Crystal growth of rare earth-transition metal borocarbides and silicides

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Journal of Crystal Growth 310 (2008) 2268–2276 www.elsevier.com/locate/jcrysgro

Crystal growth of rare earth-transition metal borocarbides and silicides Gu¨nter Behr, Wolfgang Lo¨ser, Dmitri Souptel, Gu¨nter Fuchs, Irina Mazilu, Chongde Cao, Anke Ko¨hler, Ludwig Schultz, Bernd Bu¨chner Crystal Growth Laboratory, Leibniz Institute for Solid State and Materials Research Dresden, IFW, Helmholtzstr. 20, D-01069 Dresden, Germany Available online 15 December 2007

Abstract Large high quality crystals of various classes of novel rare earth-transition metal intermetallic compounds like RNi2B2C, R11xR2xNi2B2C, R2PdSi3 and RPd2Si2 (R, R1, R2 ¼ Y, La, Sc and rare earth elements) were grown by floating zone (FZ) methods. Careful selection and handling of high-purity starting materials and the control of oxygen impurities during the whole preparation process is required. The process parameters crucially depend on the melting temperature and the solidification mode of the compound. After growth heat treatments can improve the real structure of as grown crystals. The high perfection of single crystals enabled the discovery of some novel physical and crystallographic features of these compounds. r 2008 Elsevier B.V. All rights reserved. PACS: 81.10.Fq; 81.10.Fb; 74.70.Dd Keywords: A2. Floating zone technique; A2. Single crystal growth; B1. Rare earth compounds

1. Introduction Novel rare earth-transition metal (T) intermetallic compounds exhibit a number of interesting physical properties. The R–T-borocarbides of the type RT2B2C discovered in 1994 (R ¼ rare earth elements, Sc, Y, or La; T ¼ Ni, Pd) [1–3] have attracted much attention because of superconducting transition temperatures as high as Tc=23 K and the interplay of superconductivity and magnetic ordering [4]. The RNi2B2C intermetallic compounds exhibit the tetragonal body-centered ThCr2Si2-type crystallographic structure. Soon after the discovery of this family of compounds, single crystals grown by Ni2B-flux method provided a basis for measurements of anisotropic physical properties [5,6]. But, despite the great progress of the method, they suffer from limited size and imperfections arising from the flux. As an alternative, the vertical floating zone (FZ) method was utilized to grow RT2B2C crystals of bigger dimensions, typically 6–8 mm in diameter and some centimeters in length [7–13]. Because of the high melting temperatures and the extreme reactivity of the melts with oxygen and practically all Corresponding author.

E-mail address: [email protected] (G. Behr). 0022-0248/$ - see front matter r 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.jcrysgro.2007.11.227

known crucible materials, the crucible-less FZ technique, the issue of the present review is the favored growth method. Among R–T-silicides the classes of compounds R2TSi3 with hexagonal AlB2-type crystallographic structure [14,15] and RT2Si2 with body-centered tetragonal ThCr2Si2-type crystallographic structure [16,17] are of particular interest. They exhibit a variety of interesting physical properties such as complex magnetic ordering, magnetoresistance anomalies, heavy fermion and Kondo behavior. Single crystals both of R2PdSi3 [11,18–24] and CeT2Si2 compounds [25–27] were grown by the tetra-arc Czochralski method [18,19] and preferably by FZ method. The aim of this review is a survey of crystal growth attempts of RNi2B2C, R2PdSi3 and RT2Si2 intermetallic compounds by vertical FZ methods. Challenges for feed rod preparation, principle characteristics of the FZ growth facilities, and optimum process parameters in the context of relevant phase diagram features are assessed with respect to growth of high-perfection crystals. 2. FZ crystal growth equipment The crystal growth was accomplished with various vertical FZ methods both with radio frequency (RF)

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induction heating (250 kHz or 3 MHz) and radiation heating (see Ref. [28] for details). The FZ growth with optical radiation heating was promoted because the convex melt-crystal interface facilitated the grain selection process. Most crystal growth attempts were performed in a laboratory type FZ apparatus URN-2-ZM (MPEI, Moscow) with a vertical double ellipsoid optical configuration and a 5-kW air-cooled xenon lamp positioned near the focal point of the lower mirror. The molten zone is located near the focal point of the upper mirror. A mechanical shutter, consisting of four moving horizontal sectors, made of stainless steel sheets, controls the light incident onto the molten zone within a range of 0–100% of the maximum radiation flux governed by the lamp current. The lamp power can be varied in a range from 50% to 100% of the nominal one. The original gas supply system was strongly modified for the crystal growth requirements of intermetallic compounds. A vacuum pump (down to 104 Pa), a Ti-getter system (CENTOR) at the gas inlet, and an oxygen trace measuring system (ZIROX) at the gas outlet circuit were installed to control the gas purity in the growth chamber. The ambient gas was 5 N argon at 0.1–5 MPa pressure and 6–10 l/h gas flow. After the purification the flowing argon in the growth chamber contained o0.01 vol.ppm O2. The crystals were grown with slow rod traveling rates between 1 and 10 mm/h and asymmetric counter-rotation of crystal (20–50 rpm) and feed rod (10 rpm). Details about the optical scheme of the facility were given in Ref. [29]. Alternatively, vertical FZ melting with optical heating was conducted using a four-mirror-type image furnace with four 1.5 kW-halogen lamps (CSI, Japan), which is wide spread in the laboratory practice. The crystal growth proceeds in a cylindrical quartz chamber under flowing 5 N argon ambient gas at 0.1–0.95 MPa pressure. Different from the facilities described above in that case the heating unit is moved along the rod. The growth parameters coincide with that described for URN-2-ZM. Normally, both FZ techniques with optical heating turned out to be equivalent in most practical examples of crystal growth of intermetallic compounds. However, the vertical optical configuration, although rarely employed, provides specific advantages: (i) high efficiency of the radiation flux focusing and extremely uniform azimuthal heating of the molten zone which enables melting of refractory intermetallics and oxides up to 2800 1C and (ii) a narrow bundle of illuminating rays with large incident angle on the zone [29]. This allows to use shorter quartz tubes (typically 60 mm with wall thickness up to 14 mm) for the growth chamber which are very important for application of high gas pressures (up to 10 MPa). Other advantages are the easy access to the crystal growth chamber, even during the growth process, and sufficient space for mounting of auxiliary functional components. One recent highlight was the installation of an in situ pyrometric measurement of FZ temperature using a stroboscopic method, which permits better control of the

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Fig. 1. Sketch of the floating zone apparatus with optical heating in a vertical double-mirror configuration.

growth process [30]. Normally, pyrometric temperature measurements of the zone in FZ methods with optical heating are not viable because of the inherent intense light flux. Presently, a new FZ furnace with vertical optical configuration is at work (Fig. 1), with superior performance of the mechanical and electrical components for crystal growth at elevated pressures up to 15 MPa. 3. Feed rod alloy preparation routes The use of high-purity starting materials, control of the impurity level, stoichiometry and homogeneity throughout all technological steps were crucial aspects of the feed preparation. As starting materials, metal pieces of rare earth elements with 99.9 wt% (or better) purity from different suppliers (Ames Lab, Hunan Institute of Rare Earth Metal Materials, Alfa or Goodfellow), Ni granules (99.999 wt%) or Ni powder (99.99 wt%, MaTeck), Pd lumps or shot (99.999 wt%, Heraeus), the 11B isotope (99.52 wt%, Eagle Picher), C (graphite) (99.9+ wt%, MaTeck) and Si (99.999 wt%) were utilized. For neutron diffraction studies, the crystals were grown with 11B isotope instead of natural boron to prevent the pronounced neutron absorption of 10B. The oxygen content of rare

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earth elements has been independently checked by a carrier gas hot extraction method. Typically, rare earth elements witho0.2 at% oxygen have been selected and their surfaces were mechanically ground in Ar atmosphere to remove oxide traces. Details of the influence of oxygen are given in Ref. [31]. The reactivity of elements and high melting temperatures are challenging for feed rod preparation. Therefore, crucible-less or cold crucible melting techniques are preferred. For the preparation of R–Ni-borocarbides, bulk rare earth elements and powder blends of Ni, 11B and C were weight in a desired ratio, mixed and pressed into tablets in a stainless steel die. The RNi2B2C alloys were synthesized by RF induction melting of the pressed pellets under 5 N Ar atmosphere in a Hukin-type copper cold-crucible facility. This soft melting process permits a complete reaction of the elements to form the quaternary intermetallic RNi2B2C compound and minimizes selective element evaporation. The received buttons were several times remelted in an arc-melting furnace on a water-cooled copper plate in a Zr-gettered Ar atmosphere. Polycrystalline feed rods, 6.3 mm in diameter and 60–85 mm in length, were finally cast in copper moulds utilizing Hukin-type cold-crucible facility and homogenized by annealing at temperatures up to 1400 1C. This sophisticated preparation technique results in homogeneous feed rods with controlled composition and reduced the weight loss to 0.1–0.2 wt%. Feed rods of R–T-silicides were prepared in a two-step melting process, which considerably reduced evaporation losses. First, the bulky constituents (Si, T ¼ Pd, Cu, Ni) were arc melted to form binary T–Si alloys. Then, the buttons of the binary alloy were co-melted with appropriate portions of the rare earths in Hukin-type copper cold-crucible and cast into 6.3 mm diameter feed rods.

Fig. 2. (a) Phase diagram section TbB2C2–TbNi2B2C–TbNi4B relevant for TbNi2B2C crystallization processes [28]. Superposed is a DTA heating plot. (b) Metallographic section of a stoichiometric LuNi2B2C sample illustrating the properitectic phase LuB2C2 and peritectic crystallization of LuNi2B2C on the surface of long needles of the second properitectic phase LuNiBC.

4. Main features of crystal growth processes 4.1. Phase diagram features and process parameters The process parameters crucially depend on the melting temperature and solidification mode of the desired compound. Because phase diagram data of multi-component alloy systems were scarce, we have performed own studies. It turned out that the RNi2B2C compounds studied so far (R ¼ Y, Tb, Dy, Ho, Er, Tm) formed via a peritectic reaction RNi2B2C2RB2C2+L (Fig. 2a) [8,9]. This implies that for a stoichiometric feed rod the crystallization starts with a primary phase RB2C2. The single-phase growth of RNi2B2C compounds can only be achieved in the traveling solvent FZ mode, i.e. with an off-stoichiometric FZ composition. For this purpose a small disk (of a height corresponding to FZ length) with an off-stoichiometric composition within the primary RNi2B2C solidification field is placed between seed and feed rod and molten along with the adjacent parts of these rods. This serves for RNi2B2C crystallization from the beginning. Alternatively,

one can start with a stoichiometric zone composition. In this case the precipitation of RB2C2 normally leads to a continuous shift of the element concentration of the FZ until the primary solidification range of RNi2B2C is approached. This simple method does not work for the superconducting LuNi2B2C compound where two properitectic phases, LuB2C2 and LuNiBC, compete (Fig. 2b) and the primary solidification range of LuNi2B2C is far from its stoichiometric composition. Accordingly, FZ crystal growth of large crystals was not realized so far although small LuNi2B2C crystals were already grown from the Ni2B flux [6]. Because of the finite composition difference at the crystal/melt interface the FZ growth of RNi2B2C crystals in the traveling solvent regime proceeds with comparably slow growth velocities 1–2 mm h1. For FZ growth of LuNi2B2C crystals, constitutional supercooling was still observed for 0.8 mm h1 because of a large composition difference between the crystal and the traveling zone. But 0.4 mm h1 looks more promising.

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transparency of quartz tubes and accordingly the zone stability in case of optical heating [31].

Er50-xPdxSi50

1900

Temperature [ °C]

1800

4.2. Factors determining grain selection

1700 1600

ErSi + L

L + Er 2PdSi3 Er2PdSi3 + L

1500 ErSi + Er2PdSi3

1400

Er2PdSi3

Er2PdSi3 + ErPd2Si2

1300 0 ErSi

2271

5

10

15 Pd [at.%]

20

25

30

Fig. 3. Er–Pd–Si isoplethal phase diagram section at 50 at% Si relevant for Er2PdSi3 crystal growth [31].

By contrast, congruent melting behavior at temperatures 41500 1C was revealed for several ternary silicides R2PdSi3 (R ¼ Ce, Gd, Tb, Dy, Ho, Er and Tm). A typical Er–Pd–Si phase diagram section is shown in Fig. 3 [32]. These compounds can directly grow from the stoichiometric melt and a composition difference in FZ only occurs if the actual crystal composition differs from that of the maximum melt temperature (dystectic point). This permits higher growth velocities of 3–10 mm h1. A similar congruent melting behavior of ErPd2Si2 at Tm ¼ 1420 1C was recently found, which is representative for the class of RPd2Si2 compounds. However, different from the R2PdSi3 compounds ErPd2dSi2+d exhibits a relatively extended homogeneity range of dp0.3 which can give rise to a crucial dependence of physical properties on growth conditions [32]. One fortunate effect inherent to all compounds studied is the formation of volatile oxides SiO, BO, CO of nonmetallic components in the melt at elevated temperatures 41500 1C. This is well known as silicothermic reaction (1) of metal oxides with Si [33,34]: MOx þ x  Si2M þ x  SiO " ðgasÞ;

(1)

where M is a metal. Borothermic and carbothermic reactions proceed in a similar way [9,13]. During the FZ growth, SiO evaporates from the molten zone and may condensate at the cold parts of the growth chamber. In that way, SiO is removed from ambient Ar atmosphere and the thermodynamic equilibrium is shifted toward the righthand side of Eq. (1). A steady SiO flow from the melt is induced even in the case of thermodynamically more stable rare earth oxides. This effect may lead to a self-refining of the crystal during FZ growth with respect to oxygen. Drawbacks of these effect are a slight loss of Si during the growth process which must be balanced by an excess of Si in the feed rod composition and the deposition of oxides at the walls of the growth chamber which deteriorate the

For grain selection from the feed rod in the initial stage of FZ process or maintenance of the single crystalline state after seeding, the following main factors have been elaborated: (i) the crystal/melt interface shape, (ii) impurities and (iii) foreign phases. The interface shape is basically governed by temperature distribution in the zone and determined by heating method. A convex interface shape (toward the melt) is desired, which promotes the outward growth of the grains in the starting phase and prevents the penetration of parasitic grains from the surface toward the center. If FZ growth with RF induction heating is applied, the growing interface is concave at least in a narrow range near the rod surface. This is due to the finite penetration depth of the RF fields into the melt, but the effect is further enhanced by electromagnetically driven melt flow directed from the hot surface toward the central part of the interface. Accordingly, the as-grown crystals often displayed an inner single crystalline core region covered by a narrow polycrystalline rim. This detrimental effect was more pronounced for slow growth velocities necessary for RNi2B2C. The polycrystalline rim caused other defects like thermal stresses and microcracks within the crystal. For optical heating the crystal–melt interface normally is convex over the whole cross-section and misoriented grains are avoided. However, the zone length is another decisive factor. Shorter zones enabled a convex growth front. If a critical length is exceeded a concave growth front emerges and multiple grains can grow from the surface toward the center. Another preference of FZ growth with radiation heating is the high stability of the FZ leading to a smoother crystal surface than for RF induction heating. The control of the oxygen impurity content of the melt turned out to be another key problem for growing R–T-compound crystals. As visualized in Fig. 4, the grain selection process in the initial state of Er2PdSi3 FZ crystal growth proceeds smoothly for polycrystalline feed rods (left) prepared from high-grade Er (0.01 wt% O). Already after 1.5 h growth time with 3 mm/h, the cross-section exhibits only a few columnar grains with straight grain boundaries and relatively small mutual inclination determined by X-ray Laue back scattering. Finally, one single grain covers the whole cross-section. Its /1 0 0S-axis is tilted 121 against rod axis. For feed rods prepared from low-grade Er (0.1 wt% O), an effective grain selection during the growth was suppressed by numerous small Er2O3 particles, which can pin the grain boundaries [23]. The purification of the ambient gas in the growth chamber is of similar importance as the purity of the feed rod. Secondary phases can often encounter the crystal growth process of multi-component intermetallics due to phase diagram peculiarities. Incongruent melting of RNi2B2C

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1

Inclination against crystal axis

4 2 Er2 PdSi3

3

1 mm

1 2 3 44

c

a

82.1° 88.5° 81.7° 79,1° 79.1°

10.9° 2.2° 18.0° 11,5° 11.5°

Growth direction

Fig. 4. Optical images in polarized light visualizing the grain selection process in the initial stage of FZ crystal growth of Er2PdSi3 from a polycrystalline feed rod (left) prepared from high-grade Er (0.01 wt% O). (a) Longitudinal section parallel to rod axis and (b) cross-section after about 1.5 h growth with 3 mm/h. For the grains (1–4) the orientation was determined by X-ray Laue back scattering.

compounds implies that for unseeded FZ growth from stoichiometric RNi2B2C feed rods crystallization starts with a properitectic RB2C2 phase. In the case of TbNi2B2C FZ growth (similar to other compounds of this class), a dense band of small TbB2C2 grains is formed before TbNi2B2C phase crystallizes. Further on, periodic striations of TbB2C2 grains arise within TbNi2B2C matrix until the melt composition of the traveling zone approaches the primary crystallization field of TbNi2B2C. Surprisingly, the grain selection of TbNi2B2C principle phase is virtually unaffected by small TbB2C2 second-phase particles. Typically, after 20 mm distance, a single grain with a preferred orientation close to c-axis of the tetragonal unit cell becomes dominant in the entire cross-section [13]. The off-stoichiometric composition of the traveling zone for RNi2B2C crystal growth implies severe element partition at melt–crystal interface and a tendency toward cellular instability. This is often manifested by the presence of Ni-rich phases (RNi4B, R2Ni3B6, R2Ni6B, Ni2B, Ni3B), which can form eutectics with RNi2B2C principle phase [13]. They preferably occur in cell grooves in the final part of the crystal near the ultimate zone and deteriorate the crystal perfection. Therefore, the growth process is controlled by a small pulling velocityo2 mm/h and a high stability of the heating power. Asymmetric counterrotation of feed rod and crystal supports composition homogenization by convective flow within the molten zone.

evaporation from the molten zone. This is a very unique opportunity of the vertical double-ellipsoid optical configuration utilized. Evaporation during growth is still substantial because the melting temperatures of the compounds 1733 1C (Tm2PdSi3), 1422 1C (Eu2CuSi3), 1255 1C (EuCu2Si2) far exceed those of copper and the constituent rare earths, respectively. All three compounds melt congruently, which facilitates crystal growth at reasonable high velocities 3 mm/h and with short zone lengths. As an example, the growth process of EuCu2Si2 from a Eu20.5Cu39.5Si40 feed rod is illustrated in Fig. 5. The expected evaporation loss was compensated by 0.5 at% excess Eu. Congruent melting is supposed because of the absence of secondary phases at the transition from the feed rod to the growing crystal (Fig. 5a). The single crystalline crosssection after 13 h growth is shown in Fig. 5b. The crystal contains some subgrain boundaries and macroscopic cracks, but there is not any sign of a secondary phase. The composition of the crystalline matrix is 21.8 at% Eu, 39.7 at% Cu, and 38.6 at% Si (from EDX analysis) and the growth axis is close to an orientation on the basis plane of the tetragonal unit cell (/1 0 0S) with about 151 deviation from rod axis. The final section shows the interface crystal/ quenched FZ (Fig. 5c). The presence of Cu-rich phase in EuCu2Si2 matrix of FZ indicates a slight composition shift during the growth process. Similar features were also observed for the other crystals.

4.3. Crystal growth of compounds with volatile elements

5. Crystal perfection and selected properties

Normally, FZ method is not suitable for crystal growth of refractory intermetallics containing volatile elements such as Cu, Yb, Sm, Tm and Eu. The presence of silicon in the melt reduces the partial vapor pressure of the rare earth elements significantly. However, another key point is the application of high pressures (up to 5 MPa) of Ar atmosphere in the growth chamber, which can sluggish the

5.1. Rare earth-transition metal borocarbides Crack-free RNi2B2C single crystals were received from FZ process. The X-ray Laue back-scattering patterns do not show twinning and stress distortions. For selected crystals narrow neutron diffraction rocking curves were received with a typical half-width full-maximum of

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Fig. 5. Optical images in polarized light visualizing the grain selection process in the initial stage of FZ crystal growth of EuCu2Si2. (a) Longitudinal section parallel to rod axis (left: polycrystalline seed), (b) cross-section after about 13 h growth with 3 mm/h and (c) longitudinal section at the end of the crystal (right: quenched traveling zone).

2.5

10 RRR = 24 (as grown)

YNi2B2C

2.0

HoNi2B2C

1.5

6 ρ (μΩ cm)

ρ (μΩ cm)

RRR = 12

8

RRR = 39

RRR = 62 H=0

1.0

4 RRR = 19

0.5

2

0.0

0 10

12

14 16 Temperature (K)

18

20

Fig. 6. Electrical resistivity vs. temperature plot r(T) of an YNi2B2C single crystal showing the effect of heat treatment: RRR ¼ 24 as-grown; RRR ¼ 39 after annealing 72 h at 900 1C; RRR ¼ 62 after optimum annealing (cf. text).

GE0.451, which matches the resolution of the diffractometer used [13]. Within the angular difference G, the intensity of a selected reflection of neutron diffraction decays to 50% of its maximum. The residual resistance ratio (RRR) is an appropriate measure of the crystal perfection. Typical values of R300 K/R16 KE24 for as-grown YNi2B2C (Fig. 6) and R300 K/R9 KE12 for HoNi2B2C (Fig. 7) have been achieved. The width of superconducting transition (from 90% to 10% of the normal resistivity) in YNi2B2C has been considerably improved from DTcE0.8 K for crystals grown by FZ with RF induction heating to DTcE0.1 K for large crystals grown by FZ with optical heating [12]. The perfection of the latter crystals is even superior to flux grown crystals displaying DTcE 0.25 K and RRRp20 [5]. Due to element segregation, the properties in large RNi2B2C single crystals can vary along rod axis and in the radial direction [13]. The control of the feed rod composition and growth parameters strongly reduced segregation effects. Single crystals of HoNi2B2C with specific modifica-

5

6

7

8 9 Temperature (K)

10

11

Fig. 7. Electrical resistivity vs. temperature plot r(T) of a HoNi2B2C single crystal showing the effect of heat treatment: RRR ¼ 12 as-grown; RRR ¼ 19 after optimum annealing (cf. text).

tions of superconducting behavior were grown [12,13]. One interesting feature of the HoNi2B2C compound is the co-existence of superconductivity and magnetic phenomena. By neutron diffraction experiments on large HoNi2B2C single crystals (with high signal-to-noise ratio), space division of different magnetic structures (antiferromagnetic modulations) could be identified at T=5.3 K, which provided new insight into the character of the magnetic order. 5.2. Rare earth-transition metal silicides Substitution of various rare earths within the class of R2PdSi3 single crystals revealed a systematic dependence of anisotropic magnetic properties governed by the delicate interplay of crystal-electric field (CEF) effects and magnetic two-ion interactions. In particular, the anisotropy of the magnetic properties is largely governed by the shape of the 4f-electron orbital [24]. For rare earth elements with cigar-shaped orbitals (Er, Tm), c-axis of the hexagonal unit

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cell is the easy magnetic axis, whereas for oblate-shaped orbitals (Tb, Ce) the magnetic easy axis is in (a,b)-basal plane. For Gd with nearly spherical 4f-electron orbital, the anisotropy is small. For Dy2PdSi3 and Ho2PdSi3, the magnetic easy and magnetic hard directions in the ordered state are determined by higher order CEF terms. This corresponds to a crossing of magnetic easy and hard direction in the paramagnetic state far above the Ne´el temperature in both compounds. Other interesting phenomena of this class of compounds are the anisotropic large negative magnetoresistivity observed in Ce2PdSi3, Tb2PdSi3 and Dy2PdSi3 crystals [18,21,22] and the anisotropic magnetocaloric effect discovered in Tb2PdSi3 [21]. Neutron diffraction studies at large high-quality R2PdSi3 single crystals (R ¼ Tb, Er, Tm) proved the existence of a periodic Pd, Si-superlattice on B-site instead of the hitherto-known plain hexagonal AlB2 crystallographic structure.

↑↑↓ 6

Bmeta paramagn.

Magnetic field (kOe)

2274

H II [001]

4

Bmeta'

↑↓↑↓ afm

2

H II [110] TN

Tm

0 2

4

6

8

Tc

Temperature (K)

5.3. Effects of post-growth heat treatment By heat treatments the perfection of as-grown crystals was optimized with respect to physical properties. By annealing YNi2B2C crystals for 72 h at 900 1C, the RRR increased from 24 in the as-grown state to 39. Finally, the crystals were annealed in a multi-step procedure 12 h at 1000 1C, 48 h at 750 1C and 72 h at 500 1C. The YNi2B2C crystals subjected to optimum annealing show the highest RRR r300 K/r16 K ¼ 62 (Fig. 6) and the best superconducting characteristics reported up to now, Tc ¼ 15.6 K and DTc ¼ 0.1 K. Single crystals of HoNi2B2C show a RRR r300 K/r9 K ¼ 19 and the superconducting transition temperature is shifted from Tc ¼ 7.8 K in as-grown state to Tc ¼ 8.7 K after optimum annealing (Fig. 7). The superconductivity in as-grown single crystal is believed to be strongly affected by disorder effects on C and B sites. This may explain the remarkable effect of annealing (even at relatively low temperatures) for HoNi2B2C. The temperature dependence of the upper critical field Hc2(T) of an annealed HoNi2B2C single crystal (RRR ¼ 19) for two field directions, HJ[1 1 0] and HJ[0 0 1], is shown in Fig. 8. Below the paramagnetic superconducting phase, which exists in the temperature interval TmoToTc, two incommensurate (IC) antiferromagnetic (afm) structures develop between Tm and TN, whereas the commensurate (C) afm phase appears below TN. Dashed lines denote the metamagnetic (MM) structure developing within C structure between Bmeta0 and Bmeta. The arrows indicate that the magnetic Ho moments are parallel and antiparallel to [1 1 0] direction in neighboring ab planes. The anomalous decrease of Hc2(T) with decreasing T between Tm and TN is due to suppression of superconductivity by the IC structures and/or by the MM structure which is associated with a net magnetization. The anisotropy of Hc2(T) above TN results from enhanced pair breaking for HJ[1 1 0] due to interaction of Ho moments with conduction electrons. Note that Hc2(T) becomes

Fig. 8. Temperature dependence of upper critical field, Hc2(T), in an optimal annealed HoNi2B2C single crystal (RRR ¼ 19) for the two field directions HJ[0 0 1] and HJ[1 1 0]. Magnetic structures obtained from specific heat and magnetization measurements for HJ[1 1 0] are explained in text.

isotropic below TN, which strongly suggests that superconductivity in C phase survives at a special Fermi surface sheet (FSS) which is isolated from the influence of the rare earth magnetism localized at another FSS [35]. Annealing can also improve the perfection of R2PdSi3 crystals. Various ternary R–Pd-silicides exhibited binary RSi or RSi2 precipitates within the single crystalline R2PdSi3 matrix, which are formed during the growth process due to the retrograde solubility. In some cases, Tb2PdSi3 and Ce2PdSi3 precipitates were dissolved by longterm annealing of crystals at elevated temperatures up to 1200 1C and subsequent quenching. For other compounds, Tm2PdSi3 and Ce2CoSi3, the precipitates are retained even after annealing. 5.4. Rare earth-transition metal solid solution compounds Solid solution TbxY1xNi2B2C single crystals provided unique examples for tuning the interplay of superconducting and magnetic properties. Superconductivity prevailed up to x ¼ 0.4. The anisotropy of the upper critical field Hc2(T) of these pseudo-quaternary compounds is governed both by the fraction x of Tb ions and magnetic ordering effects [36]. In solid solution crystals segregation phenomena are more severe. In TbxY1xNi2B2C this leads to a systematic increase of Tb-fraction in the crystal over the length co-ordinate and a slight increase of Ne`el temperature TN, whereas Tc was virtually constant [10]. Considerable element segregation with accumulation of metallic components in FZ was also observed for FZ crystal growth of Ce2PdxCo1xSi3 compounds. In the

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crystals the transition metal content (Pd+Co) is reduced compared to the nominal value [37]. The segregation behavior is the reason for a strong tendency towards morphological instability. The radial segregation was less severe in FZ grown crystals with RF induction heating, which permits higher growth velocities compared with FZ with optical heating due to forced convection and improved solute mixing. However, one drawback of the method was a narrow polycrystalline surface rim around the single-crystalline inner part of the rod, which probably originated from a concave melt–crystal interface area even at a high generator frequency of 3 MHz. 6. Summary FZ crystal growth methods turned out to be very convenient for growing large crystals of novel classes of rare earth-transition metal-borocarbides and-silicides because reaction of melts with crucibles is avoided. The control of oxygen impurities throughout the preparation process is a key problem for growing high-quality crystals. The process parameters crucially depend on the melting temperature and the solidification mode of the compounds. Elevated Ar pressure enabled FZ crystal growth of compounds with volatile elements. Heat treatments of asgrown crystals can improve the real structure of as-grown crystals and optimize the physical properties. From various RNi2B2C crystals, the superconducting phenomena and the interplay of superconductivity and magnetic ordering were revealed. The effect of substitution of rare earth elements R on magnetic ordering phenomena was systematically studied in R2PdSi3 single crystals. The high perfection of single crystals enabled the discovery some novel physical and crystallographic features of the compounds by neutron diffraction. Acknowledgments The authors thank S. Pichl, S. Mueller-Litvanyi, A. Teresiak, A. Ostwaldt, K. Nenkov, J. Werner, H. Bitterlich and G. Graw for experimental assistance, M. Frontzek for helpful discussions and D. Lindackers, F. Fischer, R. Voigtla¨nder and St. Ziller for providing the sketch of the new floating zone facility. We express our gratitude to the Deutsche Forschungsgemeinschaft for financial support by SFB 463 ‘‘Rare earth-transition metal compounds: Structure, Magnetism, and Transport’’. 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) 146.

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