Nov 25, 2016 - Supplementary material: Epitaxial strain-engineered self-assembly of magnetic nanostructures in FeRh thin lms. Ralf Witte, Robert Kruk, Alan ...
Supplementary material: Epitaxial strain-engineered self-assembly of magnetic nanostructures in FeRh thin lms Ralf Witte, Robert Kruk, Alan Molinari, Di Wang, Sabine Schlabach, Richard A. Brand, Virgil Provenzano, and Horst Hahn November 25, 2016
This supplementary material contains data on the in situ growth control with re ection high energy electron diraction as well as reciprocal space maps and transmission electron microscopy investigations performed on the W-V buer layers and the FeRh lms. Further information is provided on the obtained Mössbauer hyper ne parameters. Finally, results of a preliminary post-annealing study are presented.
1 In situ growth monitoring The epitaxial growth of the FeRh thin lms is monitored with re ection high energy electron diraction (RHEED). RHEED patterns of the three samples compared in the main article are shown in Fig. S1. The top row presents the diraction patterns obtained from the freshly prepared W1−y Vy layers, while the bottom row presents the nal patterns corresponding to the FeRh surfaces. The electron beam points in the [100] direction of the bcc lattice.
The streak pattern obtained in all three cases for the buer layers
as well as for the nal heterostructure con rms epitaxial growth.
Although a direct
comparison between dierent samples is di cult, due to diering experimental settings, one can observe that the diraction patterns coming from the FeRh lm grown on the intermediate buer layer with
y = 0.57
seems to be more broadened compared to the
two others. This can well be due to the distribution of in-plane lattice constants, which are a result of the two-phase nanostructure.
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y=0.57
y=0.00
Fe-Rh
W-V
y=0.75
Figure S 1: RHEED patterns for W1−y Vy (top-row) and FeRh (bottom-row) layers. The azimuthal angle is along the
y = 0.75, 0.57
and 15 kV for
h100ibcc direction. y = 0.00.
W
1 -y
V
Electron energy is 12 kV for
/ F e R h y
In te n s ity ( a . u .)
y
0 .7 5
0 .5 7 0 .0 0
2
4
6
8
1 0
2 T h e ta (° ) Figure S 2: XRR patterns measured for the three samples presented in the main paper. The thick solid line is the raw data, while the thin colored line represents the best t to the data.
2 Additional structural investigations
2.1 X-ray re ectometry The thickness
t
and roughness
R
of the individual layers was investigated with X-ray
re ectometry (XRR). Fig. S2 displays the data and the obtained t curves, which were calculated using the software package
Leptos.
2
The obtained values are summarized in
Tab. 1 of the main manuscript. Already the measured data provide evidence for the high quality and low roughness of the thin lms, given the pronounced presence of the Kiessig fringes. The shorter period oscillation corresponds to the thicker W1−y Vy buer layers, while the longer period oscillation is attributed to the FeRh layers. The obtained low roughness values
b c c
(Å
-1
)
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0 .6 5 0 .6 4 0 .6 3 0 .6 0 .6 0 .6 0 .6
4 4 2
0 .5 7
2 3 0 0
5 4
2 6 0 0
3
1 3 5 4
2 0
0 .4 4
q
||
0 .4 6
< 1 1 0 >
0 .4 8 b c c
(Å
0 .0 0
-1
)
Figure S 3: RSMs measured on the (112) re ection of the W1−y Vy buer layers. As described in the main text, an intensity line-scan in
qk h110ibcc
direction, across
the (220)/(202) re ection pair, was tted with a two peak model for the determination of the in-plane lattice constants of the
Cmcm
3
phase in sample
y = 0.57.
The data is
presented in Fig. 4. The raw data was rst smoothed, then a t with two Gaussian peaks
In te n s ity ( c p s )
was performed yielding the lattice constant values given in the main paper.
D a ta D a ta P e a k P e a k F it s u
7 .0
s m o o th e d 1 2 m
6 .5
6 .0
0 .4 3
0 .4 5
q
Figure S 4:
q -scan
in
qk h110ibcc
||
0 .4 7
< 1 1 0 > (Å
0 .4 9 -1
direction, tted with two gaussian peaks in order to
determine the in-plane lattice constants for the sample.
The
q⊥
0 .5 1
)
Cmcm phase
in the
y = 0.57
value was kept xed at the position corresponding to the
out-of-plane lattice constant. Peak one corresponds to the (220) and peak 2 to the (202) re ection respectively.
2.3 Transmission electron microscopy As stated in the main paper, it was also possible to nd areas in the lamella, where the two structural phase can be identi ed individually by high resolution transmission electron microscopy (HRTEM). This is presented in Fig. S5(a) and (b), where a detailed view of the bct and
Cmcm
structures is presented.
The out-of-plane direction of the
lms is oriented horizontally. A HRTEM investigation of the buer layer was performed in order to exclude any lateral variation of the buer layer lattice constant as possible cause of the phase separation. A high resolution micrograph is presented in Fig. S6(a), while (b) gives the result of a geometrical phase analysis (GPA) [1], which allows to compare the local lattice parameter to a reference value. Here the in-plane strain
xx
It thus represents a map of the strain state of the lm.
is presented, and the relevant reference area (average zero
strain value) is depicted by the rectangle in Fig. S6(a). The hot spots in the strain map might indicate the presence of mis t dislocations in the strained W-V layer. However, no lateral variation of the strain is observed on a scale larger than these single dislocations. Hence local structural inhomogeneity of the buer layer on the length scale of the structural features in the FeRh layer is safely ruled out as explanation for the observed phase-separation.
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[010]
[110]
(a)
(b)
bct
Cmcm
[100]
[001]
Figure S 5: Detailed view of the FeRh layer by HRTEM allowing to individually identify the bct phase shown in (a) and the
Cmcm
phase presented in (b). Both mi-
crographs are oriented with the out-of-plane direction of the lm in horizontal direction.
WV [001]
(a)
2 nm
WV [110]
WV [001]
(b)
�xx
WV [110] -0.5
�xx
0.5
Figure S 6: (a) HRTEM micrograph of W1−y Vy buer layer with rical phase analysis (GPA) of the in-plane
xx
y = 0.57.
(b) Geomet-
strain, proving lateral homo-
geneity of the in-plane lattice constants. The reference area is indicated by the white rectangle.
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3 Conversion electron Mössbauer spectroscopy Tab. S1 summarizes the hyper ne parameters obtained from the room temperature (RT) conversion electron Mössbauer spectroscopy measurements shown in Fig. 2 in the main paper. The data was tted using the WinNORMOS software package by R. A. Brand. The area ratio between second and third line A23 is xed to a value of 4, in order to avoid nonphysical values. However from the raw data it is already apparent that the second line is very pronounced, which is evidence for an in-plane orientation of the magnetization [4].
Table S 1: Mössbauer hyper ne parameters obtained from tting the CEMS spectra. Isomer shift width
Γ
δ
relative to bcc-Fe at RT, quadrupole splitting
are all given in (mm/s) while the magnetic hyper ne
∆EQ and line eld BHF is in
(T). A23 is the area ratio of second and third line. Area denotes the relative spectral fraction of the subcomponents. * labels xed parameters.
y
Phase
δ
∆EQ
0.00
Cmcm Cmcm
0.063(1)
0.17(1)
-
-
0.35(1)
1
0.064(1)
-
-
-
0.52(1)
0.59(5)
bct
0.016(3)
-
26.5(35)
4*
0.60(1)
0.41(5)
Cmcm
0.072(7)
-
-
-
1.11(2)
0.23(5)
bct
0.025(2)
-
27.2(20)
4*
0.40(1)
0.77(5)
0.57
0.75
BHF
A23
Γ
Area
4 Post-annealing study Motivated by the potential applications proposed in the main paper, we performed a preliminary post-annealing study on a sample, which showed signs of the antiferromagnetic (AF) to ferromagnetic (FM) transition already in the as-prepared state.
The partial
presence of AF order, in contrast to the samples presented in the main paper, is owing to partial B2 ordering and a chemical composition where AF order can be stabilized. The AF ground state is easily destroyed by either o-stoichiometry or anti-site defects [2]. The multilayer sample with
y = 0.50
◦
was post-annealed under vacuum at 500 C
for 90 min. Fig. S7 shows the obtained results: (a) displays HRXRD patterns and (b) temperature dependent magnetization measurements of the as prepared and the postannealed state.
The structural investigation shows the familiar presence of the two
phases. The structure is not signi cantly changed by the post-annealing procedure other than a small increase in area of the super-structure re ection, due to better B2 ordering. This indicates that the strain-induced nanostructure oers considerable thermal stability. However, the more sensitive probe for the increase of B2-ordering is the observed magnetic behavior. The magnetization curves
M (T ) for the as prepared sample show signs of
the AF-FM transition, given by an increase of magnetization with temperature. Yet, the transition appears smeared out and the sample shows some remaining magnetization at low temperature, which is due to the imperfect B2-ordering in the as prepared state eectively suppressing the subtle AF order in some areas of the sample. The post-annealing
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Intensity baxuxµ
as prepared annealed
Data as prepared annealed
55 6K 65 7K 75
2 Theta b°µ as prepared annealed
15K
Relative Emission
baµ
25 3K
Magnetization bx1K3 AAmµ
bcµ
2h
1KK µKH = 5K mT
5K
4h µKH = 4 T
bbµ
K 1KK
2KK
3KK
4KK
z9 z6 z3
Temperature bKµ
K
3
6
9
Velocity bmmAsµ
Figure S 7: In uence of post-annealing. (a) HRXRD patterns for a sample with
y = 0.50, ◦
in the as prepared and annealed state, after 90 min post-annealing at 500 C. (b) Temperature dependent magnetization curves measured between 100 400 K for the as prepared and annealed state in a constant magnetic eld
µ0 H = 50 mT
and 0.4 T, respectively.
(c) RT-CEMS spectra of the two
states, each tted with a single line (dashed) and a magnetic sextet (dotted). The arrows indicate the position of the second and fth line, see text for details.
step removes this FM contribution nearly entirely, leading to the sharp increase of the magnetization above room temperature typical for the AF-FM transition [3]. A spectroscopic view with CEMS on the two states is presented in Fig. S7(c). The as prepared sample shows a similar spectrum as in the case of the FM intermediate sample displayed in the center of Fig. 2 in the main paper, with at least two components, a paramagnetic singlet and a magnetic sextet. The latter is very broad indicating structural but possibly also magnetic disorder. The spectra are not unambiguous. A tentative t is shown with a singlet and a magnetic sextet including Gaussian broadening. However the most important change upon post-annealing is visible with the bare eye: the change in relative line intensities. The relative line intensities of the sextets drastically change from one spectrum to the other, indicating a variation of preferred magnetic orientation. For the as prepared state
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lines two and ve are present (counting left to right and denoted by the arrows) with a ratio to the inner two lines of about 3.4. This indicates a spin arrangement between nearly random (given by a ratio of 3) and in-plane (given by a ration ratio of 4) [4]. However in the spectrum of the post-annealed sample the lines two and ve are hardly observable ( tting with again a broadened sextet yields a ratio of 0.6), pointing to a spin reorientation towards an out-of-plane orientation parallel to the
γ -ray
(ratio of 0).
Bordel et al. have shown that FeRh lms with a slight tetragonal distortion (c/a=0.985) show this kind of preferred orientation of the spin arrangement in the AF phase [5]. As a tetragonal distortion with
c/a < 1
is expected in the B2 phase of the present sample,
the observed spin reorientation represents also spectroscopic evidence for the occurrence of the AF state and eectiveness of the annealing procedure. The spectral area ratios of the magnetic and paramagnetic subspectra remain about constant (50%), supporting the claim based on HRXRD results, that the nanostructure is basically unharmed by the thermal treatment. Hence, the extraordinary thermal stability of the nanostructures allows for further post-treatment of the lms, with the aim of achieving a perfect B2 ordering and the complete AF-FM transition. This result provides solid and promising evidence for the proposed potential applications of the magnetic nanostructures.
References [1] M. H¸tch, E. Snoeck, and R. Kilaas, Quantitative measurement of displacement and strain elds from HREM micrographs, Ultramicroscopy, vol. 74, pp. 131146, (1998). [2] J. B. Staunton, R. Banerjee, M. dos Santos Dias, A. Deak, and L. Szunyogh, Fluctuating local moments, itinerant electrons, and the magnetocaloric eect: Compositional hypersensitivity of FeRh, Phys. Rev. B, vol. 89, p. 054427, (2014). [3] S. Maat, J.-U. Thiele, and E. E. Fullerton, Temperature and eld hysteresis of the antiferromagnetic-to-ferromagnetic phase transition in epitaxial FeRh lms, Phys. Rev. B, vol. 72, p. 214432, (2005).
[4] P. Gütlich, E. Bill, and A. X. Trautwein, Mössbauer Spectroscopy and Transition Metal Chemistry. Berlin, Heidelberg: Springer Berlin Heidelberg, (2011).
[5] C. Bordel, J. Juraszek, D. W. Cooke, C. Baldasseroni, S. Mankovsky, J. Minár, H. Ebert, S. Moyerman, E. E. Fullerton, and F. Hellman, Fe Spin Reorientation across the Metamagnetic Transition in Strained FeRh Thin Films, Phys. Rev. Lett., vol. 109, p. 117201, (2012).
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