mapping of yttrium, (b) mapping of barium. 134 . AES image produced by compositional dot-mapping, taken from the fractured surface of a HlPed specimen reinforced with 5 volume % Ag; (0) mapping of copper, (d) mapping of oxygen. 135 . AES image produced by compositional dot-mapping, taken from f the fractured surface of a HlPed specimen reinforced with 5 volume °/o Ag; (e) mapping of silver. 136 Scanning electron micrograph of as received Cu fibers (x 250). 140 .SEM micrographs showing the oxidation of Cu fibers at the sintering temperature of 930° C for a sintering in air for 3 hours; (a) x250, (b) x1 ,500. 141 SEM micrographs showing oxidized Cu fibers in the matrix of 1-2- 3 compound after sintering at 930°C for 3 hours in air; (a) x 150, (b) x 600. 142 . SEM micrographs showing surface morphology of Ag coating on Cu fibers (x 1,200). 144 . SEM micrographs of the cross sectional area of Ag coated Cu fibers (x 500). 144 Fractured surface of 1-2- 3 compound reinforced with Ag coated Cu fibers. Samples were prepared by sintering at 930°C for 24 hours. (a) x 250, (b) x1 ,.000 145 gure Page Equilibrium phase diagram of Ag-Cu system (Ref. 113). Note that the eutectic temperature is below the sintering temperature. 146 LIST OF TABLES a Page heoretical models of energy absorption associated with mechanical aformation of fiber reinforced composites 36 est methods for fracture toughness measurements [Ref. 90] 47 ypical main auger peaks in YBa20U3O7-x composite [Ref. 97] 76 esults of transition temperature and critical current density easurements for HlPed and sintered YBfigCU307-x 91 easured and theoretical densities and porosities of YBaQCU307-x Iinforced with different volume % Ag-fibers 99 .Immary of measured mechanical properties of 1-2-3 and Ag fiber inforced and HIP consolidated 1-2-3 compound 105 xii . -.:.'r-._.—k_’_r- ---.# - ., -— .- - l. I. INTRODUCTION While the new class of high Tc (~ 93 K) ceramic superconductors repre- . sent a technological breakthrough [1], there are still a number of difficulties to Ivercome for use in practical applications. The brittle nature inherent in ceram- cs poses one of the most serious problems. It is important to recognize that the nechanical responses of these materials are similar to other advanced ceram- cs. A critical temperature of 93 K is acceptable for numerous design applica- ions provided that the devices are mechanically stable. Moreover, many appli- ations of superconducting oxides will require these materials to be fabricated 1 the form of wires or cables. The shaping of wires or cables typically require :ignificant flexibility or ductility of the material from which the wires or cables are nade. However, the oxide superconductor itself does not have enough flexibil- y and toughness to satisfy these shaping operations. The flexibility and tough- Iess can, however, be increased by making a metal-oxide superconductor :omposite. , New investigations should be focused on second phase materials such .3 filaments, whiskers or fibers which can offer a continuous path for current ow, improved mechanical strength and retained superconducting properties. 'rocessing conditions that will yield continuous, interwoven oxide and fiber- hase must be well developed. The most important criterion to be considered for successful applications 3 the understanding of processing variables which influence the microstructure nd electrical properties such as critical temperature and critical current density. is well known that critical temperature is usually influenced by the crystal :ructure and long range order, while critical current density depends predomi- antly on the microstructural features such as grain boundaries, microcracks, 1 gre p0 SU 2 grain orientation and microchemistry. The low criticalcurrent density in the bulk polycrystalline materials is thought to arise from weak-link junctions of the superconductive phase across the grain and twin boundaries, microcracks, pores and voids, inclusions of non-superconducting phase, and the lack of ap- propriate texture formation. The chemistry of impurity phases in the grain boundaries, in particular, may be the other dominant factor in determining the current carrying capability as well as the resistance to grain boundary fracture. Although the presence of fine dispersion of impurities can increase critical cur- rent density by the flux pinning mechanism, proper control of processing condi- tions and microstructure was found to be crucial in attaining a high transition temperature and improved critical current density. So far, a number of research projects have been successfully conducted in the area of superconducting composites, especially focused on particulate composites. However, none of the previous work has been undertaken in the area of fiber reinforced composites to study the modification of mechanical properties through the addition of fibers into ceramic oxide superconductors. In the present study, Ag fiber reinforced YBa20U3O7-x (1-2-3 system) su- perconductor composite has been systematically studied as a function of the volume % Ag fibers. The reason for selecting Ag fiber as the reinforcing ma- terial in the 1-2-3 matrix can be explained as follows: 1. Ag has been found to be one of the few materials that does not degrade the superconducting properties of 1-2-3 system [25]. 2. Ag remains as a pure metal, that is chemically stabilized, in the 1-2-3 system [2-5]. 3. The melting temperature of Ag, 960.500, is higher than the normal sintering or processing temperature of the 1-2-3 system. lo obi and h the IL cessi 3 4. Ag lowers the normal state resistivity and contact resistance of the composites [3]. 5. It has been shown that Ag particulates improve mechanical properties like Young's modulus and fracture strength [4,5] of the 1-2— 3 matrix. The overall objective of this research is to evaluate the feasibility in order to obtain dense superconducting composites which are mechanically reliable and have a greater fracture toughness. This practical objective is coupled with the fundamental objective, which is to understand the relationship between pro— cessing conditions, microstructure, and superconducting properties, _ .-__- 2. LITERATURE SURVEY 2. 1 High T c Superconductor: History, Structure and Properties 2. 1. 1 History of high temperature superconductors The discovery of the La-Ba-Cu-O system, with a superconducting transi- tion at 30 K, by Bednorz and Miller [6] and subsequent evolution of 90 K Y-Ba- Cu-O system [1,7] have stimulated extensive investigation in the preparation and characterization of various types of ceramic superconductors. To date, research has been conducted by a very large number of investi- gators to further increase the critical temperature. The history of the high tem- perature oxide superconductors has shown that elemental substitution is the most effective approach for raising the transition temperature. Substitution of Sr for Ba in La-Ba-Cu-O system led to discovery of 40 K La-Sr-Cu-O superconduc- tor [8, 9]. Substitution of Y for La produced the 90 K Y-Ba-Cu-O system [1], which is the first system to have its critical temperature above liquid nitrogen boiling temperature (77 K). Also the effect of various other substitutions in Y-Ba- Cu-O system, on the superconducting transition temperature, has been studied extensively [10-12]. Substitution of TI for the rare earth metal (R) in the R-Ba- Cu-O system, produced the 90 K TI—Ba-Cu-O system [13,14], and addition of Ca in the TI-Ba-Cu-O system led to the 120 K TI-Ca-Ba-Cu-O system [15]. The Tl-Ba-Cu-O system is the first non-rare earth system which reaches zero resis- tance above liquid nitrogen temperature, whereas the TI-Ca-Ba-Cu-O system is the first system which reaches zero resistance above 100 K, and has the high- est transition temperature for zero resistance. 5 On the other hand, substitution of Bi for La in La-Sr-Cu-O system led to 20 K Bi-Sr-Cu-O system superconductor [16]. But partial substitution of Ca for Sr in this system produced 110 K Bi-Ca-Sr—Cu-O superconductor [17]. A num- ber of researchers have reported that substitution of a 3-d transition metal such as Ni, Zn, or Co for Cu can sharply decrease the critical temperature [18,19,20]. The ionic size and orbital structure of the 3-d transition metals are very similar to those of Cu. The 3-d transition metals will occupy the Cu sites if transition met- als are substituted into the Y-Ba-Cu-O system. However, according to Xiao et al. [19], Tc is strongly correlated with the size of paramagnetic moments of doped elements. Addition of Zn provides an extra electron which fills up anti- bonding d-band and thereby reduces the density of state at the Fermi level. Tarascon et al. [18] have suggested and tried to explain the reduction of TC by 4 different inter-related mechanisms; structural disorder, oxygen vacancies, the dopants inducing a different oxidation state in the copper and magnetic depair— ing. However, there is no explanation as yet regarding the reduction of Tc when substitution was made for Cu. There was a report of critical temperature above 150 K in Y-Ba—Cu-O system [21]. Ovshinsky et al [21] claimed superconducting zero-state at 155 K by adding fluorine to 1-2-3 system. They synthesized five different compositions of Y882CU3FXOZ by substituting part of oxygen with fluorine. However, these results were contradicted by other researchers [2, 23]. Yan et al. [22] and Narottom et al. [23] reported they could not observe such a high transition tem- perature reported by Ovshinsky et al. [21]. Their results only indicate that su- perconducting temperature is slightly increased (up to 93.4 K) and sharpened with low levels of fluorine concentration. Yan et al. also [22] studied F and Cl effects on the 1-2-3 system. Most samples of YBazCU3FxCIyOz, showed diar sigr mo: abc ol \ me ten am to I 131 Ira Cit 6 diamagnetic susceptibility at a temperature between 89 and 94 K without any significant improvement of TC. Among high Tc oxide superconductors, Y-Ba-Cu-O is perhaps one of the most widely studied compounds because it exhibits superconducting behavior above 90 K and is relatively easy to synthesize. Fig. 1 shows the typical effect of various doping elements on the critical temperature of Y-Ba-Cu-O system as measured by ac. susceptibility [24]. It can be seen that the superconducting temperature and transition range in samples doped with Au and Cd is still high and sharp. However, doping with other elements , for example Zn or Al, Seems to destroy the superconducting transition and/or broadens the the transition range. 2. 1. 2 Crystal structure of Y8620U3O7-x YBach307-x was identified as the phase exhibiting superconducting transition at 90 K. This compound has a distorted orthorhombic oxygen defi- cient perovskitelike structure [25]. All observed x-ray powder diffraction peaks can be accounted for by choosing an orthorhombic unit cell with lattice parameters given by: a = 38.20 nm , b = 38.93 nm, and c = 116.88 nm [26]. Compounds with the perovskite structure (Fig. 2) have the formula of A803, where A is the relatively large ion in the center of the unit cell, 8 is the small ion at the corners and O is the oxygen ion at the edge of the unit cell [27]. The cell volume is 3 times the standard perovskite Cell and structure can be visualized as a stack of alternating layers in a c—axis: (BaCuOs) : (YCu03) : (BaCan) [27,28]. The crystal structure of Y-Ba-Cu-O [29] in Fig. 3 resembles three stacked cubic perovskite unit cells, however there are major differences between the two structures. While in a perovskite structure, copper ions are MAGNETIC SUSCEPTIBILITY (orb. units) Fig. 1. l l - J I I l 4O 60 80 100 120 140 160 TEMPERATURE (K) The effect of various doping elements (in weight percent) on the critical temperature of YBaZCU3O7-x as measured by a.c.susceptibility (Ref. 24). .. Oea @v ©CUII Fig. 3. The crystal structure of YBa20U307-x (Ref. 29). 10 located in the center of an oxygen octahedron, copper ions occupy two symmetry sites [30,31]. In the structure of Y-Ba-Cu-O, all oxygen atoms are missing in the (002) yttrium atom plane and oxygen vacancies occur in a strongly preferred way. In the basal plane, half of the oxygen atoms are missing; thus vacancies occur along the a-axis, leading to an orthorhombic distortion. Jorgensen et al. [32] studied Tc as a function of the fractional occupancy I of the oxygen sites. They found that YBa2Cu307-x compound can contain, on the average, from six to seven oxygen atoms per unit cell and in the orthorhombic phase, Tc decreases smoothly as x changes from 0.3 to 0.7 and reaches zero at the orthorhombic-tetragonal transition. Tetragonal YBaZCU307-x is not superconductor. Also they pointed out the structure trans- formation from the tetragonal to the orthorhombic lattice for 1-2—3 phase, strongly depends on the oxygen content and the ordered state of the oxygen vacancies in the lattice. This systematic depression of Tc is correlated with the decrease in the oxygen content which alters the electronic structure of the Cu-O network. There is a reversible phase transition from non-superconducting tetrago- nal to superconducting orthorhombic at about 750°C [29]. This means at high temperatures, oxygen and oxygen vacancy ordering are supposed to be dis- tributed at random and the structure becomes tetragonal with parameters a = b. Since the phase transition is associated with changes in oxygen content, the mechanism of high Tc superconductivity is associated with oxygen vacancies. It has been known that oxygen loss and disorder only occur on CuO basal planes and that the rest of the layers do not participate in the order-disorder transition of the oxygen and vacancies. 11 2. 1. 3 Mechanical and electrical properties in Y-Ba-Cu-O superconductor Ever since the discovery of high transition temperature superconductors [1], considerable efforts have been expended toward the investigation of this and other high temperature superconductors (HTS). A number of difficulties, however, have to be overcome before practical applications of bulk HTS can be realized. There are two main problems before these materials are successfully utilized in actual devices. One of these is the inherent brittleness of HTS. The brittleness in mechanical properties, is related not only to the intrinsic nature of ceramics with a layered structure, but also to the presence of porosity unavoid- ably present in the sintered samples. The mechanical properties of single or poly-crystalline YBach307-x are comparable to those of typical brittle ceramics, with fracture toughness values (KIC) of about 1 MPa rnl/2 [33-35]. The fracture toughness of melt processed samples, which was measured by indentation method, wasestimated to be 1.2 MPa m“2 [36]. Estimated fracture strength of pure 123, according to Nishio et. al. [37], was about 25 to 40 MPa depending on processing conditions. Also fracture strength and Young's modulus of pure 1-2-3 produced by liquid phase process- ing technique was 40 MPa and 81 GPa respectively [38]. A number of research projects have been conducted to improve the me- chanical and fracture properties of HTS and different superconducting compos- ites, especially in particulate composites[39-41]. Among these attempts, the 1- 23 system containing Ag [42-49] has been most extensively studied. Nishio et al. [37] prepared composite superconductors by sintering a mixture of YBCO 1 2 and Ag powder at various proportions, ranging up to 92 volume %, and were able to enhance the mechanical strength against fracture to about 4 - 5 times as large as that in the sintered YBCO sample. This was achieved without seriously lowering the superconducting transition temperature. A liquid phase process- ing method has been developed to manufacture bulk Ag/YBCO composites by Lee et al. [38]. In nominal 15 wt % Ag composites, the measured Young's modulus was 100 GPa and fracture strength was 70 MPa. Ag/YBCO composite, reinforced by adding Zr, has been synthesized by Oka et al. [50]. The flexural strength in Ag/ZTYYBaQCU307-x reached its maximum value of 280 MPa at the composition Y = 0.8 and 33 vol. % Ag. This value is more than 5 times as large as in the sintered YBCO. They pointed out the enhanced mechanical strength was attributed to the unique microstructure, in which dispersed Ag particles remain intact with Zr—bearing YBCO matrix, whereas CuO and BaZrOs precipitates in the 1-2-3 phase matrix act to reinforce the ceramic without dis- rupting the superconducting channel. 1-2-3/Ag composites manufactured by solid-state sintering are character- ized by randomly oriented grains of 10 - 40 um size. Typically, silver is found to reside in the voids and along the grain boundaries, with small amounts of Ag entering the 1-2-3 lattice [2]. Sintering of 1-2—3 superconductor is enhanced by Ag addition, and the percentage of theoretical density of these composites in- creases with Ag content. Consequently, mechanical properties like Young's modulus and fracture strength increase with Ag addition [4—5]. Unfortunately, even though the mechanical properties of the HTS are im- proved, the transport critical current density (Jc) of these composites remains between 50 and 600 Mom.2 [51], which is insufficient for practical applications. For practical applications, do must be on the order of 105 A/cm2 at magnetic fields on the order of 1 to 10 Tesla at 77 K [52]. However, only for epitaxially I. 13 grown thin films [53] and single crystals [54] of 1-2-3 the critical current density has reached the range of 105 - 105 A/cm2. Unfortunately do of bulk 1-2-3 super- conductors is typically three orders of magnitude lower [51]. Highly anisotropic crystal structures of all of the presently known high Tc oxides lead to strongly anisotropic electron transport properties in both the nor- mal and superconducting states, especially, in single crystals. The anisotropic behavior of single crystal YBach307-x has been studied by several investiga- tors [55, 56] by magnetization of the crystal with the field applied both parallel and perpendicular to the 00-0 planes. Fig. 4 shows magnetic hysterisis loops at 4.5 K for the Cu—O plane perpendicular to the field lines and parallel to the field line. In Fig. 4(a), actually the induced screening currents flow parallel to the Cu-O plane. Fig. 5 shows the field and temperature dependance of the criti- cal current density determined from the hysterisis curves in Fig. 4. As can be seen in Fig. 5, single crystal of YBaZCU307-x demonstrates a strong anisotropy with the good conducting directions being along the Cu-O planes. The progress toward the much expected major application of bulk high TC superconductors has been hindered by their low critical currents, especially in strong magnetic fields. In sintered superconductors, Jc decreases very rapidly with applied magnetic field. For sintered samples of YBaQCU307-x, a flux den- sity of only 0.01 T depresses Jc from its zero field value by a factor of ten. For higher fields, the Jc vs. B curve flattens and decreases slowly with a slope which depends on temperature. Thus, in bulk polycrystalline materials, it would be advantageous to have a strong crystallographic texture with the basal plane of the grains parallel to the conductor axis. 14 0.04 0.02 - A 0 3 . E O V c -0.02 .9 B .5 '8 -0.04 C O) o 2 0.002 0 _0.002 1 1 l l l l l 1 EL -40 -20 0 20 40 Applied Field (k6) Fig. 4. Magnetization hysteresis loops at 4.5 K for a single crystal of 1-2-3 with the Cu-O planes oriented (a) perpendicular to the magnetic field and (b) parallel to the applied field (Ref. 55). 15 107 -i d -i .4 d d -i .4 TIT 3 : ‘J 4.5I HIN FILM (EPITAXIAL) S R 10 W W I 5 m FLUX g.“ 10 - WITH DEFECTS PINN'NG S (IRRAD., PPT.) GAP \ < \ .- 104— \ x GRAIN-ORIENTED t; (MELT-TEXTURED) I; 103 — U RANDOM—GRAINED IOZ- BOUNDARY GAP IC— 1 L J I l 1 10 100 1000 10000 APPLIED FIELD (GAUSS) Fig. 6. Transport Jo vs. H in YBa2Cuao7-x. The shaded area represents the improved intergrain Jc values in bulk YBach307-x with added flux pinning defects (Ref. 24). 18 observed withgrespect to the grain interior [61,62]. The large anisotropy in thermal expansion and the strain associated with the phase transition, result in microcracks between grains. These microcracks also hindered supercurrent flow. However, the weak link problem in high Tc superconductors can be greatly reduced by the elimination or minimization of high angle grain boundaries in the paths of the transport current. This can be achieved by a preferential crystallographic alignment of grains parallel to the a-b conduction planes (Cqu planes). Grain alignment, through melt-processing, have been successfully used in this regard [63, 64]. Salama et al. [65] developed a liquid-phase processing technique'that results in oriented-grains of 1-2-3 superconductors. They claim that do is as high as 75,000 Mom2 at 77 K and zero applied field. Directional solidification techniques are now widely used to produce bulk 1-2—3 samples with transport critical current densities exceeding 104 A/cm2 in applied fields of several tesla. Jin et al. used the term "melt texturing" to describe their pioneer- ing work in this area [64]. In the case of 1-2-3, melt texturing describes direc- tional solidification from the melt or partially melt state. While there are a num- ber of variations of the texturing techniques, they all involve heating 1-2-3 above its perltectic decomposition temperature (1 ,010°C in air) and slow cool- ing to form aligned, generally large, grains of the superconductor. Fig. 7 Shows a schematic illustration of melt-texture processing. The optical microstructure of the melt textured material, reveals near perfect alignment of the 1-2-3 super- conductor phase in the ab direction along the length of the sample. Fig. 8 shows the effect of temperature gradient on the melt-textured microstructure. The melt-textured Y-Ba-Cu-O exhibits significantly improved Jo and greatly re- duced field dependency as shown in Fig. 8. Although a desirable microstruc- ture has been identified for producing high Jc values, the difficulty, in practice, of TEMP 1°C) 1300 ‘- 1200 1100 1000 900 § ‘\ \ \‘ \ *- _LI0.+ so - X I I I x I ‘1 .I I I \ g; V \ \ \ ‘ LIQUID \ \j'l ‘ 19 OECOMPOS I T ION CRYSTALLIZATION 800 - .. .. l. 20 L .. I Y280C005 Y80200307 BoCuOZ TIME + - CuO COMPOSITION (a) (b) Fig. 7. A schematic illustration of time, temperature and phase equilibria, Fig. 8. associated with melt-texture processing (a) A qualitative sectional phase diagram (b) Three basic processing steps, shown as the temperature-time curve (Ref. 24). LOW ANGLE 6,8, /’ LOW ANGLE 6.8. [HIGH ANGLE 6:8. - ta N0 GRADIENT ‘ GRADIENT (a) ( b) The effects of temperature gradient on the melt-textured microstructure. (a) Long-range grain alignment, obtained with a temperature gradient. (b) Locally melt-textured structure, obtained without a temperature gradient (Ref. 24). 20 achieving highly aligned microstructures over extended distances is formidable. The longest reported distances, over which "single domain" types of structures have been generated, are on the order of 5 cm or less. Also another difficulty to produce grain alignment is extremely slow travel speed (Bridgement type pro- cesses) necessary for processing. As stated above, even though progress has been made in this area, especially to increase Jc by melt-texturing techniques, there are still a number of difficulties to be overcome. The major challenge is to develop a commercially acceptable processing technique which yields high value of Jc, maintains a criti- cal temperature of ~90 K and provides sufficient fracture toughness for design purposes. 2. 2 Superconductor Processing: Techniques and Mechanisms 2. 2. 1 Sintering and Hot Isostatic Pressing Generally, in conventional Sintering process, densification occurs by capillary-driven, long range diffusion processes. The driving force for sintering is a reduction in the system free energy via the decreased surface curvatures and an elimination of surface area. A characteristic feature of sintering is that the rate is very sensitive to temperature. Consequently, full densification can be achieved only at temperatures approaching the melting temperature where dif- fusion becomes rapid. However, the achievement of high homologous sintering temperatures (~ 0.8 of melting temperature) become technically difficult and also is limited by the phase stability (see phase diagram in 8 (b)). This pre- lcludes ordinary sintering from being an effective process for achieving full density. 21 Hot Isostatic Pressing (HIP) is a process in which sintering is enhanced by the application of high isostatic pressures. The HIP brings about the consolidation of powder materials by gas pressure at elevated temperature through hot deformation processes. This means that relatively rapid densifica- tion can be carried out at a lower temperature in a relatively short time. The use of the hot isostatic press has proven to give fully dense metals and ceramics in other applications [66—69]. For polycrystalline YBa20u307-x, it has been difficult to densify to full theoretical density by HlPing, while maintain- ing the superconducting orthorhomic phase [70-73]. Basically, there are three different types of HIP techniques: (1) post-HIP, (2) Sinter-HIP and (3) encapsulation HIP. In post-HIP the material is presintered to a density at which there is little or no interconnected or open porosity and then HlPed. With sinter-HIP, the material is sintered in-Situ in the HIP equip- ment, preferably under vacuum, until a closed pore structure is formed and then HlPed without removing from the apparatus. The third method, encapsulation HIP, requires the material to be contained within a gas-tight envelope whilst still in the green state and then HlPed. The envelope must be formed from a mate- rial which is deformable at the processing temperature so that the pressure can be transmitted to the specimen to aid the sintering process. The envelope must also remain intact in order to protect the material from the gas. Among three different HIP techniques, the main attraction of encapsula- tion HIP is that the material does not need pre-sintering to give a completely closed pore system and therefore high levels of sintering aids are not required. Another advantage is that a greater degree of microstructural control is ob- tained. By encapsulation HlPing from the green state, the increased densifica- tion rate at lower sintering temperatures, clue to the application of the pressure, 22 enables fine grained microstructures to be produced by limiting the extent of grain growth. There are many methods of encapsulation using glass materials [74]: glass capsule method, glass bath method, glass powder coating method and glass powder pressing method. Also high melting point metals such as tanta- lum can be used to form deformable cans which are packed under vacuum and sealed by welding. 2. 2. 2 HIP mechanisms as applied to YBBQCU307-X consolidation Several mechanisms are considered to contribute to densification during HlPing: yielding, power law creep (PLC), Nabarro-Herring/Coble creep (NHCC) and diffusion. When pressure is applied to packed particles, it is transmitted through the powder bed as a set of forces acting across the particle contacts. If an external pressure P is applied to the compact with a density D and coordina- tion number Z, the average contact force, f can be calculated according to the following equation [75]. I: 4nPR2/ZD, (1) where R is the average powder particle radius. The contact force pro- duces a contact pressure, Pefi, on each particle contact area of a: Peff = f/a = 41tFi2P/aZD (2) The deformation at these contacts is, at first elastic but as the pressure rises, the contact forces increase, causing plastic yielding and expanding the 23 points of contact into contact areas. Yielding occurs when the contact pressure exceeds yield Strength of the material, and hot deformation occurs almost instantaneously. Once these contact areas can support the forces without fur- ther yielding, time-dependent deformation processes determine the rate of fur- ther densification. These time dependent processes are power-law creep (in the contact zones) and diffusion (from a grain boundary source to the void sur- face). Power Iaw creep, a Slower process, is governed by the usual equation; éP = A 0” eXP(-cheep/RT) (3) where, s is the creep rate, A and n are material constants, cheep is the creep activation energy, R is gas constant and T is temperature. Nabarro-Herring/Coble creep (diffusional creep) occurs by grain bound- ary diffusion and sliding. Nabarro—Herring creep results from vacancy redistri- bution from a higher vacancy concentration in the region of a material experi- encing a tensile stress to the lower vacancy concentration regions subject to compressive stresses. This results in a vacancy flux from the former to the latter areas, and a mass flux in the opposite direction. Therefore, the grain elongates in one direction and contracts in the other, that is, creep deformation occurs. Nabarro-Herring creep is expressed by following equation: on = ANH (Di/d2) (on/kn (4) where, ANH represent geometric factors, DL is self-diffusion coefficient, d is grain size, 0 is applied stress, 0 is the atomic volume, k is Boltzmanns constant and T is temperature. Nabarro-Herring creep is accomplished solely by diffusional 24 mass transport and dominates creep processes at much lower stress levels and higher temperatures than those at which creep is controlled by dislocation glide. Coble creep is closely related to Nabarro-Herring creep and is driven by the same vacancy concentration gradient that causes Nabarro-Herring creep. However, in Coble creep mass transport occurs by diffusion along grain bound- aries in a polycrystal or along the surface of a Single crystal. Coble creep mechanism is expressed by following equation: éc = Ac (Dean/d3) (on/kn (5) where, Ac represent geometric factors, D33 is grain boundary-diffusion coeffi- cient, 0' is an appropriate grain boundary thickness, d is grain size, 0 is applied stress, 0 is the atomic volume, k is Boltzmanns constant and T is temperature. From the above equations (4 and 5), it can be seen that Coble creep is more sensitive to grain size than is Nabarro-Herring creep. Even though both forms of creep are favored by high temperatures and low stresses, it is expected that Coble creep will dominate the creep rate in very fine grained materials. Also both deformation modes are effective only when the powder is polycrys- talline and the grain size is small compared to the powder particle size. During diffusional creep of a polycrystal, additional mass-transfer pro- cesses must occur at the grain boundaries to prevent the formation of internal voids or cracks. These result in grain-boundary sliding and the diffusional creep rate must be balanced exactly by the grain-boundary sliding creep rate if internal voids are not to be formed. The mass transfer can be produced by vol- ume diffusion near the grain boundary, by grain-boundary diffusion, or by both mechanisms. The mass transfer is driven by vacancy concentration gradients in the same manner that diffusional creep is driven. Diffusional flow and l_________ 25 grain-boundary sliding, therefore, can be considered sequential processes in which mass is first transported by Nabarro-Herring and/or Coble creep and a grain shape change and separation is effected. Among several mechanisms, densification by long range diffusion, or Sintering, plays only a minor role in HlPing, dominating only at very low pres- sures. The overall behavior of densifying mechanism is quite complicated be- cause each has a different dependence on particle size, on the external vari- ables P, T, and on powder properties and current geometry. One way of analyz- ing it is to construct HIP maps which identify the dominant mechanisms and pre- dict densification rates and times, as a function of pressure and temperature. The HIP process of YBa20U307-x was modeled by Tien et al. [70] in terms of pressure, temperature, powder particle size, grain size and entrapped gas. They tried to assess the feasibility of hot isostatic pressing to obtain fully dense bulk superconductors using HIP modeling and experimental verification. For the calculation, the powder was assumed to be Spherical and monosized. 2. 3 Mechanical PrOperties of Ceamic Matrix Composites 2. 3. 1 The general theory of strength of fiber reinforced composites The use of high modulus fibers in a lower modulus matrix is well known and is a routinely practiced technique for increasing the strength. The general theory of fiber reinforcement suggests that significant strengthening will occur if the elastic modulus of the fibers is greater than that of the matrix, if tensile stress can be transmitted to fibers, but less toughening will occur. Conversely, if fibers of lower modulus are employed, the ultimate failure strength will be reduced while toughening will be increased because the matrix rather than fibers will 26 carry a greater proportion of the applied load. Stress may be transmitted to the fibers by plastic or elastic deformation of the matrix. Some crack-tip energy may be absorbed by the plastic deformation of a ductile fiber. In the absence of internal stresses, the strength, 00 of continuous fiber reinforced composites may be estimated by the simple rule of mixtures criterion, assuming the strains in each component are equal: Cc = (3fo 'I' Om'Vm for Vf > Vmin (6) or 00 = OM + OmuVm for Vf > Vmin (7) where the subscripts c, f, m refer to composite, fiber, and matrix. And Of is the fracture strength of the fiber, Vi, the volume fraction of fiber, Om', the matrix stress at the fiber failure strain, Vm, the volume fraction of the matrix phase, Vmin, the minimum volume fraction of fiber and Omu is the stress carried by the matrix when the composite is strained to ultimate failure strength. The validity of both of these equations is dependent on Vi exceeding a value Vmin. Also both equa- tions depend on which component fails first, and assuming that composite fail- ure occurs immediately following failure of one of the components. If the fibers fracture first then the composite strength is followed by equation (6). Conversely, if the matrix fractures first then the composite strength is given by equation (7). A second important volume fraction is ch which is that fraction of fibers that must be exceeded for fiber strengthening to occur. These two values are illustrated in Fig. 9. In Fig. 9, plot A represents equation (6) and B repre- sents: Cc = Gmu(1 ' Vi) (8) 27 AtOc = 01 V, + Om' Vm L01“ Om. I om' {- r I I I : B=0c = Umu (1 - V1) 0 VminycrII ' -—+Vg Fig. 9. Diagram illustrating Vmin and Van and their relationship to Cf, Gm and 6m. 28 The intersection of the two curves is at Vmin, is. where fiber and matrix break simultaneously and vent is always greater than Vmin; When omu - Om' is small then little strengthening is required to exceed the matrix strength, and this can be achieved by a small volume of fiber, is. ch is small. Conversely when Omu - O'm' is large and the matrix work hardens effectively then much more strengthening is required to exceed Om“, and vent is large. Therefore, ch is of most importance where composite tensile strength is the criterion. The fibers in the composite are discontinuous, the rule of mixtures cannot be applied. The ultimate strength of a given fiber can be utilized if it lies parallel to the tensile axis, and if its length exceeds a critical fiber length. Critical fiber length is defined as the minimum fiber length in the composite which can just be loaded to its failure stress. The stress-strain state at the fiber end does not cor- respond to the condition of equal tensile strain. When the tensile load is ap- plied, the strained matrix region cannot instantaneously transfer tensile load, in an amount equivalent to that borne by continuous fibers, to the fiber at the fiber termination-matrix interface. Instead, tensile force is transmitted from the matrix to the fiber by means of shear stresses which develop at the fiber-matrix inter- face due to very different elastic modulli. The tensile force borne by the matrix deforms it and the matrix is displaced relative to the fiber along this interface. The relative displacement is zero at the fiber midpoint and a maximum at the fiber ends. The displacement results in an interfacial shear stress, greatest at the fiber ends, that is capable of transmitting tensile load to the fiber. The man- ner in which this occurs is shown schematically in Figs. 10 [76] and 11. Two consequences of this dominate the performance of discontinuous fiber compos- ites: (1), a portion of each fiber, near its ends, will not be fully loaded and will thus be ineffective in strengthening the composite, and (2) the nature of fiber/matrix bond is of critical significance in determining the Strength. 29 toodon w .— matrix QFibrc Y ,‘j__.¢5 J F’c .. "I l - I I I Matrix «Ii-- Fig. 10. Schematic illustration of the distribution of fibre tensile stress (0) and interface shear stress (I) along a short fibre embedded in a matrix (Ref. 76). I I I I I I I I I I I I I I l -------- I I I I I l 1 l I I I I I I Fibre midpoint Fig. 11. The variation of 13m and Gm with position along the fiber when the matrIx (and fiber) deforms elastically. 30 Various theories have been proposed to account for the detailed varia- tion of the fiber and interface stresses, but for most purposes these can be represented by a Simple linear function. If B is defined as the build-up of tensile stress in the fiber, the strength of a discontinuous fiber composite can be written as: 00 = 5fo I1'(1' 3) 1c/II + OmuVm. Vf > Vmin (9) In this expression, IC is the critical fiber length (a fiber shorter than this cannot be loaded to its breaking stress) and 5f is the mean fiber stress. If the length ex- ceeds a certain critical length, the fiber midpoint, and regions adjacent, experi- ence a tensile stress corresponding to the equal-strain condition. The value of Of is determined by the fiber and matrix properties through the relationship: 1C/d1 -.= O'f/21: (10) where df is the fiber diameter and 1: is the interfacial stress. Clearly if I >> to, equation (9) becomes equal to equation (6) or (7). Determination of how effec- tive discontinuous fibers are, requires knowledge of lc/CIf, where IC/df is called the critical aspect ratio. Fig. 12 schematically illustrates the effects of several variables which control the strength of the composite including: the fiber volume fraction,fiber strength and fiber length [76]. As this figure indicates, the maxi- mum Vi obtainable decreases as the fibers become shorter. If randomly oriented fibers are employed, the proportion of fibers, ca- pable of being loaded to their fracture stress will be reduced, and hence the ul- timate strength of the composite will be lower than that of an equivalent aligned 31 O'Vmox . I. 1 a) 5 O Q I— D ’ t5 I . l s 9 .. a) . :9 5 ’ “5, b c to 27) a) to ’5 Matrix UTS g o a. .E E 1 0 l O “c = 5,1; [l-(l-BIlc/II +o,',,ll -I/,I Short fibres 0 Vi Fig. 12. The effect, (shown schematically). of fiber volume fraction, V], fiber length, l, and fiber strength, or, on the strength, do, of the composite (Ref. 76). 32 system. Suitable efficiency factors have been derived to take into account the effect of fiber orientation and fiber length. 2. 3. 2 Theory of energy absorption mechanisms The incorporation of low modulus fibers into relatively high modulus ce- ramic materials was originally carried out in order to increase the fracture toughness, rather than the strength of these brittle materials. Fibers are effec- tive in providing barriers to crack propagation and thereby improve the fracture toughness of a material. Their precise role, however, is complex and difficult to describe or quantity fully. There are a number of mechanisms by which the total work of fracture may be raised. In a poorly bonded composite, the stresses near the crack tip could cause the fibers and matrix to debond from each other before the fibers break. When total debonding occurs, the elastic strain energy in debonded lengths of the fiber cannot redistribute stored energy in the rest of the compos- ite. When the fiber snaps in the debonded region, the energy is dissipated as heat I Fiber stress relaxation, a variation of the debonding model, estimates the stored elastic energy lost from a snapped fiber when the interfacial bond is not destroyed and elastic energy is lost by the fiber over its critical fiber length. It is also possible for a matrix crack to propagate past a fiber leaving it unbroken and bridging the crack. Crack bridging can result in toughness by a combina- tion of a process described above. Another mechanism which can enhance the toughness is to increase'the effective crack propagation area. This can occur by having planes of weakness within the material, in a direction parallel to the tensile axis along which a 33 propagating crack may be deflected, thereby effectively blunting crack tips. Aveston has shown [77] that very high values for the work of fracture are possible by this mechanism. If the fiber-matrix bond is sufficiently strong, a crack may propagate rela- tively unimpeded through the composite, and the work of fracture or toughness will be low. An increased bond strength would be expected to increase the in- terfacial shear strength and decrease the debonded length in the debonding model. In that case, the critical fiber length, lc would be decreased. If discontinuous fibers of length 1, where the length is less than the critical length, 1C, the fibers will not be loaded to their fracture stress and they must be withdrawn from the matrix as the fracture planes separate. It has been shown by Cottrell [78], that the contribution due to pull-out can be very significant, and the maximum value YPmax. is obtained for I = 1c, when YPmax, =Vlec/12r (11) where 1: is the interfacial shear stress resisting pull-out, and r is the fiber radius. If the fiber length 1> 1c, some pull-out is still observed, since many of the fibers will interseCt the crack plane within a distance 1cl2, and hence a fraction Id] of the fibers will not be loaded to their fracture stress, and must pull-out as the planes separate. The fact that pull-out effects have also been observed for many continuous brittle-fiber systems, has been attributed to the statistical dis- tribution of strength along these fibers, allowing fracture of many fibers to occur at positions away from the fracture plane. In the case of ductile reinforcing fibers, a large contribution to the fracture toughness of the composite may be provided by plastic flow and rupture of the fibers. For fibers of length, 1> 1c, the contribution, yr, due to fiber rupture will be 34 Yr: Vivi (l -1c/1) (12) where yr is the rupture energy. In composites with random fiber orientation, many fibers will cross the fracture plane obliquely and will be subjected to a bending moment. For the case of brittle fibers this will have the effect of lowering the applied tensile stress required to fracture the fibers. For ductile fibers, plastic bending may accommo- date the extra strain at the convex portion of the fibers during withdrawal from the matrix, and may contribute significantly to the total work of fracture. If fibers make a small enough angle to the fracture plane, they will either fall in shear, or will break through the matrix. It has been shown that the maximum work of frac- ture due to pull-out of aligned ductile fibers of critical length is significantly greater than that which can be obtained by plastic bending alone of misaligned fibers. However, the deformation of misaligned fibers can make a substantial contribution to the work of fracture when fibers of less than the critical length are employed, and would not yield very high values when aligned. Unfortunately, with randomly-oriented reinforcement, some fibers will lie approximately perpendicular to the tensile axis and may act as stress-concen- tration sites, particularly if bonding is weak, consequently reducing the Strength of the composite. In general, reinforcement of ceramics by metal or ceramic fibers, although providing a substantial improvement in fracture toughness, par- ticularly when using metal reinforcement, often leads to little improvement, or even a reduction, in the mechanical strength [79]. This is because the relatively large diameter fibers employed, normally greater than 50 pm and often as large as 200 pm, acted as stress concentration sites which weakened the matrix and toughness was therefore increased at the expense of strength. Significant 35 strengthening was possible due to a combination of factors, including the influ- ence of small size and the large difference in elastic modulus between high modulus of reinforcing fibers and the relatively low modulus of the matrix. Of the above energy absorption models, debonding and pull-out mech- anisms have been most widely accepted. Table 1 summarizes various theoreti- cal models for energy absorption mechanisms which was studied by a number of researchers. 2. 3. 3 Stress-strain curves of fiber reinforced composites For a suitable combination of the toughening mechanisms described above, a non-catastrophic mode can be obtained, as characterized by the ten- sile stress-strain curve of Fig. 13 [80]. Oualitatively, this failure mechanism is fa- vored in composites with weak interfaces, high strength fibers and tensile resid- ual stresses normal to the fiber/matrix interface. Changes in any of these pa- rameters can lead to a transition in failure mechanism to one which is catas- trophic, with a linear stress-strain curve to failure. There are generally three regions of failure in fiber reinforced compos- ites: (1) linear stress strain behavior before matrix cracking (BMC), (2) nonlinear region with increasing stress and Strain after multiple matrix cracking (AMC) oc- curs with possible crack deflections at the fiber-matrix interface, and (3) a region of decreasing stress with increasing strain where fiber pullout (FF) is occuring. For each fiber (or a bundle of fibers) pullout in region 3, a slight decrease is seen in stress locally as the general load is also decreasing. The initial departure from non-linearity in both types of stress-strain curves results from cracking of the matrix. For the non-catastrophic mode of failure, the first crack in the matrix extends indefinitely, breaking only a small 36 Table 1 Theoretical models of energy absorption associated with mechanical deformation of fiber reinforced composites Model Pull-out (short fibers of length l) Debonding Stress-relaxation Crack-bridging Matrix plastic deformation Fracture surface energy 7 Vi 0r Ic2 12I VrOr'2 12 IC Yr 01 2y 4Ef (|>IC) (I 15 %. 3. 1. 2 Sample Preparation (hot isostatic pressing) For hot isostatic pressing and specimen preparation, an lPS Eagle-6 Hot Isostatic Press system (International Pressure Service, Inc.) was used. The Ea- gle-6 uses a 6 inch internal diameter stamped monolithic pressure vessel with threaded closures, rated around 200 MPa. This equipment is capable of reach- ing a temperature 2,000°C, using a graphite furnace, which is a modular plug-in unit with a hot zone measuring 3 inches in diameter and 4 inches in length. Argon gas was used as the pressure medium. A microprocessor-based controller was utilized to control and monitor the process temperature, pressure, time, valve position, compressors, volts, amps and electrical contacts. Pressure and temperature was controlled by a programmed set point controller. A digi- tal/analog strip chart recorder was used to record system pressure and temperatures as a function of time. Temperatures was monitored by using 5 different tungsten/rhenium thermocouples located at different levels within the furnace and the pressure vessel. Assigning a true temperature to the run was difficult because neither of the thermocouples is directly attached to the specimen. Therefore temperature, in our experiments, represents a bracketed temperature. within which the upper and lower bounds. The uncertainty in temperature measurement was estimated approximately within a: 5 C of reported values. 54 Before HIP processing, given samples were conventionally sintered at 925°C for 2 hours only. The reason for this short presintering was to give some strength to the green compact so as to enable easier handling during further fabrication stages. The bar was coated with 99.99 % purity boron nitride using BN spray (Union Carbide Co.) and then it was wrapped with a 0.025 mm thick Cu foil (Aldrich Chemical Co.). There are two main reasons for using the BN spray and Cu foil. One is to reduce the oxygen partial pressure, especially at high temperatures, by induc- ing the chemical reaction between Cu foil and residual oxygen to form CuO layer. The other is to easily remove the CuO layer from the surface of the specimen separated by the BN spray. The BN sprayed and Cu foil wrapped sample was placed into a 12 mm (ID) diameter and 2 mm thickness pyrex tube. For easy decapsulation, high thermal conductivity BN powder (HCP, Union Carbide Co.) was placed between the Cu foil wrapped sample and the glass wall. The pyrex tube, with the sample, was evacuated for 20 minutes under a vacuum of 10'2 Pa and finally sealed by torch. The sealed capsule was placed into a specially designed rectangular type graphite crucible and carefully in- serted into the heating zone of the graphite furnace. In the HIP, in the first stage, the system was evacuated for 30 minutes and then purged 3 times by argon gas. The pressure was kept low to avoid glass cracking as the temperature was raised to 820°C, which is the pyrex-softening temperature. The applied heating rate was 8°C/min. Temperature and pressure were then raised up to 890°C and 170 MPa respectively and kept for 2 hours. At last, the temperature was slowly cooled to 500°C at a rate of 2.5°C/min and cooled down to room temperature. Pressure was vented. Figs. 17 and 18 show the detailed HIP processing cycle in terms of pressure, temperature and time. Specimens were removed from the glass wall of the capsule by using the grit blast method. After decapsulation, the 55 Mixing Y203, BaCOa and CuO in 1-2-3 ratio 1 Calcined and Controlled o particle 930 C, 24 hrs, 2 times less than 30 um I Composite Mixing 5, 8, 12, 15 volume % of 50 um Ag fibers (~1 cm) Cold Pressing 6 mmx 9 mmx 38 mm at 100 MPa l Presintering 925 C, 2 hrs in Air l ‘Pyrex Capsulation Cu Foiled, BN filled Vacuum and Sealing . HIPPing 890 c, 2 hr 20 min, 170 MPa l Decapsulation Grit Blast Method T Annealing . Slow heating rate (3 C/min) at 850 C for 20 hrs # HlPPed Sample (5 mmx7 mmx 33 mm) Fig. 17. Schematic diagram for the HlP processing cycle. 56 (ed w) amssaid cm 09 om? .xKOmaommmcr 82259 on: m< ac mc_n=I mezzo oan 9339993986258P .9 .9”. :5 we: was. on— wcammmcdn l l 9398.55.” l llLLl l J O O V § (3°) amteiadurel com 8 57 sample surface was found to be somewhat contaminated due to reaction with the glass wall and boron nitride. Typically, the color of the sample surface was green. This green phase was later identified as a mixture of Y2BaCu05 (2:1 :1) phase and CuO from x-ray experiments. However, this layer was usually less than 0.4 mm from the surface and disappeared during surface polishing and subsequent post-annealing processing. Annealing was conducted, at 850°C in air for 20 hours with a heating rate of 3°C/min. Typical dimensions of the H lPed - specimens after annealing and polishing was 5 mm x 7 mm x 33 mm. 58 3. 2 Sample Characterization Methods 3. 2. 1 X-ray diffraction studies For the structural analysis and phase identification, x-ray diffraction ex- periments were carried out by using a SCINTAG diffractometer (Scintac XDS 2000 diffractormeter, Scintac Co.) with a Ni filter CuKa radiation. A tube voltage of 35 kV and a tube current of 25 mA was used for all structural analysis. For diffraction studies of as-prepared samples, flat sample surfaces, which were mechanically polished, were used. 3. 2. 2 Electrical resistance measurements The electrical resistance measurement was performed by using a stan- dard four-probe technique (LR-400 four wire AC resistance bridge) on the regu- lar geometry (5 mm x 7 mm x 15 mm) specimens. Figure 19 shows the schematic diagram describing the experimental set- up for resistance measurement. Also. a schematic of the four-point probe is shown in Fig. 20. In this diagram (Fig. 20), four gold-plated pins were attached to the test specimen. The two outer pins were for current source and the inner two were for potential measurements. Four gold-plated pins and a calibrated T- type (copper-constantan) thermocouple were placed on the regular geometry of the specimen (5 mm x 7 mm x 15 mm). The specimen, thermocouple and pins were wrapped between two poly-vinyl chloride blocks (Fig. 21). All were then clamped by two plastic clamps and were immersed into the liquid nitrogen de- war for cooling. 59 AC resistance Bndge Sample Cooling System X-Y Recorder Fig. 19. Schematic diagram describing the experimental set-ups for resistance measurement. Digital Thermometer 6O LR-400 Four Wire AC Resistance Bridge Preamplifier Circuit Sample \ Current Probe Voltage Probe Current Probe Fig. 20. Schematic of the four-probe technique. 61 Stand / PC board Clamp / W \4 gold plated 4-point .. probe Sample Thermocouple — PVC plate QR QR Fig. 21. Schematic diagram of sample holder assembly for AC resistance measurement. 62 The electrical resistance was monitored by using the LR-400 AC resis- tance bridge (Linear Research Inc.). The monitored resistance was calculated by the voltage drop caused by passing a given constant current, while the tem- perature was changed from the liquid nitrogens boiling temperature, 77 K, to room temperature. The applied current used for this experiment was 3 mA. Temperature was monitored by using a digital thermometer and the uncertainty in this measurement was estimated approximately within HOC. Temperature vs. resistance (voltage) was continuously recorded by using a Houston Instrument 200, X-Y recorder. 3. 2. 3 Critical current density measurements Critical current density was measured by using a standard four-probe technique at 77 K with zero magnetic field. The temperature was held constant at 77 K while the current through the sample was varied. Fig. 22 shows a schematic diagram of the critical current density measurement set-up. A typical specimen size was 1 mm x 1.5 mm x 15 mm. The applied cur- rent range was from 0.2 to 9 Amp and the resistance was calculated from the voltage drop. The value of Jo was then deduced by dividing the critical current, at which the generated voltage was extrapolated to zero, over the effective cross section of specimen. In order to minimize the surface contact resistance, only the sample surfaces under contact were painted with Ag. The coated sample with Ag were then annealed at 450°C for about 2 hours and slowly cooled down to room temperature in air atmosphere. Cu wire was then embed- ded into the Ag dispersed surface by Indium soldering. Ag dispersion by an- nealing was believed to have created the continuous network to give a stable bonding between the superconductor and the embedded wire. 63 DC Power Supply Digital Ammeter Digital Voltmeter 1 2 3 Current probe 4 \ Test Sample Current probe Voltage probe Fig. 22. Schematic diagram describing the experimental set-ups for critical current density measurement. 64 3. 2. 4 Magnetic moment measurements The Meissner effect is a manifestation of the superconducting state. The Meissner effect is a phenomena in which the specimen expels all external magnetic fields when cooled through the superconducting transition point. Flux exclusion by a superconductor is the underlying principle associated with the Meissner effect, and is akin to a perfect diamagnetic state. For the magnetic moment measurement, a Quantum Design Magnetic Property Measurement System (MPMS, Model 2000 VHF SQUID, Quantum Design) was employed. The MPMS is a sophiscated analytical instrument specifically designed for the study of the magnetic properties of small experi- mental samples over a broad range of temperature and magnetic fields. The system hardware has two major components: (1) the MPMS dewar and probe assembly, and (2) the associated control system in the MPMS control console (Fig. 23). Automatic control and data collection are provided by an HP com- puter and two independent subsystem controllers. Most of the gas control and other ancillary functions in the system are also automated. The cryogenic probe integrates a 5.5 Tesla superconducting magnet with a SQUID detection system and a high-performance temperature control system to provide rapid precision measurements over a temperature range of 1.9 to 400 K. Liquid helium pro- vides refrigeration for the SQUID detection system and magnet, as well as pro- viding for operation down to 1.9 K. The sample handling system allows auto- matic sample measurements and position calibrations using a microstepping controller having a positioning resolution of 0.003 mm. The equipment is ca- pable of resolving variations in magnetic moments as small as 10'8 emu. The measurement is made by inserting a pair of secondary coils into the sample area. Also, external to the secondary coil, is a long primary coil. The $15!!! COMPONENTS 1. Sample Rod 2. Sample Rotator 3. Sample Transport 4. Probe S. Helium bevel Sensor 65 6. Superconducring Solenoid 7. Flow impedance 8. SQLTD Capsule 9. Dewar Cabinet 10. Dewar 11. HP Thinkiet Printer 12. Magnet Power supply 13. Model 1802 Temperature Controller 14. Console Cabinet 15. Power Distribution L‘nit 16. Model 1822 MPMS Controller 17. Gas/Magnet Control Unit 18. HP Vectra Computer 19. Monitor Fig. 23. Magnetic Property Measurement System (MPMS) components (Ref. 91 ). 66 sample is inserted into the center of one of the two secondary coils. A low-fre- quency current is passed through the primary coils. Any change in the flux linking or Meissner effect in the sample will yield a voltage (mutual inductance) in the secondary coils. A lock-in amplifier allows the magnetic susceptibility change to be detected. For the magnetic moment measurement, about 1 gram of each different volume % Ag reinforced specimens were used. The samples were cooled down to liquid He temperature in a zero field and magnetization was measured from 40 K upon warming. The applied magnetic field used for this experiment was 20 G. The measurements were taken at 5 K intervals, from 40 to 50 K, 2.5 K intervals between 50 and 80 K, and 1 K intervals from 80 to 97 K. Standard deviation in this magnetic moment measurement was around 10'8 emu. 3. 2. 5 Density and porosity measurements The density of each specimen, with different volume % Ag fibers, was measured according to ASTM B328-73 (buoyance method). About 2.5 grams of specimen was cleaned with acetone for 20 minutes and then completely dried. After cleaning, samples were weighed in air atmosphere (WA) and immersed in SAE 10 W-40 motor oil (viscosity of approximately 200 SUS at 100 F) for 4.5 hour and held at 185°C. Then, the temperature was lowered to room tempera- ture by immersion in oil at room temperature. The oil immersion introduces im- pregnation of the sample through interconnected pores (open pores). Following this, the excess oil was wiped off with a damp cloth and the samples were weighed again (WB). The samples impregnated with oil, were then tied with a‘ 0.09 mm diameter copper wire and suspended from the beam hook of a semi- micro balance. Samples were completely immersed into a beaker filled with 67 distilled water at 20°C, which was placed underneath the beam hook. The wired sample was weighed (W0) and then the wire without the sample was im- mersed again into the distilled water for measurement (WE). The density of sample can now be calculated from: D=A/(B-C+E) (20) where D = density, g/cm3 A = weight in air of oil-free specimen, 9, B = weight of oil-impregnated specimen, 9, C = weight of oil-impregnated specimen and wire in water, 9, and E = Weight of wire in water, g. The measured density was compared with the theoretical density of Ag- YBagCU307_x which was obtained from the rule of mixtures. The porosity was calculated by subtracting the ratio of the measured density to the theoretical density from 1. Theoretical density of YBa2CU3O7-x and Ag are 6.375 g/cm3 [92] and 10.491 g/cm3 [93] respectively. 3. 2. 6 Flexural strength and fracture toughness measurements For strength measurements, specimens were unnotched. For fracture toughness measurements, the specimens were machined into the single edge- notched (SENB) configuration. Typical sample dimensions were 5 mm x 7 mm x 33 mm. Seven specimens, with the same volume % of Ag, were prepared for measurements. The notch was made by 0.016 mm thickness diamond blade, and the nominal notch depth was around 0.8 mm. An instron machine, at a 68 cross head speed of 0.05 mm/min was used for all fracture tests. The testing was carried out in a 3-point bending apparatus with a span width of 28 mm. Fig. 24 shows the schematic diagram for the 3-point bending apparatus. Strength and fracture toughness values were determined from the specifications stated in ASTM D790M-82. Typical span to depth ratio was kept to 6 : 1 in order to min- imize the shear contribution to the fracture mode. From beam theory, which is based on the linear elastic fracture mechan- ics under plane strain conditions, the fracture stress (cm) of the unnotched specimen was determined as follows: 0F .-_- 3PL/2bd2 ' (13) where P is fracture load. L is the span width, and b and d are the specimen width and depth respectively. Also values for KIC were measured by the follow- ing equation: KIC = 3PL/2bd2 001/2 Y(Co') (21) where Y is the geometry factor related to the specimen, Co is the notch depth and Co' = Cold: relative notch depth. The Y(Co') relation given by Gross and Srawley [94] Is Y(Co') = 1.99 - 2.470.; + 12.97(co')2 - 23.17(co')3 + 24.8(00')4 (22) The fracture energy parameter,y, was estimated by examining the area under the load deflection curves. 69 )9 T , J'Tlfo fl." 145/T / Machlned Notch . if L .Jl. Fig. 24. Schematic diagram for notch-beam test in 3-point bending. 70 3. 2. 7 Morphological examination and EDAX analysis 3. 2. 7. 1 Optical microscopy The microstructure was observed on the polished surface of the Ag rein— forced specimens by using a Neophot-21 optical microscope (Leco 00.). Sam- ples were metallographically mounted on a Iucite block. Lucite mounted specimens were then polished on a nylon cloth with 1 to 5 um alumina powder and 0.3 - 0.5 um diamond paste. Methanol was used to prevent any possible degradation of the microstructure due to moisture pick-up. 3. 2. 7. 2 Scanning Electron Microscopy and EDAX A HITACHI 8-2500 SEM with Link energy dispersive x-ray spectroSCOpy was used to examine the microstructure and for the presence of elements on the fracture surface of the Ag reinforced specimens. Fractured specimens were mounted‘on cylindrical aluminum stubs. Silver paint was used for electrical contact to ground and also to provide a better mechanical support. 3. 2. 8 Scanning Auger Electron Microprobe Qualitative analysis was carried out to obtain information about the inter- face and elemental distributions by using a Scanning Auger Multiprobe (SAM) (PHI 660, Perkin-Elmer Corp). AES (Auger Electron Spectroscopy) is a very powerful technique often used together with XPS (X-ray Photoelectron Spectroscopy) for surface chem- istry studies. The basic principle of the Auger effect is the de-excitation of an 71 ionizedatom by a non-radiative process. When an electron is ejected from an inner shell of an atom by photon or electron bombardment, the resultant vacancy is simultaneously filled by an electron from one of the outer shells (Fig. 25). The energy released in this decay process is transferred to a third electron that is emitted outside of the sample (the Auger electron). A simplified model for calculation of Auger electron energy was developed by Burshop [95]. The basic equation in this model is: EZ(XYZ) = EZ(X) - EZ(Y) - EN) (23) where EZ(XYZ) = energy of the ejected Auger electron from element with atomic number Z EZ(X) = binding energy of the level on which initial core hole is created EZ(Y) = energy of level from which an electron falls to fill the initial hole EN) = energy appropriate to an atom already singly ionized in an inner shell. In this formula E(Y) and E(V') can be approximated as the binding energy of the Y and V' levels respectively. ‘ In Auger Electron Spectroscopy, the intensity of the emitted electron sig- nal is measured as a function of the kinetic energy of electrons. High resolution , spectra are usually obtained by using a cylindrical mirror analyzer (CMA). Auger spectra are always reported in the first derivative mode, by using the lock-in amplification technique [96]. Differentiation of the recorded signal ls done electronically by applying a modulation voltage on the incident electron beam and recording the "in-phase" component of the ejected electron beam. In the first derivative spectrum, the auger spectra appear much pronounced, and the background of inelastically scattered electrons is significantly reduced. ..... 72 or Auger electron .w—r . — —.— —.—r —.~.—..—_ I : . I}? _' x xx. ,rff/Z; 3 either photon 4. E2 Incident electron A \r e E1 . Electron 0 Electron vacancy Fig. 25. A schematic diagram illustrating the basic processes associated with Auger electron transition. 73 The surface chemistry and interface between 1-2-3 matrix and Ag fibers in a HlPed specimen, containing 5 volume % Ag fibers, were studied. Samples were prepared by two different ways, one by fracturing and the other by polishing. One of HlPed specimens with 5 volume % Ag fibers was fractured at 77 K, liquid nitrogen boiling temperature. The other specimen with 5 volume % Ag fibers was metallographically polished by using 1 - 2 um alumina powder and 0.2 - 0.5 gm diamond paste. Methanol was used to prevent possible degradation of the superconductor sample by moisture. Either fractured or pol- ished specimens were introduced into the Auger chamber which was controlled at ultra-high vacuum conditions (1.5 x.10‘9 torr). The SEM was used to exam- ine the surfaces of both specimens. Fig. 26 shows a fracture surface taken by SEM. Auger surface scanning was performed on both specimens at 10 keV electron energy and 50 nA beam current. For these conditions the beam diameter was approximately 50,000 nm. For qualitative analysis, data was acquired in N(E) mode (number of counts per energy interval) vs. kinetic energy and then peaks were recorded in the derivative mode (N(E) * E). However, Auger characterization does not give the exact answer for the quantitative analysis, overall element concentration could be roughly estimated in terms of relative intensity. In order to remove any residual surface contamination, ion sputtering was carried out on the fractured and polished specimens with a 4 keV Ar’r ion beam and a 1 mm2 raster size (the sputtering area). Sample preparation intro- duced a large amount of carbon, especially in the polished specimen due to methanol. The sputtered depth, calibrated by SiOz, was 5,000 nm/min and the controlled system pressure was 15 x 10‘3 Pa. Electron line scanning was carried out to study the chemical distribution of elements, the diffusion of Ag into matrix and the diffusion of 02 into Ag fibers 74 . . L l I ‘7 l ; >x13r92 ia.akv sao.a x 58.6mm Fig. 26. A fracture surface examined by Auger SPeCtVOSCOPY apparatus 0‘ 500)‘ 75 at the interface. Line scanning was performed on the ion sputtered fractured sample (Fig. 26) after electronically drawing the horizontal scanning line across the fractured surface. This line was focused on the interface between the Ag fiber and the matrix. The applied electron excitation voltage was 10 keV and the beam current was 50 nA. Data points of line scanning on the SEM micro- graph were taken at 2 um intervals and different element concentrations were measured as a function of scanning distance (in um). Estimated spatial resolu- tion during line scanning is approximately 50,000 nm. Typical main Auger peaks of each element (kinetic energy) [97] are summarized at Table 3. Compositional dot mapping for the distribution of chemical elements was performed on the ion sputtered fracture surface (Fig. 26). The applied electron excitation voltage was 10 keV and the beam current was 50 nA. Five different elements, Y, Ba, Cu, O and Ag are represented in terms of different contrasts in dot mapped pictures. In this experiment, Auger electrons of each element were detected around 200 to 500 nm from the surface. Table 3 Typical main auger peaks in YBa20U307-x composite [97] Elements Kinetic energy (keV) Y 127, 1746, 1821 Ba 500, 584, 600 Cu 840, 849, 920 O 468, 483, 503, A9 351,356 C 272 B 179 4. RESULTS AND DISCUSSION 4. 1 Crystal Structure versus Processing Conditions Figure 27 represents typical diffraction patterns of HlPed and unrein- forced as well as reinforced specimens with different Ag volume %. The diffrac— tion patterns shown in Fig. 27, are very similar to each other except for the extra peaks of Ag in the reinforced specimens and some variations in the relative in- tensity of prominent reflections. Relative intensity variations in prominent reflec- tions might be associated with the localized texture formation during the HlPing process [98]. To confirm the peak positions, x-ray diffraction peaks from con- ventionally sintered 15 volume % Ag reinforced samples were compared with the major peaks of pure Ag and YBagCU3O7-x phases. The YBBQCU307-X x-ray data was taken from Cava et al. [7] and pure Ag data was taken from powder diffraction index file (Fig. 28). From peak comparison, the diffraction pattern of the sintered 15 volume % Ag reinforced specimen, consisted of a superposition of diffraction patterns of pure Ag and orthorhombic YBaQCU307-x phases. Figure 29 shows the diffraction patterns of the 15 volume % Ag reinforced specimen, processed at different conditions: Fig. 29(a) sintered, Fig. 29(b) HlPed without post-annealing and Fig. 29(0) HlPed after annealing at 850°C for 20 hours. In this x-ray data, all of the diffraction peaks could be indexed. Comparing a conventionally processed sample in Fig. 29(a) with a HlPed sam- ple Fig. 29(b) or HlPed sample after annealing Fig. 29(0) it is seen that the diffraction patterns for the HlPed samples have additional peaks at a 26 value of 29, 30 (or 31) and 36.5 degrees. Composition and structure analysis [99] show that the phases responsible for these 26 values are YgBaCuOS and CUQO which are the secondary phases, resulting from the phase decomposition of 77 78 .mcoEBmam beef 80: mEouma cone—2:5 >m_-x .85.»... Km .9... lam -0 00 - a“ O I. g _ a. J . a . a . o«.i mafl «ea fi 00 I F 8 .. u< S... L. 0‘ J filial .I on I. m< O\Om I. oo.1 .Itiiisi R .. u< $3 .. 8 .. 3 $2 .. con _. _ _ Va a so.” on.u v0.0 0. 0.0«0 o.onu« o.ono« o.ovsa 0.000» 11 0.000» o.ount o.ouov O . OOVD O . Gown 010 93 new a so 9:39 9 53> 89259 :oEBm o< mxw .comtmano .8 $65 cum 9:55:50 mm; < 93 Co 96:68 cc: Eorma cozomctfi oocmcoamm .m< am 8 E0: :5qu c0309; 5 3-x .mm .55 £03039, «2% 6m /omFH/tmm>JHm \ mm>JHm n r a «a _ «a J - .mono znnmpe> mmamoo 23Hmao.s.m o . v.v«« o.om~ n.0wn o.nnv .o.unn V.ooo 0.000 u.o«o o.auo« o.vs«« ”to .wcoEccoo 9.6885 E2255 .2 Hm< xx. 0E2? 9 2:2, Soup—Em. mcméooam E0: 2553 cozomEfi >92 .mm .9“. low 00 00 DM Ow o 1 . . 1 .. a . . . . J . _ l o ‘ . C .‘ . or; C ‘ O . i 0 can m< 23a 0 d. O 1 C. S C . manna arm... 4 4 0.000 s .. $2.. Ia s. 4 4 o r a: 622:? on. m . on t O 0 I A V I 0 one as i . i 0.003 m ‘ . on J . .. . . _ f 0.33 O . k. on 4 0 an I 0.003 .82: a: e on i t 0.33 on m 0. O _ fl 4‘ Q I 0.003. . 4 4 4 4 4 4 a x C < C. ox down on 1 95355 coca. no.2: 9: A3 1 Q r 0 con . Q ”G. o.oon~ x - 35H. . so..« and V0.6 n10 81 YBach307-x. According to Wong et al. [99], Y2BaCuO5, called the "green phase", has orthorhombic perovskite type structure with lattice parameters of a = 71.31 nm, b = 56.59 nm and c = 116.80 nm. Therefore, the diffraction patterns observed after HlPing (Fig. 29) indicates the existence of a multiphase mixture: superconducting orthorhombic YBagCU307-x phase with some additional peaks of pure Ag and secondary peaks of Y2BaCu05. However, relative intensities indicate the major phase is the superconducting orthorhombic YBaQCU307-x phase. Fig. 29 also suggests that the silver remains intact with the 1-2—3 and other phases present. Another feature seen in Fig. 29 is the relative intensity decrease after HI Ping. The decrease in relative peak intensities is related to the change of atom positions within the unit cell. Based on this observation, the relative intensity variations after HlPing are presumably related to oxygen stoi- chiometry changes in YBach3O7-x system. Based on x-ray data (Fig. 29), conventionally prepared 1-2-3 after HlPing was found to be quite unstable and decomposed into other phases at the tem- perature of 890°C under 170 MPa. This phase decomposition is possibly asso- ciated with high pressure at the high temperature. The stress that causes this decomposition in the presintered sample is presumably the stress at the inter- particle contacts. The exact stress state at the inter—particle contacts varies throughout the sample. The load applied to the sample is distributed over the many inter-particle contacts, and the stress at each contact is a function of this contact load and the contact area. Since the total inter-particle contact area is less than the actual surface area of the sample, the average interparticle stress is always greater than the applied stress. For the case of isostatic pressure, the average inter-particle contact pressure, Pa, can be calculated as a function of density [75], as follows: 82 Pa = 43tPH|pR2/aND (24) for monosized powders, where PH”: is the isostatic pressure, R is the average powder particle radius, a is the average inter-particle contact area, N is the co- ordination number of the particle and D is density. From equation 25, the contact stress decreases rapidly as density in- creases because the average inter-particle contactarea increases as particles deform around each other. Possible decomposition reactions due to stress ef- fect are: YBa2Cu307-x—41/2Y2BaCu05 + ACuO + BCU20 + CBaCu02 + DMxOy (25) or 2YBa2CU307-x YgBaCuos + SCUO + 3880+ y02 (26) A recent x-ray study showed that the 211 phase (Y2BaCu05) can be easily detected after HlPing. It is known that a small amount of 211 phase is al- ways present even after careful preparation of 1-2-3 samples [100]. However, this 211 phase is quite commonly precipitated during a melt texturing process [63-65]. Based on equations 25 and 26, other second phase particles, such as CuO, or CU20 and BaCqu, should be detected in an x-ray measurement. However, the peaks related to those above mentioned phases were not strongly detected. This is possibly due to the small amounts of these phases present in the mixture. The three diffraction patterns for 1-2-3 samples processed at different conditions, are shown in Fig. 30. The bottom scan, Fig. 30(a), shows the diffrac- tion pattern for a sample presintered at 925°C for 2 hours and the center one, Fig. 30(b), shows a scan for a sample HlPed without any post-annealing. The .mcoEccoo @5880... EEotfi .0. 256.0me mm; .0 2.2.3 8.89:6 >mc-x dm .9“. law o .. no on . an mm o. 3 1 AOONV ( r. 0.00» on .. 883mg .1 0.8.» on... «-m; vocotzmoa :3 ‘85. r 0.82 3 2... , 858: .. 0.2.: cu _ on 111-1‘13? 31131511... 0.000« I .88 3.. 5.....82s 228 3-. 8...: 3. 88. r 0.8.... . 8:! on 1 89°36: 1 0.94.3 00 l ‘ I. 0.000" 3.. 8.59.5 .93 3; 8a.: .9 . _ r 0.3: 0.000n we.“ on.“ vu.o use 84 top scan, Fig. 30(0), shows a diffraction pattern for a HlPed sample after anneal- ing at 850°C for 20 hours. Comparing Figs. 30(a) and 30(0) to Fig. 30(b), an im- portant difference was found. The intensities of three pairs of diffraction peaks at around 32.5 and 32.8; 46.7 and 46.8; 58.2 and 58.3 degrees (26) are re- versed in Fig. 30(b) compared with Figs. 30(a) and 30(0). Based on the struc- ture analysis by Cava el al. [14], diffraction patterns shown in Figs. 30(a) and 30(0) were identified as those from an orthorhombic structure while that shown in Fig 30(b), is from a tetragonal structure. In an orthorhombic phase, the peaks at 32.5 and 32.8 degrees, are in— dexed as 013 and 103/110 reflections respectively and these are indexed as 103/O13 and 110 respectively in the tetragonal phase. In the orthorhombic phase, the intensity of the first peak (013) is lower than that of the 103/110 peak, but in the tetragonal phase, 013 and 103 reflections merge and give a higher intensity. Also the peaks at 46.7 and 46.8 degrees correspond to crystal planes 020/006 and 200 in the orthorhombic phase and the 006 and 200 in the tetrag- onal phase respectively. Based on these x-ray analysis, the x-ray results shown in Fig. 29 and Fig. 30, support the existence of a phase transformation under pressure, suggested by other researchers [73, 101]. Hendrix et al. [73] pointed out the YBagCU3O7-X compound was unstable with respect to pressure and it decomposes into other metal oxides (Y2BaCu05, BaCqu, CuO and others) at stresses greater than 100 MPa, and at temperatures ranging from 700 to 950°C. According to Dijken et al. [102], during HIP densification at 900°C, sintered Y862CU307-X strongly degasses and the oxygen partial pressure increases up to 24.8 MPa. This means that the transformation and decomposition reaction processes of the YBagCu307-x compound prevent full densification and possi- bly cause the glass capsule material to blow out. Therefore, it can be concluded from Fig. 29 and Fig. 30 that HlPing of the 85 YBaZCU3O7-x compound produces unstable conditions and leads to the loss of oxygen and a transformation from the high Tc orthorhombic phase to the non- superconducting tetragonal phase. This result can be explained in terms of stoichiometry of YBagCU3O7-x compound which has the highest TC orthorhom- bic phase for x close to 7.0. When it is heated, the high Tc orthorhombic phase decomposes first to a low TC superconductor for x below about 6.7 to 6.8, and fi- nally to a non-superconducting tetragonal phase for x below about 6.4. 4. 2 Superconducting Properties 4. 2. 1 Critical temperature (Tc) measurements Figure 31 shows the temperature dependance of the electrical resistance of 1-2-3 system with Ag reinforcement, ranging from zero to 15 volume % Ag. The variation of the resistance with temperature of HlPed 1-2—3 composites was similar to those observed for the sintered samples. This means the resistance decreased approximately linearly with decreasing temperature to just above the superconducting onset temperature, and then sharply dropped at the transition temperature. The onset temperature, in resistance, is defined by the smallest curvature of the resistance—temperature curve [103]. HlPed samples with differ- ent volume % Ag have onset temperatures ranging from 90 to 92 K and have zero resistance at temperatures between 88 and 90 K. The transition width, which is defined as the temperature range for 10 - 90 % of transition, is 2 - 4 K. Based on x-ray results shown in Fig. 29 and Fig. 30, HlPed samples, with Ag fiber, consisted of a multiphase mixture; superconducting orthorhombic YBa,2Cu307.x phase, pure Ag and secondary phases of YgBaCuos. Therefore, it can be concluded that as long as the superconducting phase is fully 86 0.00.. 0< 050.250. .0 90:00.0 000.0000 5.2. 00009000 9m; .0 00090.00. 00 0000000000 0.30.0902 00.. .5 .9... C: 0.20.0050... on i 0W0 .0”... 10.9 . 0.9 . 0: 2.: . 0.0 . 0.0 0.202 . “M 03.00.11 . 0< .00 0.200 0< .00 md o; m; 9N (our) aouetsiseu 87 Interconnected, the presence of second phases, suchas Y2BaCu05 and other impurity phases (BaCqu, CuO , etc), have little effect on the sharpness of the superconducting transition. It is to be noted that an improperly controlled oxygen stoichiometry, on the other hand, has a much stronger effect on the transition temperature width. As seen in Fig. 31, an unreinforced HlPed specimen shows a relatively sharp transition behavior and the highest Tc value compared with the reinforced samples. This phenomenon suggests that better interconnectivity between high TC grains can be obtained in the absence of Ag fibers in the specimen. Also the HlPed samples with 15 volume % Ag, display a larger transition temperature width. This broad transition may be indicative of the presence of weak link boundaries or minority phases along grain boundaries. Since Ag fibers are susceptible to oxidation, fiber oxidation may alter the desired oxygen stoi- chiometry of the superconducting phase near the Ag fiber. Therefore, Ag fibers might cause the formation of a low Tc phase or even a non-superconducting te- tragonal phase near the fiber-matrix interface. Figure 31 also demonstrates that the slope of the resistance vs. tempera- ture curves, in the normal state, decreases systematically as the volume % Ag fibers increases. Also resistance at the non-superconducting state, decreases with increasing volume % Ag fibers. However, the transition temperature is ap- proximately the same. Room temperature resistance of a sample with 15 vol- ume % Ag, was found to be 0.3 m0. In contrast, the resistance of an unrein- forced sample with the same geometry and at the same temperature, was 2.9 m0; a resistance almost 10 times higher than that of the 15 % reinforced speci- men. Fig. 32 shows a plot of the normal state resistance at 150 K for samples with different volume % Ag fibers. This is a derived plot from Fig. 31. As can be seen in Fig. 32, resistance decreases linearly with the increase in volume % Ag. 88 1.5 T=150K C} E d) 1.0" O C m If H .9 3’3 0.5- m 0.0 fi I j T r I .— Ififfi '7 o 2 4 6 8 1o 12 14 16 Volume%ong Fig. 32. The electrical resistance of specimens, with different volume % Ag fibers, at 150 K. 89 This can be accounted for by the Ag fiber, having a lower resistance than that of the ceramic 1-2-3 matrix. Resistivity measurements for 1-2-3 samples, processed at two different conditions, one by conventional sintering and the other by HlPing (after post- annealing at 850°C for 20 hours) are shown in Fig. 33. This figure shows that a HlPed sample after post-annealing, has a lower Tc as well as its normal state resistivity is lower than that of a presintered 1-2-3 specimen. A HlPed sample without a post-anneal was found to consist of the non-superconducting tetrago- nal crystalline structure as determined from x-ray analysis (Fig. 30(b)). However, after post-annealing the oxygen stoichiometry is restored, possibly creating a greater degree of order of oxygen in the Cu-O chain. 4. 2. 2 Critical current density (Jc) measurements The voltage-current (V-I) curves for samples with different volume % Ag, are shown in Fig. 34. Table 4 shows summarized data for critical current den- sity (Jc) and the critical temperature. The results for Jc shown in Table 4 have a greater degree of uncertainty due to the resolution of the measuring instru- ments. The voltage drop across the specimen was measured in mV scale rather than in [N scale. Though the results shown in Table 4 are less accurate, the data for JC, measured from HlPed specimens are reproducible enough to make qualitative comparisons. ° Jo values shown in Table 4 demonstrate that the current density for HlPed specimens with 15 vol. % Ag is approximately 4 times higher than that for the conventionally sintered specimens with no Ag. The observed increase in JC with increasing Ag volume %, is difficult to explain. It is possible that the Ag fiber acts as an agent which assists the formation of a denser 1-2-3 network. It 90 00.00000 000 000.1 .0 0 ”0000.008 0.0.00.0 020 .5 0.00.00 0 000 0.00.00 00.90.00 >50 0.30.0005 .w> 3300.000. .8 .90 00000090 000000000 9m; .0. 00 3.00.30.00th 09 0: 000 00 00 ‘ ‘ our L l d 03. 00000000 .000 0-0-. 000.... 30 «.04 00.0.5000 A00 (mo-0001) MIAIISISGH 91 Table 4 Results of transition temperature and critical current density measurements for HlPed and sintered YBaQCU307.x Tc (HlPed) Jc (Sintered) Jc (HlPed) 0% 90001.5 551:5 110010 5% 89.0 :1: 2.0 75 :t 5 126 a: 15 8% 88.5 :t 2.0 70 a: 8 145 a: 15 12% 88.5 :t 2.0 80 :1; 10 200 a; 15 15% 88.0 :1: 2.0 95 a: 15 205 :1: 15 .Unit of Tc : K, .Unit of Jo : Amp/cm? 92 1.25 ‘ 1 ° 0% Ag 1.00 J ° 5% Ag 1 0 8% Ag T = 77 K I o 12% Ag 9 0.75 ' , 15% Ag 5 > 0.50 2 0.25 j . 0.00 - ° . 1 1000 3000 5000 7000 I (mA) Fig. 34. The voltage-current (V-I) curves for specimens with different volume % Ag fibers. 93 is also possible that an oxygen interaction at the Ag-matrix interface might pre- cipitate non-superconducting phase (as discussed earlier) while producing flux pinning. One of the results obtained from the Jc measurements is the voltage-cur- rent (V-I) graphs for samples with different volume % Ag fibers. Beyond some critical current, a phase transformation from a superconducting to a non-super- conducting state occurs. On the non-superconducting state, ohms law (V = IR) is effective in terms of the electrical voltage and electrical current. Therefore, the slope of V-I curve as a function of volume % Ag indicates trend of resis- tance of the non-superconducting state with the variation Ag content. For the unreinforced samples in the non-superconducting state, the slope of V-I curve is steep, that is, the resistance is high. As the Ag volume % increases, the slope of the V-I curve decreases since in normal state the resistance of the Ag rein- forced composite decreases. Fig. 35 shows the slopes of the voltage—current (V-I) curves plotted as a function of volume % Ag. We note that within experi- mental errors, the slopes of the two curves shown in Figures 32 and 35 respec- tively are the same. The only difference is that the absolute value of the resis- tance as measured directly, is larger than that derived from Fig. 35. It can be speculated that the directly measured resistance sees a greater contribution from contact resistance. Also the resistance values in Fig. 35 correspond to the electrical resistance at 77 K as compared with that in Fig. 32 which is at 150 K. The effect of a 211 phase on JC was studied by McCallum et al. [100]. They pointed out the presence of the 211 phase may actually improve proper- ties by producing flux pinning sites if they are distributed properly. However, in this experiment, the amount of 211 phase was very small and the determination of the presence of the 211 phase was quite difficult. 94 0.5 (aV/aI) ' —l 8 1O 12 14 16 I fi 1 0.1 1 10 0 2 4 6 Volume °/o of Ag Fig. 35. The slopes of the voltage-current (V-I) curves in Fig. 34, are plotted as a function of volume % Ag fibers. 95 4. 2. 3 DC magnetic moment measurements The temperature dependance of the DC magnetization for HlPed speci- mens with different volume % Ag is shown in Fig. 36. As can be seen in this fig- ure, the magnetic moment is systematically changed with increasing tempera- ture. There is a strong tendency toward a diamagnetic behavior which be- comes systematically less dominant when the transition width broadens with in- creasing the volume % Ag fibers. HlPed samples with 15 volume % Ag dis- played a wider transition temperature region and a lesser diamagnetic behavior compared to other samples. The decrease of diamagnetic behavior and the broadening of the transition width, with increasing volume % Ag, presumably arise from the decrease of the volume fraction of the superconducting phase in the HlPed specimen with increasing volume % Ag. The magnitude of the Meissner effect is proportional to the volume frac- tion of the superconducting phase in the specimen. If the sample has a large amount of superconducting phase, the Meissner effect will be strong. Electrical resistivity on the other hand dependents on both the amount of the supercon- ducting phase present, as well as the grain boundary properties. Even if a specimen has a large volume fraction of the superconducting phase, it does not show a sharp superconducting transition in the electrical resistivity, if its grain boundary is degraded. Conversely, if the specimen has a low volume % of the superconducting phase, it can show a sharp superconducting transition in the electrical resistivity if the superconducting phases are fully interconnected. To determine the transition temperature from magnetic moment measure- ments, the transition area (10 - 90 % of the transition temperature), in the temperature range from 85 to 95 K, was enlarged. From this magnified area, the superconducting transition temperature was found to be approximately 88 i Magnetic Moment (emu/g) 0.000“ I I I I I I I IIIIICI‘gzv... . g p I I O . - " ' 388 38 3 o o o o o o o o o 0 ° ° : ozoo°§o° o o 0 0.0 .. 9 0 0° -0.020 6 8 o o o 8 ° 0 0 o 8 .1 0 -0.040 o 2 o o o 0%Ag O . o O . SWOAg O 0060 ° °°°°° °8%Ag ° 12%Ag ’ - 15%Ag '0o080 ' T I ' l I l v f , I _r 40 50 60 70 80 90 100 Temperature (K) o I o W ’6: 00010 0 3 . . 8 9 E 00030 0 d) o 0 V o *5 -0.0050 0 o 0 0 E 00070 0 . ° 5 000904 ° .2 5 H “0’ 00110-3. 8 00130 2 - j 1 85.0 87.5 90.0 92.5 95.0 Temperature (K) lFig. 36 The temperature dependance of the DC magnetization for HlPed specimens. 97 3 K, which is consistent with the value obtained from the electrical resistance measurement carried out at various temperatures. Based on the x-ray data, most Ag reinforced specimens have a lower vol- ume fraction of the superconducting phase. Therefore, it can be concluded that the relatively sharp superconducting transition in electrical resistivity is due to the interconnected network of superconducting phases as mentioned earlier. 4. 3 Variation of Density and Porosity of Samples with Different Types of Processing Porosity was measured as a function of the volume % Ag fibers and these results are plotted in Fig. 37. The porosity in the conventionally sintered specimens increased almost linearly with volume % Ag. However, HlPed samples were considerably more dense and did not show any increase of porosity with A9 volume %. After HlPing, the samples had densities between 92 and 96 % of the theoretical density. Table 5 summarizes density and porosity data. 4. 4 Mechanical Properties of Ag-Fiber Reinforced Composites Figure 38 shows the flexural strength of both HlPed and conventionally sintered specimens. Due to the optimization of the microstructure, the HlPed samples show significantly improved flexural strength over that of the conven- tionally sintered samples. The large amount of porosity (20 - 30 %) in the sin- tered specimens accounted for the low strength. Measured flexural strength of unreinforced 1-2-3 after HlPing and sintering, were found to be 56 MPa and 34 MPa respectively. However, as shown in Fig. 38, the flexural strength Porosity (%) 98 30 (a) the sintered (b) HlPed 10‘ i I r + 41. l I f 0 I I I 0 5 10 15 Volume % of Ag Fig. 37. Variation of porosity with volume % Ag fibers (a) conventionally sintered samples; (b) HlPed samples. 20 "v 99 Table 5 Measured and theoretical densities and porosities of YBaZCu307-x reinforced with different volume % Ag-fibers Theoretical density Measured density (g/cm3) (g/Cm3) (a) sintered (b) HlPed Porosity (%) (a) sintered (b)HlPed 0% 5% 8 °/o 1 2% 15% 6.3750 6.5808 6.7043 6.8689 6.9924 5.2593 :1: 0.1393 5.3963 a; 0.1481 5.3634 50 0.1611 5.1516 a: 0.1651 5.1045 a; 0.1660 5.9925 :1: 0.1015 6.1860 :1: 0.1130 6.3691 0 0.1208 6.4224 :1: 0.1215 6.5029 :1: 0.1253 175015 60:16 18.0 a; 2.3 6.0 a; 1.7 20.0 :1: 2.4 5.0 :1; 1.8 250024 65:18 210024 70:18 100 70 ’ 60 (b) HlPed 50 40 307 (a) the sintered Flexural Strength (MPa) 10 . . - 1 . . 0 0 s 10 15 20 Volume % of A9 Fig. 38. A plot of flexural strength vs. volume % Ag (a) conventionally sintered samples; (b) HlPed samples. 101 decreased systematically with increasing volume % Ag. There are several possible explanations for flexural strength decrease with increasing volume °/o Ag: (1) The general theory of the fiber reinforcement suggests that strength- ening will occur if the elastic modulus of the fibers is greater than that of the ma- trix. The elastic modulii of Ag fiber and 1-2-3 matrix were found to be around 71 GPa [104] and 140 GPa [105] respectively. (2) Inhomogeneous mixing of the fibers presumably causes nonuniform distribution of fibers in the fracture plane. This will also reduce the strengthening efficiency. (3) If the interfacial bonding between the matrix and the fibers is weak then strength will decrease with increasing volume % Ag fiber. (4) The diameter (50 pm) of Ag fibers might be relatively large. Fiber di- ameter influences strongly on the strength of the composite; usually the com- posite strength increases as fiber diameter decreases [106]. Also the large di- ameter fibers sometimes act as stress concentration sites which weaken the matrix. (5) If the fiber length is shorter than the critical fiber length (1c), then the matrix deforms around the fiber such that there is virtually no stress transfer and thereby the proportion of fibers capable of being loaded to the fracture stress will be reduced. Another important property of fiber reinforced composites is their tough— ness. A useful way of quantifying toughness of a material is by measuring the fracture surface energy (y) as has been discussed earlier. This is a measure of the resistance of the material to crack propagation and is defined empirically as the minimum amount of energy required to create a unit area of fracture surface. Increasing the fracture surface energy. by inhibiting crack propagation, 1 02 increases the resistance to fracture under static or shock loading. KIC as a function of Ag volume fraction is shown in Fig. 39. As can be seen, KIC increases significantly with increasing volume % Ag. With 15 volume % Ag fibers, the measured KIC value increased up to 4.5 MPa m"2 which is al- most 4 times higher than that for the unreinforced HlPed specimens (1.2 MPa m1l2)_ The nature of load-deflection curves as a function of Ag fiber volume fraction are also shown in Fig. 40. Typical fracture behavior of a 15 % Ag rein- forced HlPed specimen along the load-deflection curve, is shown by inserting typical microstructures associated with the various stages of deformation. The increasing area under load-deflection curves in Fig. 40, indicates a remarkable improvement in the average value of fracture energy (y). From these curves (Fig. 40), it is clear that the unreinforced 1-2-3 sample failed catastrophically, and simultaneously the fracture load decreased immediately to zero. This frac- ture behavior is typical of a brittle ceramic material. In contrast, as seen in the optical micrographs of the sample (Fig. 40), the reinforced specimen fractured in a very different way. The crack deviated from original fracture plane because of the bridging fibers and displayed additional displacement before the ultimate failure. Different load-deflection curves can be explained by the type of energy absorption mechanisms. Some of these mechanisms may be fiber debonding, crack-deflecion, fiber pull-out and fiber plastic deformation. However, in this ex- periment, the energy of fracture is found to be due to a combination of the work needed to debond the fibers, and the work done against friction in pulling the fibers out of the matrix. A summary of the mechanical property data obtained for samples rein- forced with different volume fraction Ag fiber is shown in Table 6. Although Km (MPa m1/2) 103 5 l 4 _ (b) HlPed 3 0 i i 2 J + (a) the sintered .l , I l + .l J + 1 0 . I “r r fi f 0 5 l 0 1 5 2 0 Volume % of Ag .Fig. 39. Km versus volume % Ag (a) conventionally sintered samples; (b) HlPed samples. 50- LOAD (N) 40' 30- 20- 104 3% A9 15% Ag 0% Ag \ / 0.5 1.0 135 2:0 DISPLACEMENT (mm) Fig. 40. Typical load-deflection curves of YBa20U3O7-x samples reinforced with various volume % Ag fibers. Also shown are corresponding typical fractographs. “VI-13‘ 105 Table 6 Summary of measured mechanical properties of 1-2-3 and Ag fiber reinforced and HlP consolidated 1-2-3 compound OF(MPa) KIC(MPa m1’2) (a) sintered (b) HlPed (a) sintered (b) HlPed O % 34.3 :1: 6.0 56.2 :1: 7.3 0.80 :1: 0.3 1.2 :1: 0.3 5 % 25.4 :1: 5.5 60.3 :1: 7.0 0.83 :1: 0.4 1.7 :l: 0.4 8 % 24.9 :1. 5.2 51.3 :1: 6.1 0.91 :1: 0.3 2.5 :1: 0.3 12 % 18.0 :1: 6.3 47.9 :1: 7.2 0.87 :1: 0.3 2.8 :1: 0.4 15 % 21.7 :1: 6.2 50.4 :1: 4.2 1.23 :1: 0.3 4.5 :1: 0.3 106 there is some scatter in the data, the results exhibit a significant improvement in K10 and fracture energy. Flexural strength values, however, are seen to de- crease systematically as a function of volume % Ag. 4. 5 Microstructure and Interfaces 4. 5. 1 Optical microscope analysis Figure 41 shows the microstructure of the as-received Ag fibers. For comparison, microstructures of the unreinforced and Ag fiber reinforced speci- mens, produced by two different processing techniques, are shown in Fig. 42 and Fig. 43 respectively. The conventionally processed specimen shows extensive porosity and large grain size, while the HlP processed specimen have relatively less porosity and a very fine grain size (Figs. 42, 43(a) and 43(b)). Optical microscopy of a HlPed specimen (observed under polarized light) revealed a bimodal grain size; the smaller grains are 1 - 5 um in diameter, and the larger grains are ap- proximately 10 - 15 pm in diameter. The HlPed sample contained a small amount of uniformly distributed porosity. Examination of the microstructure of the HlPed sample indicates an improved inter-grain integrity resulting from the elimination of residual pores. As seen from Fig. 43, the Ag fibers bond well with the 1-2-3 phase by forming a continuous network. This continuous network is believed to be responsible for improved Jc as discussed earlier. 107 Fig. 41. Microstructure of as-received pure Ag fibers (x120). ;' a . " “ti. ‘0 “" 3“: .‘f ' WV“:_z. ‘} s. ;? Fig. 42. Optical microstructures taken from the polished surface of 1-2-3 specimens at different magnifications (x 200 and x 500) (a) sintered, (b) HlPed and annealed 109 (b) Fig. 43(a). Optical microstructures of the sintered specimen reinforced by 15 volume % Ag fiber, taken at different magnifications (a) x 200,. (b) x 500. 110 (b) Fig. 43(b). Optical microstructure of HlPed specimen reinforced by 15 volume % Ag fiber, taken at different magnifications (a) x 200, (b) x 500. 111 4. 5. 2. Scanning Electron Microscopy (SEM) analysis The SEM micrographs in Fig. 44 were taken on the fractured section of a conventionally sintered, and HlPed samples as shown in Fig. 44(a) and 44(b) respectively. As observed by using the optical microscope, the sintered speci- men contains copious amounts of pores and the grain size is large. The HlPed specimen on the other hand shows fewer pores and the grain size is smaller. ' To confirm the presence of the homogeneous distribution of fibers in the matrix, fracture surfaces with increasing volume % Ag fiber were systematically studied using a SEM (Fig. 45). As can be seen in Fig. 45, the area fraction of Ag fiber increases with increasing percentage of Ag fibers added. These figures also indicate a relatively uniform distribution of fibers in the matrix. Fracture surface observations give useful information concerning energy absorption or toughening mechanisms. Some of these mechanisms may be fiber debonding, crack-deflecion, fiber pull-out and fiber plastic deformation. As can be seen in Fig. 46(a), different aspects of fiber pull-out from the matrix can be clearly seen. Also debonding phenomena between the fiber and matrix are easily seen from Fig. 46(b). One significant phenomena, observed in most micrographs, is that the topography of the Ag fibers are altered during HlP processing. As can be seen from Fig. 47, the surface of Ag fiber was visibly faceted by attached 1-2—3 grains. The faceted Ag fiber can not be fully explained by a typical ductile fracture be- havior which usually occurs in Ag after neck formation. To determine surface microhardness, vickers indentation tests were conducted on the surface of high purity annealed Ag plates and HlPed 1-2-3 samples. lndentations were made in air at a load of 100 Kg by using a 1/16 inch diameter ball indenter. This combination gives the Rockell B scale. Calculation of the micro-hardness was 112 ... :I .D\. 4 1...“. . . 1.0.209... . 1 .)x. at different magnifications ( x 500 and x 1,000 ) Fig. 44. SEM micrographs, taken from the fractured surface of HlPed 1-2-3 specimens, (a) sintered, (b) HlPed. 113 Fig. 45(a). SEM micrographs taken from the fractured surface of HlPed 1-2—3 specimens reinforced by different Ag volume % (a) 1-2-3 matrix, (b) 5 % Ag, (c) 8 % Ag. . . . _ 114 .0< .0 0. .0. .3 .x. m. .0. «a 0E:.o> m< 000.0...0 >0 000.250. 000E.0000 0&4 000.... .0 000.50 005.00.. 0... E9. 000. 0000.00.20. 2mm 3.00 .9”. 115 (b) Fig. 46. SEM_ micrographs of fracture surfaces of HlPed and Ag reinforced specrmen showrng: (a) fiber pull-out, and (b) fiber debonding. 116 (b) Fig 47. SEM micrographs showing damaged Ag fiber in HlPed 1 2--3 matrix (a) x 400, (b) x 1,500. 1 1 7 based on an average of 5 indentation measurements. Measured microhard- ness value in HRB scale was 32 for Ag and 68 for 1-2-3, respectively. These re- sults indicate that the hardness of 1-2—3 samples is more than 2 times greater than that of annealed Ag. Thus, It is possible that 1-2-3 phase particles are mechanically embedded on the softer Ag fibers during the pressurization stage of the isostatic pressing. These embedded particles then react with Ag fibers at the sintering temperature to produce a strong adherent layer. Fig. 46 shows in- deed such in the case. Therefore, these coated fibers might not necessarily have the same ductility or strength as the virgin Ag fibers. Another interesting result from both SEM and optical micrographs is the fine grained materials obtained from the HlPing process. Even after prolonged post-annealing at 850°C for 20 hours, the grain size remained almost the same as in the unannealed sample (Fig. 48). This result correlates well with the re- sults by Sadananda et al. [71]. Since the final grain size was attained during the HlPing and not during the annealing process, it can be suggested that grain re- finement occured at the interparticle contact areas due to high localized strain produced by the applied pressure. Since H lP processing is conducted at a lower temperature (890°C) than for conventional sintering (925°C), and the HIP processing time (2.5 hours) is compared with the conventional sintering time (24 hours), these factors combine to produce a smaller grain size. A fine grained materials has higher strength according to the Hall-Petch relationships [107]: Of = 00 + kd'“2 (27) where Of is the fracture strength, d is the average grain diameter and 00 and k are obtained from the intercept and slope respectively of the flow stress versus 118 (b) (a) ured surfaces (x 500) t-annealin 0 , at 850 C for 20 hours. Fig. 48. SEM micrographs taken from the fract 9 y DOS (a) HlPed specimen without an (b) HlPed with post-annealing 119 d“2 plot. Therefore, the fine grain structure, as discussed above, will have a beneficial effort on the mechanical properties (i.e., toughness and tensile strength) of the HlPed superconductor composites. Returning to superconducting properties, it is well known that the grain size and grain boundary chemistry play important roles in affecting the Jc. A sample with a large grain size, favorably oriented and non-equiaxed grains, and with low porosity, can have a higher critical current density. A fine grain size results in a decrease of current transport capabilities because of the in- crease in the number of grain boundaries along a conducting path. HlPing is one of effective ways of reducing the porosity and we know that porosity causes the lack of continuity of the superconductive phase across the grain boundaries in sample. If there is no texture formation between neighboring grain bound- aries, the proportion of high-angle grain boundaries is increased in a fine grained material. In HlPed composite superconductors, the grains are small and randomly oriented, as has been discussed earlier. Thus, .JC values are still relatively low for the HlPed samples. The secondary product, such as, Y2BaCu05, due to the phase decomposition of 1-2-3, is known to improve Jc through flux pinning. However, it was difficult to detect this phase (2—1-1) by using SEM or optical microscopy. Micro—cracking in the specimens is one of the detrimental effects which occurs in reinforced composites. It is important to choose materials which have a similar difference between their thermal expansion coefficients in order to avoid micro—crack formation. The coefficient of thermal expansion of the matrix and the fibers are 16.9 x 10'5/K [108] and 19 x 10'5/K [109], respectively. Thus the fibers contract slightly more than the matrix during cooling. It seems that the fibers do not introduce microcracks into the 1-2—3 based matrix due to the ther- mal expansion mismatch. Because porosity exists between the fiber-matrix 120 interface and the difference in thermal-expansion coefficients is quite low, stresses generated by the thermal-expansion mismatch could be relieved and the residual stresses may be small. The best composites have a differential ex- pansion coefficient of about 1 to 2 x 10'5/K [106], with 01f > 01m (where 01f and 01m are thermal expansion coefficients of fiber and matrix respectively) for the fibers to have residual tensile stress and the matrix to have residual compres- sion stress during heating. For our present case this condition is satisfied. 4. 6 Microchemistry of Bulk and Interface Regions 4. 6. 1 Energy dispersive x-ray analysis Energy dispersive X-ray analysis was conducted for regions of 1-2-3 grains and Ag fiber (Fig. 49). Fig. 50(a) shows the elemental energy dispersive x-ray data (EDAX) for the common elements in the 1-2—3 superconductors. The bottom portion of this figure shows a semi-quantitative (without correction fac- tors) analysis of the EDAX data. Fig. 50(b) and 50(c) show similar results for Ag fiber reinforced regions. Fig. 50(b) is that from a large area of the fracture sur- face and the 50(c) is that from a Ag fiber region. 4. 6. 2 Auger Electron Spectroscopy (AES) analysis After Auger surface scanning, it was found that the sample surfaces were contaminated by carbon. Sample surface were polished with diamond paste and alumina powder dispersed in methanol (Fig. 51). Even after a 5,000 nm layer was ion-sputtered away, the Auger data (Fig. 52) indicated that the carbon content remained almost the same as before. Due to this carbon contamination, 121 1|117|1J .60. x 000 00.. x. 8.00 .80 3.00 .0. . 00 00.“. 0. 00.00005 00 0.0.. 000 .0.00E0.0 x F 8:01.14 00: "'8 C: 11 “1’99: 1 0 ct. S HEH1= NEHI: wINbow START END 01979 GROSS NET EFF. 1990 LABEL keV keV cunus INTEGRAL INTEGRAL FACTOR TOTAL v 1.92 2.96 13 1925 949 1.99 1.93 an 4.36 4.59 12 21327 14655 1.99 29.32 RR 4.72 4.98 14 14082 8487 1.66 '16.46 BA 5.99 5.26 19 5557 2257 1.99 4.36 BA 5.44 5.62 16 3183 973 1.66 1.88 CU 7.96 8.26 16 25689 21661 1.66 46.59 cu 9.79 9.96 15 5132 2777 1.99 5.37 Y 14.84 15.68 13 2264 646 1.66 1.25 Fig. 50(a). Elemental analysis by EDAX of a 1-2-3 superconductor. 123 10.0064 keU 20.2 > FS=511 as: -9 ch 510: 11 cts l'lEl‘ll ‘-' HEHI: . WINDOW START END HIDTH GROSS NET EFF. SAGE ZAGE LABEL kev kev CHANS INTEGRAL INTEGRAL FACTOR TOTAL RATIO Ag 2.88 3.39 22 1292 698 1.99 7.23 292.31 A9 3.72 3.82 6 149 14 1.99 .17 6.73 Ba 3.84 4.98 13 499 56 1.99 .66 26.68 Ba 4.28 5.32 53 5633 3919 1.99 46.52 1889.95 89 5.59 5.64 8 335 47 1.99 .56 22.69 Ba 5.72 5.86- 8 261 13 1.99 .15 6.25 Cu 7.99 8.29 16 3742 3158 1.99 37.57 1518.27 Cu 8.78 9.92 13 684 392 1.99 4.66 188.22 Y 14.82 15.98 14 439 298 1.99 2.47 199.99 Fig. 50(b). Elemental analysis by EDAX on a fractured surface (with overall area scanning). 124 V Preset: 1005 Remaining: is 10% Dead. L (C) E e a a E Y —H AuLLAa—L A 10.000 keU 20.2 > FS= 1K OS= ~16 ch S10= 9% 915 MEMl: HEH1: WINDOW START END HIDTH GROSS NET EFF. SAGE LABEL ch keV CHANS INTEGRAL INTEGRAL FACTOR TOTAL Ag 2.89 3.44 33 15997 12639 1.99 91.69 Ba 4.38 4.56 19 719 149 1.99 1.98 8: 4.78 4.98 11 699 151 1.99 1.19 Cu 7.96 8.18 12 1798 694 1.99 5.93 Cu 8.86 9.99 8 464 56 1.99 .41 Y 14.82 15.96 13 362 96 1.99 .69 Fig.50(c). Elemental analysis by EDAX of a region of Ag fiber. 1255 oomfi _ coofi .m< .\. mE:_o> m 55, 98.259 cméooam 999.1 9 Lo momtzm 9959.06 95 So: :6me 85on mw< .5 .9“. >6 .>ummzm qumsz can I. 1! VI _ —1—- 41‘ —— :nv 30 30 2...... 5.2.2355g - ~ fa. — db 1 Li I fi—I I l 1 I L I h T I If T czoooe.anHa >xeo.ofin>n ‘1— mxo x coo.o .ZHz Hm.o~nmsz coc fl cmmc mmon\H —u— npmmumc _ — .mxo x Ham.mmm nachocm mgcom Louosccooamasm osomm> c_ ac N m fimfiosumn> ”mAHL oi; >m>z=m mmc ofi SJJIP‘II°"S'3¥(3)N 126 1‘1. Ill IIYI. ll 1T 11' H coma _ _ so -9:3.$2;6.5.,§=E.§._.%.qééieégifixé a}... .929th :2 5:9 m< £6 55689 89:1 6 Lo momtnm 995:8 m: diu- >6 .52sz 3525. 82 :0 0 8 p _ O 0829 m 55> 98.859 6 E0: :33 28on mm... .mm .9“. m 2., ‘1' _ 1 «:88 6an 2.28 .3”; q- — — mxo x coed ”53.8 .mxo 6. mmodm ”229E ”:ch 6386:8682...“ :32; 5 ma N m 983$; ”BE .22 NHAHHHEE 8c H 5% mmxmfixfi Qin— E>~5m wmc S SJJIP 'IIWS '3¥(3)N 127 all elemental intensities were reduced. All samples fractured at room temperature showed that the Ag fiber was debonded from matrix during fracturing, and fibers were covered with very fine grains of 1—2—3 phase. Specimens fractured at cryogenic temperatures, were used to study the fiber-matrix interface. Figures. 53 and 54 show the typical Auger electron spectra in terms of the derivative (N(E) * E) and the kinetic energy modes before and after ion sputtering of the fractured surface (Fig. 26). All of the peaks in the spectra can be accounted for by the components of the compound. After ion sputtering, car- bon peaks remained even though the intensity of the carbon peak was some- what reduced. The presence of a small quantity of carbon is most likely due to the contamination of the high vacuum system (please note that this system is also used to study polymeric materials) [110]. Small amounts of boron were also detected. This suggests that there was a possible surface reaction be- tween the boron nitride used during HlPing and the specimen. From the line scan data (Fig. 55) taken on the same fracture surface, the oxygen concentration, which is based on the molar ratio of the 1-2—3 system, was somewhat lower than expected. This might indicate a loss of surface oxy- gen in the high vacuum environment [111]. At the same time this data indicates that oxygen did not diffuse very far into the Ag fibers. The line scanning was carried out from the interface between matrix and Ag fiber well into the Ag fiber. Thus, the result suggests that at the HlPing temperature, the oxygen diffusion rate into Ag is quite low. However, diffusion of Ag into 1-2—3 matrix phase could not be detected conclusively. The line scan seems to show that the interface between the matrix and the Ag fiber is not sharp. One possible explanation is that elements of the matrix phase has diffused into the Ag fiber during process— ing. .95. 6\. oEEo> 55 £55 556955an 99:50me 995.: 9 5c 8959 905599: 95 So: :39 950on mw< .mm .9“. >5 .5525 “5:5sz Sam 85,; 82 8m —1 _ _ _ 5 —u— _. fi— ~r- J _] CD :0 l . «m -r v 555 128 -r :0 5. 555 H5...........Z.S_. {55.55.555.59..__.§...<55.5..52...zs5>.5.555545%}: . 2.5/.35. BJJIP 'IIOUS '3¥(3)N _ . 3 _ —— _. q q— -- .— 555555.5u55 5555.55n55 5\6 5 555.5 “555555 .5\6 5 555.555 ”555555 55555 56565556655555 555555 55 55 N 5 555555555 H55 .255 55.5Nu5555 555 5 5555 N5\55\5 5-5 555555 555 ‘N29 .mccmtzam :0. 56:56 95. 6\6 mE:_o> m £55 68505565 559.5550me 995.1 m 50 momtzm 5595:5095.— m£ E055 559.55 55550on ww< .vm .95 >5 .>Qmmzm lomzwx oom5 ooo5 com _ LT. Ir 5 _ _ . 5 5‘ 5 5 5 —1.— any >. “v 1 . 3 o :0 Lg m< 5: m. -m...5.5.5.5.3525.<_._.5_3,.>6_5\(5..55.3593. 5.335.555.6555:.3..555515s...ff..$$<..5.5._(t5é:54 fig}; >51§3é>> 5 5 L l 4 IA .5 — 51 T v' 5 I 5. T t _ 5 1 5 1 _ q 5555555155 56.55.55.155 56 6. 555.5 ”55555 .56 .5 555.555 4.555555 55555 56565666656655 555655 65 55 N 5 555555655 H5555 .sz em.m5umzH5 auc 5 «mm: mmxm5x5 91¢ >w>m=m mmc o5 GJJIP'II°"S'3¥(3)N 13C) .955 ..\o mE:_o> m £555 55852559 5595550me 55005:... m 50 83:6 85305555 9: 5:95 5595555 .8955 9555555558 ma: 9: 5.55 55550wa wm< .mm .95 555555: -moz¢pmmc 58 555 555 55 55 -- -- . t L L (al.-g .. \\ o -5 5 -- 5. 4 mm .55. .5 .5- -. .5 -- 555 55 -5 \ f 5 I... IT -5 -- 5- -- 15.. m o J] 5- < 55 -- -5 1.. .55. ll 1‘.- -5 55 1.. 1T II 5 Ll _ L 5 5f 5 5 5 4 c:aoee.enHm paeo.e~u>a mxo x cao.e upmm55o .mxo x moe.mm5~ nuc~oc5 m5¢um 55555555555555 555555 55 55 N 5 555555555 55555 .sz 5c.55um=H~ ooc H mzH5 m cum 59u5m memfixfi Qua mzH5 mmc 95 Ilous‘lHOIJH 363d 131 It is known that yttrium has a relatively low sensitivity for Auger transition [97]. Therefore, the excitation of the Auger transition was difficult with a 5 keV beam voltage (Fig. 56). However, when a 15 keV electron beam was applied to increase the Auger sensitivity for yttrium, significant radiation damage occured on the fracture surface. Fig. 57 shows one of the fractured surfaces damaged by electron radiation upon application of 15 keV for 3 hours. Please note the glassy appearance of the surface which strongly suggests melting and vitrifica- tion. Figure 58 shows dot-mapped AES images for different elements (Y, Ba, Cu, O and Ag) with different contrasts. Yttrium element (Fig. 58(a)) was lightly and uniformly scattered on the surface (Fig. 26) while barium (Fig. 58(b))and oxygen elements (Fig. 58(d)) were heavily distributed on the surface. In a dot- maped image, representing barium and oxygen elements, dark background indicates the presence of deep porosity or other elements such as silver. However, in some sections of the fractured surface (in as Fig 26), the copper rich phases were observed as clusters, based on the copper dot-mapped image (Fig. 58(c)). Silver fibers were detected clearly as large bright spots and silver fibers can be seen in several other areas. Only one silver fiber was selected by SEM in Fig. 26, on which dot-mapping was performed. However, as pointed out in discussing the line scan data, interfaces between Ag fibers and matrix are not well defined because some of the 1-2—3 grains (shown as dark image in silver fiber itself) are deposited on or attached to the silver fibers. in this experiment, Auger electrons of each elements were detected around 500 nm from the sur- face. 132 .m< o\o oE:_o> m £5, 88299 5&6me no .. . . of m 0 wow 3 8589.. 9: E0: :38 >995 cozmzoxm >9. m 55, Emcwam wtmw .mm .9“. 83 -— >o .52sz— “:5sz 83 com cup- lT % 3.» s. 2% a a Ind- MF- _ L 17 :0 . mm .- J‘Jiij I_IL_1 $598 an: 3.8 dua— m\o v. 38.9 ”$ch .mxo V. mmmdm ”meat 38m 83389.83 95m? 5 ac N m «@839; ”5H... .22 mm .amnmzfl 8c a «mac NQNE .._\> 53.3 $5 3 SJJ! P 'IIOUS '3¥(3)N 133 Fig. 57. Fracture surface of a HlPed specimen reinforced with 5 volume % Ag which exhibits electron beam radiation damage at 15 keV excitation energy. Rounded projections indicate melting. (b) Fig. 58(a). AES image produced by compositional dot-mapping, taken from the fractured surface of a HlPed specimen reinforced with 5 volume % Ag; (a) mapping of yttrium, (b) mapping of barium. Fig. 58(b). AES image produced by compositional dot-mapping, taken from the fractured surface of a HlPed specimen reinforced with 5 volume % Ag; (0) mapping of copper, (d) mapping of oxygen. 136 (9) Fig. 58(0). AES image produced by compositional dot-mapping, taken from the fractured surface of a HlPed specimen reinforced with 5 volume % Ag; (e) mapping of silver. 5. CONCLUSIONS The objective of this research is to evaluate the feasibility of obtaining dense superconducting composites, which have a critical temperature > 77 K, a reasonable critical current density, and which have a greater fracture tough- ness. After investigating several types of reinforcement materials, Ag fibers were chosen because Ag fibers produced the minimum degradation of super- conducting properties. To meet research objectives, specimens were fabricated by using two different processing techniques; one by conventional sintering and the other by HlPing. In this research, physical, electrical and mechanical properties of Ag fiber reinforced YBa2Cuao7-x superconductor composites have been systematically investigated as a function of volume % fibers and processing techniques. Following are the main conclusions of this research: 1. In the non-superconducting state, the temperature dependance of the electrical resistance decreased systematically as the Ag volume % increased. increasing volume % of Ag, however, did not have any significant negative ef- fects on Tc and JC. in fact, Jc increased by a factor of almost 2 for the reinforced and HlPed samples compared with unreinforced samples. 2. HlPed specimens, with different volume percents of Ag, have a su- perconducting onset temperature range of 90 to 92 K. In all of the cases zero resistance state occurred at approximately 88 to 90 K. The critical temperature data agree well with transition temperature determined from magnetic sus- ceptibility studies. 3. HlPed sample shows a higher densification with approximately 5 % porosity as compared with the conventioally processed samples. Since porosity adversely affects both the electrical and mechanical properties, HlP densification method for processing is recommended. 137 138 4. A significant improvement of fracture toughness of YBa20U3O7-x su- perconductor can be obtained by Ag fiber reinforcement and HIP processing. For example, the value of KIC, (4.5 MPa mm), in a specimen with 15 % Ag- fiber, increased by a factor of 4 over that for an unreinforced specimens, (1.2 MPa m1/2). 5. With the addition of Ag fibers, the fracture energy, yrz, was found to in- crease linearly but the fracture strength,oi:, decreased systematically. 6. Some decomposition of the 1-2-3 phase occurred in the HlPed samples at a temperature of 890°C under 170 MPa. 7. A fine-grained consolidated superconductor composite was obtained after HlPing. It is believed that this fine grain size is a limitation to a dramatic improvement in Jc. Jc could be improved by texturing and/or grain growth. Our initial efforts to produce grain growth in these samples were not successful. It was found that increased HlPing temperature, which might promote grain growth, had an adverse effect on superconducting properties due to phase decompositions. 8. Although Jc is not very large in samples processed in this research, the gain in mechanical stability, retention of a relatively high critical temperature and a modest improvement in Jc will make this processing techniques useful for developing some engineering devices (eg. low current carrying electrical busses, switches etc). APPENDIX 1. Research history Prior to using Ag fibers, 3 different types of reinforcing materials were in- vestigated. First, carbon fibers with a 5 um diameter were consolidated into a 1- 23 matrix. However, the carbon fibers reacted with oxygen, decomposed and did not reinforce the 1-2-3 matrix. Second, copper fibers with a 15 um diameter was introduced into the matrix and processed. The copper fibers similarly oxidized during sintering. Fig. 59 shows the as received Cu fiber and Figs. 60(a) and 60(b), different aspects of the oxidized COpper fiber after heating to 930°C for 3 hours in air. Third, the copper fiber was cladded with about 0.2 mm of silver by using an electroless plating technique [112]. The silver plating material was chosen because the element silver was known to have little or no effect on the super- conducting properties of YBach307-x. This technique produced a considerably thick plating in a short time. This process actually is a controlled reduction pro- cess of a metal ion species by a chemical reducing agent. A general equation for this type of reaction for deposition on a metallic substrate (Ms) is M+2 + reducer __.._.. M0 (on Ms) + reaction products (29) In this experiment, the reducer was 50 g/I of potassium iodide, which re- acted with 500 g/l of silver nitrate. Basic procedures for this technique were as follows. First to remove the oxidized layer, the copper surface was dipped into a dilute nitric acid for few seconds at room temperature and then immediately 139 Fig. 59. Scanning electron micrograph of as received Cu fibers (x 250). Fig. 60(a). SEM micrographs showing the oxidation of Cu fibers at the sintering temperature of 930°C for a sintering in air for 3 hours; (a) x 250, (b) x 1,500. Fig. 60(b). (1)) SEM micrographs showing oxidized Cou fibers in the matrix of 1-2-3 compound after sintering at 930 C for 3 hours in air; (a) x 150, (b) x 600. 143 rinsed with distilled water. The plating solution was made by mixing two chemi- cals; potassium iodide and silver nitrate, at 100°C. Clean copper fibers (as de- scribed above) were immersed in the hot solution. Plating thickness increased with time and with increasing solution temperature. Although, exhibiting good adhesion, the silver deposited was somewhat porous and may not be protected from oxidation and mechanical damage. Fig. 61(a) and Fig. 61(b) show several SEM micrographs of copper fibers after the silver deposition. After cladding with silver, copper fibers were embedded into 1-2-3 matrix for synthesis. The detrimental effect during this process was the Cu-Ag eutectic phase formation at approximately 800°C. The liquid eutectic alloy drained out, leaving behind a sponge and much weakened Cu fiber skeleton. These remaining fibers did not bond well with the matrix and they did not contribute to the mechanical properties of the composite in a positive way. Fig. 62 shows typical SEM micrographs taken at different magnifications after sintering. Fig. 63 is the phase diagram of Ag-Cu system [113]. Since Ag metal did not react, per se, with the 1-2-3 matrix, it was then decided to conduct our experiments with pure Ag fibers. 144 Fig. 61(a). SEM micrographs showing surface morphology of Ag coating on Cu fibers (x 1,200). Fig. 61 (b). SEM micrographs of the cross sectional area of Ag coated Cu fibers (x 500). Fig. 62. Fractured surface of 1-2-3 compound reinforced with Ag coated Cu fibers. Samples were prepared by sintering at 930° C for 24 hours. (a) x 250, (b) x 1,000. 1.4(5 WEIGHT PER CENT COPPER 5 10 15 20 25 30 35 4O 45 $0 60 l l ‘ l TO 80 90 "00 i I l A J_ g l l 1083° ‘ 2”//27 1000 . _ 960.S° I / 900 '. aAo\ / “cu ortAqI I\ \ / I orlCulw 300 f \7 2,79. i 3’ 14.1 39.9 95.1 - (8.8) ms il 192.0l u I t: 700 3 : [I C 2‘: 600 3 Z w 500 1 400 I 300 I >99.94 200 0'35 “121’ (>99.9l O 10 20 30 40 50 60 70 80 90 100 Ag nomc PER cam COPPER Cu Fig. 11. Ag—Cu Fig. 63. Equilibrium phase diagram of Ag-Cu system (Ref. 113). Note that the eutectic temperature is below the sintering temperature. REFERENCES 1. C. W. Chu, P. H. Hor, R. L. Meng, L. Cao and Z. J. Huang, Science, 885, 567 (1987) Y. H. Kao, Y. D. Yao, L. Y. Jang, F. Xu, A. Krol, L W. Song, C. J. Sher, A Darovsky and J. C. Phillips, J. Appl. Phys, 87, 353 (1990) F DesIandes, B. Raveau, P. Dubots and D. Legat, Solid State Commun, H, 407 (1989) G. Xiao, F. H. Streitz, M. Z. Cieplak, A. Bakhshrai, A. Gavrin and C. L. Chien, Phys. Rev., B 38, 776 (1989) V. Plechacek, V. Landa, Z. 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