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Single-crystal growth of Tb0.4Y0.6Ni2B2C compounds by the floating-zone method
Journal of Crystal Growth 213 (2000) 319}327
Single-crystal growth of Tb Y Ni B C compounds by the #oating-zone method H. Bitterlich *, W. LoK ser , G. Behr , G. Graw , W. Yang-Bitterlich, U. KraK mer, L. Schultz Institut fu( r Festko( rper- und Werkstoworschung Dresden, Helmholzstr. 20, Postfach 270016, D-01171 Dresden, Germany Institut fu( r Kristallographie und Festko( rperphysik, Technische Universita( t Dresden, D-01062 Dresden, Germany Received 27 September 1999; accepted 18 February 2000 Communicated by L.F. Schneemeyer
Abstract Single crystals of a Tb Y Ni B C intermetallic compound, which exhibit the coexistence of superconductivity and low-temperature magnetic ordering, have been grown by vertical #oating-zone melting. The growth process of these intermetallic compound crystals is analysed from in situ tracking of the zone temperature, segregation behaviour along the axis, and the constitution of the ultimate zone. The slight shift of composition along the crystal axis, increase in Tb/(Tb#Y)-fraction but decrease in Ni/B ratio, were revealed by electron probe microanalysis. The superconducting transition temperature along the axis co-ordinate is nearly constant at ¹ "4.0$0.1 K with a transition range of *¹ "0.8 K. Stacking faults resulting from a local deviation of the stoichiometric composition were found near the end of the crystal by TEM investigations. 2000 Elsevier Science B.V. All rights reserved. PACS: 74.72.Ny; 81.10.A; 61.72 Keywords: Crystal growth; Intermetallic compounds; Superconductivity; Lattice defects
1. Introduction The class of RETM B C (RE"rare earth or Sc, Y, La; TM"transition metal) borocarbide intermetallic compounds [1}3] has attracted much attention because of relatively high superconducting transition temperatures up to 23 K and the coexistence of superconductivity and magnetic ordering phenomena [4]. Single crystals grown by the Ni B * Corresponding author. Tel.: #49-351-4659/422/451. E-mail address: [email protected] (H. Bitterlich).
#ux method provide a basis for determination of anisotropic physical properties [3,5]. Bulk single crystals of selected RENi B C (RE"Y, Ho, Tb) compounds were recently grown by #oating-zone methods with optical [6,7] and radio frequency (RF) heating [8], respectively. Alloying with both, magnetic and nonmagnetic rare earth elements, in pseudoquaternary Tb Y Ni B C intermetallic V \V compounds can give rise to a gradual shift between magnetic ordering and superconducting properties. This has been revealed by susceptibility and resistivity measurements on polycrystalline samples [9]. However, elucidating of the anisotropy of some
0022-0248/00/$ - see front matter 2000 Elsevier Science B.V. All rights reserved. PII: S 0 0 2 2 - 0 2 4 8 ( 0 0 ) 0 0 3 6 7 - 5
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physical phenomena and details of the magnetic order by neutron di!raction, e.g., require studies on bulk single crystalline specimen. Among the possible compositions Tb Y Ni B C has been selected because of the co-existence of superconductivity and antiferromagnetic order. Therefore, our main objective was the crystal growth of a superconducting pseudoquaternary Tb Y Ni B C intermetallic compound by a computer-aided vertical #oating-zone melting technique with RF heating. Directionally solidi"ed crystals o!er the possibility to investigate the constitution and the physical properties along the growth direction leading to a better understanding of the crystallization process. Therefore, segregation behaviour and homogeneity of physical properties along the crystal axis have been extensively studied in order to establish links between process parameters and features of the crystal perfection which has never been reported before. The lattice defects, arising from di!erent thermal histories and local stoichiometry deviations, were investigated by TEM. The growth process was analysed from in situ tracking of the zone temperature and features of the as-grown crystal making use of relevant parts of phase diagrams, which have recently been determined [8}11] in order to optimise process parameters and crystals perfection.
power, and zone temperature, which was determined by a two-colour pyrometer near the middle of the zone surface, were tracked during the growth process. To investigate the homogeneity of the polycrystalline material small pieces were subjected to a special high-temperature homogenisation treatment up to 14353C within a resistance furnace under puri"ed Ar atmosphere. Composition, microstructure and crystal perfection of the samples were investigated by chemical analysis, optical metallography, scanning electron microscopy and electron probe microanalysis (EPMA). The reduced resistance ratio (R ! ) R )/R has been determined from electrical ) ) measurements down to 20 K. In addition, electrical conductivity measurements and AC-susceptibility measurements have been accomplished in a Hecryostat down to 1.2 K in order to reveal the superconducting properties. The orientation of single crystals was determined by the X-ray Laue backscattering method. Thin foils were prepared by ion milling for transmission electron microscopy (TEM) investigation of lattice defects. The specimens were examined at 200 kV in a Philips CM200 FEG transmission electron microscope.
3. Experimental results 2. Experimental methods Polycrystalline feed rods were prepared from boron, carbon and nickel powder and bulk Y- and Tb-pieces (purity 99.9% or better) which were mixed and pressed to pellets. After melting in a Hukin-type cold crucible rods of 6 mm in diameter and 55 mm in length were cast from the RF levitated melt [12]. Bulk crystals of 6 mm diameter and 40 mm in length were grown by a #oating-zone facility with RF induction heating (250 kHz, 30 kW). The crystals were grown from the o!-stoichiometric initial zone with rod travelling rates of 1.5}10 mm/h (details of the method are given below). Asymmetric counter-rotation of crystal (10/min) and feed rod (20}40/min) was employed. Growth parameters like rod travelling rate, radio frequency, input
3.1. Constitution and growth regime of the (Tb,Y)Ni2 B2 C crystals The quaternary compounds RENi B C (RE"Y, Tb) are formed by peritectic reactions of the melt with properitectic REB C phases [6]. The solidi"cation process can be discussed on the relevant pseudobinary sections of the quaternary phase diagram REB C -RENi B C and RENi B C-RENi B, respectively [8,11]. The Tb Y Ni B C pseudoquaternary compounds V \V are formed by the same peritectic reaction mechanism from a properitectic Tb Y B C phase. VY \VY As reported in a previous paper [9], the onset temperature of the peritectic reaction is gradually reduced with increasing Tb content from ¹ " 15703C for YNi B C (x"0) to ¹ "15183C for TbNi B C (x"1). The primary crystallisation of
H. Bitterlich et al. / Journal of Crystal Growth 213 (2000) 319}327
the Tb Y Ni B C compound is approached below the peritectic temperature ¹ "15563C. Therefore, a travelling solvent method was applied. This was accomplished by placing a piece of Tb Y Ni B onto the polycrystalline slightly hyperstoichiometric Tb Y Ni B C feed rod. This piece was co-melted with an appropriate portion of the feed rod until a molten zone with excess Ni was formed. That is, the nominal composition of the molten zone (c(Ni)/c(B)+1.4) approached the primary solidi"cation range of the Tb Y Ni B C phase. The hyperstoichiometric feed rod com position was chosen to ensure crystallization below ¹ during the whole growth process. Finally the Tb Y Ni B C seed was dipped into the o! stoichiometric initial zone and the crystal growth process was initiated by moving seed and feed rod. In this way the crystallization of the properitectic Tb Y B C phase from the seed can be e!ecVY \VY tively circumvented. The operating point of the growth process, primarily "xed by the average composition of the #oating zone, may drift continuously by any small composition deviation between feed rod and growing crystal. The temperature of the #oating zone can provide some information on the growth process and the actual melt composition. As shown in Fig. 1 the zone temperature starts below the peritectic temperature 15563C, i.e. in the primary crystallization range of the Tb Y Ni B C phase, and slopes down with processing time t, which is represented by the crystal axis co-ordinate z"vt, where v"2 mm/h is the zone travelling rate. The course of the average zone temperature suggests a slight drift and possible temporal #uctuations of the melt composition. The large composition di!erence between the growing crystal and the melt renders the process sensitive to #uctuation of the growth parameters, namely the zone geometry. Moreover, composition gradients in the boundary layer adjacent to the growing interface imply the tendency to morphological instability of the crystal}melt interface. Therefore, the growth process was controlled by small zone travelling rates &2 mm/h and by asymmetric counter-rotation of feed rod and seed rod which serves for mixing of elements within the molten zone.
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Fig. 1. Temperature of the #oating zone during the growth process with rod travelling velocity v"2 mm/h. The axial coordinate of the crystal z"tv is related to the processing time t. (dashed line) Recorded temperature; (solid line) averaged temperature. For comparison: peritectic temperature ¹ "15563C.
3.2. Crystal characterization The metallographic inspection of as-grown rods revealed a central single-crystalline core region with 5 mm in diameter surrounded by a narrow polycrystalline rim [8]. Single-crystalline samples were prepared by cutting o! the polycrystalline parts. By the Laue back scattering from di!erent sections of a cross-sectional plane of the crystal it was revealed that the 10 0 12 c-axis of the tetragonal unit cell of the single-crystalline part of the rod is tilted 133 with respect to the rod axis that means the rod axis is near to the 11 1 22 direction. The macrosegregation along the crystal length co-ordinate z was revealed by EPMA in the WDX mode. In the polycrystalline feed rod and the ultimate zone we must distinguish between the composition of the principal (Tb,Y)Ni B C phase and the averaged composition comprising the minority phases, too. The latter estimates are important for understanding the mass balance during the growth process. The rare-earth element contents were determined with su$cient accuracy, e.g. c(Tb)"6.6$0.3 at%
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Fig. 2. Segregation of the Tb-fraction x along the rod axis z of the as-grown Tb Y Ni B C crystal. (For comparison) quen ched ultimate zone (hatched) and the polycrystalline feed rod. Full symbols: Tb Y Ni B C principal phase. (Open sym bols) averaged composition of the section.
and c(Y)"10.3$0.3 at% at the top of the crystal. As shown in Fig. 2 there is a slight drift of the Tb fraction x"c(Tb)/c(Tb)#c(Y) from x"0.39$ 0.003 at the top towards x"0.415$0.003 at the end of the crystal (z"38 mm). An apparent discontinuity of this trend with a local maximum (x+0.41) near z"14 mm is correlated with a zone temperature anomaly. The Tb fraction of Tb Y Ni B C grains in the polycrystalline feed V \V rod exhibits considerable scatter, but the average x"0.405$0.003 is de"nitely higher than that at the top of the crystal and exceeds the nominal Tb content of the feed rod x +0.40. The Ni content displays a transient behaviour in the initial part of the crystal with an increase from c(Ni)"32.9$0.2 at% to 34.8$0.2 at%. The Ni concentration is descending with axis co-ordinate z throughout the crystal, except for one region near z+10 mm. However, except for the initial part the Ni concentration of the crystal exceeds the rod composition 33.6 at%. Within the ultimate zone and the polycrystalline feed rod the Tb Y V \V
Ni B C phase displays the stoichiometric (33.2$ 0.2 at%) and a hypostoichiometric (32.9$0.2%) Ni-concentration, respectively (Fig. 3a). Besides of the principal Tb Y Ni B C phase the quenched V \V ultimate zone contains Ni-rich phases Tb Y V \V Ni B and Tb Y Ni B. The average Ni-con V \V centration c (Ni)+44.7$0.5 at% and Tb-fraction x +0.44$0.005 of the ultimate zone were estimated. The errors of the EPMA determination of light elements in the Tb Y Ni B C crystal, V \V c(B)"34.2$1.7 at% and c(C)"16.1$1.2 at%, are too large to derive de"nite tendencies along the crystal axis. The ratio c(Ni)/c(B) versus length co-ordinate z, which may a!ect the superconducting properties [9}11], is presented in Fig. 3b. There is a decreasing Ni/B ratio from 0.98 at the top towards 0.92 at the end of the crystal. The average c (Ni)/c (B)"0.92$ 0.04 of the Tb Y Ni B C phase in the polyV \V crystalline feed rod di!ers from the nominal rod composition c (Ni)/c (B)"1.01. The average c (Ni)/ c (B)"1.46$0.07 of the ultimate zone was estimated, which comprises Ni-rich phase fractions, too. The EPMA investigations have con"rmed small composition deviations along the axis coordinate and the vast composition di!erence between crystal and molten zone. 3.3. Lattice defects The one- and two-dimensional lattice defects near the top and the end of the crystal di!er signi"cantly. At the end of the crystal (near the quenched ultimate zone) there are numerous stacking faults and perfect dislocations with Burgers vectors b of the 11 1 02 and 11 0 02 type, respectively (see Fig. 4). In this area the 11 1 02 dislocations are dominant. Near the top of the crystal no stacking faults have been detected. The majority of dislocations exhibit Burgers vectors of 11 0 02 type (see Fig. 5). The Burgers vectors were determined by contrast experiments. There is no or only residual contrast of dislocations with Burgers vector b"[1 0 0] on imaging with di!raction vector u"(0 2 0) and (0 0 4) and of dislocations with b"[1 1 0] at u"(1 1 2) and (0 0 4). The stacking faults are of the same type like in YNi B C specimens [13]. The stacking fault plane
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Fig. 3. Ni-segregation (a), Ni/B ratio (b) and reduced resistance ratio (R !R )/R (c) along the rod axis z of the ) ) ) as-grown Tb Y Ni B C crystal. (For comparison) quen ched ultimate zone (hatched) and the polycrystalline feed rod. (Full symbols) Tb Y Ni B C principal phase. (Open sym bols) averaged composition of the section.
is (0 0 1) and the displacement vector is R+[0 0 1]. The contrast of the stacking faults occurs for imaging with di!raction vector u" (1 1 2) (phase factor a"2puR+p) (see Fig. 5(a)), the contrast disappears with u"(2 0 0) (see Fig. 5(b)). If stacking faults are bounded by partial dislocations those exhibit a Burgers vector of bP +[0 0 1]. 3.4. Physical properties The onset temperature of superconducting transition which has been determined by non-destructive resistivity measurements is nearly constant at ¹ +4.0$0.1 K along the crystal axis and ! reduced to ¹ "3.9 K in the ultimate zone. From ! the susceptibility measurements generally smaller onset temperatures are derived than from resistivity
Fig. 4. TEM image of stacking faults and dislocations near the end of the Tb Y Ni B C crystal with foil orienta tion+[1 1 1]. (a) beam direction+[1 1 1], di!raction vector u"(1 1 2); (b) beam direction+[0 0 1], di!raction vector u"(2 0 0). (A) [1 1 0] dislocations; (B) [0 1 0] dislocation, (P) bounding partial dislocation.
measurements. As shown in Fig. 6 the superconducting onset temperature of polycrystalline samples is decreased by annealing up to 14353C from ¹ "4.2 down to 3.0 K and the transition width is reduced from *¹ +1.8 K to *¹ +0.9 K [11]. ! ! For a single-crystalline sample susceptibility measurement revealed ¹ "3.6 K and a transition width *¹ +0.8 K is superior to the annealed ! polycrystalline state. This is not surprising because as-grown crystal sections are subjected to postsolidi"cation annealing during the slow growth process. It is also interesting that the single crystal
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Fig. 6. Comparison of the superconducting transition in asgrown and annealed samples (inductive method). As-solidi"ed (䊐) and annealed (䊏) polycrystalline rod compared with an as-grown crystal section (/). Inset: detail indicating the antiferromagnetic transition at ¹ "5.8 K. ,
Fig. 5. TEM image of the dislocation structure near the tip of the Tb Y Ni B C crystal. Foil orientation between [0 2 1] and [1 1 1]. (a) beam direction+[0 2 1], di!raction vector u"(1 1 2); (b) beam direction+[2 0 1], di!raction vector u"(0 2 0). (A) [1 0 0] dislocations; (B) [0 1 0] dislocations.
exhibits the susceptibility maximum at ¹ "5.8 K , caused by the antiferromagnetic transition. The smooth peak cannot be detected in polycrystalline samples because of statistical grain orientations. This gives some evidence of single-crystalline Tb Y Ni B C samples for revealing the co existence of superconductivity and magnetic ordering phenomena. The reduced resistance ratio (RRR) is known as one gross measure for crystal perfection, which is sensitive to deviations from the stoichiometric composition of intermetallic compounds. As shown in Fig. 3c there is an increase from RRR+3.5 towards 5.0 along the axis coordinate z. The RRR+5 and 5.8 in both ultimate
zone and the feed rod exceed that of the crystal. After annealing there is only a slight reduction to RRR+5.2 in polycrystalline samples of the feed rod.
4. Discussion and conclusions 4.1. Comparison with existing crystal growth methods Since the discovery of RE}TM-borocarbides the #ux method is well-established in crystal growth of this class of compounds [3]. Single crystals of relatively high perfection manifested e.g. by small *¹ +0.3 K [14] have been grown from Ni B-#uxes up to considerable size of 10 mm; 10 mm;1 mm [5]. One drawback is the relatively small extension of the plate-like single crystals into the c-axis direction due to growth kinetic reasons. Moreover, it is di$cult to control process parameters. In the present investigation the #oatingzone method with RF inductive heating has been
H. Bitterlich et al. / Journal of Crystal Growth 213 (2000) 319}327
employed to pseudoquaternary Tb Y Ni B C compounds, which has the advantage that rodshaped crystals several mm in diameter and with arbitrary axis orientation can be grown by using single-crystalline seeds. Process parameters such as zone travelling rate and zone length can be well controlled. On this basis, process models can be developed which link growth characteristics with crystal perfection. Compared with optical heating [6] vigorous #ow caused by RF inductive heating o!ers the great advantage of convective mixing of the multi-component alloys in the molten zone. The zone temperature is accessible to measurements in contrast to optical heating where light re#ection from the zone surface practically hampers the temperature determination. It turned out that a serious de"ciency of the crystals was a narrow circumferential polycrystalline rim which was not reported for #oating-zone growth with optical heating [6]. This de"ciency seems to be caused by concave melt}crystal interface regions near the outer surface of the molten zone arising from the "nite penetration depth of RF electromagnetic "elds and has been observed for related substances, too [8,15]. By optimization of process parameters, coil geometry and melt #ow conditions, this detrimental e!ect was only partially circumvented so far. Therefore, single-crystalline specimen, e.g. 4 mm;4 mm;8 mm size with near 10 0 12 c-axis orientation of the tetragonal cell, were cut from the central part of 6 mm diameter as-grown rods. The perfection of bulk Tb Y Ni B C crystals, which manifests in a narrow superconducting transition width *¹ +0.8 K is nearly identical with #oating zone grown of quaternary YNi B C compound crystals both, with optical (*¹ +0.7 K [6]) and inductive heating (*¹ +0.8 K [8]). Plate-like YNi B C crystals grown by an im proved #ux method display a superior *¹ + ! 0.3 K [5] compared with the bulk crystals. Di!erent from the #oating-zone method with optical heating of Takeya et al. [6] we have employed a Ni-rich solvent as initial #oating zone. This permits a continuous growth of the seed with given crystallographic orientation into the o!stoichiometric zone and prevents a partial decay by a peritectic reaction into the REB C phase which might induce parasitic grain grow.
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4.2. Ewect of constitution on crystal growth and perfection Because of incongruent melting of the Tb Y Ni B C phase and the o!-stoichiometric zone composition during the growth process the crystal is prone to segregation e!ects, which may be understood in terms of existing equilibrium phase diagrams [8}11]. The local composition of the crystal re#ects the solidus line at the operation point, given by the composition of the molten zone. Although the solidus composition of the Tb Y Ni B C phase is not known precisely, it may be inferred that segregation along the crystal axis means a drift of the operation point due to a slight composition di!erence between solidifying crystal and feed rod. This becomes also apparent from the course of the temperature of the #oating zone. The Tb and Y elements are randomly distributed on the rare-earth lattice sites with unlimited mutually solubility [9]. The Tb segregation in the crystal is attributed to a slope of the melting point of Tb Y Ni B C compounds with increasing V \V Tb content x [9]. Therefore, crystallization starts with a Y-enriched compound and the partition coe$cient k is less than unity. The ratio of Tb 2 concentration in the initial part of the crystal x "0.39 (compare Fig. 2) and the nominal fraction of the initial zone x "0.4 leads to an e!ective partition coe$cient k "0.975. The Tb fraction at 2 the end of the crystal x "0.415 subdivided by x "0.44 of the quenched zone results in k "0.94. 2 A partition coe$cient k (1 is consistent with the 2 gradual increase of the Tb fraction along the axis, and Tb is accumulated in the travelling zone during the growth. However, the Tb fraction in the crystal does not level o! as expected for constant k . 2 A partition coe$cient k "0.76 is estima, ted from the Ni concentration ratio in the crystal and the quenched ultimate zone, respectively. For crystallization of the incongruent melting Tb Y Ni B C compound within the primary solidi"cation range k (1 is mandatory according , to the phase diagram sections [11]. The gradual decrease of Ni content with z is consistent with retrograde solidus concentration throughout the primary solidi"cation range of Tb Y Ni B C. V \V The growth process is sensitive to any parameter
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#uctuations, in particular, the zone geometry. We suppose that local composition changes (at z'7 mm) may originate from #uctuations of the zone length due to sizeable composition di!erences between crystal/feed rod and the travelling zone. Rising zone lengths would e.g. reduce the Ni excess in the molten zone and accordingly decrease the local Ni content in the grown crystal. A shrinking zone accompanied by a temperature drop (through reduced power input), however, would raise the Ni and Tb content in the molten zone. The stacking faults near the end of the crystal are due to local deviations from the stoichiometric composition. They may be caused by introducing additional RE}C- or B}Ni-layers into the crystal lattice. In accordance with results for quaternary LuNi B C, which exhibit a weak tendency of form ing stacking faults [16], a high density of stacking faults was only detected in as-cast specimen which were homogenized only at relatively low temperature (11003C). There were 11 0 02- and 11 1 02dislocations but only sparse or no stacking faults in polycrystalline YNi B C and TbNi B C samples after annealing at 14353C and in single crystals [13]. Some micro-inhomogeneity causing stacking faults in the last part of the Tb Y Ni B C crystal may arise from quenching of the ultimate zone implying enhanced cooling rates of order of 10 K/s between ¹ and 13003C compared with
35 K/h during the steady-state growth process. For the YNi B C and HoNi B C intermetallic compounds a strong e!ect of composition on superconducting and magnetic properties has been reported [11,17]. As reported previously [9] for a series of annealed Tb Y Ni B C polycrystals V \V an increasing fraction of magnetic Tb-ions x not only reduced ¹ but also the reduced resistance ratios from RRR+9.5 (x"0) to RRR+4 (x"0.3). For x'0.3 a nearly constant RRR+4 is approached. Therefore, we attribute the increase of RRR with axis co-ordinate z of the Tb Y Ni B C crystal rather to the decreasing c(Ni)/c(B) ratio (see Fig. 4). Di!erent solidi"cation pathways in multi-component alloys may result in di!erent compositions of intermetallic compounds within their homogeneity range. The Tb Y Ni B C crystals are grown from a Ni-rich #oating zone. Therefore, stoichiometry deviations (Ni
excess e.g.) might explain the lower level (compared with RRR+5.8 of the polycrystalline feed rod) and the increase from RRR+3.5 to 5.0 with z of the crystals grown by the #oating-zone method. We mention that comparison of RRR between di!erent authors may lead to some confusion and misleading conclusions on crystal perfection because of di!erent de"nitions. From the literature data we inferred that e.g. R /R +19 reported for ) ) #ux-grown TbNi B C crystals [18] will be equi valent to (R !R )/R +5.8 in terms ) ) ) of the de"nition employed in the paper which is of the order of the pseudo-quaternary Tb Y Ni B C-compound. The superconducting transition temperature ¹ neither re#ect much scatter along the crystal axis. We suppose that the e!ect of increasing Tb content with axis co-ordinate z, which tends to reduce ¹ [9], is compensated by the decreasing c(Ni)/c(B) ratio which tends to raise ¹ [10]. ¹ is also independent on the density of planar defects, since the average distance between stacking faults 500 nm is large compared with the coherence length of Cooper pairs of the order of 10 nm. We conclude that the present work must be understood as a "rst attempt to analyse physical properties, chemical composition and perfection of intermetallic compound crystals grown by travelling zone method. There is a lack of more accurate chemical analysis methods (especially for the light elements B and C), which are highly desirable in order to get a more quantitative correlation between local element concentration, intrinsic properties of the crystal and growth process features of this class of multi-component intermetallic compounds.
Acknowledgements The authors express their gratitude to M. Doerr, H. Vinzelberg, A. Teresiak, M. FroK mmel, and S. Pichl for useful discussions and experimental assistance and to R. Goldberg for providing the TEM facilities. The "nancial support of the Deutsche Forschungsgemeinschaft within the SFB 463 `Rare Earth-Transition Metal-Intermetallics: Structure, Magnetism and Transporta is gratefully acknowledged.
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References [1] G. Nagarajan, G. Mazumdar, Z. Hossain, S.K. Dhar, K.V. Gopalakrishnan, L.C. Gupta, C. Godart, B.D. Padalia, R. Vijayaghavan, Phys. Rev. Lett. 72 (1994) 274. [2] R.J. Cava, H. Takagi, B. Batlogg, H.W. Zandbergen, J.J. Krajewski, W.F. Peck Jr., R.B. van Dover, R.J. Felder, T. Siegrist, M. Mizuhashi, J.O. Lee, H. Eisaki, S.A. Carter, S. Uchida, Nature 367 (1994) 252. [3] M. Xu, P.C. Can"eld, J.E. Ostenson, D.K. Finnemore, B.K. Cho, Z.R. Wang, D.C. Johnston, Physica C 227 (1994) 321. [4] R.J. Cava, H. Takagi, B. Batlogg, H.W. Zandbergen, J.J. Krajewski, W.F. Peck Jr., R.B. van Dover, R.J. Felder, T. Siegrist, M. Mizuhashi, J.O. Lee, H. Eisaki, S.A. Carter, S. Uchida, Nature 367 (1994) 146. [5] P.C. Can"eld, P.L. Gammel, D.J. Bishop, Phys. Today 51 (10) (1998) 40. [6] H. Takeya, T. Hirano, K. Kadowaki, Physica C 256 (1996) 220. [7] H. Takeya, K. Kadowaki, K. Mirata, T. Mirano, J. Alloys Compounds 245 (1996) 94.
327
[8] G. Behr, W. LoK ser, G. Graw, H. Bitterlich, J. Freudenberger, J. Fink, L. Schultz, J. Crystal Growth 198(199) (1999) 642. [9] H. Bitterlich, W. LoK ser, G. Behr, K. Nenkow, G. Fuchs, A. Belger, L. Schultz, Physica C 308 (1998) 243. [10] H. Szillat, P. Majewski, F. Aldinger, J. Alloys Compounds 261 (1997) 242. [11] G. Behr, W. LoK ser, G. Graw, K. Nenkow, U. KraK mer, A. Belger, B. Wehner, J. Mater. Res. 14 (1999) 16. [12] O. Thomas, J.P. Senateur, R. Modar, O. Laborde, E. Rosender, Solid State Commun. 55 (1985) 629. [13] W. Yang, U. KraK mer, G. Behr, G. Graw, Crystal Res. Technol. 34 (1999) 1213. [14] B.K. Cho, P.C. Can"eld, L.L. Miller, D.K. Finnemore, D.C. Johnston, W.P. Beyermann, A. Yatskar, Phys. Rev. B 52 (1995) 3684. [15] S. Otani, T. Tanaka, J. Crystal Growth 51 (1981) 381. [16] H.W. Zandbergen, R.J. Cava, J.J. Krajevski, W.F. Peck Jr., Physica C 224 (1994) 6. [17] H. Schmidt, M. Weber, H.F. Braun, Physica C 246 (1995) 177. [18] B.K. Cho, P.C. Can"eld, D.C. Johnston, Phys. Rev. B 53 (1996) 8499.
Single-crystal growth of Tb Y Ni B C compounds by the #oating-zone method H. Bitterlich *, W. LoK ser , G. Behr , G. Graw , W. Yang-Bitterlich, U. KraK mer, L. Schultz Institut fu( r Festko( rper- und Werkstoworschung Dresden, Helmholzstr. 20, Postfach 270016, D-01171 Dresden, Germany Institut fu( r Kristallographie und Festko( rperphysik, Technische Universita( t Dresden, D-01062 Dresden, Germany Received 27 September 1999; accepted 18 February 2000 Communicated by L.F. Schneemeyer
Abstract Single crystals of a Tb Y Ni B C intermetallic compound, which exhibit the coexistence of superconductivity and low-temperature magnetic ordering, have been grown by vertical #oating-zone melting. The growth process of these intermetallic compound crystals is analysed from in situ tracking of the zone temperature, segregation behaviour along the axis, and the constitution of the ultimate zone. The slight shift of composition along the crystal axis, increase in Tb/(Tb#Y)-fraction but decrease in Ni/B ratio, were revealed by electron probe microanalysis. The superconducting transition temperature along the axis co-ordinate is nearly constant at ¹ "4.0$0.1 K with a transition range of *¹ "0.8 K. Stacking faults resulting from a local deviation of the stoichiometric composition were found near the end of the crystal by TEM investigations. 2000 Elsevier Science B.V. All rights reserved. PACS: 74.72.Ny; 81.10.A; 61.72 Keywords: Crystal growth; Intermetallic compounds; Superconductivity; Lattice defects
1. Introduction The class of RETM B C (RE"rare earth or Sc, Y, La; TM"transition metal) borocarbide intermetallic compounds [1}3] has attracted much attention because of relatively high superconducting transition temperatures up to 23 K and the coexistence of superconductivity and magnetic ordering phenomena [4]. Single crystals grown by the Ni B * Corresponding author. Tel.: #49-351-4659/422/451. E-mail address: [email protected] (H. Bitterlich).
#ux method provide a basis for determination of anisotropic physical properties [3,5]. Bulk single crystals of selected RENi B C (RE"Y, Ho, Tb) compounds were recently grown by #oating-zone methods with optical [6,7] and radio frequency (RF) heating [8], respectively. Alloying with both, magnetic and nonmagnetic rare earth elements, in pseudoquaternary Tb Y Ni B C intermetallic V \V compounds can give rise to a gradual shift between magnetic ordering and superconducting properties. This has been revealed by susceptibility and resistivity measurements on polycrystalline samples [9]. However, elucidating of the anisotropy of some
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physical phenomena and details of the magnetic order by neutron di!raction, e.g., require studies on bulk single crystalline specimen. Among the possible compositions Tb Y Ni B C has been selected because of the co-existence of superconductivity and antiferromagnetic order. Therefore, our main objective was the crystal growth of a superconducting pseudoquaternary Tb Y Ni B C intermetallic compound by a computer-aided vertical #oating-zone melting technique with RF heating. Directionally solidi"ed crystals o!er the possibility to investigate the constitution and the physical properties along the growth direction leading to a better understanding of the crystallization process. Therefore, segregation behaviour and homogeneity of physical properties along the crystal axis have been extensively studied in order to establish links between process parameters and features of the crystal perfection which has never been reported before. The lattice defects, arising from di!erent thermal histories and local stoichiometry deviations, were investigated by TEM. The growth process was analysed from in situ tracking of the zone temperature and features of the as-grown crystal making use of relevant parts of phase diagrams, which have recently been determined [8}11] in order to optimise process parameters and crystals perfection.
power, and zone temperature, which was determined by a two-colour pyrometer near the middle of the zone surface, were tracked during the growth process. To investigate the homogeneity of the polycrystalline material small pieces were subjected to a special high-temperature homogenisation treatment up to 14353C within a resistance furnace under puri"ed Ar atmosphere. Composition, microstructure and crystal perfection of the samples were investigated by chemical analysis, optical metallography, scanning electron microscopy and electron probe microanalysis (EPMA). The reduced resistance ratio (R ! ) R )/R has been determined from electrical ) ) measurements down to 20 K. In addition, electrical conductivity measurements and AC-susceptibility measurements have been accomplished in a Hecryostat down to 1.2 K in order to reveal the superconducting properties. The orientation of single crystals was determined by the X-ray Laue backscattering method. Thin foils were prepared by ion milling for transmission electron microscopy (TEM) investigation of lattice defects. The specimens were examined at 200 kV in a Philips CM200 FEG transmission electron microscope.
3. Experimental results 2. Experimental methods Polycrystalline feed rods were prepared from boron, carbon and nickel powder and bulk Y- and Tb-pieces (purity 99.9% or better) which were mixed and pressed to pellets. After melting in a Hukin-type cold crucible rods of 6 mm in diameter and 55 mm in length were cast from the RF levitated melt [12]. Bulk crystals of 6 mm diameter and 40 mm in length were grown by a #oating-zone facility with RF induction heating (250 kHz, 30 kW). The crystals were grown from the o!-stoichiometric initial zone with rod travelling rates of 1.5}10 mm/h (details of the method are given below). Asymmetric counter-rotation of crystal (10/min) and feed rod (20}40/min) was employed. Growth parameters like rod travelling rate, radio frequency, input
3.1. Constitution and growth regime of the (Tb,Y)Ni2 B2 C crystals The quaternary compounds RENi B C (RE"Y, Tb) are formed by peritectic reactions of the melt with properitectic REB C phases [6]. The solidi"cation process can be discussed on the relevant pseudobinary sections of the quaternary phase diagram REB C -RENi B C and RENi B C-RENi B, respectively [8,11]. The Tb Y Ni B C pseudoquaternary compounds V \V are formed by the same peritectic reaction mechanism from a properitectic Tb Y B C phase. VY \VY As reported in a previous paper [9], the onset temperature of the peritectic reaction is gradually reduced with increasing Tb content from ¹ " 15703C for YNi B C (x"0) to ¹ "15183C for TbNi B C (x"1). The primary crystallisation of
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the Tb Y Ni B C compound is approached below the peritectic temperature ¹ "15563C. Therefore, a travelling solvent method was applied. This was accomplished by placing a piece of Tb Y Ni B onto the polycrystalline slightly hyperstoichiometric Tb Y Ni B C feed rod. This piece was co-melted with an appropriate portion of the feed rod until a molten zone with excess Ni was formed. That is, the nominal composition of the molten zone (c(Ni)/c(B)+1.4) approached the primary solidi"cation range of the Tb Y Ni B C phase. The hyperstoichiometric feed rod com position was chosen to ensure crystallization below ¹ during the whole growth process. Finally the Tb Y Ni B C seed was dipped into the o! stoichiometric initial zone and the crystal growth process was initiated by moving seed and feed rod. In this way the crystallization of the properitectic Tb Y B C phase from the seed can be e!ecVY \VY tively circumvented. The operating point of the growth process, primarily "xed by the average composition of the #oating zone, may drift continuously by any small composition deviation between feed rod and growing crystal. The temperature of the #oating zone can provide some information on the growth process and the actual melt composition. As shown in Fig. 1 the zone temperature starts below the peritectic temperature 15563C, i.e. in the primary crystallization range of the Tb Y Ni B C phase, and slopes down with processing time t, which is represented by the crystal axis co-ordinate z"vt, where v"2 mm/h is the zone travelling rate. The course of the average zone temperature suggests a slight drift and possible temporal #uctuations of the melt composition. The large composition di!erence between the growing crystal and the melt renders the process sensitive to #uctuation of the growth parameters, namely the zone geometry. Moreover, composition gradients in the boundary layer adjacent to the growing interface imply the tendency to morphological instability of the crystal}melt interface. Therefore, the growth process was controlled by small zone travelling rates &2 mm/h and by asymmetric counter-rotation of feed rod and seed rod which serves for mixing of elements within the molten zone.
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Fig. 1. Temperature of the #oating zone during the growth process with rod travelling velocity v"2 mm/h. The axial coordinate of the crystal z"tv is related to the processing time t. (dashed line) Recorded temperature; (solid line) averaged temperature. For comparison: peritectic temperature ¹ "15563C.
3.2. Crystal characterization The metallographic inspection of as-grown rods revealed a central single-crystalline core region with 5 mm in diameter surrounded by a narrow polycrystalline rim [8]. Single-crystalline samples were prepared by cutting o! the polycrystalline parts. By the Laue back scattering from di!erent sections of a cross-sectional plane of the crystal it was revealed that the 10 0 12 c-axis of the tetragonal unit cell of the single-crystalline part of the rod is tilted 133 with respect to the rod axis that means the rod axis is near to the 11 1 22 direction. The macrosegregation along the crystal length co-ordinate z was revealed by EPMA in the WDX mode. In the polycrystalline feed rod and the ultimate zone we must distinguish between the composition of the principal (Tb,Y)Ni B C phase and the averaged composition comprising the minority phases, too. The latter estimates are important for understanding the mass balance during the growth process. The rare-earth element contents were determined with su$cient accuracy, e.g. c(Tb)"6.6$0.3 at%
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Fig. 2. Segregation of the Tb-fraction x along the rod axis z of the as-grown Tb Y Ni B C crystal. (For comparison) quen ched ultimate zone (hatched) and the polycrystalline feed rod. Full symbols: Tb Y Ni B C principal phase. (Open sym bols) averaged composition of the section.
and c(Y)"10.3$0.3 at% at the top of the crystal. As shown in Fig. 2 there is a slight drift of the Tb fraction x"c(Tb)/c(Tb)#c(Y) from x"0.39$ 0.003 at the top towards x"0.415$0.003 at the end of the crystal (z"38 mm). An apparent discontinuity of this trend with a local maximum (x+0.41) near z"14 mm is correlated with a zone temperature anomaly. The Tb fraction of Tb Y Ni B C grains in the polycrystalline feed V \V rod exhibits considerable scatter, but the average x"0.405$0.003 is de"nitely higher than that at the top of the crystal and exceeds the nominal Tb content of the feed rod x +0.40. The Ni content displays a transient behaviour in the initial part of the crystal with an increase from c(Ni)"32.9$0.2 at% to 34.8$0.2 at%. The Ni concentration is descending with axis co-ordinate z throughout the crystal, except for one region near z+10 mm. However, except for the initial part the Ni concentration of the crystal exceeds the rod composition 33.6 at%. Within the ultimate zone and the polycrystalline feed rod the Tb Y V \V
Ni B C phase displays the stoichiometric (33.2$ 0.2 at%) and a hypostoichiometric (32.9$0.2%) Ni-concentration, respectively (Fig. 3a). Besides of the principal Tb Y Ni B C phase the quenched V \V ultimate zone contains Ni-rich phases Tb Y V \V Ni B and Tb Y Ni B. The average Ni-con V \V centration c (Ni)+44.7$0.5 at% and Tb-fraction x +0.44$0.005 of the ultimate zone were estimated. The errors of the EPMA determination of light elements in the Tb Y Ni B C crystal, V \V c(B)"34.2$1.7 at% and c(C)"16.1$1.2 at%, are too large to derive de"nite tendencies along the crystal axis. The ratio c(Ni)/c(B) versus length co-ordinate z, which may a!ect the superconducting properties [9}11], is presented in Fig. 3b. There is a decreasing Ni/B ratio from 0.98 at the top towards 0.92 at the end of the crystal. The average c (Ni)/c (B)"0.92$ 0.04 of the Tb Y Ni B C phase in the polyV \V crystalline feed rod di!ers from the nominal rod composition c (Ni)/c (B)"1.01. The average c (Ni)/ c (B)"1.46$0.07 of the ultimate zone was estimated, which comprises Ni-rich phase fractions, too. The EPMA investigations have con"rmed small composition deviations along the axis coordinate and the vast composition di!erence between crystal and molten zone. 3.3. Lattice defects The one- and two-dimensional lattice defects near the top and the end of the crystal di!er signi"cantly. At the end of the crystal (near the quenched ultimate zone) there are numerous stacking faults and perfect dislocations with Burgers vectors b of the 11 1 02 and 11 0 02 type, respectively (see Fig. 4). In this area the 11 1 02 dislocations are dominant. Near the top of the crystal no stacking faults have been detected. The majority of dislocations exhibit Burgers vectors of 11 0 02 type (see Fig. 5). The Burgers vectors were determined by contrast experiments. There is no or only residual contrast of dislocations with Burgers vector b"[1 0 0] on imaging with di!raction vector u"(0 2 0) and (0 0 4) and of dislocations with b"[1 1 0] at u"(1 1 2) and (0 0 4). The stacking faults are of the same type like in YNi B C specimens [13]. The stacking fault plane
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Fig. 3. Ni-segregation (a), Ni/B ratio (b) and reduced resistance ratio (R !R )/R (c) along the rod axis z of the ) ) ) as-grown Tb Y Ni B C crystal. (For comparison) quen ched ultimate zone (hatched) and the polycrystalline feed rod. (Full symbols) Tb Y Ni B C principal phase. (Open sym bols) averaged composition of the section.
is (0 0 1) and the displacement vector is R+[0 0 1]. The contrast of the stacking faults occurs for imaging with di!raction vector u" (1 1 2) (phase factor a"2puR+p) (see Fig. 5(a)), the contrast disappears with u"(2 0 0) (see Fig. 5(b)). If stacking faults are bounded by partial dislocations those exhibit a Burgers vector of bP +[0 0 1]. 3.4. Physical properties The onset temperature of superconducting transition which has been determined by non-destructive resistivity measurements is nearly constant at ¹ +4.0$0.1 K along the crystal axis and ! reduced to ¹ "3.9 K in the ultimate zone. From ! the susceptibility measurements generally smaller onset temperatures are derived than from resistivity
Fig. 4. TEM image of stacking faults and dislocations near the end of the Tb Y Ni B C crystal with foil orienta tion+[1 1 1]. (a) beam direction+[1 1 1], di!raction vector u"(1 1 2); (b) beam direction+[0 0 1], di!raction vector u"(2 0 0). (A) [1 1 0] dislocations; (B) [0 1 0] dislocation, (P) bounding partial dislocation.
measurements. As shown in Fig. 6 the superconducting onset temperature of polycrystalline samples is decreased by annealing up to 14353C from ¹ "4.2 down to 3.0 K and the transition width is reduced from *¹ +1.8 K to *¹ +0.9 K [11]. ! ! For a single-crystalline sample susceptibility measurement revealed ¹ "3.6 K and a transition width *¹ +0.8 K is superior to the annealed ! polycrystalline state. This is not surprising because as-grown crystal sections are subjected to postsolidi"cation annealing during the slow growth process. It is also interesting that the single crystal
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Fig. 6. Comparison of the superconducting transition in asgrown and annealed samples (inductive method). As-solidi"ed (䊐) and annealed (䊏) polycrystalline rod compared with an as-grown crystal section (/). Inset: detail indicating the antiferromagnetic transition at ¹ "5.8 K. ,
Fig. 5. TEM image of the dislocation structure near the tip of the Tb Y Ni B C crystal. Foil orientation between [0 2 1] and [1 1 1]. (a) beam direction+[0 2 1], di!raction vector u"(1 1 2); (b) beam direction+[2 0 1], di!raction vector u"(0 2 0). (A) [1 0 0] dislocations; (B) [0 1 0] dislocations.
exhibits the susceptibility maximum at ¹ "5.8 K , caused by the antiferromagnetic transition. The smooth peak cannot be detected in polycrystalline samples because of statistical grain orientations. This gives some evidence of single-crystalline Tb Y Ni B C samples for revealing the co existence of superconductivity and magnetic ordering phenomena. The reduced resistance ratio (RRR) is known as one gross measure for crystal perfection, which is sensitive to deviations from the stoichiometric composition of intermetallic compounds. As shown in Fig. 3c there is an increase from RRR+3.5 towards 5.0 along the axis coordinate z. The RRR+5 and 5.8 in both ultimate
zone and the feed rod exceed that of the crystal. After annealing there is only a slight reduction to RRR+5.2 in polycrystalline samples of the feed rod.
4. Discussion and conclusions 4.1. Comparison with existing crystal growth methods Since the discovery of RE}TM-borocarbides the #ux method is well-established in crystal growth of this class of compounds [3]. Single crystals of relatively high perfection manifested e.g. by small *¹ +0.3 K [14] have been grown from Ni B-#uxes up to considerable size of 10 mm; 10 mm;1 mm [5]. One drawback is the relatively small extension of the plate-like single crystals into the c-axis direction due to growth kinetic reasons. Moreover, it is di$cult to control process parameters. In the present investigation the #oatingzone method with RF inductive heating has been
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employed to pseudoquaternary Tb Y Ni B C compounds, which has the advantage that rodshaped crystals several mm in diameter and with arbitrary axis orientation can be grown by using single-crystalline seeds. Process parameters such as zone travelling rate and zone length can be well controlled. On this basis, process models can be developed which link growth characteristics with crystal perfection. Compared with optical heating [6] vigorous #ow caused by RF inductive heating o!ers the great advantage of convective mixing of the multi-component alloys in the molten zone. The zone temperature is accessible to measurements in contrast to optical heating where light re#ection from the zone surface practically hampers the temperature determination. It turned out that a serious de"ciency of the crystals was a narrow circumferential polycrystalline rim which was not reported for #oating-zone growth with optical heating [6]. This de"ciency seems to be caused by concave melt}crystal interface regions near the outer surface of the molten zone arising from the "nite penetration depth of RF electromagnetic "elds and has been observed for related substances, too [8,15]. By optimization of process parameters, coil geometry and melt #ow conditions, this detrimental e!ect was only partially circumvented so far. Therefore, single-crystalline specimen, e.g. 4 mm;4 mm;8 mm size with near 10 0 12 c-axis orientation of the tetragonal cell, were cut from the central part of 6 mm diameter as-grown rods. The perfection of bulk Tb Y Ni B C crystals, which manifests in a narrow superconducting transition width *¹ +0.8 K is nearly identical with #oating zone grown of quaternary YNi B C compound crystals both, with optical (*¹ +0.7 K [6]) and inductive heating (*¹ +0.8 K [8]). Plate-like YNi B C crystals grown by an im proved #ux method display a superior *¹ + ! 0.3 K [5] compared with the bulk crystals. Di!erent from the #oating-zone method with optical heating of Takeya et al. [6] we have employed a Ni-rich solvent as initial #oating zone. This permits a continuous growth of the seed with given crystallographic orientation into the o!stoichiometric zone and prevents a partial decay by a peritectic reaction into the REB C phase which might induce parasitic grain grow.
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4.2. Ewect of constitution on crystal growth and perfection Because of incongruent melting of the Tb Y Ni B C phase and the o!-stoichiometric zone composition during the growth process the crystal is prone to segregation e!ects, which may be understood in terms of existing equilibrium phase diagrams [8}11]. The local composition of the crystal re#ects the solidus line at the operation point, given by the composition of the molten zone. Although the solidus composition of the Tb Y Ni B C phase is not known precisely, it may be inferred that segregation along the crystal axis means a drift of the operation point due to a slight composition di!erence between solidifying crystal and feed rod. This becomes also apparent from the course of the temperature of the #oating zone. The Tb and Y elements are randomly distributed on the rare-earth lattice sites with unlimited mutually solubility [9]. The Tb segregation in the crystal is attributed to a slope of the melting point of Tb Y Ni B C compounds with increasing V \V Tb content x [9]. Therefore, crystallization starts with a Y-enriched compound and the partition coe$cient k is less than unity. The ratio of Tb 2 concentration in the initial part of the crystal x "0.39 (compare Fig. 2) and the nominal fraction of the initial zone x "0.4 leads to an e!ective partition coe$cient k "0.975. The Tb fraction at 2 the end of the crystal x "0.415 subdivided by x "0.44 of the quenched zone results in k "0.94. 2 A partition coe$cient k (1 is consistent with the 2 gradual increase of the Tb fraction along the axis, and Tb is accumulated in the travelling zone during the growth. However, the Tb fraction in the crystal does not level o! as expected for constant k . 2 A partition coe$cient k "0.76 is estima, ted from the Ni concentration ratio in the crystal and the quenched ultimate zone, respectively. For crystallization of the incongruent melting Tb Y Ni B C compound within the primary solidi"cation range k (1 is mandatory according , to the phase diagram sections [11]. The gradual decrease of Ni content with z is consistent with retrograde solidus concentration throughout the primary solidi"cation range of Tb Y Ni B C. V \V The growth process is sensitive to any parameter
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#uctuations, in particular, the zone geometry. We suppose that local composition changes (at z'7 mm) may originate from #uctuations of the zone length due to sizeable composition di!erences between crystal/feed rod and the travelling zone. Rising zone lengths would e.g. reduce the Ni excess in the molten zone and accordingly decrease the local Ni content in the grown crystal. A shrinking zone accompanied by a temperature drop (through reduced power input), however, would raise the Ni and Tb content in the molten zone. The stacking faults near the end of the crystal are due to local deviations from the stoichiometric composition. They may be caused by introducing additional RE}C- or B}Ni-layers into the crystal lattice. In accordance with results for quaternary LuNi B C, which exhibit a weak tendency of form ing stacking faults [16], a high density of stacking faults was only detected in as-cast specimen which were homogenized only at relatively low temperature (11003C). There were 11 0 02- and 11 1 02dislocations but only sparse or no stacking faults in polycrystalline YNi B C and TbNi B C samples after annealing at 14353C and in single crystals [13]. Some micro-inhomogeneity causing stacking faults in the last part of the Tb Y Ni B C crystal may arise from quenching of the ultimate zone implying enhanced cooling rates of order of 10 K/s between ¹ and 13003C compared with
35 K/h during the steady-state growth process. For the YNi B C and HoNi B C intermetallic compounds a strong e!ect of composition on superconducting and magnetic properties has been reported [11,17]. As reported previously [9] for a series of annealed Tb Y Ni B C polycrystals V \V an increasing fraction of magnetic Tb-ions x not only reduced ¹ but also the reduced resistance ratios from RRR+9.5 (x"0) to RRR+4 (x"0.3). For x'0.3 a nearly constant RRR+4 is approached. Therefore, we attribute the increase of RRR with axis co-ordinate z of the Tb Y Ni B C crystal rather to the decreasing c(Ni)/c(B) ratio (see Fig. 4). Di!erent solidi"cation pathways in multi-component alloys may result in di!erent compositions of intermetallic compounds within their homogeneity range. The Tb Y Ni B C crystals are grown from a Ni-rich #oating zone. Therefore, stoichiometry deviations (Ni
excess e.g.) might explain the lower level (compared with RRR+5.8 of the polycrystalline feed rod) and the increase from RRR+3.5 to 5.0 with z of the crystals grown by the #oating-zone method. We mention that comparison of RRR between di!erent authors may lead to some confusion and misleading conclusions on crystal perfection because of di!erent de"nitions. From the literature data we inferred that e.g. R /R +19 reported for ) ) #ux-grown TbNi B C crystals [18] will be equi valent to (R !R )/R +5.8 in terms ) ) ) of the de"nition employed in the paper which is of the order of the pseudo-quaternary Tb Y Ni B C-compound. The superconducting transition temperature ¹ neither re#ect much scatter along the crystal axis. We suppose that the e!ect of increasing Tb content with axis co-ordinate z, which tends to reduce ¹ [9], is compensated by the decreasing c(Ni)/c(B) ratio which tends to raise ¹ [10]. ¹ is also independent on the density of planar defects, since the average distance between stacking faults 500 nm is large compared with the coherence length of Cooper pairs of the order of 10 nm. We conclude that the present work must be understood as a "rst attempt to analyse physical properties, chemical composition and perfection of intermetallic compound crystals grown by travelling zone method. There is a lack of more accurate chemical analysis methods (especially for the light elements B and C), which are highly desirable in order to get a more quantitative correlation between local element concentration, intrinsic properties of the crystal and growth process features of this class of multi-component intermetallic compounds.
Acknowledgements The authors express their gratitude to M. Doerr, H. Vinzelberg, A. Teresiak, M. FroK mmel, and S. Pichl for useful discussions and experimental assistance and to R. Goldberg for providing the TEM facilities. The "nancial support of the Deutsche Forschungsgemeinschaft within the SFB 463 `Rare Earth-Transition Metal-Intermetallics: Structure, Magnetism and Transporta is gratefully acknowledged.
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References [1] G. Nagarajan, G. Mazumdar, Z. Hossain, S.K. Dhar, K.V. Gopalakrishnan, L.C. Gupta, C. Godart, B.D. Padalia, R. Vijayaghavan, Phys. Rev. Lett. 72 (1994) 274. [2] R.J. Cava, H. Takagi, B. Batlogg, H.W. Zandbergen, J.J. Krajewski, W.F. Peck Jr., R.B. van Dover, R.J. Felder, T. Siegrist, M. Mizuhashi, J.O. Lee, H. Eisaki, S.A. Carter, S. Uchida, Nature 367 (1994) 252. [3] M. Xu, P.C. Can"eld, J.E. Ostenson, D.K. Finnemore, B.K. Cho, Z.R. Wang, D.C. Johnston, Physica C 227 (1994) 321. [4] R.J. Cava, H. Takagi, B. Batlogg, H.W. Zandbergen, J.J. Krajewski, W.F. Peck Jr., R.B. van Dover, R.J. Felder, T. Siegrist, M. Mizuhashi, J.O. Lee, H. Eisaki, S.A. Carter, S. Uchida, Nature 367 (1994) 146. [5] P.C. Can"eld, P.L. Gammel, D.J. Bishop, Phys. Today 51 (10) (1998) 40. [6] H. Takeya, T. Hirano, K. Kadowaki, Physica C 256 (1996) 220. [7] H. Takeya, K. Kadowaki, K. Mirata, T. Mirano, J. Alloys Compounds 245 (1996) 94.
327
[8] G. Behr, W. LoK ser, G. Graw, H. Bitterlich, J. Freudenberger, J. Fink, L. Schultz, J. Crystal Growth 198(199) (1999) 642. [9] H. Bitterlich, W. LoK ser, G. Behr, K. Nenkow, G. Fuchs, A. Belger, L. Schultz, Physica C 308 (1998) 243. [10] H. Szillat, P. Majewski, F. Aldinger, J. Alloys Compounds 261 (1997) 242. [11] G. Behr, W. LoK ser, G. Graw, K. Nenkow, U. KraK mer, A. Belger, B. Wehner, J. Mater. Res. 14 (1999) 16. [12] O. Thomas, J.P. Senateur, R. Modar, O. Laborde, E. Rosender, Solid State Commun. 55 (1985) 629. [13] W. Yang, U. KraK mer, G. Behr, G. Graw, Crystal Res. Technol. 34 (1999) 1213. [14] B.K. Cho, P.C. Can"eld, L.L. Miller, D.K. Finnemore, D.C. Johnston, W.P. Beyermann, A. Yatskar, Phys. Rev. B 52 (1995) 3684. [15] S. Otani, T. Tanaka, J. Crystal Growth 51 (1981) 381. [16] H.W. Zandbergen, R.J. Cava, J.J. Krajevski, W.F. Peck Jr., Physica C 224 (1994) 6. [17] H. Schmidt, M. Weber, H.F. Braun, Physica C 246 (1995) 177. [18] B.K. Cho, P.C. Can"eld, D.C. Johnston, Phys. Rev. B 53 (1996) 8499.