Mechanical alloying of Fe100−x−yNbxBy (x = 5, 10; y = 10, 15): from pure powder mixture to amorphous phase

Intermetallics. Vol. 16. Núm. 9. 2008. Pag. 1073-1082 http://dx.doi.org/10.1016/j.intermet.2008.06.006 1 Mechanical alloying of Fe100-x-yNbxBy (x=5, 10; y=10, 15): From pure powder mixture to amorphous phase J.J. Ipus, J.S. Blázquez, V. Franco, A. Conde* Dpto. Física de la Materia Condensada, ICMSE-CSIC, Universidad de Sevilla, P.O. Box 1065, 41080, Sevilla, Spain. Abstract The mechanical alloying process of Fe75Nb10B15 and Fe85Nb5B10 systems has been studied from an initial mixture of elemental powders. The amorphization process is monitored by X-ray diffraction, Mössbauer spectroscopy and magnetization measurements. An amorphous phase (with a Curie temperature of ~250 K) is formed after 400 h milling only for Fe75Nb10B15 alloy, whereas a bcc supersatured solid solution is the final product after milling Fe85Nb5B10 alloy. For both cases, a dispersion of ~10 % in the Fe content of the powder particles persists after 400 h milling. Powder particle size, Cr content and lattice parameter of bcc phase are larger for the alloy with the highest Nb content. Keywords: A. Nanostructured intermetallics; A. Magnetic intermetallics; C. Mechanical alloying and milling. *Corresponding author: Prof. A. Conde Departamento de Física de la Materia Condensada. Universidad de Sevilla. Apartado 1065, 41080 Sevilla (Spain). Phone : (34) 95 455 28 85/ Fax : (34) 95 461 20 97 E-mail: [email protected] Intermetallics. Vol. 16. Núm. 9. 2008. Pag. 1073-1082 http://dx.doi.org/10.1016/j.intermet.2008.06.006 2 1 Introduction Nanocrystalline materials are defined by a crystal size below 100 nm. As a limit, amorphous materials are solid systems where the structural long range order is lost. These materials have received much attention due to their physical properties (mechanical and magnetic), which are clearly different from those exhibited by conventional microstructures and, in some cases, improve their technological applicability [1,2,3]. One way to obtain nanocrystalline materials is by partial devitrification of a precursor amorphous alloy during controlled thermal annealing. This technique controls the microstructure of materials and thus optimizes the properties of the final product. Another possibility is mechanical alloying, which has become a very versatile technique to directly produce metastable microstructures (amorphous, nanocrystallines, supersaturate solid solution, etc) [1] from elemental powders or alloys. During this milling process the material is submitted to fracture and cold welding phenomena, as well as intensive plastic deformation, which define the powder morphology, microstruture and properties. The continuous storing of defects in the crystalline phase during milling process unstabilizes it, leading to nanocrystalline and/or amorphous structures [1]. Nanocrystalline Fe-M-B type alloys (M= Zr, Nb, etc), so-called Nanoperm [4], are attractive due to their soft magnetic properties after optimum thermal treatment and are used in commercial applications such as telecommunications, micro devices and power electronics [5,6]. Although these systems are generally obtained by rapid quenching and subsequent annealing, nanocrystalline alloys of these compositions can be directly obtained by mechanical alloying of elemental powders. As soft magnetic properties depend on the structure of the material [7,8,9], its structural characterization Intermetallics. Vol. 16. Núm. 9. 2008. Pag. 1073-1082 http://dx.doi.org/10.1016/j.intermet.2008.06.006 3 is a very important task to understand the system behavior and to predict its possible technological capabilities. In this study, two Fe100-x-yNbxBy (x=5, y=10 and x=10, y=15) alloys were produced by mechanical alloying from a mixture of pure elements and their morphological, compositional and microstructural evolution, as well as their thermal stability, were studied as a function of milling time. The amorphization of the ternary FeNbB system by rapid quenching methods has been studied previously [10]. Whereas compositions similar to Nb10 can be obtained in amorphous structure, those similar to Nb5 can not be obtained as amorphous. 2 Experimental Fe100-x-yNbxBy (x=5, y=10 and x=10, y=15) compositions were prepared by ball milling in a planetary mill Fritsch Pulverisette 4 Vario from elemental powders ( 99 % purity), with particle size d <200 m for Fe and Nb and d <1 mm for B. For simplicity, the studied alloys will be named in the following by their Nb content: Nb10 for Fe75Nb10B15 and Nb5 for Fe85Nb5B10. The initial powder mass was 30 g and the ball to powder ratio was 10:1. The rotational speed of the disk which supports the vials was 150 rpm and that of the vials was 300 rpm in opposite direction. After selected times, some powder was taken out from the vials to characterize the morphology, composition, microstructure and thermal evolution. The opening and closing of the vials was done under argon atmosphere in a Saffron Omega glove box to avoid oxygen and humidity contamination. Particle size distribution and morphology were studied by scanning electron microscopy (SEM) using secondary electrons (SE) and backscattered electrons (BSE) modes in a Jeol JSM-6460 LV operated at 30 kV. Compositional evolution was studied Intermetallics. Vol. 16. Núm. 9. 2008. Pag. 1073-1082 http://dx.doi.org/10.1016/j.intermet.2008.06.006 4 by energy dispersive X-ray (EDX) analysis using an Incax-sight of Oxford Instruments. Phase composition and structure were studied from X-ray diffraction (XRD), using Cu K radiation, and Mössbauer spectrometry (MS). Mössbauer spectra were recorded at room temperature in a transmission geometry using a 57Co(Rh) source. Values of the hyperfine parameters were obtained by fitting with NORMOS program [11]. The isomer shift, I, was quoted relative to that of -Fe at room temperature. Magnetization was measured using a maximum applied field of 1.5 T, in the temperature range from 77 to 440 K, every 15 K, in a vibrating sample magnetometer (VSM). The values of magnetization were obtained by extrapolation to zero field of the linear fitting of the high field magnetization. Thermal characterization of the samples was studied by differential scanning calorimetry (DSC) using a Perkin-Elmer DSC7 in Ar atmosphere. 3 Results 3.1 Morphology and composition Figure 1 shows SEM images of both alloys obtained after different milling times. For both alloys at short milling times, t <20 h, the particles are formed by joint layers and inclusions, which present different composition. In order to appreciate this heterogeneity in more detail, figure 2 shows SE (figure 2a) and BSE (figure 2b) images taken on a typical particle after 2 h milling for Nb10 alloy. Figure 2c shows EDX spectra taken on the different points marked in figure 2a. The A spectrum, on a dark inclusion, only shows the emission line of boron besides the typical background at low energy, so this zone is rich in this light element. The B and C spectra show zones rich in Fe and Nb, respectively. Therefore, the heterogeneity of the individual powder particles is evidenced for short milling times in the conditions used in this study. Intermetallics. Vol. 16. Núm. 9. 2008. Pag. 1073-1082 http://dx.doi.org/10.1016/j.intermet.2008.06.006 5 After 20 h milling, the powders show pores and cracks on their surface for Nb10 and Nb5 alloys but no clear layers and, for longer milling times, t 100 h, agglomeration of powder is observed. This process seems to depend on milling conditions, as other author have detected agglomeration of particles much earlier for similar compositions (e.g. after 5 and 15 h for Fe84Nb7B9 alloy with more energetic milling conditions [12]). The particles that form these agglomerates are larger for Nb10 than for Nb5 alloy. After 400 h milling, the agglomerates are not observed for Nb10 alloy but, for Nb5 alloy, the particles are still agglomerated. From SEM images, a statistical analysis of the average particle size evolution, <d>, has been performed over ~100 particles per sample (Figure 3) for both alloys. After 20 h milling, <d> increases over <d> > 200 m for both alloys. As milling time increases, a decrease in <d> is observed and, about 100 h milling, this value is stabilized, being smaller for Nb5 alloy (~ 25 m) than for Nb10 alloy (~ 50 m). This indicates that a stationary situation between cold welding and fracture has been achieved. The compositional evolution was studied by EDX from a statistical set of ~20 particles of each sample. As B content cannot be quantitatively measured by EDX, compositional analysis of the systems is referred to the relative amounts of Fe and Nb. In figure 4, histograms of Fe content of the powder particles are presented for Nb10 and Nb5 alloys after different milling times. For short milling times, t ≤5 h, it is possible to find Nb rich particles and a high fraction of Fe rich particles in both alloys. After 10 h milling, in Nb5 alloy a strong reduction in the broadening of the compositional distribution, ΔCFe, is observed (calculated as the difference in Fe/Fe+Nb ratio between the Fe richest and the Fe poorest powder particles found). However, this reduction occurs for 20 h in the Nb10 alloy, being a wider compositional distribution for the alloy Intermetallics. Vol. 16. Núm. 9. 2008. Pag. 1073-1082 http://dx.doi.org/10.1016/j.intermet.2008.06.006 6 with a higher Nb content. For longer milling times, t > 20 h, the broadening of the distribution of Fe content is almost constant and, at the end of the studied range, there is not a unique composition but a certain broadening is found (ΔCFe =16 at. % of Fe for Nb10 and 8 at. % of Fe for Nb5). Small Fe contamination is hard to measure in such Fe rich compositions. However, Cr is easily quantified as the initial powder mixture of this study is Cr free. Therefore, figure 5 shows the Cr concentration as a function of the milling time. For both alloys, a linear increase with the milling time is observed in the explored range, being the amount of Cr higher for Nb10 (~2 at. %) alloy than for Nb5 (~1 at. %) alloy after 400 h milling, in agreement with the expected increase of hardness as Nb increases in Fe based alloys [13]. Similar Cr contamination has been found in other ball milled systems [14]. 3.2 Structural evolution 3.2.1 X-ray diffraction Figure 6 shows the XRD patterns of both alloys as a function of milling time. For short milling times, t  20 h, a slight broadening of (110) diffraction peak of the αFe phase can be observed for both alloys. The full width at half maximum (FWHM) is approximately the same for both alloys at this stage; FWHM increases from 0.26 to 0.52 ± 0.10º from 1 to 20 h, respectively. This effect can be related with the decrease of the crystalline size and an increase of microstrains. After 50 h milling, a strong broadening of (110) peak is observed, as well as a shift to lower values of 2 position of this peak. Moreover, the different maxima of bccNb are no longer detected. These effects are related to the Nb incorporation into the bccFe lattice and to the formation of the supersaturated solid solution. For longer milling Intermetallics. Vol. 16. Núm. 9. 2008. Pag. 1073-1082 http://dx.doi.org/10.1016/j.intermet.2008.06.006 7 times, t >200 h, the XRD pattern of Nb10 alloy shows a halo centered at 2 ~ 44º and the -Fe (200) peak is reduced, being unappreciable for t =400h. This is related to the formation of the amorphous phase in the powder particles. 3.2.2 Mössbauer spectrometry Mössbauer spectra along with the hyperfine magnetic field distributions of the different samples are shown in figures 7 and 8 for Nb10 and Nb5, respectively. Mössbauer spectra were fitted using a ferromagnetic site contribution with HF =33 T for pure α-Fe phase (site-F) and two hyperfine magnetic field distributions; one for low field contributions, D1 (from 0 to 10 T) and other for high field contributions, D2 (>8 T). Furthermore, other site contribution but paramagnetic with quadrupolar splitting ~0.5 mm/s (site-P) is necessary to fit the spectra of Nb10 for long milling times. It is worth mentioning that, for such complex systems as the studied here, there is ambiguity between low field ferromagnetic sites (< 5 T) and paramagnetic ones. For t 10 h milling, the spectra only show a sextet with narrow absorption peak (width ~0.30 mm/s) and hyperfine magnetic field 33 T for both alloys, indicating that the α-Fe lattice has not been significantly affected by milling, in agreement with EDX (heterogeneous powder particles, not really alloyed) and XRD (lattice parameter close to pure α-Fe) results. For Nb10 alloy spectra at t  20 h, two new contributions appear, D1 and D2. This shows the existence of Fe atoms in three different main environments: first (siteF), pure bcc-Fe phase environment; second (D1), Nb rich environments; and third (D2), Fe rich environments. This latter contribution can be related with Fe atoms in the α-Fe phase but in the presence of impurities of Nb, B and/or Cr (due to contamination by the grinding media) or to Fe atoms at the interface region of nanocrystals [15]. However, in the Nb5 alloy, along with the crystalline contribution, site-F, only the D2 distribution is Intermetallics. Vol. 16. Núm. 9. 2008. Pag. 1073-1082 http://dx.doi.org/10.1016/j.intermet.2008.06.006 8 observed (D1 is negligible). This means that, in Nb5 alloy, Nb rich environments are not as significant as in Nb10 alloy, as it would be expected. 4 Discussion 4.1 Simulation of powder size evolution As it was previously mentioned, for 20 h milling, cracks are observed in the particles surface, which may cause the fracture of the particles when the milling continues. The change in the evolution trend of <d> could be ascribed to changes in the mechanical properties of the system due to rapid accumulation of defects into the particles [16]. This changes the balance between the two processes responsible for the evolution of <d>: fracture and cold welding. In fact, a qualitative change in the mechanical behavior of the powder is clearly observed for milling times longer than 50 h. Until 50 h milling, the powder sticks on the vial wall and ball surface. However, for longer milling times the powder detaches from milling media surfaces, noticeably increasing the amount of loose powder. Based on these two processes, a basic simulation algorithm was performed to obtain a first approximation of the tendency followed by <d> during milling. The initial system consist of 500 equally sized particles. Once a random particle is chosen, its probability to be cold welded with another particle is defined by Pcw =exp(-di/dc). This expression depends on the size of the i particle, di, and a critical size, dc. Particles with di>dc will tend to fracture, while if di <dc, the particle will tend to join with another. As can be inferred from the experimental evolution of <d>, dc is not constant during the milling process. In our simple simulation, only two different values of dc were used (supported by the abrupt change in the mechanical properties detected between 20 and 50 h. A dc value 100 times the initial size was used to reproduce the rapid increase in <d> at short milling times. For long milling times, the constant experimental value of Intermetallics. Vol. 16. Núm. 9. 2008. Pag. 1073-1082 http://dx.doi.org/10.1016/j.intermet.2008.06.006 9 <d> was used as dc. Simulation results reproduce the experimental ones after properly rescaling the iteration steps with the milling time (see figure 3). 4.2 Simulation of compositional evolution EDX results indicate that the composition of particles is not unique, even at long milling time but compositional distributions achieve a stationary situation. The parameter ΔCFe was simulated considered a system of particles defined by a 100 components binary array, being 0 for Nb and 1 for Fe (1 % in compositional resolution). Two random particles will interact changing a portion of their arrays with the same aleatory size. Therefore, powder particles size and number remain unchanged and only compositional evolution is simulated in this very simple simulation. Results are shown in figure 9a along with experimental ones for comparison. Experimental and simulation results agree even quantitatively (once iteration steps are conveniently rescaled), showing a stationary situation of the system with a compositional distribution and, therefore, the existence of a certain degree of heterogeneity (ΔCFe =18 at. % of Fe for Nb10 and 14 at. % of Fe for Nb5). Other parameter which enables to follow the compositional evolution is the most probable Fe content in the particles, CFe, (shown in figure 9b). This parameter rapidly decreases with the milling time close to the nominal composition for both alloys. After 20 h milling, CFe= 89 and 95 ± 2 at. % for Nb10 and Nb5, respectively. An increase with respect to the nominal compositions (88 and 94 at. % for Nb10 and Nb5, respectively), although into the experimental error, would not be surprising and could be expected due to Fe contamination from milling media. The evolution of CFe was also obtained from the simple simulation described above (also shown in figure 9b, using the Intermetallics. Vol. 16. Núm. 9. 2008. Pag. 1073-1082 http://dx.doi.org/10.1016/j.intermet.2008.06.006 16  For milling times t ≥50 h, the only phase present in the samples is a bcc Fe(Nb,B) supersaturated solid solution for both alloys. After 400 h, only for the alloy with the highest Nb content, the presence of amorphous phase is evidenced by an amorphous halo observed by XRD, as well as the detection of an exothermic process in DSC ascribed to crystallization. This amorphous phase is stable up to 850 K.  For the alloy with the lowest Nb content, after 50 h milling no new Fe atomic sites appear, being the more important contribution at 33 T. However, for the alloy with the highest Nb content, the Fe atomic environment continuously evolves increasing the fraction of paramagnetic environments.  The Curie transition of amorphous phase in the alloy with the highest Nb content is below room temperature (~250 K). This phase, detected by magnetization measurement for milling times t ≥200 h, was also detected by X-ray diffraction and Mössbauer spectroscopy techniques. Acknowledgments This work was supported by the Spanish Government and EU FEDER (Project MAT 2004-04618) and by the PAI of the Regional Government of Andalucía (Project P06-FQM-01823). J.J.I. acknowledges a fellowship from the Spanish Ministry of Education and Science. J.S.B. acknowledges a research contract from the Regional Government. References [1] C. Suryanarayana, Prog. Mater. 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Huang, B. Lu, J. All. Comp. 413 (2006) 85-89. [13] J.R. Stephens, W.R. Witzke, J. Less-Common Met. 48 (1976) 285-308. [14] J.J. Ipus, J.S. Blázquez, V. Franco, A. Conde, Intermetallics 15 (2007) 11321138. [15] J. M. Greneche, A. Slawska-Waniewska, J. Magn. Magn. Mater. 215-216 (2000) 264-267. [16] B.Q. Zhang, L. Lu, M.O. Mai, Phys. B 325 (2003) 120-129. [17] E. Jartych, D. Oleszak and J.K. Zurawicz, Hyp. Int. 136 (2001) 25-33. [18] Z. Caamaño, G. Pérez, L.E. Zamora, S. Suriñach, J.S. Muñoz, M.D. Baró, J. Non-cryst. Solids. 287 (2001) 15-19. Intermetallics. Vol. 16. Núm. 9. 2008. Pag. 1073-1082 http://dx.doi.org/10.1016/j.intermet.2008.06.006 18 [19] J.Y. Yang, T.J. Zhang, K. Cui, X.G. Li, J. Zhang, J. All. Comp. 242 (1996) 153156. [20] J.J. Suñol, A. González, J. Saurina, Ll. Escoda, P. Bruna, Mater. Sci. Eng. A 375-377 (2004) 874-880 [21] J.S. Blázquez, V. Franco, C.F. Conde, A. Conde, Intermetallics 15 (2007) 13511360. [22] J. Torrens-Serra, J. Rodríguez-Viejo, M.T. Clavaguera-Mora, J. Non-Cryst. Sol. 353 (2007) 842-844. [23] V. Franco, C.F. Conde, A. Conde, J. Magn. Magn. Mater. 303 (1999) 60-62. [24] D. Oleszak, P.H. Shingu, J. Appl. Phys. 79 (1996) 2975-2980. [25] D. Henderson, J. Non-cryst. Solids 30 (1979) 301-315. [26] A. Makino, K. Suzuki, A. Inoue, T. Masumoto, Mater. Sci. Eng. A 179-180 (1994) 127-131. Intermetallics. Vol. 16. Núm. 9. 2008. Pag. 1073-1082 http://dx.doi.org/10.1016/j.intermet.2008.06.006 19 Figure captions Figure 1. SEM images of both alloys after selected milling times. Figure 2. a) SE image, b) BSE image of a typical particle of Nb10 alloy after 2 h milling and c) EDX spectra of selected points indicated in a). Figure 3. Experimental and simulated data of the average powder particles size as a function of the milling time for both alloys. Figure 4. Histograms of the percentage of Fe content in total Fe+Nb content for both alloys at selected milling times. Figure 5. Cr content as a function of milling time for both alloys. The slopes of the different fitted lines are also indicated. Figure 6. XRD patterns of both alloys after selected milling times. Figure 7. Mössbauer spectra and hyperfine magnetic field distributions for Nb10 alloy after selected milling times. Figure 8. Mössbauer spectra and hyperfine magnetic field distributions for Nb5 alloy after selected milling times. Figure 9. Experimental and simulated values of a) ΔCFe and b) CFe as a function of the milling time (experimental) and iteration steps (simulation). Horizontal lines at 88 and 94 at.% Fe correspond to the nominal values of Nb10 and Nb5, respectively. Iteration steps axes have been conveniently rescaled to show the agreement with the experimental data. Figure 10. a) Lattice parameter, b) minimum crystal size and c) maximum microstrain as a function of milling time for both alloys. Figure 11. Area fraction of the different Mössbauer contributions to the total fitting as a function of milling time for both alloys. Intermetallics. Vol. 16. Núm. 9. 2008. Pag. 1073-1082 http://dx.doi.org/10.1016/j.intermet.2008.06.006 20 Figure 12. Magnetization curves as a function of temperature for Nb10 alloy after 100, 200, 300 and 400 h milling. Figure 13. DSC scans at 40 K/min for both alloys after selected milling times. Symbols are superimposed over the curves to distinguish them. Figure 14. XRD patterns for Nb5 alloy after different heated temperatures. Figure 15. XRD patterns for Nb10 alloy after different heated temperatures. Figure 16. Mössbauer spectra and hyperfine magnetic field distribution for Nb5 alloy after annealed treatment. Figure 17. Mössbauer spectra and hyperfine magnetic field distribution for Nb10 alloy after annealed treatment. Intermetallics. Vol. 16. Núm. 9. 2008. Pag. 1073-1082 http://dx.doi.org/10.1016/j.intermet.2008.06.006 21 Figure 1 2 h 500  m 20 h 500  m 200  m 100 h 100  m 400 h Nb10 500  m 2 h 20 h 500  m 100  m 100 h 50  m 400 h Nb5 Intermetallics. Vol. 16. Núm. 9. 2008. Pag. 1073-1082 http://dx.doi.org/10.1016/j.intermet.2008.06.006 22 Figure 2 50 m A B C B Fe. Nb 0 2 keV c ) a ) b) Fe. Nb Intermetallics. Vol. 16. Núm. 9. 2008. Pag. 1073-1082 http://dx.doi.org/10.1016/j.intermet.2008.06.006 23 Figure 3 0 100 200 300 400 0 100 200 300 Nb10 Nb5 iteration step (x103) <d> [m] milling time [h] 0 5 10 15 20 25 30 35 simulated Intermetallics. Vol. 16. Núm. 9. 2008. Pag. 1073-1082 http://dx.doi.org/10.1016/j.intermet.2008.06.006 24 Figure 4 4 8 12 2 4 6 2 4 2 4 6 Nb10 0.2 0.4 0.6 0.8 1.0 2 4 % Fe 4 8 12 Nb5 2 h 4 8 10 h 2 4 6 20 h 4 8 12 50 h 0.2 0.4 0.6 0.8 1.0 4 8 400 h % Fe Intermetallics. Vol. 16. Núm. 9. 2008. Pag. 1073-1082 http://dx.doi.org/10.1016/j.intermet.2008.06.006 25 Figure 5 0 100 200 300 400 0.0 0.6 1.2 1.8 Cr [%] milling time [h] N5B N10B (7 ±2)*10-4 at. % Cr/h (3.6±0.3)*10-3 at. % Cr/h Intermetallics. Vol. 16. Núm. 9. 2008. Pag. 1073-1082 http://dx.doi.org/10.1016/j.intermet.2008.06.006 32 Figure 12 100 200 300 400 0 30 60 90 120 400 h 300 h 200 h 100 h M 0 [emu/g] T [K] Intermetallics. Vol. 16. Núm. 9. 2008. Pag. 1073-1082 http://dx.doi.org/10.1016/j.intermet.2008.06.006 33 Figure 13 400 500 600 700 800 900 1000 temperature [K] 400 h 400 h 50 h 50 h 20 h 20 h 5 h 5 h Nb10 Nb5 0.2 W/g (exo) dH/dt Intermetallics. Vol. 16. Núm. 9. 2008. Pag. 1073-1082 http://dx.doi.org/10.1016/j.intermet.2008.06.006 34 Figure 14 40 50 60 600 K 800 K as-milled 2  [degree] 1000 K 500 1000 0,286 0,288 T [K] a [nm] Intermetallics. Vol. 16. Núm. 9. 2008. Pag. 1073-1082 http://dx.doi.org/10.1016/j.intermet.2008.06.006 35 Figure 15 40 50 60 600 K 800 K as-milled 2  [degree] 1000 K 500 1000 0 50 100 T [K] X C [%] Intermetallics. Vol. 16. Núm. 9. 2008. Pag. 1073-1082 http://dx.doi.org/10.1016/j.intermet.2008.06.006 36 Figure 16 0,1 0,2 0,08 0,16 0,24 0,2 0,4 -8 -6 -4 -2 0 2 4 6 8 0 5 10 15 20 25 30 35 0,2 0,4 0,6 as-milled 600 K relative transmition probability 800 K velocty [mm/s] B hyp [T] 1000 K Intermetallics. Vol. 16. Núm. 9. 2008. Pag. 1073-1082 http://dx.doi.org/10.1016/j.intermet.2008.06.006 37 Figure 17 0,1 0,2 -8 -6 -4 -2 0 2 4 6 8 0,1 0,2 0,03 0,06 0 5 10 15 20 25 30 35 0,1 as-milled probability velocity [mm/s] relative transmition 600 K 800 K B hyp [T] 1000 K