Thin films having giant magnetostriction at low fields and an ordering temperature high enough for the applications as magnetostrictive microactuators have been the goal of magnetostriction researches. Various attempts have been focused mainly on amorphous Terfenol and Terfenol-D (TbDyFe2) alloys. However, the magnetostriction of these amorphous alloys has been found to be one order of magnitude lower than that of its crystalline counterpart. Magnetic investigations have shown that such amorphous R-Fe alloys could never offer promising magnetostrictive performances, due to the distribution of the signs of the Fe-Fe interactions and the Fe- and R-sperimagnetic characters. Another great disadvantage of the amorphous R-Fe thin films, from the point of view of application in microsystems is the comparatively low Curie temperature (To ~< 400 K for amorphous Terfenol-D films). The polycrystalline Terfenol-D film processes a room temperature magnetostriction of ~s ~ 1500x 10 -6 and Tc = 650K. This film is, however, not suitable for micro-mechanical applications due to its comparatively high coercive fields. Material designs always require a combination between the advantages of the crystalline films (giant magnetostriction and high Tc) and the good magnetic softness of the amorphous films (low coercivity). In this context, the nanocrystalline structure is expected to be able to realise this objective (Schatz et al. 1993, 1994, Williams et al. 1994, Miyazaki et al. 1997, Wada et al. 1997c,d, Ried et al. 1998, Winzek et al. 1999, Farber and Kronmtiller 2000a,b). In these alloys, the dimensions of crystallites are sufficiently large enough to allow the exchange coupling to be effective but small enough to prevent the appearance of any macroscopic anisotropy energy. Thus one may consider the material as an isotropic ferromagnet, in which the magnetostriction is expected to be the same but the magnetocrystalline anisotropy has disappeared. By
GIANT MAGNETOSTRICTION IN R - T THIN FILMS 47
combining the annealing temperature and time, the magnetostriction value can be doubled.
Indeed, the best magnetic properties were obtained on the 'cold grown' (Tb0.3Dy0.7)l xFex (with x ~ 0 . 3 ) films annealed at TA=600°C for 10 minutes with a grain size around 10 nm: a saturation magnetostriction )~ = 860 x 10 -6 and #0Hc = 120 mT (Ried et al. 1998).
In addition, high-temperature annealing is not desirable because of the formation of iron- rich phases, like RFe3 and R6Fe23, with even less magnetostriction and higher magnetic coercivity.
According to the random anisotropy model (Herzer 1990, Hofmann et al. 1992), a further reduction of the crystal size will reduce the coercive field if the exchange length is larger than the average grain size. Hence, technologies which at the same time enhance grain nucleation and limit grain size must be applied to achieve a fine nanocrystalline structure. For this purpose Zr and Mo have been chosen as additives in order to enhance grain nucleation and to limit grain growth, respectively. Practically, the influence of the Zr and Mo additives on the crystallisation and the magnetic properties was studied by Winzek et al. (1999). Films with 3 at% Zr, crystallised at 973 K for 10 minutes, showed )~lJ = 430 x 10 .6 at #0H = 1.0 T and kt0Hc = 120 mT, which was a large improvement compared to the starting alloys: ;tll= 230x 10 .6 and/~0Hc = 300 mT. These additives were also thought to enhance the growth of the RFe2 grains and to hinder the formation of RFe3 ones. It is assumed that the latter was responsible for the high coercivity values above 150 mT.
In general, it should be noted that a reduction of the average grain size of the cubic Laves phases below 10 nm and, therefore, the coercivity values below 100 mT could not be achieved in single layer films. Grain growth, however, enabled one to control nanometer- scaled multilayers with interlayers of Nb (Fischer et al. 1999, Winzek et al. 1999). These authors fabricated a multilayer system containing TbDyFe + Zr with a thickness of 5 nm separated by Nb-layers with an average thickness of 0.25 nm. The introduction of a partial Nb layer in this multilayer structure was thought to cause reduced dimensions and increase the amount of inner surfaces, while Zr is assumed to play a dominant role in forming nucleation centers for the nanograins. After 10 min. annealing at temperatures from 873 K to 973 K the magnetic phase transition temperature increased from Tc = 333 K to 592 K accompanied by an increase of the magnetostriction from 265 x 10 -6 to 520 x 10 6, while the coercive fields increased from kt0Hc = 5 to 75 mT, which lies distinctively below 100mT. Research is in progress to prevent not only the growth of crystallites but also aging effects.
In the amorphous state, however, it is preferable to replace the iron by cobalt because the amorphous alloys near to the composition a-RCo2 (named a-TercoN6el, where N6el identifies the laboratory where the magnetostriction of this composition was studied first; Duc 2001) have higher ordering temperatures and higher magnetostrictions than the equivalent iron-based alloys. In fact, the magnetostriction has been optimised in a series of thin films of the type a-(Tb,DY)2(Fe,Co)n: the material Tb(Fe0.4sCo0.s5)2.~
showed a high record magnetoelasticity o f ~' = 63.5 Mpa and )~v, 2= 1020× 10 -6, which is almost fully developed at low fields. The giant magnetostrictions obtained in this series of alloys has been explained in terms of an increase in the 3d-3d exchange
48 NGUYEN HUU DUC
coupling within (Fe,Co) sublattice. Recently, Danh et al. (2000) and Duc et al.
(2000b) have reported a value of the magnetoelastic susceptibility Zzl] = 1.8 × 10 -2 T -1 (or ZbN = 1100MPa) for a-Tb(Fe0.55Co0.45)1.5. (denoted a-TerfecoHan, where Han stands for Hanoi, the capital where studies on this composition have been carried out; Duc 2001).
Values of magnetostriction and magnetostrictive susceptibility which are interesting from the applications viewpoint are summarised in table 4. It is clearly seen that this magnetostrictive susceptibility is comparable to that of the magnetostrictive multilayers.
Even better performances were obtained from magnetostrictive mulfilayers of the novel type of "Spring Magnets", where the saturation field of the magnetostrictive, amorphous Tb-FeCo is lowered by increasing the average magnetisation through exchange coupling with the soft-magnetic Fe-Co layers. In addition to this, increasing magnetisation by closing the cone angle in a sperimagnefic structure is also an efficient means of increasing low field magnetostrictive susceptibility, since this reduces the saturation field, and the magnetoelastic coupling is strongly correlated to the mean value of the magnetic moments correlation function. This process can be achieved by increasing the molecular field. Experimentally, the molecular field of R-T may be strengthening in the neighbourhood of the strongly magnetic FeCo layer. In this case, layers must be sufficiently thin, then, a noticeable volume of the R-T layer may be submitted to the large molecular field. In the last two years, important progress has been achieved in magnetostrictive multilayer research. Ludwig and Quandt (2000) reported the possibility of controlling the orientation of the magnetic easy axis by magnetic annealing and, thus, were able to enhance the magnetostriction in the desired direction. An induced uniaxial anisotropic multilayer TbFe/Fe can also be created by deposition under a magnetic polarisation field. High amplitude flexural and torsional oscillation modes were observed for these films (Le Gall et al. 2000). Duc et al. (2001a) reported a large magnetostricitive susceptibility and discussed the so-called working point for the magnetostrictive multilayers in microactuator devices.
In attempts to prepare the magnetostrictive multilayers consisting of nanocrystalline magnetostrictive layers with soft magnetic interlayers, Farber and Kromntiller (2000b) have studied TbDyFe/Finemet multilayers (Finemet is a nanocrystalline FeSiBNbCu soft magnetic alloy). The contribution of the individual layers to the magnetoelastic couplings was deduced from their magnetoelastic data, where an opposite contribution of the Finemet (b v,2= 15) with respect to that of the TbDyFe (b v,2 =-18) was found.
In the TbDyFe/Fe multilayers, however, the magnetoelastic coupling has the same sign:
b v, 2 = -22 and - 6 for TbDyFe and Fe layers, respectively (Farber and Kronmfiller 2000b).
Attempts to reduce /20Hc also imply the possibility of shifting the working point to lower fields. For this purpose, Quandt and Ludwig (1999) have prepared the TbFe/FeCoBSi multilayers. It was shown that the FeCoBSi layers improved the magnetic softness of the multilayer. Recently, Duc et al. (2001b) have succeeded in preparing Tb(Fe0.55Co0.45)1.5/(Y0.2Fe0.8) (i.e., a-TerfecoHardn-YFe, Duc 2001) multilayers with
#0Hc=0.5mT. Initially, this multilayer consists of the amorphous TbFeCo and not- well crystallised FeCo layers. In this state, a soft magnetic and magnetostrictive character with a coercivity /20He = 3 mT and a parallel magnetostfictive susceptibility
GIANT MAGNETOSTRICTION IN R T THIN FILMS 49 Table 4
Comparison of the magnetoelastic data for magnetostrictive bulk and thin-film materials
Materials b v,z (MPa) )~,2 (10 6) ObJOB (Mpa/T) References
Bulk crystalline
Terfenol-D Tbo.27Dyo.73 Fe 2 -101 2400 568 [1]
Single layer films
a-TbFe 2 19.4 321 20 [2]
a-Tbo.27 Dyo.73 Fe 2 - 17.2 300 50 [3]
n-Tbo.27 Dyo.73 Fe 2 49.0 800 [4]
a-TbCo 2 -24.5 400 155 [5]
a-Tbo.27Dyo.7~ Co 2 -15.1 260 190 [3]
a-Tb(Feo.45Coo.ss)2 -63.5 1040 300 [6,7]
a-Tb(Feo.55 Coo,45 )1.5 65.9 1080 1100 [8]
a-Tbo.27 Dyo.73 (Feo.~5 Coo.55 )2 -20.15 330 430 [6,7]
a-SmFel 6 25.9 -380 [9]
a-(SmFe2)99.26Bo,74 45.6 -670 [10]
a-SmCo z 11.0 -161 40 [7]
a-Sm(Feo.ss Coo.42)1.54 27.4 -320 76 [ 11 ]
Multilayers
Tbo.4Feo.6/Fe 20 300 [12]
Tbo.27 Dyo.73 FejFe - 1 2 650 [13]
Tbo.27 Dyo.73 Fee/Finemet 300 [13]
Tb(Feo 55 Coo.45 )1.5/Fe - 3 9 3040 [14]
Tbo.4Feo.6/Fe0.5 Coo 5 5 8 410 1000 [12]
Wbo.37 F%. 63/][:~%.65 C00,35 --31.1 600 4000 [12]
Tbo.27F%.T3/Feo.75 C00.25 - 2 7 348 4800 [12]
Tbo.i8 Feo.82/Feo.75 Coo.25 44.5 890 [12]
Yb(Feo.55 Coo.4s)l.s/Feo.85 Coo.15 -32 530 7850 [15]
Sandwich system
Ndo.25 Coo.75/Tbo.28 Coo.72/Ndo.25 Coo.75 -15.2 248 560 [8]
Tbo.zs Coo.72/Ndo.25 Coo.75/Tbo.2s Co0,72 16.5 270 117 [7,8,16]
References [1] Clark (1980) [2] Hayashi et al. (1993) [3] Duc et al. (2000a) [4] Ried et al. (1998) [5] Betz et al. (1999) [6] Duc et al. (1996)
[7] Betz (1997) [8] Duc et al. (2000b) [9] Honda et al. (1994) [10] Kim (1993)
[11] S. David, unpublished data
[12] Quandt and Ludwig (1997) [13] Farber and Krorlmfiller (2000b)
[14] Duc et al. (2001a) [15] Duc et al. (2001b) [16] Givord et al. (1996)
X~.II, max = 3.8 X 10 2 T-1 h a d b e e n a c h i e v e d . T h i s m a g n e t o s t r i c t i v e s o f t n e s s w a s s t r o n g l y i m p r o v e d b y h e a t t r e a t m e n t s : g o H c = 0 . 5 m T a n d X~ll = 1 3 x 10 -2 T -1 i n a f i e l d o f 1.8 m T . T h e s e n o v e l p r o p e r t i e s a r e a s s o c i a t e d w i t h t h e d e v e l o p m e n t o f Y F e n a n o s t r u c m r e l a y e r s .
50 NGUYEN HUU DUC
Aside from technical considerations, we must stress that studies on magnetoelastic effects in thin films also present a fimdamental interest. The magnetoelastic coupling is quite dependent on the symmetry of the material. At the surface of any alloy, the symmetry is broken since there are no neighbouring atoms. The surface contribution to the magnetoelastic coupling, thus, must be considered. This contribution is usually negligible in bulk materials, but it is no longer true for thin films where surface effects are essential. In particular, for magnetostrictive multilayers the nature of the interfaces is critical; their contributions to the magnetic and magnetostrictive properties needs to be considered. In this context, a general consideration of the magnetoelasticity in the heterogeneous magnetic materials (including nanoscale thin films, multilayers, superlattices, nanocrystalline magnetic alloys and magnetic granular films) may be useful. As noted above the contributions of interfaces to the magnetic and magnetostrictive properties have been experimentally verified in heterogeneous magnetic systems. However, the magnetoelastic interactions as well as magnetic properties of these materials are complex and need further experimental as well as theoretical studies. At the moment, only phenomenological models can describe their properties.
Finally, significant improvements in the field of giant magnetostrictions have allowed the design and test of some prototypes of microactuators and motors taking advantage of a wireless magnetic excitation.
Acknowledgements
The support of the Vietnam National University, Hanoi under projet QG.99.08 is useful for the completion of this work. The collaborations with Dr. D. Givord, Dr. K. Mackay, Dr. J. Betz and Dr. E. du Tr6molet de Lacheisserie at the Laboratoire Louis N6el, CNRS, Grenoble (France) are very useful, not only to become familiar with the field, but also to expand my research in this area. Special thanks are due to J. Betz for permission to include his work prior to this overview. I am particularly grateful to Prof. Dr.
T.D. Hien, Prof. Dr. N.H. Luong, Dr. N.H. Chau and co-workers, who have started together with me to build equipment for measuring the thermal expansion and magnetostriction using the capacitance method, and to develop research on magnetovolume effects and magnetostriction of the lanthanide-transition metal intermetallics at the Cryogenic Laboratory in Hanoi since 1981. Dr. N.T. Hien and Prof. Dr. N.P. Thuy have greatly encouraged me to write this chapter.
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Handbook on the Physics and Chemistry of Rare Earths VoI. 32
edited by K.A. Gschneidner, Jr., L. Eyring and G.H. Lander
© 2001 Elsevier Science B.V. All rights reserved
Chapter 206
;tSR STUDIES OF RARE-EARTH AND ACTINIDE MAGNETIC MATERIALS*
G . M . K A L V I U S
Physics Department, Technical University Munich, 85747 Garching, Germany
D . R . N O A K E S
Department o f Physics, Virginia State University, Petersburg, VA 23806, USA
O. H A R T M A N N
Department o f Physics, University o f Uppsala, 75121 Uppsala, Sweden
C o ~ e ~ s List of symbols List o f acronyms 1. Introduction 2. The ~tSR technique
2.1. General aspects 2.2. Properties of the muon 2.3. Muon generation and decay 2.4. Muon implantation 2.5. Muon beam characteristics 2.6. The recording of [xSR spectra 2.7. Special applications
2.7.1. Extreme temperatures 2.7.2. High magnetic fields 2.7.3. High pressure 2.7.4. Small samples 2.7.5. Low-energy muons 3. ~tSR magnetic response
3.1. The local field at the muon site 0 3.2. ~tSR spectroscopy of dia- and
paramagnets
3.2.1. Transverse field measurements (Knight shift)
3.2.2. Zero field measurements
3.2.3. Longitudinal field measurements 102 58
59 3.3. poSR spectroscopy of ordered magnets 105
60 3.4. Double relaxation 110
62 3.5. Data analysis 111
62 3.6. Determination o f the muon site 114 66 3.7. Modeling of the internal field 116 67 3.8. Muon diffusion and magnetism (site
69 averaging) 121
71 4. Elemental metals 122
76 4.1. Anisotropic spin fluctuations 124
81 4.2. Gadolinium 126
81 4.2.1. Paramagnetic region and critical
82 behavior 127
83 4.2.2. Ferromagnetic region 129
85 4.2.3. High-pressure measurements 131 86 4.3. Holmium, dysprosium and erbium 132
88 4.3.1. Ambient pressure 132
88 4.3.1.1. Holmium 132
4.3.1.2. Dysprosium 134
92 4.3.1.3. Erbium 136
4.3.2. High-pressure measurements 138
93 5. Intermetallics 140
97 5.1. Singlet ground state systems 140
* This work has been supported by the BMBF, Federal Republic of Germany under contract 03-KA4-TU1-9, by US DOD grants F-49620-97-0297 (AFOSR) and -0532 (BMDO), and by the Swedish National Research Council.
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