Structural and magnetic investigation of nonstoichiometric YFe10V2 and its interstitial carbide prepared by arc-melting

Journal of Alloys and Compounds 299 (2000) 45–54

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Structural and magnetic investigation of nonstoichiometric YFe 10 V2 and its interstitial carbide prepared by arc-melting a a, a,b a b b c N. Plugaru , D.P. Lazar *, A. Galatanu , M. Morariu , G. Wiesinger , A. Kottar , E. Vasile a

National Institute of Material Physics, P.O. Box MG-07, 76900 Bucharest–Magurele, Romania b Technical University of Vienna, Wiedner Hauptstraße 8 – 10, A-1040 Vienna, Austria c METAV-S. A., P.O. Box 18 /3, Bucharest, Romania Received 8 November 1999; accepted 18 November 1999

Abstract A far-from-equilibrium preparation route and using a particular combination of starting materials led us to obtain nonstoichiometric Y(Fe,V ) 12 and its interstitial carbide. Structural analysis and EDAX results reveal the formation of the main 1:12 phase with the ThMn 12 -type structure, with 13–15% vacancies at the 3d metal sites. Insertion of interstitial carbon increases the T C and magnetization of ¨ the 1:12 phase, whereas the unit cell volume remains unchanged. The effects of vacancies and interstitial carbon on the Mossbauer hyperfine fields and isomer shifts are discussed with reference to the calculated quantities.  2000 Elsevier Science S.A. All rights reserved. 57 ¨ spectroscopy; Local environment analysis Keywords: Rare earth intermetallic compounds; Structural analysis, Rietveld refinement, Fe Mossbauer

1. Introduction There is considerable interest in the investigation of interstitial RFe 122x M x (A y ) compounds with ThMn 12 -type structure (M standing for a stabilising element and A for carbon or nitrogen) for both their intrinsic magnetic properties and permanent magnet applications [1–6]. The carbon occupancy of the interstitial (2b) sites in the compounds prepared by gas–solid phase reaction has well been established, as well as its effects on increasing the Curie temperature, iron sublattice magnetization and reinforcing the uniaxial anisotropy at the rare earth site. However, controversial results as regards the site occupancy of carbon and its effects on the structural and magnetic properties of the compounds prepared by arcmelting have been reported. Thus, Drzazga et al. [7] found that carbon increases the Curie temperature and iron magnetization in RFe 122xVx (C y ), with R5Y or Dy, while decreasing the lattice parameters, and inferred from their results that carbon occupies the interstitial (4d) sites and also some Fe sites. In a neutron diffraction study of RFe 10 V2 C y , Lin et al. [8] determined that carbon occupies

*Corresponding author. Fax: 140-1-4930-267. E-mail address: [email protected] (D.P. Lazar)

the interstitial (2b) sites and that the lattice parameters, T C and iron magnetization increase slightly for a carbon uptake of y50.3. Recently, Mao et al. (see Ref. [9] and the references cited therein) have determined by neutron diffraction investigation of YFe 10.5 Mo 1.5 Cy with y#0.6, that carbon occupies almost exclusively the (2b) sites and expands significantly the crystal lattice. Also, these authors have found that the presence of the interstitial carbon enhances considerably the hard magnetic properties of RFe 10.5 Mo 1.5 C y , with R5Y and Dy. Usually, the crystal structure analyzes performed by Rietveld refinement of the diffraction patterns of R(Fe,M) 12 Ay were carried out under the assumption of the ideal occupancies of the iron sites or very close to those values [8–13]. It is worth noting that, occasionally, the presence of significant amounts of a-Fe or a-(Fe,M) or some other secondary phases was indicated. Actually, the formation of secondary phases may be consistent with a deficit of elements in the main 1:12 phase. Much attention has been paid to the study of the ironrich corner of A–Fe–Mo systems, where A is a 4f or 5f element and several authors have found that the ternary ThMn 12 -type phase may form within rather wide composition ranges (see Refs. [14,15] and references cited therein). The investigation of the structure of nonstoichiometric Gd(Fe,Mo) 12 solid solution led Zinkevich et al.

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N. Plugaru et al. / Journal of Alloys and Compounds 299 (2000) 45 – 54

[14], to conclude that the Fe-sublattice may contain as much as 6–8% constitutional vacancies. Obviously, the effects due to the existence of vacancies and non-stoichiometry superimposes to those of interstitials and solute additions. Thus, Garcia et al. [16], proposed a model based on the diffusion of Fe vacancies and interstitials and its influence on the mobility of domain walls to describe the behaviour of the magnetic susceptibility in the presence of disaccommodation processes in several classes of rare earth–iron compounds. Welch [17] has pointed out the importance of using thermodynamic and kinetic models in understanding and optimizing processing methods for rare earth–transition metal magnetic materials for technical applications. The estimates of defect energetics and free energies in the frame of such models are dependent on the realism of the description of metallic bonding. The experimental evidence on the properties of point defects and non-stoichiometry should provide useful data for the theoretical calculations. Therefore, we found it worthwhile to investigate the structural and magnetic properties of nonstoichiometric Y(Fe,V ) 12 (C y ) produced by arc-melting and to tentatively assert the effects contributed by the Fe vacancies and carbon atoms, respectively.

2. Experimental Alloys with the nominal compositions YFe 10 V2 and YFe 10 V2 C 0.5 were prepared by arc-melting stoichiometric quantities of yttrium (99.8%), iron (99.9%), a pre-alloy (V514.1 wt% Fe) and carbon (99.99%), in a water-cooled copper crucible, under argon atmosphere. In the case of the carbon-containing sample, yttrium and carbon were melted separately and then remelted with the other constituents. An excess of yttrium was added in order to compensate for the mass losses by evaporation during repeated melting of the alloys, to attain better homogeneity. Only samples showing mass variations smaller than 0.5% of the mass of the nominal compound were selected. The ingots were wrapped in Ta foil, sealed in evacuated silica tubes and annealed at 9008C for 8 days. Then, the samples were water quenched. Secondary electron images of the samples polished and etched with Nital 2% solution were obtained utilizing a scanning electron microscope Philips 515, also equipped with an Energy Dispersive Analysis of X-ray (EDAX) facility for chemical composition determination. Elements lighter than Na are not detectable by EDAX, thus the carbon content in the YFe 10 V2 C 0.5 alloy was estimated by refinement of the X-ray diffraction data, as discussed later. The crystal structure was investigated by X-ray diffraction on powder samples with grain size less than 20 mm. The diffraction patterns were recorded with high statistics utilizing a Siemens D500 diffractometer with Co–K a ˚ selected by a graphite monoradiation (1.789007 A)

chromator in the secondary beam, at room temperature (RT). The patterns were analyzed using the FULLPROF.98 Version 0.2 software for Win. NT / 95 [18]. Thermal variations of magnetisation curves were measured by means of a Weiss-type balance in the temperature range 293–950 K in a magnetic field of 200 Oe. Magnetisation isotherms at 4 K were measured utilizing an extraction magnetometer in fields up to 10 T. The ¨ Mossbauer spectra were recorded with a Promeda spectrometer at room temperature (RT) and by means of a conventional constant acceleration type spectrometer with a helium flow cryostat at 5 K. A 57 Co in rhodium 20 mCi source was used and the velocity scale was calibrated to ¨ the spectrum of a-Fe foil, at RT. Homogeneous Mossbauer absorbers of about 35 mg / cm 2 effective thickness were prepared from fine powders mixed with polystyrene dissolved in toluene and fixed on iron-free mica foil. The spectra were analyzed in the approximation that treats the quadrupole interaction as a perturbation to the magnetic hyperfine interaction, with the MOSSFIT programme, utilizing a least squares minimization subroutine [19]. As fit parameters for a sextet the hyperfine field, Hhf , the isomer shift, d, the quadrupole shift, e, and a single linewidth, G, were used. We used Lorentzian lines and constrained their relative intensities to the ratio 3:(26x):1:1:(26x):3, with x in the range 0–0.15, to allow for a slight texture of the samples. The estimated errors are 63 kOe for Hhf , 60.015 mm / s for d and 60.03 mm / s for e. The d values are reported relatively to a-Fe foil at RT.

3. Results and discussion

3.1. Phase and structural analysis The SEM images of the YFe 10 V2 and YFe 10 V2 C 0.5 alloys revealed that they mainly consist of a matrix with the ThMn 12 -type structure. Two types of impurity phases were also observed, the first one identified as a-(Fe,V ), present in both samples, and a cubic phase, identified as vanadium carbide (S.G. Fm-3 m), present only in the interstitial YFe 10 V2 C 0.5 solid solution. The chemical compositions of these phases, as given by EDAX, are presented in Table 1. Although the EDAX results at different spots on the Table 1 Results of EDAX measurements Phase

YFe 10 V2 , matrix a-(Fe,V ) YFe 10 V2 C, matrix a-(Fe,V ) Vanadium carbide

Element content (%) Y

Fe

V

8.7(5) 0.7(4) 9.0(5) 0.9(4) 0.8(4)

76.5(7) 79.7(7) 78.4(7) 87.7(8) 3.5(4)

14.7(5) 19.6(6) 12.6(5) 11.4(5) 95.7(8)

N. Plugaru et al. / Journal of Alloys and Compounds 299 (2000) 45 – 54

same phase showed a statistic dispersion of the atomic percentages x?Y, y?Fe and z?V, which is expressed by standard deviations between 0.5 and 0.7 for the main 1:12 phase, we remarked that the values of the ratios y?Fe /z?V were approximately constant, 5.260.2 in the case of YFe 10 V2 and 6.260.2 for YFe 10 V2 C 0.5 . Consequently, these ratios were further used in the refinements of the X-ray spectra of the compounds. In the tetragonal unit cell (S.G. I4 /mmm, Z52) the rare earth atoms occupy the (2a) sites and the iron atoms and structure-stabilising element, M, share the (8i), (8j) and (8f ) sites. Two types of interstitial sites also exist, the octahedral (2b) sites and the tetrahedral (4d) sites. Firstly, we refined the diffraction pattern of YFe 10 V2 sample, assuming a stoichiometric 1:12 main phase with vanadium atoms preferentially distributed over the (8i) sites and the atomic positions given in [20]. There were refined the scale factors, zero point of the detector, lattice constants, atomic positions, overall Debye–Waller factor, profile parameters for pseudo-Voight line shape and occupation numbers (chemical occupancy multiplied by site multiplicity). We carried out the fit by applying the constraint n?V(8i)1m?Fe(8i)58 for the (8i) site occupation numbers. As a result, we obtained n51.20(2) and m56.80(2). The reliability factors of the fit for the main phase were R wp 50.126, x 2 59.33 and R B 50.108. Also, an a-Fe secondary phase was introduced in the fitting. The FULLPROF analysis of the refinement and the significant differences between observed and calculated intensities in

47

the low angle region (see Fig. 1) strongly demanded for an improvement of the structural model. Moreover, the refined composition, YFe 11.40 V0.60 , seemed rather unrealistic when comparing it to both the nominal composition YFe 10 V2 and the EDAX composition of YFe 8.77 V1.69 for this sample. As a next step in the refinement procedure, we tried to follow the strategy described by Zinkevitch et al. [14] for the case of nonstoichiometric solid solutions with ThMn 12 type structure in the Fe–Gd–Mo system. These authors concluded that the (8i) sites show the tendency of being fully occupied, whereas the (8f ) and (8j) sites can be partially vacant. Under these assumptions our refinement of the occupation numbers for the (8f ) and (8j) sites produced the values nhFe(8f )j56.78(4) and nhFe(8j)j5 6.79(3), which indeed supports the existence of iron vacancies on these sites. However, the constraint of full occupation imposed to the (8i) site has led to a negative value of the iron occupation number for this site, which may be an indication that it is partially vacant, also. In order to check this assumption, one should refine the occupation numbers of the Fe and V atoms on the (8i) sites as independent parameters. Still, they cannot be refined simultaneously because of the strong correlation between them. Also, one could expect that the refined composition is not unique (different combinations of site occupation numbers may result in the same final R values) all the more that the X-ray scattering factors of iron and vanadium are almost the same. Thus, we refined only the (8i)

Fig. 1. Ritveld refinement of the XRD pattern assuming a stoichiometric composition for the YFe 10 V2 sample.

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N. Plugaru et al. / Journal of Alloys and Compounds 299 (2000) 45 – 54

occupation number of the Fe atom for different fixed values of the occupation number of the V atoms on (8i) sites and we chose that solution for which the ratio of the total number of Fe atoms in (8i), (8j) and (8f ) positions to the number of V atoms is close to that in the EDAX chemical composition, i.e. 5.2. Finally, the best agreement between the observed and calculated patterns was obtained for a refined composition YFe 8.70 V1.67 and a-(Fe,V ) as the secondary phase (see Fig. 2). The cell parameters, occupation numbers and R-factors are given in Table 2. This solution is also satisfactory as regards the total atomic balance between the nominal composition and the refined composition yielded by the fitting: YFe 10 V2 →YFe 8.70 V1.67 1Fe 1.30 V0.33 , where the composition of the remaining a-(Fe,V ) alloy is close to Fe 79.7 V19.6 measured by EDAX (see Table 1). In the case of the YFe 10 V2 C 0.5 alloy, three phases were considered in the refinement of the diffraction pattern: the main ThMn 12 -type phase, an a-(Fe,V ) phase and a vanadium carbide phase. The initial atom coordinates for the 1:12 phase were those given for YFe 11 TiN x in [21]. We found that the fitting in compliance to the stoichiometric composition and placing the interstitial carbon atoms at the (2b) sites led to unsatisfactory reliability factors, R wp 5 0.133, x 2 56.55 and R B 50.120 and also to significant differences between the observed and the calculated intensities in the low angle region (see Fig. 3). At the same time, the refined composition YFe 9.25 V2.75 C 0.5 was different to YFe 8.67 V1.4 C y , measured by EDAX (see Table 1).

Table 2 Lattice constants, atomic positions, occupancy numbers, overall isotropic temperature factors, as well as the fit quality factors: x 2 , R wp and R B Parameter

YFe 8.70 V1.67 h 1.63

YFe 8.74 V1.41 h 1.85 C 0.3

a (nm) c (nm) x(8i) x(8j) Occ. No. Fe(8i) Occ. No. V(8i) Occ. No. Fe(8j) Occ. No. Fe(8f ) Occ. No. C(2b) ˚ 2) B (A R wp x2 RB

0.848362(18) 0.476556(15) 0.3625(3) 0.2755(3) 3.87(4) 3.35 6.76(3) 6.77(3) – 0.86(3) 0.111 7.15 0.060

0.848026(13) 0.476708(11) 0.3604(3) 0.2761(3) 4.06(4) 2.82 6.76(4) 6.66(3) 0.60(7) 0.73(3) 0.107 4.29 0.067

Following the same refinement strategy as previously described, in the third step the refined composition of the main phase, YFe 8.74 V1.41 C 0.3 matched the EDAX composition and the quality of the fit was rather good. The experimental and calculated patterns are shown in Fig. 4 and the cell parameters, atomic occupation numbers and reliability factors are also listed in Table 2. Similarly to the case of YFe 10 V2 , one can obtain a reasonable atomic balance between the nominal and refined compositions: YFe 10 V2 C 0.5 →YFe 8.74 V1.41 C 0.3 1Fe 1.26 V0.16 1V0.43 C 0.2 , in which the derived stoichiometry of the (Fe,V ) alloy is close to the measured one by EDAX. Also, according to these results, the refined content in interstitial carbon in the

Fig. 2. Final Ritveld refinement of the XRD pattern according to the strategy described in the text for the YFe 10 V2 sample.

N. Plugaru et al. / Journal of Alloys and Compounds 299 (2000) 45 – 54

Fig. 3. Ritveld refinement of the XRD pattern assuming a stoichiometric composition for the YFe 10 V2 C 0.5 sample.

Fig. 4. Final Ritveld refinement of the XRD pattern according to the strategy described in the text for the YFe 10 V2 C 0.5 sample.

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N. Plugaru et al. / Journal of Alloys and Compounds 299 (2000) 45 – 54

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Table 3 Distribution of iron vacancies, average numbers of Fe NN and vacancies NN at the three Fe sites and the refined phase compositions Iron site

YFe 8.70 V1.67 h 1.63 h Fe

(8i) (8j) (8f )

0.78 1.24 1.23

h Fe ThMn 12 FeV VC

YFe 8.74 V1.41 h 1.85 C 0.3 Av. Fe–NN 9.18 7.01 7.01

Av. h–NN

h Fe

1.72 1.32 1.32

1.12 1.24 1.34

13.6% 92.3 % 7.7% –

Av. Fe–NN 9.25 7.05 7.08

Av. h–NN 1.85 1.54 1.52

15.4% 86.3% 5.1% 8.6%

compound prepared by arc-melting is the same as previously determined by refinement of neutron diffraction data [8]. Table 3 displays the distribution of the Fe vacancies on the iron sites in the elementary cell. One may observe that while the numbers of vacancies are almost the same for the (8j) and (8f ) sites in the two 1:12 compounds, the (8i) site is less filled in YFe 10 V2 C 0.5 because of the vanadium carbide formation and the reluctance of iron atoms to occupy the (8i) site. The refined phase contents of the investigated alloys are also presented in Table 3. The refined cell parameters of the a-(Fe,V ) alloys, a5 0.288591(12) nm for the first sample and 0.287557(14) nm for the second one, are situated between that of pure a-Fe (0.28665 nm) and pure a-V (0.303 nm) and reflect correctly the relative contents in Fe and V of the alloys.

3.2. Magnetic properties The Curie temperature increases from 542 K in the YFe 8.70 V1.67 h 1.63 compound to 568 K in the interstitial YFe 8.74 V1.41 h 1.94 C 0.3 , in spite of the slight contraction in ˚ 3 for y50 to 342.8 the unit cell volume, from V5343.0 A 3 ˚ for y50.3. The 4 K saturation magnetization also A

increases from 15.3 m B / f.u. to 17.9 m B / f.u. by carbonation. It should be remarked that the magnetization values were obtained by using the masses of the nominal formulae of the compounds and therefore are affected by some unknown error, due to the contribution of the (Fe,V ) secondary phase to magnetization. Nevertheless, taking into account that this impurity phase is present in both samples in similar fractions (see Table 3), one may assume that the increase in the saturation magnetisation, |17%, primarily reflects the increase in the mean iron magnetic moment in the 1:12 carbide. The experimental Curie temperature and magnetization variations may be analyzed on the grounds of the nearest neighbour environments at the different iron sites (see Tables 3 and 4) in conjunction with the results of first-principles calculations of the magnetic properties in 1:12 compounds and their interstitial counterparts. Deniszczyk and Borgiel calculated the iron magnetic moments [22] and iron magnetic moments ¨ and Mossbauer hyperfine fields [23] in Y(Fe,V ) 12 , interstitial YFe 8 V4 C with the carbon atoms located at the octahedral (2b) or tetrahedral (4d) voids and YFe 8 (V3 C) with one carbon replacing vanadium at the (8i) sites. According to their results for the hypothetical interstitial YFe 8 V4 C, the effect of carbon introduced at the (2b) voids is to decrease

Table 4 Number of nearest neighbours and interatomic distances in YFe 8.70 V1.67 and YFe 8.74 V1.41 C 0.3 compounds YFe 8.70 V1.67

Y(2a)

Fe / V(8i)

Fe(8j)

Fe(8f )

(2b)

˚ 3] VW– S [A

Y(2a) Fe / V(8i)

– ˚ 133.08 A

– –

29.3 12.8

˚ 233.05 A

˚ 833.05 A ˚ 232.61 A ˚ 232.66 A ˚ 232.69 A

˚ 833.23 A ˚ 432.61 A

Fe(8j)

˚ 432.44 A

–

11.6

Fe(8f )

˚ 233.23 A

˚ 433.08 A ˚ 132.33 A ˚ 432.90 A ˚ 232.61 A ˚ 232.66 A ˚ 432.61 A

˚ 432.44 A

˚ 232.38 A

–

11.1

˚ 833.05 A ˚ 232.62 A ˚ 232.65 A ˚ 232.68 A

˚ 833.23 A ˚ 432.61 A

˚ 232.38 A ˚ 033.88 A

28.3 12.9

˚ 432.44 A

˚ 131.90 A

11.0

˚ 432.44 A ˚ 431.90 A

˚ 232.38 A ˚ 033.23 A

˚ 033.23 A

11.1 3.2

YFe 8.74 V1.41 C 0.3 Y(2a) Fe / V(8i)

– ˚ 133.06 A

Fe(8j)

˚ 233.05 A

Fe(8f ) C(2b)

˚ 233.23 A ˚ 232.38 A

˚ 433.06 A ˚ 132.37 A ˚ 432.91 A ˚ 232.62 A ˚ 232.65 A ˚ 432.61 A ˚ 033.88 A

N. Plugaru et al. / Journal of Alloys and Compounds 299 (2000) 45 – 54

drastically the T C and moderately the iron magnetization. Also, they found for the Fermi contact contribution to the hyperfine field the relation Hhf (8f ).Hhf (8j) for all compounds studied [23] and derive the conclusion that carbon atoms should be distributed over the (4d) and (8i) sites in order to justify the experimental data reported in [7]. At variance with their results are those obtained by Yang et al. [24] for YFe 10 Mo 2 X, with X5H, B, C, N, O, F, and Asano et al. [25] for RFe 12 X, with R5Y, Ce, Gd and X5N, C. These latter calculations show that the magnetovolume effect results in an increase of magnetic moments at all Fe sites, while the chemical bonding effect of C atoms increases the magnetic moments at the Fe(8f ) and Fe(8i) sites but decreases that at the Fe(8j) sites. Since the structural parameters obtained in our work indicate that carbon atoms are located at the (2b) voids, we discuss our results with reference to [24,25].

3.3. Effect of vacancies Inspection of the Fe–Fe interatomic distances, d Fe – Fe , listed in Table 4 shows that the presence of interstitial C leads to anisotropic variations in d Fe – Fe and enhances the structural disorder compared to the base compound. This is also evidenced by the variations in the atomic volumes approximated by the Wigner–Seitz cell volumes, VW– S , at the different lattice sites, calculated irrespective of the number of vacant sites nearest neighbours. One may observe a small increase in VW– S (8i), decreases in VW– S (2a) and VW– S (8j), while VW– S (8f ) remains constant as a result of carbonation. Actually, due to the presence of vacancy nearest neighbours (see Table 3) the VW– S values should be larger as compared to the calculated ones, in the base compound and its carbide. Furthermore, from Table 3 one may observe that the average numbers of vacancies NN increase in the carbide relative to the base compound at the different Fe sites due to the vanadium carbide phase formation. The existence of the supplementary 3d metal vacancies in the carbide contributes to increasing the VW– S at the (8i) and (8f ) sites and therefore increasing the Fe(8i) and Fe(8f ) magnetic moments compared to those in the base compound.

3.4. Effect of interstitial C It should be remembered that due to the C content of less than one atom per formula unit in our sample and assuming a random distribution of the interstitial on the (2b) sites, there may exist unit cells with zero, one or two interstitial carbon atoms. In the following discussion we consider the case of the stoichiometric YFe 10 V2 C, as the chemical bonding effects of the interstitial are maximal in that case. One may observe from Table 4 that Fe(8j) has ˚ and the Y atom has one C atom NN at a distance of 1.90 A ˚ Thus, according two C atoms NN at a distance of 2.38 A. to the calculations [24,25] the iron moment at the (8j) site

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is strongly reduced due to a considerable overlap between the valence states of the Fe(8j) and C atoms. In contrast, the iron moment at the (8f ) site, situated at a large distance ˚ from the C atom, increases because of d ( 8f ) – (2b ) 53.23 A the depressed bonding to the 2 yttrium NN, the latter forming tight covalent bonds with the C atoms NN. The iron moment at the (8i) site is the least affected by the chemical bonding effect, having both the largest distance ˚ and number of Fe NN to the C atom, d ( 8i ) – (2b ) 53.88 A, (see Table 3). Thus, the predictions of band structure calculations [24,25] taking into account the structural data displayed in Tables 3 and 4 support an increase in the mean iron magnetic moment, as experimentally observed, as a result of the competitive magnetovolume and chemical bonding effects associated with the interstitial C atom accommodation in the crystal lattice. The increased mean iron magnetic moment and therefore enhanced Fe–Fe overall exchange are responsible for the observed 4.8% increase in T C in the carbide. ¨ 3.5. Mossbauer spectral results ¨ The experimental Mossbauer spectra and their fittings are displayed in Fig. 5. The experimental patterns represent superpositions of two subspectra, originating from the main 1:12 phase and the minoritary a-(Fe,V ) phase. The contribution of the a-(Fe,V ) in a spectrum was accounted for by using a sextet with a broader linewidth, G 50.37 mm / s, compared to the linewidth in the calibration spectrum, G 50.26 mm / s, in order to approximate the (rather narrow) distribution of the hyperfine field caused by the different Fe–V environments possible at the Fe site. Also, we used the values of 280 and 304 kOe for the magnetic hyperfine splittings of Fe 1.30 V0.33 and Fe 1.26 V0.16 alloys, respectively, which we measured on separate samples with these compositions, at RT. Even if these values were somewhat different from those of the minoritary phase dispersed in the 1:12 matrix, there is no better approximation. The isomer shift value was set more negative than that of a-Fe, about 20.03 mm / s for Fe 1.30 V0.33 and 20.02 mm / s for Fe 1.26 V0.16 , having in view the average numbers of V atoms NN at an iron site, 1.6 and 0.9, respectively [26]. Also, the relative intensities of the 1:12 phase and a-(Fe,V ) phase in the spectra of the base alloy and its carbide were constrained to the ratios of 12:1 and 17:1, in agreement with the refined phase contents (see Table 3). In the fitting of the 1:12 phase we followed the procedure described in Refs. [27,28], assuming a binomial distribution of environments at the Fe sites and hence a binomial distribution of the relative intensities of the sextets, calculated using the refined iron occupancies at the three different crystallographic sites. Eleven sextets were used to account for the absorbtion spectrum of the 1:12 phase, with linewidths of about 0.32 mm / s at 5 K and 0.36 mm / s at RT.

N. Plugaru et al. / Journal of Alloys and Compounds 299 (2000) 45 – 54

52

¨ Fig. 5. Experimental Mossbauer spectra and fit results using the deduced relative sextets area (see text).

The site-averaged values of the hyperfine fields and isomer shifts, as well as their compound averages are presented in Table 5. For the YFe 8.70 V1.67 h 1.63 compound, the highest hyperfine field was obviously assigned to the Fe(8i) sites, on the basis of their relative intensity and biggest average number of Fe NN (see Table 3). In the assignment of the next Hhf value to the Fe(8j) sites and the lowest one to the Fe(8f ) sites, both having the same relative intensity and average number of Fe NN, we followed the guidelines provided by the calculations of the local iron magnetic moments [29] and local iron magnetic moments and hyperfine fields [24,25] which, independently, indicate the order Hhf (8j).Hhf (8f ) in the 1:12 yttrium compounds. In the case of the interstitial carbides, the calculations of Asano et al. [25] for the hypothetical YFe 12 C show that Hhf (8i) decreases moderately, but still preserves the highest value since no dramatic change takes place at these

sites. At the same time, Hhf (8j) strongly decreases because of the hybridization between the iron 3d states and the valence states of the carbon nearest neighbour, while Hhf (8f ) remains about the same as in YFe 12 . Consequently, the calculated hyperfine fields order in the sequence Hhf (8i).Hhf (8f ).Hhf (8j) and the compound-averaged Hhf value slightly decreases. These results are consistent with the calculations carried out by Yang et al. [24] for the YFe 10 Mo 2 C compound, with the difference that in the latter case a considerable increase in the hyperfine field at the Fe(8f ) sites was obtained. A survey of the Hhf data presented in Table 5 shows that all hyperfine fields increase in the interstitial YFe 8.74 V1.41 h 1.94 C 0.3 compared to their values in the base compound. The observed Hhf increases at the 8i and 8j sites are in contrast to the predictions of the calculations [24,25]. However, for the same reasons as in the base compound and in agreement with the calculations we

Table 5 Hyperfine parameters Compound

YFe 8.70 V1.67 YFe 8.74 V1.41 C 0.3

T (K)

5 293 5 293

d (mm / s)

Hhf (kOe) (8i)

(8j)

(8f )

Wt.av.

(8i)

(8j)

(8f )

Wt.av.

297 281 334 298

239 223 260 222

219 199 276 248

244 226 285 250

20.008 20.134 0.041 20.074

20.034 20.171 20.086 20.227

20.060 20.241 20.045 20.185

20.038 20.190 20.041 20.175

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assigned the highest Hhf value to the Fe(8i) site. In the assignment of the two lower hyperfine fields to the Fe(8j) and Fe(8f ) sites, we have considered it reasonable to attribute the Hhf value giving the smallest increase to the (8j) sites and the Hhf value giving the largest increase to the (8f ) site, which is the only choice consistent with the predictions of the calculations [24,25]. The 5 K values of the hyperfine fields for YFe 8.70 V1.67 h 1.63 compare well to the 4.2 K values reported in Ref. [27] and 10 K values reported in Ref. [28] for the stoichiometric YFe 10 V2 . We mention that in the case of Ref. [27] we refer to a reversed assignment of the hyperfine fields to the (8j) and (8f ) sites. We suggest that the consistency of the low temperature Hhf data results from the contrasting effects of vacancies and magnetic nearest neighbours on the hyperfine field. Thus, iron atoms having vacant iron sites as nearest neighbours have larger atomic volumes and hence larger magnetic moments and hyperfine fields. However, the smaller number of magnetic atom nearest neighbours reduces the magnetic moment and hyperfine fields at that site. For small concentrations in vacancies these effects may balance each other. One may also observe that the 5 K hyperfine fields increase 12.5, 8.8 and 26% for the (8i), (8j) and (8f ) sites, respectively, as a result of the interstitial insertion of carbon. The smallest increase takes place at the (8j) site, which may be accounted for by the competing effects of the atomic volume variation and the chemical bonding to the C atom nearest neighbour. This assignment of the hyperfine fields to the three different iron sites is also supported by the outcoming correlations between the isomer shifts and Wigner–Seitz cell volumes and their variations due to the interstitial C insertion (see Tables 4 and 5). Thus, the isomer shift values scale to the VW– S values according to the sequences d (8i).d (8j).d (8f ) in the base compound and d (8i). d (8f ).d (8j) in the carbide. Also, in the carbide the isomer shift increases at the Fe(8i) site because of the increased Wigner–Seitz cell volume at that site. The isomer shift at the Fe(8f ) site is constant in the stated error and also the VW– S (8f ) is rather constant. Finally, the decrease of the isomer shift at the Fe(8j) site suggests that the VW– S (8j) decrease due to a local lattice distortion which dominates the size effect of the vacancies nearest neighbours.

4. Conclusions Structural and phase analyzes by X-ray diffraction and EDAX suggest that nonstoichiometric YFe 10 V2 (C 0.52 e ) compounds may be prepared by arc melting. We have found that the formation of secondary phases is consistent with the non-stoichiometry as well as with the presence of 13–15% vacancies at the 3d metal sites. The use of a V–Fe pre-alloy as starting material is decisive in controlling to some extent the formation of vacancies. Also, Rietveld

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refinements of XRD data have shown that up to 0.3 C at. / f.u. may occupy the interstitial (2b) sites. The combined effects of interstitial carbon and 3d metal vacancies lead to increased Curie temperature and magnetization in the YFe 8.74 V1.41 h 1.94 C 0.3 compound. The analysis of the effects induced by the vacancies and the interstitial carbon ¨ on the Mossbauer hyperfine fields and isomer shifts is possible only using the predictions of band structure calculations. The present results need further investigation and may be of some significance in the preparation route of the materials for technical applications.

Acknowledgements The authors are indebted to Mr. Gh. Cotiga for his help in performing SEM and EDAX measurements. This work was partially supported by a Grant from National Agency for Science, Technology and Innovation, Romania.

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