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Large lattice mismatch effects on the epitaxial growth and magnetic properties of FePt films
Accepted Manuscript Large lattice mismatch effects on the epitaxial growth and magnetic properties of FePt films Jinyu Deng, Kaifeng Dong, Ping Yang, Yingguo Peng, Ganping Ju, Jiangfeng Hu, Gan Moog Chow, Jingsheng Chen PII: DOI: Reference:
S0304-8853(17)32420-4 http://dx.doi.org/10.1016/j.jmmm.2017.09.014 MAGMA 63143
To appear in:
Journal of Magnetism and Magnetic Materials
Received Date: Revised Date: Accepted Date:
1 August 2017 5 September 2017 6 September 2017
Please cite this article as: J. Deng, K. Dong, P. Yang, Y. Peng, G. Ju, J. Hu, G. Moog Chow, J. Chen, Large lattice mismatch effects on the epitaxial growth and magnetic properties of FePt films, Journal of Magnetism and Magnetic Materials (2017), doi: http://dx.doi.org/10.1016/j.jmmm.2017.09.014
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Large lattice mismatch effects on the epitaxial growth and magnetic properties of FePt films Jinyu Deng1, Kaifeng Dong2, Ping Yang3, Yingguo Peng4, Ganping Ju4, Jiangfeng Hu5, Gan Moog Chow1 and Jingsheng Chen1,* 1
Department of Materials Science and Engineering, National University of Singapore, Singapore
117575, Singapore 2
School of Automation, China University of Geoscience, Wuhan 430074, China
3
Singapore Synchrotron Light Source, National University of Singapore, Singapore 117603,
Singapore 4
Seagate Technology, Fremont, California 94538, USA
5
Data Storage Institute (DSI), Singapore 117608, Singapore
Heteroepitaxial film growth is crucial for magnetic and electronic devices. In this work, we reported the effects of the large lattice mismatch and film thickness on the epitaxial growth and magnetic properties of FePt films on ZrxTi1-xN (001) intermediate layer. FePt films with different thickness were deposited on ZrTiN intermediate layers with various doping concentration of TiN in ZrN. The increase in doping concentration of TiN caused a decrease in the lattice parameters of ZrTiN intermediate layer. It was found that (001) epitaxy of FePt 10 nm films was only achieved on ZrTiN intermediate layer when the TiN composition was ≥25 vol. %, while (001) texture of 5 nm films was achieved on ZrTiN intermediate layer with a minimum of 50 vol. % TiN composition. The in-plane lattice constants of FePt and Zr0.70Ti0.30N (25 vol. % TiN) were 3.870 Å and 4.476 Å, respectively, which resulted in a lattice mismatch as large as 15.7%. These large lattice mismatch heterostructures adopted 7/6 domain matching epitaxy. The magneto-
1
crystalline anisotropy of FePt films was improved with the increase in lattice mismatch. Intrinsic magnetic properties were extrapolated for FePt (30 nm)/ Zr0.70Ti0.30N (30 nm)/TaN (30 nm)/MgO, and the Ms(0 K) and K1(0 K) were 1042 emu/cc and 5.10×107 erg/cc, respectively, which is comparable to that of bulk L10 FePt.
*Author to whom correspondence should be addressed. Electronic mail: [email protected]. Tel.: 65-65167574. FAX: 65-67763604
2
I.
INTRODUCTION Over the last few decades, much research has been done and major advances have
occurred in both understanding and practicing with regard to heat assisted magnetic recording (HAMR). For developing a magnetic storage media fulfilling industrial requirements, there is still room for further improvement in decreasing grain size, narrowing size distribution, enhancing perpendicular magnetic anisotropy, etc. Engineering the intermediate layer is one of the methods to modify the microstructure of FePt films. For this reason, extensive studies have been carried out on the intermediate layer through tuning its surface chemistry [1-5] and controlling the mismatch with respect to FePt epilayer [6-10]. The effects of the lattice mismatch between 2.23% and 8.86% on the epitaxial growth and magnetic properties of FePt films have been investigated [6, 11]. However, to our knowledge, there is no report on how the lattice mismatch affects the epitaxial growth and magnetic properties of FePt films when the lattice mismatch is larger than 8.86%. Conventionally, it is believed that lattice-matching epitaxy (LME) during thin film growth is possible as long as the lattice misfit between the film and substrate is small (less than 7% to 9%). Above this limit, it is assumed that the film will be textured or polycrystalline [12, 13]. However, it was reported that in some heteroepitxial system films with lattice mismatch larger than 9% were epitaxially grown by domain-matching epitaxy (DME), where integral multiples of unit cells of film and substrate match across the interface [13, 14]. With DME, the total strain energy caused by mismatch in FePt heteroepitaxy could be released greatly by forming interfacial dislocations. The energy minima occurs when grain size is integral times of the domain size. Therefore, the stable grain size is defined by the domain size and any further change of the grain size needs to overcome an energy barrier imposed by misfit strain. Larger 3
lattice mismatch could result in a smaller domain size which might be beneficial in reducing FePt grain size for ultra-high density HAMR. In the present work, the effects of lattice mismatch larger than 9% on the epitaxial growth, microstructure and magnetic properties of FePt films were investigated. The films were deposited on ZrxTi1-xN intermediate layers. ZrN with lattice constant of 4.570 Å (18.0% mismatch w.r.t. FePt) and TiN with lattice constant of 4.231 Å (9.3% mismatch w.r.t. FePt) were used. The lattice constants of ZrxTi1-xN solid solution with different compositions changed from 4.570 Å to 4.231 Å, depending on the molar concentration of each component according to Vegard’s Law [15]. The lattice mismatch with regard to (w.r.t.) the substrate or the film mentioned in this work is defined as
or
, where af and as are lattice constants a of the substrate and the film, respectively.
II.
EXPERIMENT FePt (5, 10, 20 and 30 nm) / ZrTiN (30 nm ZrN with 0, 15, 25, 30, 40, and 50 vol. % TiN)
/ TaN (30 nm) / MgO (001) single crystal substrate films were prepared by using magnetron sputtering system with a base pressure lower than 2×10-8 Torr. The ZrTiN intermediate layers were prepared by co-sputtering ZrN and TiN targets. The deposition temperature of FePt, ZrTiN and TaN was 380 ˚C. The deposition rate for FePt was fixed at 24 Å/min. The TaN layer was employed to promote the (001) texture of ZrTiN intermediate layer. The crystallographic structure was examined by X-ray diffraction (XRD) using Cu Kα radiation (λ=0.1542 nm), and synchrotron radiation X-ray diffraction. The microstructure of the films was observed by transmission electron microscopy (TEM). The magnetic properties were measured using the vibrating sample magnetometer (VSM), and the anomalous Hall effect (AHE) measurement. The
4
Hall effect measurement was carried out using a four point ac resistance bridge with a small bias current of 100 μA.
III.
RESULTS AND DISCUSSIONS
A. Crystallographic Structure Fig. 1 (a) shows the -2 XRD spectra of 10 nm FePt films grown on ZrTiN intermediate layer with TiN doping concentration varied from 0 to 40 vol. %. ZrTiN solid solution was formed and a slight shift of ZrTiN (002) peak towards higher angle was observed. The right shift indicated that with the increase in TiN doping concentration, the lattice constant c decreased. According to Vegard’s Law [15], the lattice constant
,
where x was the molar fraction of TiN in solid solution. Both FePt (002) and (200) peaks appeared when TiN concentration was below 25 vol. %, which was not desirable for perpendicular magnetic recording. When the doping concentration of TiN reached 25 vol. %, FePt (200) peak disappeared, indicating an excellent epitaxial growth of FePt film. FePt films of different thickness were then deposited on ZrTiN intermediate layer with 25 vol. % TiN doping concentration, and the XRD spectra is shown in Fig. 1 (b). For the 5 nm FePt sample, only a very weak (002) FePt bump could be observed, while the other FePt films with thicknesses of 10 nm, 20 nm and 30 nm all exhibited good (001) epitaxy. The observation of FePt (003) peaks in thicker films clearly indicated a good long range chemical ordering. As seen in Fig. 1 (c), 5 nm epitaxial FePt films were achieved when the TiN doping concentration in ZrTiN was ≥ 50 vol. %. For those epitaxial FePt films, the (001) and (002) peak positions did not change with the varying interlayer mismatch, showing that the lattice constant c of FePt remained unchanged. It has been reported earlier that when the lattice 5
mismatch is 6.33%, the lattice constant c increases and lattice constant a decrease with lattice mismatch [6, 11], due to the tensile strain that expands the lattice constant a and shrink the lattice constant c [16-18]. Further increase in lattice mismatch to 8.86%, the lattice constant c increased and lattice constant a decreased due to strain relaxation. The FePt lattice was fully relaxed and the lattice parameters kept constant with further increase in lattice mismatch, based on the findings from current work. The rocking curves of the FePt films grown on ZrTiN intermediate layer were carried out to examine the quality of (001) orientation. Summaries of the full width at half maximum (FWHM) Δθ50 of the rocking curves of FePt (001) peak are shown in Fig. 1 (d). All the Δθ50 were larger than 2˚, which is mainly due to the grain tilting induced by the large lattice mismatch. Moreover, the increase in the FePt thickness led to a decrease in the Δθ50 of FePt (001) peak. This could be due to the formation of continuous film when the FePt layer grew thicker [7]. The XRD θ-2θ scan could only reveal the information of crystalline structure along film normal i.e. lattice constant c. In order to acquire in-plane information such as in-plane lattice plane spacing and strain states for structural analysis, asymmetrical
reciprocal space
mapping (RSM) was done by using synchrotron radiation and is shown in Fig. 2. The vertical axis was along [00L] direction while the horizontal axis was along [HH0] direction. Strong diffraction peaks were observed for
planes of MgO substrate, TaN and ZrTiN layers,
which were all labelled accordingly. As seen from Figs. 2 (a), (b) and (c), with the increase of TiN doping concentration from 0 vol. % to 25 vol. %, the ZrTiN towards the TaN
peak position shifted
peak, indicating that both lattice constant a and c of ZrTiN decreased.
The peak positions needed for precise calculation of the lattice constant and strain analysis were determined by fitting Gaussian functions to the intensity integrated along and perpendicular to 6
the
dicrections. The results of lattice constant a of ZrTiN intermediate layers with various
TiN doping concentration determined by RSM measurement were displayed in Table 1, together with the results calculated by using Vegard’s Law. It was found that, with the increase in TiN doping concentration, the lattice constant a of ZrTiN intermediate layer decreased and the experimental values matched with the calculated values. Further increasing the TiN doping concentration, the
diffraction peaks of TaN and ZrTiN merged, which made it very
difficult to determine the precise lattice constants a from RSM. In these cases, Vegard’s Law was used for lattice constant determination. Table 1. Both experimental and calculation values of lattice parameter a of ZrTiN intermediate layer with various TiN doping concentration TiN doping concentration
ZrTiN (HH0) index
Lattice Parameter a (Å)
Lattice Parameter a (Å)
Experimental results
Calculated results
0 vol. %
-0.920
4.578
4.570
15 vol. %
-0.932
4.517
4.508
25 vol. %
-0.942
4.471
4.469
As seen from Figs. 2(d) and 2(e), strong FePt
diffraction spots were observed for both 10
nm and 30 nm films grown on ZrTiN intermediate layer with 25 vol. % TiN doping. With increasing FePt thickness, the diffraction spots of FePt
plane became stronger. The
distribution of the diffuse scattered intensity around the reciprocal lattice point (RLP) maxima represented the degree of imperfect texture. The broadening of the FePt
plane in
horizontal direction could be observed, implying a large variation in lattice constant a. This may
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be caused by the relaxation by forming misfit dislocation due to a large mismatch [7]. The [00L] dash line passing through the ZrTiN RLP maxima is a reference line to determine the strain status of the FePt epilayer [19]. The peak position of FePt
plane was far away from the
dash line, implying that the FePt films were fully relaxed. The lattice constant a of FePt film derived from the RSM peak position was 3.87 Å, which was very close to the value of bulk L10 FePt. The relaxation of FePt films might imply that the film adopted DME to relieve large strain energy by forming misfit dislocation. By adopting DME, every m planes of the FePt film matched with every n planes of ZrTiN intermediate layer [13]. The residue misfit strain in FePt film could be thus reduced to (1)
ε
where ds and df are the interplanar spacing of matching lattice planes of the substrate and the film, respectively, and coefficient m and n are integer numbers. When FePt was deposited on ZrTiN with 25 vol. % TiN doping, the system was expected to follow 7/6 DME which possessed the lowest strain energy. Table 2 shows a complete summary of in-plane information including lattice constant a of both ZrTiN and FePt, lattice misfit between the two layers, and as well as the predicted DME m/n ratio. Table 2. Lattice constant a of both ZrTiN and FePt, lattice misfit between the two layers, m/n ratio as function of TiN concentrations in ZrTiN. ZrTiN intermediate layer TiN composition Lattice Parameter a (Å) 25 vol. %
4.476
FePt films Lattice Parameter a (Å) 3.870
Lattice Misfit
Predicted DME m/n ratio
ε 15.7%
7/6
8
30 vol. %
4.451
3.870
15.0%
7/6 or 8/7
40 vol. %
4.415
3.870
14.1%
8/7
50 vol. %
4.380
3.870
13.2%
8/7 or 9/8
B. Microstructure Low magnification and high resolution cross-sectional TEM images were taken for further investigations on microstructure. 10 nm FePt film grown on ZrTiN with 25 vol. % TiN doping concentration had an island structure, as shown in Fig. 3 (a). The formation of islands was mainly attributed to large lattice mismatch based on Chen’s [20] study. Chen et.al. proposed that the mismatched film thermodynamically favored island structure when the size of the epilayer exceeded the critical thickness that was inversely proportional to the mismatch strain. The decrease of strain energy caused by misfit dislocation (or island edge) formation could outweigh the increase in the surface energy during the formation of islands. The FePt grain on the right shown in fig. 3 (b) was tilted 2.8˚. The tilting was one of the lattice relaxation mechanisms to reduce the large misfit strain energy brought by the lattice mismatch [21]. The grain tilt resulted in a relatively large Δθ50 of FePt film, and it might cause the opening-up in its in-plane hysteresis loop as well. The insert was the diffraction pattern with <010> zone axis, which demonstrated the excellent epitaxial growth of FePt (001) <100>//ZrTiN (001) <100> layers. Fig. 3 (c) is the enlarged image of selected area with (200) lattice fringes and ½[200] Burger vector labelled, from which, 7/6 domain match with ½[200] edge dislocations could be easily observed. Fig 3 (e) is the EDX elemental mapping which showed a rough FePt/ZrTiN interface. The elemental inter diffusion with a 2 nm-thick mixing zone was observed.
9
The cross-sectional images of textured 5 nm FePt film grown on ZrTiN intermediate layer with 25 vol. % TiN doping concentration were shown in Fig. 4. A layer of FePt nano-grains were observed, and the morphology was very different from that of 10 nm FePt film. Statistically, 65% of FePt grains had a preferred (001) orientation, within which some tilted grains were found as well. As shown in Fig. 4 (c), except the majority (001) textured L10 FePt grains, grains with (100) and (110) orientations were observed as well. Furthermore, some of the grains were in disordered A1 phase. From the EDX mapping show in Fig. 4 (b), a rough interface could also be observed. Ti, Zr and N diffused upwards and acted like spacing material between FePt grains. Interestingly, (001) oriented 5nm FePt film was achieved when the TiN doping concentration in ZrTiN layer increased to 50 vol. %. Compared with the (001) epitaxial 10 nm film, Fig. 5 (a) showed that FePt islands had a similar truncated square pyramid configuration. However the size of the islands were much smaller than that of the 10 nm film. The DME model was observed and shown in Fig. 5 (e), and the majority of domains was confirmed to have an 8/7 ratio which matched with previous prediction. The change from 7/6 to 8/7 was due to the decrease in lattice mismatch. A less rough interface and a thinner elemental mixing zone was shown in the EDX mapping. It could be due to the presence of less defects and build-up of strain at the interface.
C. Magnetic Properties The out-of-plane and in-plane M-H loops of FePt films with different thickness deposited on ZrTiN intermediate layer with different TiN doping concentration were measured by VSM. All FePt films with thicknesses of 10, 20 and 30 nm grown on ZrTiN intermediate layer exhibited good perpendicular magnetic anisotropy. A typical out-of-plane and in-plane M-H loops of 10 nm films are shown in Fig. 6 (a). The opening up of in-plane M-H loop was observed which was possibly due to the grain tilting caused by large interface mismatch [21]. The out-of-plane 10
coercivities of FePt films as function of thickness and TiN concentration in ZrTiN is shown in Fig. 6 (c). The thinner FePt films had larger coercivity, which might be attributed to the StonerWohlfarth rotation or domain wall pinning dominated magnetization reversal mechanism since the grain boundaries and the gaps between the islands in thinner film could acted as obstacles to domain wall motion [22].With the increase in FePt film thickness, the out-of-plane coercivity gradually decreased. This was due to the change in magnetization reversal mechanism from S-W coherent rotation/pinning mode to domain nucleation mode, as a result of forming continuous film with the increase in film thickness. The out-of-plane coercivities did not show strong dependence on the TiN concentrations i.e. lattice mismatch. As for 5 nm FePt films, the epitaxy was not achieved until the TiN doping concentration reached 50 vol. %. It caused poor perpendicular magnetic anisotropy of the samples with 25, 30 and 40 vol. % TiN doping concentration in ZrTiN intermediate layer. However, for the sample with 5 nm FePt deposited on ZrTiN interlayer with 50 vol. % TiN doping concentration, it had the best magnetic properties among all of the samples. It had the highest out-of-plane coercivity of 23.4 kOe, the highest squareness and the smallest in-plane opening up as shown in Fig. 6 (b). The magnetic anisotropy constant K as function of TiN concentration, as illustrated in Fig.6 (d), were measured by Anomalous Hall effect with analyzing the normalized hall voltage curves using generalized Sucksmith-Thompson (GST) method [23]. It was found that with increase in TiN doping concentration the magnetic anisotropy decreased, implying that large lattice mismatch could improve the magnetic anisotropy. The change in magnetic anisotropy with TiN doping concentration coincide with that of chemical ordering parameter S which was estimated using the integrated peak intensity ratio (I001/I002) [24]. Ding et al. reported that the in-plane tensile strain might increase with an increase of the lattice mismatch from 2.23% to 6.33% (w.r.t. 11
intermediate layer). The strain helped in expansion of the a axis and contraction of the c axis of FePt, thus assisted the phase transformation and chemical ordering. With a mismatch larger than 6.33%, dislocations formed at the interface to release the strain, which might result in a lower degree of chemical ordering [6]. However in our experiment, the lattice constants of FePt films did not change with the increase in the lattice misfit (w.r.t. film) from 13.2% to 15.7%. There was no in-plane tensile strain since the FePt lattice was relaxed. Obviously, there was discrepancy as compared to Ding’s paper. The previous studies were conducted based on different type of substrates or intermediate layers, such as Pt, Cr, MgO etc. which may introduce more uncertainties to the system such as the interfacial interaction between these materials and FePt film. In this work, all of the FePt films were deposited on Zr xTi1-xN intermediate layer. Zr and Ti belong to the same group of element and minor change in atomic ratio of these 2 elements does not greatly affect the surface chemistry. However, the reason to its increasing chemical ordering parameter S and magnetocrystalline anisotropy with the lattice misfit beyond 13.2% still requires further investigation. The temperature dependence of saturation magnetization Ms and magnetic anisotropy constant K for 30nm thick FePt film deposited on ZrTiN intermediate layer with 25 vol. % TiN doping concentration was measured. The M-T and K-T curves with least square fit were shown in Fig. 6 (e) and (f). The M-T curve of ferromagnetic materials could be described by the general analytical function proposed by Kuz’min [25]: (2)
where m is the normalized moment; is the reduced temperature which
; s and p are
parameters, p>3/2, s>0. This equation obey Bloch’s 3/2 power law at low temperature, and when 12
approaching another limit
, m(τ) is proportional to (1- τ)1/3, as described by Heisenberg
model [26]. In our experiment, as an inter-diffusion zone was observed at the interface, the total magnetic moment μtot was considered as the summation of the body moment μbody and the deadlayer moment μdeadlayer. Each component was fitted using equation (2). Tc (body) obtained from the fitting was 733 K, and Ms body (0) was obtained to be 1042 emu/cc. The K-T curve was fitted by the theory within the framework of thermal fluctuations of magnetic moments [27], which can be expressed as: (3)
The m was obtained to be equal to 2.4, which was smaller than 3, close to 2 as predicted by single-ion model for L10 magnets [28]. The K1(0) obtained from the fitting was about 5.10 × 107 erg/cc which was comparable to magneto-crystalline anisotropy of bulk FePt (K1 = 7 × 107 erg/cc measured at room temperature [29]).
D. Discussions Similar to LME, for DME there is a critical point of the lattice misfit between the film and substrate, beyond which the film will grow textured or largely polycrystalline. The limit in FePt (10 nm)/ZrTiN system was about 15.7% (corresponding to 25 vol. % TiN). Beyond this mismatch limit, the epitaxy could not be maintained and part of the FePt film was (100) oriented. For FePt (5 nm)/ ZrTiN system, the TiN concentration in ZrTiN must be more than 50 vol. % (13.2% lattice mismatch correspondingly), otherwise the FePt film grew textured. The limit in lattice mismatch was thermodynamically determined. Based on the experimental observations, a model of FePt (fixed thickness)/ ZrxTi(1-x)N system is proposed. Energy terms for both DME and polycrystalline configuration with incoherent interface are given and compared. Because Zr and 13
Ti are the same group elements, they have the same number of valance electron. It is reasonable to assume a constant surface chemistry of ZrTiN which is independent from the TiN composition. Moreover, the ZrTiN layer is assumed rigid in this model. For a polycrystalline film with an incoherent interface, there will be no long-range misfit strain energy, but only short-range distortions caused by chemical binding at the interface [30]. Other than the chemical interfacial energy, surface energy is another term for consideration. However, these two terms are both independent from the TiN doping concentration in ZrTiN and considered as constant,
. Unfortunately, it is
very difficult to precisely calculate the total energy. What we know is that DME will occur in FePt/ZrTiN system as long as its total energy is smaller than this critical value of
.
As for a DME FePt/ZrTiN system, a model is built based on the generalized schematic configuration shown in Fig. 7 (a). In order to fulfill the traction-free boundary conditions on the free surface, a series of image dislocations with opposite sign to misfit dislocations were assumed [31]. The total energy of the system includes residual strain energy, surface energy, interfacial energy and dislocation energy. The coherent residual strain energy density in FePt film can be expressed as (4)
where E is the Young’s modulus and v is the Poisson ration of FePt [32]. h is the film thickness. For a perfect DME system, the residual strain is zero, resulting in a zero strain energy density. Even if εr is finite meaning nds cannot totally cancel out mdf, the value of strain energy is still very small and insignificant. The interfacial energy is determined by the surface chemistry. With the assumption stated earlier, the FePt/ZrTiN chemical interfacial energy should stay constant. 14
The surface energy is determined by the morphology of the film surface and its equilibrium crystal shape. It can be treated as a constant if the film thickness and interfacial energy have been fixed. The energy terms change with lattice mismatch are mainly related with the formation of dislocation. For the purpose of simplifying the calculation, the dislocation-related energy density is considered per domain. Within one domain, the self-energy density per area for edge dislocations is (5)
where (6)
G is the shear modulus of FePt; b is the Burger vector which equals to a/2[100]; d is the average domain size; R is the crystal size; and r0 is the cut-off core radius of the dislocation, which is close to one Burger vector [33]. The interaction between misfit dislocation and its image dislocation can be expressed as (7)
To sum up equations (5) and (7) we can get the total dislocation energy (8)
The final dislocation related energy term is the interaction energy between dislocations from other domains. The interaction energy between two closest neighbors can be expressed as [33]
15
(9)
Since interaction energy decreases quickly with increasing distance, first order approximation is used for estimation (10)
The insertion of the extra half planes also introduce the ledge surface energy (11)
where
is the surface energy of FePt (001) plane. Therefore, according to the equations, the
total energy arises from dislocation formation
is a function of the domain
size d. In order to correlate the total dislocation-related energy with lattice misfit, the domain size
can be expressed as (12)
(13)
By substituting (12) into (13): (14)
16
(15)
For the calculation, the Young’s Modulus of 180 GPa and Poisson ratio of 0.33 for bulk FePt were used [34]. The edge dislocation energy density as a function of misfit strain regarding domain matching FePt/ZrTiN is plotted in Fig. 7 (b). The blue triangle represents the edge dislocation energy density corresponding to each perfect m/n domain match system. The full energy profile is a continuous function instead of a discrete step function. This is because, in case of imperfect domain matching, the m/n and (m+1)/(n+1) domains may alternate with a frequency to provide a perfect matching according to (16)
where α is the frequency factor [13]. Therefore, the average domain size changes continuously with the misfit strain. As shown in the plot, with a large misfit strain, the dislocation energy is largely contributed by the interaction energy. This is because when the dislocations are distributed closer, the repulsion among the parallel misfit dislocations become stronger. With the decrease in misfit strain caused by the increase in TiN doping concentration, the total dislocation energy density drops dramatically. Once the doping concentration reaches the critical value, the DME will be thermodynamically favored. In present work, when lattice mismatch was larger than 15.7%, the DME could not be retained. However, it has been reported that DME was observed in ZnO/α-Al2O3 (0001) (18.3% misfit w.r.t. ZnO) [14] and TiN/Si (100) (28.4% misfit w.r.t. TiN) [35]. To our knowledge, in order to achieve DME, it is important to keep the total interfacial energy low. The interfacial energy comprises dislocation-related energy and interfacial chemical binding. A large lattice mismatch can cause a high density of dislocation, and induce a huge dislocation-related energy. In order to 17
stabilize the DME system with a large lattice mismatch, the chemical interfacial energy must be decreased by choosing similar type of materials for these 2 layers. For example, the ZnO/αAl2O3 (0001) system already has a very large lattice mismatch of 18.3%. The strong ionic bond also causes a rigid structure which can induce an even greater dislocation energy. The DME in ZnO/α-Al2O3 (0001) system is still retained despite the large dislocation energy since it has a low oxide/oxide chemical interfacial energy. On the opposite, perpendicular anisotropy of FePt films could also been obtained non-epitaxially on amorphous substrates [36]. However, the FePt (001)/SiO2/Si film obtained by post-annealing might be at a thermodynamic metastable state, and eventually the FePt would become (111) oriented. These findings may offer a new way to deposit epitaxial films on substrates with large lattice mismatch for next-generation solid state technology and industrial applications.
IV.
SUMMARY Domain matching epitaxial (DME) FePt (001) films were deposited on a ZrxTi(1-x)N
intermediate layer with a lattice misfit no greater than 15.7% (w.r.t. film). Strain energy was released through the formation of edge dislocations at the interface of layers with mismatched lattice constants. The density of dislocations increases with the interlayer misfit, causing an increase in total energy and destabilizing the DME. The effect of large lattice mismatch on magnetic properties of FePt was investigated as well. The experiments and analysis revealed that the FePt films had good perpendicular magnetic anisotropy with large out-of-plane coercivity. The in-plane magnetic component was caused by grain tilting which is a strain relieving mechanism. The large lattice mismatch could improve both chemical ordering and as well as the magnetic anisotropy constant Ku, which is desirable for magnetic recording applications.
18
ACKNOWLEDGEMENTS The research is supported in part by the Singapore National Research Foundation under CRP Award No. NRF-CRP10-2012-02 and Seagate Technology.
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[10] A. Hotta, T. Ono, M. Hatayama, K. Tsumura, N. Kikuchi, S. Okamoto, O. Kitakami, T. Shimatsu, J. Appl. Phys. 115 (2014) 17B712. [11] C.J. Aas, L. Szunyogh, J.S. Chen, R.W. Chantrell, Appl. Phys. Lett. 99 (2011) 132501. [12] J.W. Mathews, A.E. Blakeslee, J. Cryst. Growth 27 (1974) 118-125. [13] J. Narayan, B.C. Larson, J. Appl. Phys. 93 (2003) 278. [14] J. Narayan, K. Dovidenko, A.K. Sharma, S. Oktyabrsky, J. Appl. Phys. 84 (1998) 25972601. [15] A.R. Denton, N.W. Ashcroft, Phys. Rev. A 43 (1991) 3161-3164. [16] Y.-N. Hsu, S. Jeong, D.N. Lambeth, D.E. Laughlin, IEEE Trans. Magn. 36 (2000) 29452947. [17] Y.-N. Hsu, S. Jeong, D.E. Laughlin, D.N. Lambeth, J. Appl. Phys. 89 (2001) 7068-7070. [18] Y. Xu, J.S. Chen, J.P. Wang, Appl. Phys. Lett. 80 (2002) 3325-3327. [19] E. Koppensteiner, G. Bauer, H. Kibbel, E. Kasper, J. Appl. Phys. 76 (1994) 3489. [20] Y. Chen, J. Washburn, Phys. Rev. Lett., 77 (1996) 4046. [21] X. Jiang, R.Q. Zhang, G. Yu, S.T. Lee, Phys. Rev. B 58 (1998) 15351-15354. [22] K.F. Dong, H.H. Li, J.F. Hu, Y.G. Peng, G. Ju, G.M. Chow, J.S. Chen, IEEE Trans. Magn. 49 (2013) 668-674. [23] S. Okamoto, K. Nishiyama, O. Kitakami, Y. Shimada, J. Appl. Phys. 90 (2001) 4085. [24] S. Okamoto, N. Kikuchi, O. Kitakami, T. Miyazaki, Y. Shimada, K. Fukamichi, Phys. Rev. B 66 (2002) 024413. [25] M.D. Kuz’min, Phys. Rev. Lett. 94 (2005) 107204. [26] E. Callen, H. Callen, J. Appl. Phys. 36 (1965) 1140-1142. [27] C. Zener, Phys. Rev. 96 (1954) 1335-1337.
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[28] R. Skomski, O.N. Mryasov, J. Zhou, D.J. Sellmyer, J. Appl. Phys. 99 (2006) 08E916. [29] O.A. Ivanov, L.V. Solina, V.A. Demshina, L.M. Magat, Phys. Met. Metallogr. 35 (1973) 81. [30] I.A. Ovid’Ko, Rev. Adv. Mater. Sci. 1 (2000) 61-107. [31] C. Wei, W.D. Nix, Imperfections in Crystalline Solids, 1st ed., Cambridge University Press, United Kindom, 2016. [32] M. Ohring, Chapter 8 - Epitaxy, in: Materials Science of Thin Films (Second Edition), Academic Press, San Diego, 2002, pp. 429. [33] J.P. Hirth, J. Lothe, Theory of Dislocations, 2nd ed., McGraw-Hill, New York, 1968. [34] P. Rasmussen, X. Rui, J.E. Shield, Appl. Phys. Lett. 86 (2005) 191915. [35] T. Zheleva, K. Jagannadham, J. Narayan, J. Appl. Phys. 75 (1994) 860-871. [36] H. Zeng, M. L. Yan, N. Powers, D. J. Sellmyer, Appl. Phys. Lett. 80, (2002) 2350.
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Figures:
FIG. 1. (a) XRD spectra of 10 nm FePt films grown on ZrTiN intermediate layers with various TiN doping concentration, (b) XRD spectra of FePt films with various thickness grown on ZrTiN intermediate layer with 25 vol. % TiN doping concentration, (c) XRD spectra of 5 nm FePt films grown on ZrTiN intermediate layers with various TiN doping concentration, (d) FWHM Δθ50 of the rocking curves of FePt (001) peak as a function of film thickness and TiN concentration.
22
FIG. 2. Reciprocal space maps around ( 3) planes of FePt grown on ZrTiN intermediate layer with (a) 0 vol. %, (b) 15 vol. % and (c) 25 vol. % TiN doping concentration, and maps of (d) 10nm and (e) 30nm FePt grown on ZrTiN intermediate layer with 25 vol. % TiN doping concentration.
23
FIG. 3. (a)Low magnification cross-sectional TEM image of FePt (10nm) / ZrTiN (25 vol.% TiN doping), (b) high resolution cross-sectional TEM image showing grain tilting, the insert is the FFT image of selected area, (c) enlarged selected area showing 7/6 domain matching epitaxy, with (200) lattice fringes and ½[200] Burger vector labelled, (d) schematic drawing of 7/6 domain matching epitaxy and (e) EDX elemental mapping at the interface. 24
FIG. 4. (a)Low magnification cross-sectional TEM image, (b) & (c) EDX elemental mapping of selected area of FePt (5nm) / ZrTiN (25 vol.% TiN doping), and (c) high resolution crosssectional TEM images of selected grains with corresponding FFT image, showing polycrystalline characteristic of the film.
25
FIG. 5. (a)Low magnification cross-sectional TEM image of FePt (5nm) / ZrTiN (50 vol.% TiN doping), (b) EDX elemental mapping of selected area, (c)high resolution cross-sectional TEM image at interface, (d) corresponding FFT image and (e) enlarged image of selected area with (200) lattice fringes and ½ [200] Burger vector labelled, and (f) schematic drawing of 8/7 domain matching epitaxy.
26
FIG. 6. (a) The hysteresis loop of FePt (10nm) / ZrTiN (25 vol.% TiN), (b) the hysteresis loop of FePt (5nm) / ZrTiN (50 vol.% TiN), (c) summary of out-of-plane coercivity of all samples, (d) the dependence of magnetic anisotropy constant Ku and integrated peak intensity ratio I001/I002 of 10nm FePt films on the TiN doping concentration in ZrTiN intermediate layer, (e) M-T and (f) K-T curves with least square fit of FePt (30nm) / ZrTiN (25 vol.% TiN).
27
FIG. 6 (a) Schematic drawing of the dislocation configuration in FePt / ZrTiN domain matching epitaxial system, and (b) the dependence of various energy densities in FePt (10nm) / ZrTiN (001) system on the misfit strain between the two layers.
28
Highlights:
L10 FePt films were epitaxially deposited on ZrxTi1-xN intermediate layer with a lattice mismatch up to 15.7%.
Detailed TEM studies showed that the domain matching epitaxy (DME) of FePt films was achieved by the formation of dislocations to reduce strain energy.
The dislocation-related energy increased with the lattice mismatch and destabilize the DME system.
The large lattice mismatch could improve both chemical ordering and as well as the magnetic anisotropy constant Ku, which is desirable for magnetic recording applications.
29
S0304-8853(17)32420-4 http://dx.doi.org/10.1016/j.jmmm.2017.09.014 MAGMA 63143
To appear in:
Journal of Magnetism and Magnetic Materials
Received Date: Revised Date: Accepted Date:
1 August 2017 5 September 2017 6 September 2017
Please cite this article as: J. Deng, K. Dong, P. Yang, Y. Peng, G. Ju, J. Hu, G. Moog Chow, J. Chen, Large lattice mismatch effects on the epitaxial growth and magnetic properties of FePt films, Journal of Magnetism and Magnetic Materials (2017), doi: http://dx.doi.org/10.1016/j.jmmm.2017.09.014
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Large lattice mismatch effects on the epitaxial growth and magnetic properties of FePt films Jinyu Deng1, Kaifeng Dong2, Ping Yang3, Yingguo Peng4, Ganping Ju4, Jiangfeng Hu5, Gan Moog Chow1 and Jingsheng Chen1,* 1
Department of Materials Science and Engineering, National University of Singapore, Singapore
117575, Singapore 2
School of Automation, China University of Geoscience, Wuhan 430074, China
3
Singapore Synchrotron Light Source, National University of Singapore, Singapore 117603,
Singapore 4
Seagate Technology, Fremont, California 94538, USA
5
Data Storage Institute (DSI), Singapore 117608, Singapore
Heteroepitaxial film growth is crucial for magnetic and electronic devices. In this work, we reported the effects of the large lattice mismatch and film thickness on the epitaxial growth and magnetic properties of FePt films on ZrxTi1-xN (001) intermediate layer. FePt films with different thickness were deposited on ZrTiN intermediate layers with various doping concentration of TiN in ZrN. The increase in doping concentration of TiN caused a decrease in the lattice parameters of ZrTiN intermediate layer. It was found that (001) epitaxy of FePt 10 nm films was only achieved on ZrTiN intermediate layer when the TiN composition was ≥25 vol. %, while (001) texture of 5 nm films was achieved on ZrTiN intermediate layer with a minimum of 50 vol. % TiN composition. The in-plane lattice constants of FePt and Zr0.70Ti0.30N (25 vol. % TiN) were 3.870 Å and 4.476 Å, respectively, which resulted in a lattice mismatch as large as 15.7%. These large lattice mismatch heterostructures adopted 7/6 domain matching epitaxy. The magneto-
1
crystalline anisotropy of FePt films was improved with the increase in lattice mismatch. Intrinsic magnetic properties were extrapolated for FePt (30 nm)/ Zr0.70Ti0.30N (30 nm)/TaN (30 nm)/MgO, and the Ms(0 K) and K1(0 K) were 1042 emu/cc and 5.10×107 erg/cc, respectively, which is comparable to that of bulk L10 FePt.
*Author to whom correspondence should be addressed. Electronic mail: [email protected]. Tel.: 65-65167574. FAX: 65-67763604
2
I.
INTRODUCTION Over the last few decades, much research has been done and major advances have
occurred in both understanding and practicing with regard to heat assisted magnetic recording (HAMR). For developing a magnetic storage media fulfilling industrial requirements, there is still room for further improvement in decreasing grain size, narrowing size distribution, enhancing perpendicular magnetic anisotropy, etc. Engineering the intermediate layer is one of the methods to modify the microstructure of FePt films. For this reason, extensive studies have been carried out on the intermediate layer through tuning its surface chemistry [1-5] and controlling the mismatch with respect to FePt epilayer [6-10]. The effects of the lattice mismatch between 2.23% and 8.86% on the epitaxial growth and magnetic properties of FePt films have been investigated [6, 11]. However, to our knowledge, there is no report on how the lattice mismatch affects the epitaxial growth and magnetic properties of FePt films when the lattice mismatch is larger than 8.86%. Conventionally, it is believed that lattice-matching epitaxy (LME) during thin film growth is possible as long as the lattice misfit between the film and substrate is small (less than 7% to 9%). Above this limit, it is assumed that the film will be textured or polycrystalline [12, 13]. However, it was reported that in some heteroepitxial system films with lattice mismatch larger than 9% were epitaxially grown by domain-matching epitaxy (DME), where integral multiples of unit cells of film and substrate match across the interface [13, 14]. With DME, the total strain energy caused by mismatch in FePt heteroepitaxy could be released greatly by forming interfacial dislocations. The energy minima occurs when grain size is integral times of the domain size. Therefore, the stable grain size is defined by the domain size and any further change of the grain size needs to overcome an energy barrier imposed by misfit strain. Larger 3
lattice mismatch could result in a smaller domain size which might be beneficial in reducing FePt grain size for ultra-high density HAMR. In the present work, the effects of lattice mismatch larger than 9% on the epitaxial growth, microstructure and magnetic properties of FePt films were investigated. The films were deposited on ZrxTi1-xN intermediate layers. ZrN with lattice constant of 4.570 Å (18.0% mismatch w.r.t. FePt) and TiN with lattice constant of 4.231 Å (9.3% mismatch w.r.t. FePt) were used. The lattice constants of ZrxTi1-xN solid solution with different compositions changed from 4.570 Å to 4.231 Å, depending on the molar concentration of each component according to Vegard’s Law [15]. The lattice mismatch with regard to (w.r.t.) the substrate or the film mentioned in this work is defined as
or
, where af and as are lattice constants a of the substrate and the film, respectively.
II.
EXPERIMENT FePt (5, 10, 20 and 30 nm) / ZrTiN (30 nm ZrN with 0, 15, 25, 30, 40, and 50 vol. % TiN)
/ TaN (30 nm) / MgO (001) single crystal substrate films were prepared by using magnetron sputtering system with a base pressure lower than 2×10-8 Torr. The ZrTiN intermediate layers were prepared by co-sputtering ZrN and TiN targets. The deposition temperature of FePt, ZrTiN and TaN was 380 ˚C. The deposition rate for FePt was fixed at 24 Å/min. The TaN layer was employed to promote the (001) texture of ZrTiN intermediate layer. The crystallographic structure was examined by X-ray diffraction (XRD) using Cu Kα radiation (λ=0.1542 nm), and synchrotron radiation X-ray diffraction. The microstructure of the films was observed by transmission electron microscopy (TEM). The magnetic properties were measured using the vibrating sample magnetometer (VSM), and the anomalous Hall effect (AHE) measurement. The
4
Hall effect measurement was carried out using a four point ac resistance bridge with a small bias current of 100 μA.
III.
RESULTS AND DISCUSSIONS
A. Crystallographic Structure Fig. 1 (a) shows the -2 XRD spectra of 10 nm FePt films grown on ZrTiN intermediate layer with TiN doping concentration varied from 0 to 40 vol. %. ZrTiN solid solution was formed and a slight shift of ZrTiN (002) peak towards higher angle was observed. The right shift indicated that with the increase in TiN doping concentration, the lattice constant c decreased. According to Vegard’s Law [15], the lattice constant
,
where x was the molar fraction of TiN in solid solution. Both FePt (002) and (200) peaks appeared when TiN concentration was below 25 vol. %, which was not desirable for perpendicular magnetic recording. When the doping concentration of TiN reached 25 vol. %, FePt (200) peak disappeared, indicating an excellent epitaxial growth of FePt film. FePt films of different thickness were then deposited on ZrTiN intermediate layer with 25 vol. % TiN doping concentration, and the XRD spectra is shown in Fig. 1 (b). For the 5 nm FePt sample, only a very weak (002) FePt bump could be observed, while the other FePt films with thicknesses of 10 nm, 20 nm and 30 nm all exhibited good (001) epitaxy. The observation of FePt (003) peaks in thicker films clearly indicated a good long range chemical ordering. As seen in Fig. 1 (c), 5 nm epitaxial FePt films were achieved when the TiN doping concentration in ZrTiN was ≥ 50 vol. %. For those epitaxial FePt films, the (001) and (002) peak positions did not change with the varying interlayer mismatch, showing that the lattice constant c of FePt remained unchanged. It has been reported earlier that when the lattice 5
mismatch is 6.33%, the lattice constant c increases and lattice constant a decrease with lattice mismatch [6, 11], due to the tensile strain that expands the lattice constant a and shrink the lattice constant c [16-18]. Further increase in lattice mismatch to 8.86%, the lattice constant c increased and lattice constant a decreased due to strain relaxation. The FePt lattice was fully relaxed and the lattice parameters kept constant with further increase in lattice mismatch, based on the findings from current work. The rocking curves of the FePt films grown on ZrTiN intermediate layer were carried out to examine the quality of (001) orientation. Summaries of the full width at half maximum (FWHM) Δθ50 of the rocking curves of FePt (001) peak are shown in Fig. 1 (d). All the Δθ50 were larger than 2˚, which is mainly due to the grain tilting induced by the large lattice mismatch. Moreover, the increase in the FePt thickness led to a decrease in the Δθ50 of FePt (001) peak. This could be due to the formation of continuous film when the FePt layer grew thicker [7]. The XRD θ-2θ scan could only reveal the information of crystalline structure along film normal i.e. lattice constant c. In order to acquire in-plane information such as in-plane lattice plane spacing and strain states for structural analysis, asymmetrical
reciprocal space
mapping (RSM) was done by using synchrotron radiation and is shown in Fig. 2. The vertical axis was along [00L] direction while the horizontal axis was along [HH0] direction. Strong diffraction peaks were observed for
planes of MgO substrate, TaN and ZrTiN layers,
which were all labelled accordingly. As seen from Figs. 2 (a), (b) and (c), with the increase of TiN doping concentration from 0 vol. % to 25 vol. %, the ZrTiN towards the TaN
peak position shifted
peak, indicating that both lattice constant a and c of ZrTiN decreased.
The peak positions needed for precise calculation of the lattice constant and strain analysis were determined by fitting Gaussian functions to the intensity integrated along and perpendicular to 6
the
dicrections. The results of lattice constant a of ZrTiN intermediate layers with various
TiN doping concentration determined by RSM measurement were displayed in Table 1, together with the results calculated by using Vegard’s Law. It was found that, with the increase in TiN doping concentration, the lattice constant a of ZrTiN intermediate layer decreased and the experimental values matched with the calculated values. Further increasing the TiN doping concentration, the
diffraction peaks of TaN and ZrTiN merged, which made it very
difficult to determine the precise lattice constants a from RSM. In these cases, Vegard’s Law was used for lattice constant determination. Table 1. Both experimental and calculation values of lattice parameter a of ZrTiN intermediate layer with various TiN doping concentration TiN doping concentration
ZrTiN (HH0) index
Lattice Parameter a (Å)
Lattice Parameter a (Å)
Experimental results
Calculated results
0 vol. %
-0.920
4.578
4.570
15 vol. %
-0.932
4.517
4.508
25 vol. %
-0.942
4.471
4.469
As seen from Figs. 2(d) and 2(e), strong FePt
diffraction spots were observed for both 10
nm and 30 nm films grown on ZrTiN intermediate layer with 25 vol. % TiN doping. With increasing FePt thickness, the diffraction spots of FePt
plane became stronger. The
distribution of the diffuse scattered intensity around the reciprocal lattice point (RLP) maxima represented the degree of imperfect texture. The broadening of the FePt
plane in
horizontal direction could be observed, implying a large variation in lattice constant a. This may
7
be caused by the relaxation by forming misfit dislocation due to a large mismatch [7]. The [00L] dash line passing through the ZrTiN RLP maxima is a reference line to determine the strain status of the FePt epilayer [19]. The peak position of FePt
plane was far away from the
dash line, implying that the FePt films were fully relaxed. The lattice constant a of FePt film derived from the RSM peak position was 3.87 Å, which was very close to the value of bulk L10 FePt. The relaxation of FePt films might imply that the film adopted DME to relieve large strain energy by forming misfit dislocation. By adopting DME, every m planes of the FePt film matched with every n planes of ZrTiN intermediate layer [13]. The residue misfit strain in FePt film could be thus reduced to (1)
ε
where ds and df are the interplanar spacing of matching lattice planes of the substrate and the film, respectively, and coefficient m and n are integer numbers. When FePt was deposited on ZrTiN with 25 vol. % TiN doping, the system was expected to follow 7/6 DME which possessed the lowest strain energy. Table 2 shows a complete summary of in-plane information including lattice constant a of both ZrTiN and FePt, lattice misfit between the two layers, and as well as the predicted DME m/n ratio. Table 2. Lattice constant a of both ZrTiN and FePt, lattice misfit between the two layers, m/n ratio as function of TiN concentrations in ZrTiN. ZrTiN intermediate layer TiN composition Lattice Parameter a (Å) 25 vol. %
4.476
FePt films Lattice Parameter a (Å) 3.870
Lattice Misfit
Predicted DME m/n ratio
ε 15.7%
7/6
8
30 vol. %
4.451
3.870
15.0%
7/6 or 8/7
40 vol. %
4.415
3.870
14.1%
8/7
50 vol. %
4.380
3.870
13.2%
8/7 or 9/8
B. Microstructure Low magnification and high resolution cross-sectional TEM images were taken for further investigations on microstructure. 10 nm FePt film grown on ZrTiN with 25 vol. % TiN doping concentration had an island structure, as shown in Fig. 3 (a). The formation of islands was mainly attributed to large lattice mismatch based on Chen’s [20] study. Chen et.al. proposed that the mismatched film thermodynamically favored island structure when the size of the epilayer exceeded the critical thickness that was inversely proportional to the mismatch strain. The decrease of strain energy caused by misfit dislocation (or island edge) formation could outweigh the increase in the surface energy during the formation of islands. The FePt grain on the right shown in fig. 3 (b) was tilted 2.8˚. The tilting was one of the lattice relaxation mechanisms to reduce the large misfit strain energy brought by the lattice mismatch [21]. The grain tilt resulted in a relatively large Δθ50 of FePt film, and it might cause the opening-up in its in-plane hysteresis loop as well. The insert was the diffraction pattern with <010> zone axis, which demonstrated the excellent epitaxial growth of FePt (001) <100>//ZrTiN (001) <100> layers. Fig. 3 (c) is the enlarged image of selected area with (200) lattice fringes and ½[200] Burger vector labelled, from which, 7/6 domain match with ½[200] edge dislocations could be easily observed. Fig 3 (e) is the EDX elemental mapping which showed a rough FePt/ZrTiN interface. The elemental inter diffusion with a 2 nm-thick mixing zone was observed.
9
The cross-sectional images of textured 5 nm FePt film grown on ZrTiN intermediate layer with 25 vol. % TiN doping concentration were shown in Fig. 4. A layer of FePt nano-grains were observed, and the morphology was very different from that of 10 nm FePt film. Statistically, 65% of FePt grains had a preferred (001) orientation, within which some tilted grains were found as well. As shown in Fig. 4 (c), except the majority (001) textured L10 FePt grains, grains with (100) and (110) orientations were observed as well. Furthermore, some of the grains were in disordered A1 phase. From the EDX mapping show in Fig. 4 (b), a rough interface could also be observed. Ti, Zr and N diffused upwards and acted like spacing material between FePt grains. Interestingly, (001) oriented 5nm FePt film was achieved when the TiN doping concentration in ZrTiN layer increased to 50 vol. %. Compared with the (001) epitaxial 10 nm film, Fig. 5 (a) showed that FePt islands had a similar truncated square pyramid configuration. However the size of the islands were much smaller than that of the 10 nm film. The DME model was observed and shown in Fig. 5 (e), and the majority of domains was confirmed to have an 8/7 ratio which matched with previous prediction. The change from 7/6 to 8/7 was due to the decrease in lattice mismatch. A less rough interface and a thinner elemental mixing zone was shown in the EDX mapping. It could be due to the presence of less defects and build-up of strain at the interface.
C. Magnetic Properties The out-of-plane and in-plane M-H loops of FePt films with different thickness deposited on ZrTiN intermediate layer with different TiN doping concentration were measured by VSM. All FePt films with thicknesses of 10, 20 and 30 nm grown on ZrTiN intermediate layer exhibited good perpendicular magnetic anisotropy. A typical out-of-plane and in-plane M-H loops of 10 nm films are shown in Fig. 6 (a). The opening up of in-plane M-H loop was observed which was possibly due to the grain tilting caused by large interface mismatch [21]. The out-of-plane 10
coercivities of FePt films as function of thickness and TiN concentration in ZrTiN is shown in Fig. 6 (c). The thinner FePt films had larger coercivity, which might be attributed to the StonerWohlfarth rotation or domain wall pinning dominated magnetization reversal mechanism since the grain boundaries and the gaps between the islands in thinner film could acted as obstacles to domain wall motion [22].With the increase in FePt film thickness, the out-of-plane coercivity gradually decreased. This was due to the change in magnetization reversal mechanism from S-W coherent rotation/pinning mode to domain nucleation mode, as a result of forming continuous film with the increase in film thickness. The out-of-plane coercivities did not show strong dependence on the TiN concentrations i.e. lattice mismatch. As for 5 nm FePt films, the epitaxy was not achieved until the TiN doping concentration reached 50 vol. %. It caused poor perpendicular magnetic anisotropy of the samples with 25, 30 and 40 vol. % TiN doping concentration in ZrTiN intermediate layer. However, for the sample with 5 nm FePt deposited on ZrTiN interlayer with 50 vol. % TiN doping concentration, it had the best magnetic properties among all of the samples. It had the highest out-of-plane coercivity of 23.4 kOe, the highest squareness and the smallest in-plane opening up as shown in Fig. 6 (b). The magnetic anisotropy constant K as function of TiN concentration, as illustrated in Fig.6 (d), were measured by Anomalous Hall effect with analyzing the normalized hall voltage curves using generalized Sucksmith-Thompson (GST) method [23]. It was found that with increase in TiN doping concentration the magnetic anisotropy decreased, implying that large lattice mismatch could improve the magnetic anisotropy. The change in magnetic anisotropy with TiN doping concentration coincide with that of chemical ordering parameter S which was estimated using the integrated peak intensity ratio (I001/I002) [24]. Ding et al. reported that the in-plane tensile strain might increase with an increase of the lattice mismatch from 2.23% to 6.33% (w.r.t. 11
intermediate layer). The strain helped in expansion of the a axis and contraction of the c axis of FePt, thus assisted the phase transformation and chemical ordering. With a mismatch larger than 6.33%, dislocations formed at the interface to release the strain, which might result in a lower degree of chemical ordering [6]. However in our experiment, the lattice constants of FePt films did not change with the increase in the lattice misfit (w.r.t. film) from 13.2% to 15.7%. There was no in-plane tensile strain since the FePt lattice was relaxed. Obviously, there was discrepancy as compared to Ding’s paper. The previous studies were conducted based on different type of substrates or intermediate layers, such as Pt, Cr, MgO etc. which may introduce more uncertainties to the system such as the interfacial interaction between these materials and FePt film. In this work, all of the FePt films were deposited on Zr xTi1-xN intermediate layer. Zr and Ti belong to the same group of element and minor change in atomic ratio of these 2 elements does not greatly affect the surface chemistry. However, the reason to its increasing chemical ordering parameter S and magnetocrystalline anisotropy with the lattice misfit beyond 13.2% still requires further investigation. The temperature dependence of saturation magnetization Ms and magnetic anisotropy constant K for 30nm thick FePt film deposited on ZrTiN intermediate layer with 25 vol. % TiN doping concentration was measured. The M-T and K-T curves with least square fit were shown in Fig. 6 (e) and (f). The M-T curve of ferromagnetic materials could be described by the general analytical function proposed by Kuz’min [25]: (2)
where m is the normalized moment; is the reduced temperature which
; s and p are
parameters, p>3/2, s>0. This equation obey Bloch’s 3/2 power law at low temperature, and when 12
approaching another limit
, m(τ) is proportional to (1- τ)1/3, as described by Heisenberg
model [26]. In our experiment, as an inter-diffusion zone was observed at the interface, the total magnetic moment μtot was considered as the summation of the body moment μbody and the deadlayer moment μdeadlayer. Each component was fitted using equation (2). Tc (body) obtained from the fitting was 733 K, and Ms body (0) was obtained to be 1042 emu/cc. The K-T curve was fitted by the theory within the framework of thermal fluctuations of magnetic moments [27], which can be expressed as: (3)
The m was obtained to be equal to 2.4, which was smaller than 3, close to 2 as predicted by single-ion model for L10 magnets [28]. The K1(0) obtained from the fitting was about 5.10 × 107 erg/cc which was comparable to magneto-crystalline anisotropy of bulk FePt (K1 = 7 × 107 erg/cc measured at room temperature [29]).
D. Discussions Similar to LME, for DME there is a critical point of the lattice misfit between the film and substrate, beyond which the film will grow textured or largely polycrystalline. The limit in FePt (10 nm)/ZrTiN system was about 15.7% (corresponding to 25 vol. % TiN). Beyond this mismatch limit, the epitaxy could not be maintained and part of the FePt film was (100) oriented. For FePt (5 nm)/ ZrTiN system, the TiN concentration in ZrTiN must be more than 50 vol. % (13.2% lattice mismatch correspondingly), otherwise the FePt film grew textured. The limit in lattice mismatch was thermodynamically determined. Based on the experimental observations, a model of FePt (fixed thickness)/ ZrxTi(1-x)N system is proposed. Energy terms for both DME and polycrystalline configuration with incoherent interface are given and compared. Because Zr and 13
Ti are the same group elements, they have the same number of valance electron. It is reasonable to assume a constant surface chemistry of ZrTiN which is independent from the TiN composition. Moreover, the ZrTiN layer is assumed rigid in this model. For a polycrystalline film with an incoherent interface, there will be no long-range misfit strain energy, but only short-range distortions caused by chemical binding at the interface [30]. Other than the chemical interfacial energy, surface energy is another term for consideration. However, these two terms are both independent from the TiN doping concentration in ZrTiN and considered as constant,
. Unfortunately, it is
very difficult to precisely calculate the total energy. What we know is that DME will occur in FePt/ZrTiN system as long as its total energy is smaller than this critical value of
.
As for a DME FePt/ZrTiN system, a model is built based on the generalized schematic configuration shown in Fig. 7 (a). In order to fulfill the traction-free boundary conditions on the free surface, a series of image dislocations with opposite sign to misfit dislocations were assumed [31]. The total energy of the system includes residual strain energy, surface energy, interfacial energy and dislocation energy. The coherent residual strain energy density in FePt film can be expressed as (4)
where E is the Young’s modulus and v is the Poisson ration of FePt [32]. h is the film thickness. For a perfect DME system, the residual strain is zero, resulting in a zero strain energy density. Even if εr is finite meaning nds cannot totally cancel out mdf, the value of strain energy is still very small and insignificant. The interfacial energy is determined by the surface chemistry. With the assumption stated earlier, the FePt/ZrTiN chemical interfacial energy should stay constant. 14
The surface energy is determined by the morphology of the film surface and its equilibrium crystal shape. It can be treated as a constant if the film thickness and interfacial energy have been fixed. The energy terms change with lattice mismatch are mainly related with the formation of dislocation. For the purpose of simplifying the calculation, the dislocation-related energy density is considered per domain. Within one domain, the self-energy density per area for edge dislocations is (5)
where (6)
G is the shear modulus of FePt; b is the Burger vector which equals to a/2[100]; d is the average domain size; R is the crystal size; and r0 is the cut-off core radius of the dislocation, which is close to one Burger vector [33]. The interaction between misfit dislocation and its image dislocation can be expressed as (7)
To sum up equations (5) and (7) we can get the total dislocation energy (8)
The final dislocation related energy term is the interaction energy between dislocations from other domains. The interaction energy between two closest neighbors can be expressed as [33]
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(9)
Since interaction energy decreases quickly with increasing distance, first order approximation is used for estimation (10)
The insertion of the extra half planes also introduce the ledge surface energy (11)
where
is the surface energy of FePt (001) plane. Therefore, according to the equations, the
total energy arises from dislocation formation
is a function of the domain
size d. In order to correlate the total dislocation-related energy with lattice misfit, the domain size
can be expressed as (12)
(13)
By substituting (12) into (13): (14)
16
(15)
For the calculation, the Young’s Modulus of 180 GPa and Poisson ratio of 0.33 for bulk FePt were used [34]. The edge dislocation energy density as a function of misfit strain regarding domain matching FePt/ZrTiN is plotted in Fig. 7 (b). The blue triangle represents the edge dislocation energy density corresponding to each perfect m/n domain match system. The full energy profile is a continuous function instead of a discrete step function. This is because, in case of imperfect domain matching, the m/n and (m+1)/(n+1) domains may alternate with a frequency to provide a perfect matching according to (16)
where α is the frequency factor [13]. Therefore, the average domain size changes continuously with the misfit strain. As shown in the plot, with a large misfit strain, the dislocation energy is largely contributed by the interaction energy. This is because when the dislocations are distributed closer, the repulsion among the parallel misfit dislocations become stronger. With the decrease in misfit strain caused by the increase in TiN doping concentration, the total dislocation energy density drops dramatically. Once the doping concentration reaches the critical value, the DME will be thermodynamically favored. In present work, when lattice mismatch was larger than 15.7%, the DME could not be retained. However, it has been reported that DME was observed in ZnO/α-Al2O3 (0001) (18.3% misfit w.r.t. ZnO) [14] and TiN/Si (100) (28.4% misfit w.r.t. TiN) [35]. To our knowledge, in order to achieve DME, it is important to keep the total interfacial energy low. The interfacial energy comprises dislocation-related energy and interfacial chemical binding. A large lattice mismatch can cause a high density of dislocation, and induce a huge dislocation-related energy. In order to 17
stabilize the DME system with a large lattice mismatch, the chemical interfacial energy must be decreased by choosing similar type of materials for these 2 layers. For example, the ZnO/αAl2O3 (0001) system already has a very large lattice mismatch of 18.3%. The strong ionic bond also causes a rigid structure which can induce an even greater dislocation energy. The DME in ZnO/α-Al2O3 (0001) system is still retained despite the large dislocation energy since it has a low oxide/oxide chemical interfacial energy. On the opposite, perpendicular anisotropy of FePt films could also been obtained non-epitaxially on amorphous substrates [36]. However, the FePt (001)/SiO2/Si film obtained by post-annealing might be at a thermodynamic metastable state, and eventually the FePt would become (111) oriented. These findings may offer a new way to deposit epitaxial films on substrates with large lattice mismatch for next-generation solid state technology and industrial applications.
IV.
SUMMARY Domain matching epitaxial (DME) FePt (001) films were deposited on a ZrxTi(1-x)N
intermediate layer with a lattice misfit no greater than 15.7% (w.r.t. film). Strain energy was released through the formation of edge dislocations at the interface of layers with mismatched lattice constants. The density of dislocations increases with the interlayer misfit, causing an increase in total energy and destabilizing the DME. The effect of large lattice mismatch on magnetic properties of FePt was investigated as well. The experiments and analysis revealed that the FePt films had good perpendicular magnetic anisotropy with large out-of-plane coercivity. The in-plane magnetic component was caused by grain tilting which is a strain relieving mechanism. The large lattice mismatch could improve both chemical ordering and as well as the magnetic anisotropy constant Ku, which is desirable for magnetic recording applications.
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ACKNOWLEDGEMENTS The research is supported in part by the Singapore National Research Foundation under CRP Award No. NRF-CRP10-2012-02 and Seagate Technology.
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Figures:
FIG. 1. (a) XRD spectra of 10 nm FePt films grown on ZrTiN intermediate layers with various TiN doping concentration, (b) XRD spectra of FePt films with various thickness grown on ZrTiN intermediate layer with 25 vol. % TiN doping concentration, (c) XRD spectra of 5 nm FePt films grown on ZrTiN intermediate layers with various TiN doping concentration, (d) FWHM Δθ50 of the rocking curves of FePt (001) peak as a function of film thickness and TiN concentration.
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FIG. 2. Reciprocal space maps around ( 3) planes of FePt grown on ZrTiN intermediate layer with (a) 0 vol. %, (b) 15 vol. % and (c) 25 vol. % TiN doping concentration, and maps of (d) 10nm and (e) 30nm FePt grown on ZrTiN intermediate layer with 25 vol. % TiN doping concentration.
23
FIG. 3. (a)Low magnification cross-sectional TEM image of FePt (10nm) / ZrTiN (25 vol.% TiN doping), (b) high resolution cross-sectional TEM image showing grain tilting, the insert is the FFT image of selected area, (c) enlarged selected area showing 7/6 domain matching epitaxy, with (200) lattice fringes and ½[200] Burger vector labelled, (d) schematic drawing of 7/6 domain matching epitaxy and (e) EDX elemental mapping at the interface. 24
FIG. 4. (a)Low magnification cross-sectional TEM image, (b) & (c) EDX elemental mapping of selected area of FePt (5nm) / ZrTiN (25 vol.% TiN doping), and (c) high resolution crosssectional TEM images of selected grains with corresponding FFT image, showing polycrystalline characteristic of the film.
25
FIG. 5. (a)Low magnification cross-sectional TEM image of FePt (5nm) / ZrTiN (50 vol.% TiN doping), (b) EDX elemental mapping of selected area, (c)high resolution cross-sectional TEM image at interface, (d) corresponding FFT image and (e) enlarged image of selected area with (200) lattice fringes and ½ [200] Burger vector labelled, and (f) schematic drawing of 8/7 domain matching epitaxy.
26
FIG. 6. (a) The hysteresis loop of FePt (10nm) / ZrTiN (25 vol.% TiN), (b) the hysteresis loop of FePt (5nm) / ZrTiN (50 vol.% TiN), (c) summary of out-of-plane coercivity of all samples, (d) the dependence of magnetic anisotropy constant Ku and integrated peak intensity ratio I001/I002 of 10nm FePt films on the TiN doping concentration in ZrTiN intermediate layer, (e) M-T and (f) K-T curves with least square fit of FePt (30nm) / ZrTiN (25 vol.% TiN).
27
FIG. 6 (a) Schematic drawing of the dislocation configuration in FePt / ZrTiN domain matching epitaxial system, and (b) the dependence of various energy densities in FePt (10nm) / ZrTiN (001) system on the misfit strain between the two layers.
28
Highlights:
L10 FePt films were epitaxially deposited on ZrxTi1-xN intermediate layer with a lattice mismatch up to 15.7%.
Detailed TEM studies showed that the domain matching epitaxy (DME) of FePt films was achieved by the formation of dislocations to reduce strain energy.
The dislocation-related energy increased with the lattice mismatch and destabilize the DME system.
The large lattice mismatch could improve both chemical ordering and as well as the magnetic anisotropy constant Ku, which is desirable for magnetic recording applications.
29