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MATERIALS RESEARCH

Nanostructured high-temperature superconductors:Creation of strong-pinning columnar defects innanorod/superconductor compositesPeidong Yang and Charles M. LieberDivision of Engineering and Applied Sciences, and Department of Chemistry, Harvard University,Cambridge, Massachusetts 02138

(Received 14 April 1997; accepted 10 June 1997)

A chemical approach to the formation of columnar defects involving the growth andincorporation of MgO nanorods into high temperature superconductors (HTS’s) hasdeveloped. MgO nanorods were incorporated into Bi2Sr2CaCu2Oz, Bi2Sr2Ca2Cu3Oz, andTl2Ba2Ca2Cu3Oz superconductors at areal densities up to2 3 1010ycm2. Microstructuralanalyses of the composites demonstrate that the MgO nanorods create a columnarstructure in the HTS matrices, form a compositionally sharp interface with the matriand self-organize into orientations perpendicular and parallel to the copper oxide plMeasurements of the critical current density demonstrate significant enhancementsMgO nanorod/HTS composites at elevated temperatures and magnetic fields compwith reference samples.

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I. INTRODUCTION

Large critical current densities (Jc ’s) are essentialto many proposed applications of HTS’s, such as wirfor power transmission cables and solenoids.1,2 Conse-quently, enhancingJc in HTS materials has been anremains a topic of great scientific and technologicalterest. In general,Jc is limited by both intrinsic andextrinsic factors.1,3,4 For example,Jc is limited by ther-mally activated flux flow (TAFF); that is,Jc vanisheswell below the upper critical field lineHc2sTd as a re-sult of the motion of magnetic flux lines. This limitationis intrinsic, and arises from the short coherence lengand large anisotropies of the HTS materials that leto weak pinning of flux lines.3–6 In the polycrystallinematerials used for large-scale applications,Jc is also lim-ited by extrinsic factors such as the interfaces betwesuperconducting grains and the development of miccracks during processing.3,7 Poor alignment of crystallinegrains and chemical heterogeneity at their boundarieswell as microcracks, produce weak links with low valueof Jc. Over the past several years, significant progrehas been made in improving the alignment of graiand consequently improvingJc ’s in wires and tapesbased on Ag–Bi2Sr2Can–1CunO2n+4 (n 2, BSCCO-2212; n 3, BSCCO-2223), and YBa2Cu3Oz (YBCO)by identifying how material processing conditions affemicrostructure and transport currents.3,8–10 Despite thisprogress in processing, which is now yielding BSCCObased wires withJc values sufficiently high for somecommercial applications10 and exciting performance inaligned YBCO tapes,9 the intrinsic problem of TAFF

J. Mater. Res., Vol. 12, No. 11, Nov 1997

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remains a limitation to the performance of HTS materiaat high temperatures and magnetic fields.11–13

The traditional method of flux line pinning to reduceTAFF in conventional type-II superconductors, whicinvolves introducing point-like defects into the supeconductor matrix, has been ineffective in the HTS’Theoretical and experimental work has shown, howevthat the problem of TAFF in HTS’s can be reduced significantly by creating correlated defects, such as colunar defects, in the crystal lattice.14–17 The interactionof flux lines with columnar defects results in a largpinning energy that effectively resists thermally activatemotion, and thus increasesJc significantly at high tem-peratures and magnetic fields (Fig. 1).14,18 Significantly,columnar defects improveJc in both BSCCO and YBCOsuperconductors,16,18 and thus have the potential to become a key technological tool in the fabrication of higJc wires. These line-like defects are usually created byradiating samples with heavy ions having energies on torder of a billion electron volts (GeV). This proceduryields amorphous damage tracks 10 to 20 nm in diameand tens of micrometers in length.19 The small defectdiameter is essential to maximize the flux line corepinning interaction without destroying an unnecessavolume fraction of the superconductor. Although theffectiveness of heavy-ion generated columnar defein enhancingJc has been clearly demonstrated,16–18 thesuitability of this technique for large-scale applicationis limited by the short penetration depth of heavy ionand the low thermal stability of the defects.20 Thesefactors eliminate the possibility of producing columna

1997 Materials Research Society 2981

http://www.mrs.org/publications/jmr/comments.html

P. Yang et al.: Nanostructured high-temperature superconductors

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FIG. 1. Schematic diagram illustrating magnetic flux lines in a HTbefore and after the creation of columnar defects. The linear defewhich can be made by heavy ion irradiation, nuclear fission, ananorod composite formation, strongly pin the flux lines in thsuperconductor.

defects in processed wires or of forming wires froirradiated BSCCO powders, respectively. However, Gprotons can be used to create isotropic columnar daage tracks in BSCCO materials through the fissionBi nuclei.21,22 The advantage of this approach is thlarge, 0.5 m penetration depth of the energetic protothat makes possible the production of columnar defein completed wire cables. The fission process dohowever, produce an isotropic distribution of defecthat reduces the pinning efficiency of the columndefects,15 and produces radioactive daughter nuclei thmay make this approach unattractive to the publiclarge. Furthermore, it is unclear whether high-energion-irradiation techniques are economically viable flarge-scale commercial applications.

The large enhancement ofJc produced by columnardefect structures argues that other strategies shouldconsidered for their creation in HTS materials.23 Forexample, the incorporation of rod-like nanostructurinto a HTS matrix should yield the same beneficieffect as ion-generated columnar defects. There ahowever, several critical constraints to achieving a use

2982 J. Mater. Res., Vol. 1

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columnar defect structure in a composite. First, the rolike structures must have diameters comparable todamage tracks or the coherence length of the HTS’s (inanorods) so that they effectively pin the flux lines whiat the same time minimizing the nonsuperconductiphase in the HTS matrix. Second, the nanorods shobe chemically inert in the aggressive metal oxide enronment used to prepare HTS’s; otherwise, the nanorwill not create columnar defects and may have adveeffects on the superconducting properties due toincorporation of impurity ions in HTS lattice. Third, thenanorods should be oriented within the superconduing matrix to maximize the pinning efficiency. Fromthe perspective of size, carbon nanotubes and carbnanorods are materials that could be consideredcreating columnar defects.24,25 Carbon nanotubes arehowever, oxidized to carbon dioxide gas in the oxygemetal oxide environment used to prepare high-qualsamples, and thus are not useful for introducing a hidensity of well-defined columnar defects. Likewise, cabide nanorods25 also react with the HTS matrix andseverely degrade the superconducting properties.

This key issue of chemical reactivity has led uto focus on the growth of nanorods of simple ineoxide materials such as MgO, and their subsequincorporation into HTS’s.26 MgO has been found to bechemically compatible with most HTS’s and is widelused as a substrate for HTS thin films and as a crucimaterial for the growth of HTS single crystals. Prevous studies of micrometer-diameter MgO whiskers hashown that this material can be incorporated into BSCCsuperconductors with little chemical reaction.27,28 Theprincipal effect of these micron size whiskers is aimprovement in mechanical properties;Jc does not ap-pear to be enhanced and in some cases is loweby their addition.28 Presumably the reduction ofJc isdue to a decrease in the cross-sectional superconducarea in the composite, the inhibition of grain growtof the BSCCO phase, and the interference of graalignment in the presence of the large whiskers.contrast, we have recently found that large enhancemein Jc can be achieved by using nanometer size Mgsingle crystal whiskers termed nanorods.26,29,30 Hereinwe describe in detail the growth and structural propeties of MgO nanorods and nanorod composites formwith thick film and bulk BSCCO-2212, BSCCO-2223and Tl2Ba2Ca2Cu3Oz (TBCCO-2223) HTS materials.Magnetization measurements made on these samdemonstrate that the nanorod/HTS composites exhenhancements inJc similar to those produced previ-ously by high-energy ion irradiation. Our approachovercoming the intrinsic limitation of thermally activateflux flow can be implemented in processes currently usto produce large-scale HTS samples and thus mightuseful in applications such as power transmission cab

2, No. 11, Nov 1997


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II. EXPERIMENTAL SECTION

A. Growth of MgO nanorods

The MgO nanorods were prepared using a modifivapor-solid crystal growth process. Briefly, a mixtuof MgO powder (200 mesh, Alfa AESAR) and carbopowder (300 mesh, Johnson Matthey) in a 1 : 3 weigratio was placed at the center of a tube furnace withingraphite boat. The mixture was heated to approximat1200 ±C under a flow of high purity argon to generamagnesium vapor which is then transported to nucleatsites located downstream from the MgOycarbon mixture.The nanorod growth time was typically between 5 a30 min. Two types of nucleation sites have been usedproduce nanorods: (1) nanoscale etch pits/hillocks a(2) MgO nanocrystallites. The nanoscale pits/hillocwere created by etching MgO (001) substrates in 0.5NiCl2 solutions for 1 to 30 min. Alternatively, MgOnanocrystallites were prepared using a previouslyported method31 or depositedin situ. Additional detailsabout the growth and structure of the MgO nanorodsdescribed in Sec. III. A.

B. Preparation of nanorod/HTS composites

To prepare textured thick film BSCCO-2212 composites, pulsed laser deposition (PLD) was first usto deposit 1–5mm BSCCO films onto substrates owhich oriented MgO nanorods had been previousgrown (Fig. 2). The BSCCO-2212 PLD targets wemade by grinding a stoichiometric ratio of Bi2O3, SrCO3,CaCO3, and CuO, compacting this mixture into a pelland then sintering to form a dense target. A Nd:YAlaser (532 nm) was used to ablate this target usingfluence of approximately 1 Jycm2.32 The BSCCO filmswere deposited in 50–200 mTorr of oxygen with sustrate temperatures of 25 and 650±C. Next, the thickfilms produced in this way were partially melted intube furnace at 880±C for 15 min, and then slowlycooled to form textured superconducting thick films.33

Films prepared by deposition at both temperatures exhited similar texture, although microcracks were reducin samples deposited at 650±C.34 Nanorod/BSCCO-2223 composite films were prepared by laser depositof Bi1.8Pb0.3Sr2Ca2Cu3Oz targets at 680±C followedby annealing at 800–810±C in air.35 The as-preparedBSCCO-2212 and 2223 samples were then post-annein flowing Ar at 350±C for 3 h to optimize the superconducting transition temperature,Tc. Reference sampleswithout nanorods were prepared using similar procedufor the purpose of comparison.

Textured Tl-2223 film composites were prepareda two-stage process that involves (1) PLD of 1–2mmthick amorphous Ba2Ca2Cu3Oz thick films at 200±Cfollowed by (2) ex situ thallination.36 During the thal-lination process, the amorphous films were encapsula

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FIG. 2. Schematic diagram illustrating the preparation of a nanorHTS composite starting from a nanorod forest oriented on a substsurface.

2, No. 11, Nov 1997 2983


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with TBCCO-2223 powder in Ag foils and placed ihalf-sealed alumina tubes with Tl2O3. The Tl source tem-perature was kept at 890±C, while the film temperaturewas maintained at 790±C. Reference samples withounanorods were also prepared using similar procedure

In addition, bulk BSCCO-2212 composites havbeen prepared using a two-step process. First, 5–1by weight bulk MgO nanorods were mixed homogeneously with prereacted BSCCO-2212 powdSecond, the nanorod/BSCCO-2212 mixture was partiamelted on silver foil at 880±C, and then slowly cooledto form a textured superconducting tape. Referensamples without nanorods were prepared using simprocedures.

C. Characterization of the nanorods andcomposites

The MgO nanorods and nanorod/HTS compositwere analyzed using a variety of techniques to charterize their microstructure andJc ’s. Scanning electronmicroscopy (JEOL, JSM 6400) was used to study toriented growth of MgO nanorods on MgO single crystsubstrates and the microstructure of as-deposited filtextured films, and bulk samples. Backscattered eltron images were also obtained to map qualitativeelemental distributions on the surfaces of the composiTransmission electron microscopy (Philips, EM420was used to study the growth behavior and size ofMgO nanorods and the local structure, orientation, ainterfaces of the nanorod/HTS composites. In additiospatially resolved energy dispersive x-ray spectrosco(EDAX) was used to determine the local compositioacross the nanorod/HTS interface. The overall textuand crystallinity of the films were characterized usinx-ray (Rigaku, D/Max)u-2u scans and rocking curvemeasurements.

Magnetization measurements were made usingsuperconducting quantum interference device (SQUQuantum Design MPMS1) to determineTc andJc. TheBean model, which relates a measured hysteresismagnetization,DMsH, T d, to the supercurrents flowingwithin the sample, was used to determineJc : Jc 15DMyR, where R is the radius of the sample.37

Although this approach could underemphasize the limtation of grain boundaries onJc, it is appropriate forprobing the intrinsic effects of defect pinning that athe focus of our studies.11,21

III. RESULTS AND DISCUSSION

A. Structure of MgO nanorods

Figure 3 shows cross-sectional SEM micrographsMgO nanorods grown on etched MgO (100) substratMost of the rods are observed to grow normal to tsubstrate, and thus form what we have termed a nano


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forest.26 High densities of nanorods can be grown on tsubstrates by this approach. For example, the nanodensity in Fig. 3(a) is3 3 109ycm2 and that in Fig. 3(b)is nearly a factor of ten larger,2 3 1010ycm2. Thesedensities correspond to effective matching fields of 6and 4000 G, respectively. The true density of pinnisites is, however, likely to be much larger, as discusspreviously.26 Analyses of TEM micrographs recorded onanorods cleaved from substrate surfaces and dispeonto holely carbon grids (e.g., Fig. 4) show that thdiameters of these rods range from 7 to 40 nm withaverage of 15 nm. The lengths of the rods range fro1 to 3 mm. The rods thus have quite high aspect ratioThe nanorods also appear to have smooth surfacesmay be gently curved.

The composition and structure of the nanorods wconfirmed by EDAX and diffraction studies. EDAXdata recorded on nanorods show no evidence for heelements other than Mg. In addition, convergent beelectron diffraction on isolated nanorods [Fig. 4(b)] anx-ray diffraction measurements on bulk samples (Fig.

2, No. 11, Nov 1997


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FIG. 4. (a) TEM micrograph of MgO nanorods cleaved from a MgO substrate. (b) TEM image of a single MgO nanorod. Thek100l growthdirection was determined from convergent beam electron diffraction (inset) recorded with the electron beam perpendicular to theaxis. (c) Growth habits of the MgO nanorods. The initial nucleation site is shown on the left and a view of the square cross seca fractured rod is shown on the right.

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Since MgO has an isotropic, cubic structure, itsomewhat surprising to observe the growth of highanisotropic MgO nanorods. Detailed TEM experimenwere carried out on the nanorods to elucidate possigrowth mechanisms leading to these highly anisotropstructures. The growth axis of the MgO nanorods, whiwas determined using convergent beam electron diffrtion [Fig. 4(b), inset], was found to be along thek001ldirection in all samples studied. In addition, we founthat one end of the nanorods typically terminatedan octahedrally shaped cluster and that the nanorexhibited a zigzag morphology for several hundred nan


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meters past these sites before changing to the smomorphology seen over the major portion of their length[Fig. 4(c)]. Furthermore, tilting experiments showed ththe other ends of the nanorods are square [Fig. 4(inset]. Because the nanorod axis corresponds to thek001ldirection, these square ends can be assigned to (1surfaces. We believe that the octahedrally shaped clusand zigzag morphology correspond to the nucleatisites and initial growth stage, respectively.

The above results suggest that nanorod growth pceeds through a vapor-solid mechanism analogous to tdescribed previously for micron diameter whiskers.38

In this mechanism, the nucleation site size is critical fdefining the rod diameter when the vapor supersaturat

, No. 11, Nov 1997 2985


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is appropriately controlled. The pits and hillocks creaton MgO (100) substrates by etching thus appear to dethe critical nanoscale nucleation sites. The observatof rod-like structures with cubic MgO also requires ththe axial growth rate be greater than the overgrowth raIn MgO, previous studies found that the growtrates increased fromk111l to k110l to k001l; these ratedifferences were suggested to reflect the decreassolid-vapor interfacial free energy of the correspondisurfaces.39 A model that explains qualitatively nanorogrowth in light of these observations is proposedFig. 6. The nanorod growth can be divided into twstages after the initial nucleation. In the first stabeyond nucleation, growth along thek001l directionoccurs with (111) side planes. In the second stathe side planes smooth into (110) or (001) planes dto the smaller free energy associated with the lowindex faces. These growth habits are also obserexperimentally, as shown in Fig. 4(c). The, 55± anglemeasured between facets for this nanorod, which wimaged along the [110] zone axis, clearly indicates th

FIG. 6. Schematic diagram illustrating nucleation and growth of tMgO nanorods. The relevant crystallographic planes are highlight


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the corrugation corresponds to (111) planes at the inistage of growth. Figure 4(b) shows the second stagethe crystal growth in which faces are smooth and parato the growth direction. As discussed below, the growhabits of the MgO nanorods also serve an importafunction in the development of the local microstructuin the nanorod/superconductor composites.

The growth of unsupported MgO nanorods appeato proceed in a similar way with MgO nanocrystalliteserving as the nucleation sites. These nanocrystallitesaggregates of polyhedral particles with kinks, steps, atruncated corners that expose high index crystal plansuch as (111).40 The nanocrystallites thus exhibit featuresimilar to the pits and hillocks created on the Mgsubstrates by etching. Figure 7(a) shows a TEM imaof nanorods nucleated and grown on 5–8 nm diame

FIG. 7. TEM micrographs of unsupported MgO nanorods growusing (a) 5–8 nm and (b) 100 nm diameter nucleation sites.

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MgO nanocrystallites. The diameters of the corresponing nanorods range from 8 to 50 nm with an averaof 25 nm. These samples also exhibit a similar growbehavior (i.e., growth direction and habits) as thogrown on the etched MgO (100) substrates. We haalso used the growth of unsupported MgO nanorodstest the importance of nucleation in determining nanordiameter. Nanorod growth using larger, 100 nm MgOcrystallites for nucleation but otherwise identical reacticonditions produced nanorods with an average diameof 80 nm [Fig. 7(b)]. Hence, it is clear that nucleatiorepresents a critical step in the growth of the Mgnanorods.

B. Microstructure of the nanorod/HTS composites

In order to elucidate the phase formation anchemistry of nanorod/HTS composites, we have invetigated composites prepared using the oriented nanoforests grown on MgO substrates and unsupported Mnanorods. The nanorod forests provide a particulawell-controlled environment for studying HTS composite formation since the positions and orientations of tnanorods are fixed by the substrate.

In general, high-quality nanorod/BSCCO-221composites and reference BSCCO-2212 samples hbeen made through the use of standard melt-texturprocedures20,33 that produce polycrystalline samplewith well-oriented (textured) grains. X-ray diffractionstudies of the composites show that the partial meltincooling cycle leads to crystalline, textured BSCCO-22[Fig. 8(a)] in which the crystallographicc-axis isnormal to the substrate surface. Rocking curve analy[Fig. 8(b)] of the composite and reference samplfurther shows that the nanorods do not adversely affthe development of a well-textured (FWHM, 1±)polycrystalline product. The columnar defect structuin the nanorod/BSCCO-2212 composite prepared o3 3 109ycm2 MgO nanorod forest is observed clearin SEM images (Fig. 9). First, low-magnification sideview images of nanorod/BSCCO composites preparedPLD at low temperature exhibit a columnar morphologthat reflects deposition around the forest of nanoro[Fig. 9(a)]. In contrast, reference samples prepared isimilar manner do not exhibit a columnar morphologSecond, images of melt-textured samples showcharacteristic plate-like morphology of BSCCO anexhibit a density of nanorods protruding from thBSCCO surface (1 3 109ycm2) that is similar tothe one measured on the substrate before compoformation [Fig. 9(b)]. Similar results were also obtainefor BSCCO-2212 composites formed using a muhigher 2 3 1010ycm2 density nanorod forest, as showin Fig. 10. These results show that dense columdefect arrays needed for effectively pinning flux line


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at high fields are accessible using our approach.addition, extended annealing studies at temperatubetween 800 and 850±C suggest that the nanorods dnot react significantly with the BSCCO matrix on a timscale of 3 days (below the liquidus point).

We have also investigated the formationnanorod/HTS composites using TBCCO-2223 aBSCCO-2223 materials. Structural studies of thecomposites show that in both systems the nanorare incorporated into the superconductor crystal grain a manner similar to the BSCCO-2212 system. Fexample, x-ray diffraction measurements made onnanorod/TBCCO-2223 composite (Fig. 11) show thTBCCO is textured with thec-axis perpendicular tothe substrate surface and that the sample is. 95%phase pure. Similar diffraction patterns were recordon composites formed using high-density nanorforests as well. Cross-sectional SEM images of t

, No. 11, Nov 1997 2987


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FIG. 9. (a) Cross-sectional SEM image of a nanorod/BSCCO-22composite before melt-texturing. (b) SEM image of the nanoroBSCCO composite after texturing; the image was recorded a45± angle to the surface normal. The nanorods protruding fromBSCCO basal plane appear light gray.

Ba2Ca2Cu3Oz films show that the nanorod array ifilled with amorphous Ba2Ca2Cu3Oz, and that manynanorods are protruding from the free surface of tmaterial [Fig. 12(a)]. After thallination and texturingthe superconducting composite films have a smoplate-like morphology with a high density of nanorodprotruding from the basala-b plane of crystal grains[Fig. 12(b)]. The observed density of nanorod1 3 109ycm2, corresponds well to the initial density onanorods.

In addition, we have obtained similar results fonanorod/BSCCO-2223 composites. Figure 13 shoSEM images of reference and nanorod compossamples obtained after annealing BSCCO-2223 films815 ±C for 48 h. These images exhibit similar BSCCgrain structure, and importantly, the presence ofhigh-density of nanorods protruding from individuagrains in the composite. The 110 KTc ’s determinedfor the reference and nanorod composite sampsuggest strongly that BSCCO-2223 phase is formhowever, x-ray diffraction measurements do indicate tpresence of a 5–10% impurity phase in the samples.believe that these results are quite promising, althouclearly additional work will be needed to optimize thmicrostructure in our nanorod/BSCCO-2223 composit


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FIG. 11. X-ray diffraction pattern recorded on a textured nanoroTBCCO-2223 composite.

, No. 11, Nov 1997


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Lastly, we have studied the formation of bulnanorod/BSCCO-2212 composites. Composites prepaon silver tapes exhibited features similar to those dscribed above for the film composites. X-ray diffractiomeasurements made on both composite and referesamples exhibited only the (001) reflections of BSCC2212, thus showing a good texture of the BSCCgrains within these samples. More importantly, TEstudies of these samples showed that a high densityMgO nanorods were incorporated into the crystalligrains (Fig. 14). Images recorded with the electrbeam along thec-axis showed roughly circular, darkspots with diameters of 30 nm that contained primarMg. Selected area electron diffraction data recordedthe same time confirm the orientation of the observMgO nanorods and BSCCO crystal grain [Fig. 14(bThese features are similar to those observed forsamples containing substrate-oriented nanorods, and


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suggest that many nanorods preferentially orient alothe BSCCO-2212c-axis in these silver tape sampleThe columnar structure and nearc-axis orientationof the nanorods were further confirmed by samptilting experiments [Fig. 14(c)] and high-resolutiocross-sectional images [Fig. 15(a)]. The high-resolutiimages show clearly the BSCCO layered structu(spacing cy2) with the nanorod perpendicular to thlayers. In addition, these TEM studies demonstrate ta large number of nanorods also orient preferentiallythe BSCCO a-b plane [Fig. 14(d)]. The preferentiaorientation of nanorods is summarized in Fig. 1Because oriented columnar defects are expectedproduce improved values ofJc compared with anisotropic distribution,15 we believe that the observeself-organization of MgO nanorods during composprocessing is significant for applications. The originself-organization is discussed below.

, No. 11, Nov 1997 2989


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FIG. 14. (a) In-plane TEM micrograph of a bulk nanorod/BSCCO-2212 composite. The roughly circular dark spots correspond tnanorods oriented perpendicular to the copper oxide planes. The in-plane density of rods is ca.1 3 1010ycm2. (b) Selected area electrondiffraction recorded on the sample shown in (a). Diffraction peaks corresponding to the in-plane BSCCO and MgO lattices are(c) TEM image of the sample after tilting the crystal off-axis. (d) TEM image showing several nanorods oriented parallel to theoxide planes; these nanorods are highlighted with arrows.

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We have also directly assessed the stability of tMgO nanorods by making spatially resolved EDAmeasurements across the nanorod/BSCCO-2212 inface. EDAX measurements indicate that there is a retively sharp interface between the MgO nanorods andBSCCO-2212 matrix with little interdiffusion and reaction. Figure 15(a) shows a TEM bright-field image oftypical, sharp nanorod/BSCCO interface. Figures 15and 15(c) are typical EDAX spectra recorded acrothe nanorod/BSCCO-2212 interface using a 10 nm besize. No Mg signal can be detected after moving seeral nm beyond the interface into the BSCCO mat[Fig. 15(c)]. Backscattered electron imaging on compoite surfaces further indicates there is little diffusion oMg into the BSCCO-2212 lattice. These results demostrate that the MgO nanorods are chemically compati


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with BSCCO and maintain their structure after incorpration into this matrix.

C. Entrapment and self-organization of nanorodsin the HTS matrix

The observed entrapment of MgO nanorods withBSCCO-2212 crystallites can be understood within tframework of pushing/trapping behavior of foreigparticles at an advancing solid/liquid interface durinsolidification and crystallization. According to theUhlmann–Chalmers–Jackson (UCJ) model,41 the rela-tive magnitude of the particle/solid (sps), particle/liquid(spl), and solid/liquid (ssl) interfacial free energiesdetermines whether a particle is trapped or pushed outhe growth interface. At equilibrium, particle interactio

, No. 11, Nov 1997


c

ioC

Oxy.

d

les1

3itbeds)

erynotialeststal

FIG. 15. (a) High-resolution TEM image obtained from the interfabetween a MgO nanorod and BSCCO-2212 bulk composite sam(b) Typical EDAX spectrum recorded on the MgO nanorod regof the sample. (c) Typical EDAX spectrum recorded on the BSCregion of the sample. The electron beam size was 10 nm in thexperiments.

at the solid/liquid interface is given by

sps spl 1 ssl . (1)

When sps is greater than the sum ofspl and ssl, theparticle is pushed from the interface toward the liqu


eple.nO

ese

id

FIG. 16. Schematic diagram illustrating the self-organization of Mgnanorods in HTS grains. Self-organization is driven by lattice epita

phase. In contrast, whensps is less than the sum ofspl

andssl, the particle is trapped within the growing soli[Figs. 17(a) and 17(b)].

It has also been demonstrated that smaller particare more easily pushed outside the solid in the Y-21yY-123 system42–45where it is likely thatsps . spl 1 ssl.Small Y-211 particles can be still be trapped in Y-12if sufficiently high growth rates are used. Similarly,is expected that very high crystal growth rates wouldrequired to trap nanometer scale particles (i.e., nanoroinside BSCCO grains ifsps . spl 1 ssl. The fact thata large number of nanorods are incorporated under vmodest growth rates suggest that this inequality doeshold for our system. Indeed, the relatively good epitaxrelationships between MgO and BSCCO-2212 suggthat the nanorods can serve as nucleation sites for cry

2, No. 11, Nov 1997 2991


nditions:
FIG. 17. Schematic diagrams showing the interaction of MgO nanorods and the HTS crystal growth front under three hypothetical co(a) repulsion, (b) entrapment, and (c) nucleation and entrapment.
e

g0w

n

yno

(ts

2

n

rted

t an

urport

re.tededisve

s.al-

er

itiesx

tialonles

s-

s,d

r

growth. In this case, we can rewrite Eq. (1) as

sps spl 1 ssl cos u , (2)

whereu is the wetting angle of the liquid, and thus wexpect the MgO nanorods to be trapped regardlesstheir size [Fig. 17(c)].

The preferred orientation or self-organization of thMgO nanorods after melt-texturing is not surprisinwhen we consider that the exposed MgO (110) and (0surfaces are lattice-matched substrates for the groof BSCCO with either itsa- and b-axes or thec-axisoriented parallel to the nanorod axis. Indeed, (110) a(001) epitaxial growth has been widely documentedthe literature.46,47Lastly, we note that the high anisotropof the nanorods could also affect the trapping aorganization within the HTS matrix and will need tbe considered in the future.

D. Critical current densities of nanorod/HTScomposites

We have systematically characterizedJc in thenanorod/HTS composites and HTS reference sampthrough measurements of the sample magnetizationas a function of magnetic field and temperature usingBean model. In general, we see large enhancementJc for composite samples compared with the referensamples. Field-dependentJc data determined at severatemperatures for reference and composite BSCCO-2samples with nanorod densities of3 3 109ycm2 and2 3

1010ycm2 are shown in Fig. 18. The nanorod/BSCCO2212 composites exhibit large increases inJc comparedwith the reference samples, and these increasesespecially significant at higher magnetic fields a


of

e

1)th

din

d

lesM)he

ofcel12

-

ared

temperatures. Because theJc values of our BSCCO-2212reference samples are comparable to the best repofor polycrystalline thick films,17,21 we believe that theincreases observed in the nanorod composites are noartifact of our chosen reference point.

The large increases inJc are attributed to enhancedpinning arising from nanorod columnar defects in ocomposite samples. Several pieces of evidence supthis proposal. First, the increases inJcsH, Td becomemore pronounced with increasing field and temperatuThese results are characteristic of the behavior expecfor pinning by columnar defects and cannot be explainby point-defect pinning, since point defect pinningweak above 30 K. Second, our previous studies hashown that the enhancement inJc exhibits an inversedependence on the average diameter of the nanorod26

These results are consistent with expectations for optimpinning,14 and also show the importance of developing small nanorods for the HTS composites in ordto achieve the optimal pinning efficiency.45 Third, thepresent results demonstrate that higher nanorod densyield largerJc’s at elevated fields, as expected for fluline pinning by columnar defects.

We examined the temperature dependence ofJc forselected magnetic field values to evaluate the potenof these composites for applications. Results obtainednanorod/BSCCO-2212 composite and reference sampat H 0.3 and 0.5 T are shown in Fig. 19; the compoite sample has a nanorod density of2 3 1010ycm2. In thereference samples,Jc fell off rapidly with temperatureas a result of TAFF. In the nanorod/BSCCO compositehowever, this rapid fall-off was dramatically reduced anled to improvements inJc that could exceed one ordeof magnitude. The large increases inJcsH, Td for the

, No. 11, Nov 1997


m

hulf

tte

as

ed.ity

nforof

d/he-ral,

in-

FIG. 18. Comparison of the magnetic field dependence ofJc fornanorod/BSCCO-2212 composites formed with nanorod densities3 3 109ycm2 (≤) and 2 3 1010ycm2 (m), and a BSCCO references.d. The data were recorded at (a) 20, (b) 40, and (c) 60 K. TheTc ’sof all three samples were 80 K.

nanorod/BSCCO-2212 composites can also be sumrized by a plot of the irreversibility line21 in the magneticfield-temperature plane (Fig. 20). From the standpointapplications, the operating field and temperature mustchosen to be below this line because dissipation (tis, loss of superconductivity) occurs above it. In ostudies, the irreversibility point was taken as the fieat which the magnetization hysteresis loop closedeach temperature. The closing criterion was takenbe 5 3 1025 electromagnetic units and was appliedsimilar-sized composite and reference samples inanalysis. Large upward shifts in the irreversibility lin


of

a-

ofbeatrdortoohe

FIG. 19. Comparison of the temperature dependence ofJc in typicalnanorod/BSCCO-2212 compositesmd and references.d samples at(a) 0.3 and (b) 0.5 T. The nanorod density in the composite w2 3 1010ycm2.

for the nanorod/BSCCO-2212 composites are observIndeed, the composite containing the highest dens(2 3 1010ycm2) of nanorods exhibits shifts of 23 and30 K at fields of 0.5 and 1.0 T, respectively. We cathus conclude that the potential operating regimeBSCCO-2212 is extended significantly by formationthe nanorod composite.

Similar results have been obtained on nanoroTBCCO-2223 systems as well. Figure 21 shows tfield dependence ofJc determined at several temperatures for reference and composite samples. In genethe nanorod/TBCCO-2223 composites exhibit large

, No. 11, Nov 1997 2993


e

es

r

t

a

m

uy

of

eegtothe

reS

FIG. 20. Plots of the irreversibility lines for nanorod/BSCCO-221composites with nanorod densities of3 3 109ycm2 (≤) and 2 3

1010ycm2 smd, and a BSCCO references.d sample.

creases inJc compared with the reference sampleand these increases are especially significant atvated fields and temperatures. The explicit tempeture dependencies ofJc obtained on nanorod/TBCCOcomposite and reference samples at 0.5 and 0.8 Tshown in Fig. 22.Jc typically was found to drop morerapidly with temperature in the reference comparto the composite samples. Lastly, the large increain JcsH, Td for the nanorod/TBCCO-2223 compositehave been summarized by a plot of the irreversibililine in Fig. 23. Significantly, we observe large upwashifts in the irreversibility line for the nanorod/TBCCOcomposites: 30 and 35 K at fields of 0.5 and 1.0respectively. These upward shifts in the irreversibililine are comparable to our observations in BSCCO ademonstrate the generality of our approach to differeHTS materials.

IV. CONCLUSIONS

A chemical approach to the formation of columndefects has been developed involving the growthMgO nanorods and subsequent incorporationthese rod-like nanostructures into HTS’s to fornanorod/superconductor composites. MgO nanorohave been successfully incorporated into Bi2Sr2CaCu2-Oz, Bi2Sr2Ca2Cu3Oz, and Tl2Ba2Ca2Cu3Oz supercon-ductors at densities up to2 3 1010ycm2. Microstructuralanalyses of the composites demonstrate thatMgO nanorods (1) create a columnar defect structin the HTS matrices, (2) form a compositionall


2

s,le-

ra-

are

des

styd

T,yndnt

rofof

ds

there

FIG. 21. Comparison of the field dependence ofJc for a typicalnanorod/TBCCO-2223 composite with a nanorod density3 3 109ycm2 sjd and a typical TBCCO referenceshd. The datawere recorded at (a) 50, (b) 60, and (c) 90 K. TheTc ’s of bothsamples were 120 K.

sharp interface with the matrix, and (3) self-organizinto orientations perpendicular and parallel to thHTS copper oxide planes. A model incorporatinentrapment and lattice epitaxy has been usedunderstand these observations. Measurements ofcritical current density as a function of temperatuand field demonstrate that the MgO nanorod/HT

2, No. 11, Nov 1997


r

ta

a

3

y,

d

.

n-d

.. E.

a,

ar,

s.

d

FIG. 22. Comparison of the temperature dependence ofJc in atypical nanorod/TBCCO-2223 composite with a nanorod density3 3 109ycm2 sjd and a TBCCO reference sampleshd at (a) 0.5 and(b) 0.8 T.

composites exhibit significant enhancements at elevatemperatures and magnetic fields compared wreference samples. These enhancements inJc leadto large upward shifts in the irreversibility line fothe composites, and thus show that the operatregime of these HTS materials can be extendsignificantly through the incorporation of nanorodBecause MgO nanorods can be incorporated into HTusing current processing schemes, these results sugthat the nanorod/HTS composites may representechnologically viable strategy for increasing criticcurrents in large-scale applications.

ACKNOWLEDGMENTS

We acknowledge helpful discussions with D. R. Neson, F. Spaepen, and Y. Lu. This work was supportedpart by the Office of Naval Research and the MateriResearch Science and Engineering Center Programthe National Science Foundation.


of

tedith

ingeds.S’sgest

al

l-inlsof

FIG. 23. Plots of the irreversibility lines for nanorod/TBCCO-222composites with nanorod densities of3 3 109ycm2 srd and TBCCOreference samplessed.

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, No. 11, Nov 1997

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Journal of MATERIALS RESEARCH Welcome Comments Helpincorporation of MgO nanorods into high temperature superconductors (HTS’s) has been ... metal oxide environment used to prepare