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Au multilayers
Journal of Alloys and Compounds 680 (2016) 722e728
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Ion irradiation induced phase transition of Co in Co/Au multilayers Vantari Siva a, Siddharth S. Sahu a, D.P. Datta a, P.C. Pradhan b, M. Nayak b, V. Solanki c, D. Topwal c, Kartik Senapati a, Pratap K. Sahoo a, * a
School of Physical Sciences, National Institute of Science Education and Research, Bhubaneswar, Jatni 752050, India Indus Synchrotrons Utilization Division, Raja Ramanna Centre for Advanced Technology, Indore 452013, India c Institute of Physics, Sachivalaya Marg, Bhubaneswar 751005, India b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 5 January 2016 Received in revised form 14 April 2016 Accepted 16 April 2016 Available online 20 April 2016
We have studied the structural phase transformation in thin Co films using moderate energy Au ion irradiation. We found that local energy transfer along the track of the ion beam is an effective method for transforming fcc phase of Co into technologically appealing hcp phase. As a function of the irradiation fluence, a gradual crossover to the hcp phase was observed in X-ray diffraction data. Such phase transition was explained on the basis of evolution of grains under increasing ion fluence. Bloch-Grüneisen simulations to the four contact resistivity data showed a decrease in Debye temperature with evolution of hcp-phase Cobalt, consistent with the reduced modal freedom for phonons in the hcp structure compared to the fcc structure. © 2016 Elsevier B.V. All rights reserved.
Keywords: Phase transformation Ion irradiation RBS AFM Resistivity
1. Introduction In last few decades, Cobalt, one of the strongest ferromagnetic elements with very high Curie temperature (1388 K) [1], has received much attention [2e5] in the context of various magnetoactive devices. Magnetoresistive sensors, magneto-optic recording media [3], and thin film spintronic devices [6] are of major interest. Among the various possible phases of Co, hcp and fcc are the most stable ones in thin films. However, Bulk Co exists in hcp phase at room temperature and transforms into high temperature fcc phase above ~693 K [7]. These phase transformations can be induced by the post deposition treatments such as thermal annealing [7] and ion irradiation [8e11]. Ion irradiation has the advantage of being a low temperature process along with possibility of spatial selectivity, compared to usual thermal annealing processes. Structural phase transition in Co films has been studied by Zhang et al. [9] and Yan et al. [11] using 200 keV Xe ion irradiation in Co/Fe bilayers and Ag/Co multilayers, respectively. Similarly, Zhang et al. have studied this transition by irradiating 200 keV Xe ions on Co films [10]. Gupta et al. [8] have irradiated Co thin films using 320 KeV and 120 MeV
* Corresponding author. E-mail addresses: [email protected] (K. Senapati), [email protected] (P.K. Sahoo). http://dx.doi.org/10.1016/j.jallcom.2016.04.155 0925-8388/© 2016 Elsevier B.V. All rights reserved.
Au ions, to understand the effects of nuclear and electronic energy loss on this phase transition. Johansen et al. [12] has also reported the same transition in Co single crystals by 150 keV Pb ion implantation. All these reports have found that phase transformation occurs from hcp to fcc phase of Cobalt. In contrast, here we show phase transformation of Co from fcc to hcp phase using ion irradiation. We note here that among all the stable phases of Co, hcp phase has the highest coercive field, which is desirable for ultra-high density storage applications [13]. Unlike previous studies, we have used 1.5 MeV Au ion for irradiation of Co/Au multilayer films. It is well known [14] that irradiated ions deposit energy in material primarily through two channels i.e. electronic energy loss and nuclear energy loss. While electronic energy loss is dominant at higher energies, nuclear energy loss dominates at low energies. All previous experiments on ion irradiation induced phase transformation in Co have focused on either the low energy or swift heavy ions. Here we have used moderate energy (1.5 MeV) ions which have similar contribution towards nuclear energy loss (Sn) and electronic energy loss (Se). The moderate energy range also has the added advantage of relatively lower sputtering rate of material compared to the lower energy ions. However, one can not completely rule out irradiation induced sputtering at this moderate ion energy. Therefore, we have used Au layers on the surface to minimize damage to the Co layer. The reason behind the specific choice of Au as a damage protective layer
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for Co is the positive heat of mixing (DHmix ¼ þ11 kJ/mol) [16] in the Co/Au system. This implies that Au and Co are immiscible, thereby ensuring no ion-beam induced AueCo phase formation happens at the interface of Co and Au. The overall effects of the ion irradiation on the multilayer Co/Au films, including the phase transition of cobalt from fcc to hcp phase, has been studied using glancing angle X-ray diffraction, scanning electron microscopy, atomic force microscopy, and electrical transport measurements.
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A series of multilayers of Co and Au thin films were deposited on Si substrates by e-beam evaporation technique in a high vacuum chamber with a base pressure below 1 107 mbar. The thickness of each layer was 18 nm, shown by quartz crystal monitor (QCM). The rates of deposition, 0.6 nm/min and 0.4 nm/min for Co and Au respectively, were maintained throughout the deposition for all samples. The samples were irradiated at room temperature by Au ions of 1.5 MeV energy at different fluences in the range of 5 1014 to 1 1016 ions/cm2. For 1.5 MeV Au ions, the calculated value of Se (from SRIM [15]) is ~3.3 keV/nm and Sn is ~8 keV/nm in Co. The energy of these ions was chosen in such a way that the range of these ions (~220 nm) is more than the total thickness of the films, so that the phase transition of Co would be uniform throughout the layers. The structural properties of Co and Au were investigated using glancing angle X-ray diffraction (GAXRD) with the synchrotron radiation X-ray source at Indus-2, RRCAT, Indore. The X-ray reflectivity (XRR) measurements were also performed in the pristine sample in order to verify the thickness and interface roughness of the samples. The modifications in surface topography were observed by Field emission scanning electron microscopy (FESEM). To understand the morphological changes in detail, we have performed atomic force microscopy (AFM) on the samples. The information of interlayer diffusion, sputtering and thickness was obtained from Rutherford back scattering spectroscopy (RBS). Resistivity of all the samples in the temperature range of 2 Ke300 K was measured in the standard four-probe geometry in a liquid cryogen free cryostat.
along with the simulated curve. The fitting to the reflectivity data (indicated by the solid line in Fig. 1) was performed using Bruker software. The thickness and interface roughness of Au and Co layers obtained from XRR fitting are shown in Table 1. The simulated thickness was found to be very close to the quartz crystal monitored thickness. Maximum interface roughness was found to be less than 1 nm, indicating reasonably sharp interfaces, which can be attributed to the immiscibility of Co and Au. Fig. 2(a) shows X-ray diffraction patterns of the pristine sample and irradiated ones using synchrotron radiation at different fluences. From Fig. 2(a), we find that the native phase of as deposited Co is fcc as seen for the pristine sample with only the fcc (111) peak of Co. We note here that (200) peak of cubic Au also falls at the same angle as the (111) fcc peak of Co. With the lowest irradiation dose of 5 1014 ions/cm2 we observe a small hump near 2q ¼ 22.2 , corresponding to the (1010) peak of hcp Co. In the sample irradiated at an intermediate dose of 3 1015 ions/cm2, we observe a broad and prominent peak in between the positions of hcp and fcc phases of Co, which is separately shown in Fig. 2(b). This may correspond to an intermediate phase of Co. On further increasing the fluence, a clear hcp peak segregates out of the broad intermediate peak and transforms into hcp and fcc phases of Co. However, the possibility of Co fcc-phase presence can not completely be ruled out based on the presented data only. The remaining intensity at the position of fcc (111) Co is due to the Au (200) planes. In all the samples, we have observed the cubic phase of Au with space group Fm3m. No alloy phases or compounds of CoeAu has been found after ion irradiation, due to positive heat of mixing (as mentioned earlier). Moreover, we have noticed a shift in the peak position of all the peaks towards the higher angles due to the compressive stress developed in the irradiated samples. We have extracted FWHM values for the different Au peaks and plotted in Fig. 2(c). As shown in Fig. 2(c), the FWHM values increase upto a fluence of 7 1015 ions/cm2 followed by a sharp decrease at higher fluences. An initial increase in FWHM represents random amorphization while the sharp decrease in FWHM values at higher fluences indicates re-crystallization through ion induced local melting.
3. Results and discussions
3.2. Scanning electron microscopy (SEM)
3.1. X-ray diffraction (XRD)
Surface evolution of the films with increasing ion fluence, obtained from field emission SEM, is shown in Fig. 3. Typically bulk cobalt remains in hcp phase along with a small fraction of fcc phase down to particle sizes of about 40 nm. For grain sizes below ~ 20 nm, it has been reported that Co remains almost completely in the fcc phase [7]. In the range of 20e40 nm grains, a mixed phase of Co is observed with almost equal percentage of fcc and hcp phases [7,8]. The pristine sample consists of small grains distributed uniformly over the surface. The rate of evaporation of the Co layer in the present study compared to the other reports in the literature is much smaller (at least 5 times), which is a possible reason for formation of fcc phase. In fact, it is shown in Ref. [7] that the crystal structure of evaporated Co grains is dependent upon size
2. Experimental methods
Fig. 1 shows the X-ray reflectivity data of as-deposited sample,
Table 1 Thickness and interlayer roughness values in nm obtained from XRR and RBS. Layer
Thickness XRR
Fig. 1. X-Ray reflectivity curves of as deposited bilayers. The red solid line is fitting to the data. A schematic of the bilayer is shown in the inset. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
Au Co Au Co
16.1 18.4 16.0 17.3
Roughness RBS
± ± ± ±
0.09 0.10 0.79 0.40
16.3 17.8 17.1 17.8
± ± ± ±
0.5 0.6 0.7 0.6
XRR
RBS
0.86 ± 0.03 1.0 ± 0.2 0.79 ± 0.09 0.9 ± 0.31
0.7 0.6 0.7 0.5
± ± ± ±
0.02 0.03 0.04 0.02
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Fig. 2. (a) Glancing angle XRD patterns of pristine and irradiated samples, (b) Deconvolution of the peak in 2q range of 21.5 e23.4 from the sample irradiated at a fluence of 3 1015 ions/cm2 and (c) The FWHM of the different Au peaks as a function of the irradiation fluence.
distribution of the grains. We observe that the low rate of evaporation in the present study compared to Refs. [8e11], leads to a reduction in the particle size distribution, resulting in the formation of fcc phase in pristine sample. The size of distributed particles on the surface is seen to increase (up to ~ 50 nm) at higher fluences, possibly due to ion induced local melting. Local melting takes place in the material due to energy deposition by ions which gives rise to a local annealing effect, leading to the agglomeration of small grains into the bigger ones. Consistent with this general feature of
Co grains, we observe the evolution of hcp-phase Co for increased grain size as observed from electron microscopy. An additional feature of the SEM images is the black and grey colour contrast observed on the surface of all the irradiated samples (as shown in Fig. 3(b) and (c)). The darker contrast corresponds to regions where the top Au layer is sputtered out due to irradiation. A Cross-sectional SEM image of a sample irradiated to a fluence of 1 1016 ions/cm2 is shown in Fig. 3(d), which explains that the colour contrast is indeed due to the formation of hillocks and
Fig. 3. Panels (a), (b) and (c) show the planar SEM images of pristine, irradiated films at 5 1014 and 1 1016 ions/cm2. The panel (d) represents a cross sectional view of the sample irradiated at 1 1016 ions/cm2 (Inset is the magnified image of panel d).
V. Siva et al. / Journal of Alloys and Compounds 680 (2016) 722e728
craters. This can be due to the sputtering at lower fluences, while at higher fluences, it can be attributed to a combined effect of ion induced local melting and sputtering. 3.3. Rutherford back scattering (RBS) Fig. 4 shows RBS spectra of the pristine and irradiated samples. The fitting of the data was performed using RUMP software [17,18]. As shown in Table 1, the thickness and interface roughness values obtained from RBS analysis are in very good agreement with those found from XRR data fittings. The integral intensity (which is a measure of the amount of material) of a peak corresponding to top Au layer was observed to decrease with the increasing ion fluence systematically, which is a clear indication of the sputtering (consistent with the observation from SEM). Such sputtering, though much less in amount, can also be observed in the top Co layer beyond a fluence 3 1015 ions/cm2. To understand this phenomenon in more detail, we have estimated the sputtering yields from TRIDYN [19] simulations and results are shown in Fig. 4(b). The sputtering yield of Au drops sharply beyond a fluence of 2 1015 ions/cm2. It is well known that the TRIDYN allow dynamic simulations, which takes preferential sputtering as well as stoichiometry changes into account within the collision cascade [20]. Therefore, at the higher fluences, as most of the top Au layer is sputtered, the collision cascade is expected to be altered. This explains the
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reduction in sputtering yield with increasing fluence, which is due to decrease in amount of Au (as observed from SEM and RBS studies discussed above). The sputtering yield in case of Co is almost zero at lower fluences since the Co layer is buried under the top Au layer. On the other hand, it increases at higher fluences as the Au layer is systematically sputtered out and Co layer underneath gets exposed. Moreover, a tailing effect of the peaks can be noticed in Fig. 4(a), which is not due to the interfacial mixing but due to the increase in interface roughness with the ion fluence. The possibility of interfacial mixing of CoeAu system can be safely excluded as it is an immiscible system. This is the reason why we did not notice any trace of alloy phase formation in XRD data. In fact, a tailing effect at the high energy side of Si is also observed in the RBS data at higher fluences, but the presence of alloy phases of Co and Si peaks are not observed from XRD analysis. Therefore, we can conclude that the slope noticed in the RBS spectra [21] is not due to mixing, but it is related to increased rms roughness, as found from AFM studies (discussed below). 3.4. Atomic force microscopy (AFM) To further understand the ion induced modifications of the films, we performed atomic force microscopy (AFM) on all the samples (shown in Fig. 5). As-deposited films consist of plain and uniform surface throughout the scanned region as observed from SEM. As soon as ion fluence reaches 5 1014 ions/cm2, formation of hillocks and craters can be clearly seen. The RMS roughness of the as-deposited sample is 4.6 nm, which is in good agreement with the total roughness (3.55 nm) obtained from XRR data. The rms roughness of the films irradiated at 5 1014 ions/cm2 has reached 4 times that of the pristine sample and it increases gradually beyond this ion fluence. The main reason for the increase in the roughness is the sputtering of the material from the surface, which is confirmed from RBS results. A difference in the size of the objects has been observed in the AFM images compared to the ones obtained from SEM. This can be explained as the AFM images show larger 3D features and the SEM images show details of surface morphology within such features at respective fluences. Another important aspect of AFM study is that the underlying mechanism of the surface modification can be extracted from AFM images by using power spectral density (PSD). The PSD can be mathematically defined by Ref. [22].
2 1 d2 r iq$r CðqÞ ¼ e < hðrÞ > t ∬ area 2p
(1)
where q is the spatial frequency and h(r) is the height of the surface at a point r ¼ (x, y). The PSD functions, determined using WSxM software [23], is plotted against the spatial frequency for all samples in Fig. 6. This PSD function corresponds to a surface that can be divided into two different regimes depending upon surface corrugation. In the low frequency regime (q
Fig. 4. (a) RBS spectra of the pristine and irradiated samples at the fluences 5 1014 ions/cm2, 3 1015 ions/cm2 and 1 1016 ions/cm2. (b) Variation of sputtering yields of Co and Au with ion fluence.
CðqÞ ¼
4 X 1
!1 ag qg
(2)
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V. Siva et al. / Journal of Alloys and Compounds 680 (2016) 722e728
Fig. 5. AFM images for (a) pristine and after 1.5 MeV Au ion irradiation at fluences of (b) 5 1014 ions/cm2, (c) 3 1015 ions/cm2 and (d) 1 1016 ions/cm2. Creation of Hillocks and craters after irradiation can be seen in (b). The scales are shown at the right hand side of each image and length of the scales are provided in the top-right corner of the respective images.
where ag is a constant. The exponent (g) of spatial frequency in the second regime provides the useful information of the underlying mechanism of the surface evolution. Herring et al. [25] reported using a linear dimensional analysis that the values of g are 1, 2, 3 and 4 corresponds to bulk viscous flow, evaporationcondensation, volume diffusion, and surface diffusion respectively [24]. In the pristine sample, we found the value of g to be 2, which is expected because the films are formed through the process of evaporation and condensation. For the sample irradiated to a fluence of 5 1014 ions/cm2, the value of g is 3, which indicates volume diffusion from the surface. This can be understood as the removal of material from surface which is basically the ion induced sputtering as is also observed from RBS spectra. Beyond this ion fluence, the value of g is 4, which means that the surface diffusion was driven by ion induced local melting (as noticed from XRD and SEM analysis). This is consistent with the findings from SEM and RBS.
layers as the parallel resistors. The solid lines represent the simulations to the experimental data. The Bloch-Grüneisen (BG) formula for metals is given by Ref. [26].
3.5. Electrical transport Fig. 7 shows the measured resistivity (r) as a function of temperature in the pristine sample and irradiated ones at the fluences of 5 1014 ions/cm2, 3 1015 ions/cm2 and 1 1016 ions/cm2. The temperature dependent resistivity of all the samples were simulated using Bloch-Grüneisen formula by considering the Au and Co
Fig. 6. Power spectral density (PSD) of all samples as a function of spatial frequency. The vertical arrows indicate the position of qo at respective fluences. The q2 and q4 dependences are also shown for comparison with variation of PSD.
V. Siva et al. / Journal of Alloys and Compounds 680 (2016) 722e728
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Fig. 7. Temperature dependent resistivity (r) of pristine sample and irradiated ones at 5 1014 ions/cm2, 3 1015 ions/cm2 and 1 1016 ions/cm2.The red solid lines represent the simulations of the respective experimental data. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
rðTÞ ¼ r0 þ A
T
QD
n QZD =2T 0
xn
dx ðsinhxÞ2
(3)
where r0 is the residual resistivity, A is a pre-factor of the BlochGrüneisen formula and qD is the Debye temperature of the metal. In general, the constant n takes different integral values depending on the nature of electronic interaction. In magnetic metals, in addition to the usual electron-phonon interactions, one expects electronmagnon interactions as well. Therefore, a third term was introduced to the above formula to account for resistivity due to magnon scattering in Co. In the Bloch-Grueisen formula used for simulating the resistivity data, the three exponents i.e., nCo1, nCo2 and nAu were actually constrained very close to the theoretical values because they arise from well established models of electron-phonon interaction and electron-magnon interactions. Therefore, essentially there were 4 parameters which determined the simulated curves. The simulated parameters obtained from the data are tabulated in Table 2. The values of resistivities of both Co and Au layers are seen to be increased monotonically with the ion fluence, which is expected for ion induced defect formation. From Table 2, it can be seen that Co has two different values of n i.e., 2, and 5 corresponding to electronmagnon and electron-phonon interactions respectively. Since, Au is a non-magnetic metal, it has only one value of n (equal to 5), which is due to electron-phonon interaction. The Debye temperature of Co(QCo ) for pristine sample is found to be 445 K, which is same as that of bulk Co. The values of QCo reduces continuously with the increasing ion fluence. In contrast, an increment in the Debye temperatures of Au is found, upon increasing the fluence. The Debye temperature (QD) can be expressed as the following [1].
QD ¼
1=3 Zn 6p2 N V kB
(4)
where Z, n, kB and N are reduced Planck’s constant, velocity of sound in the sample, Boltzmann constant and total number of phonons in volume V respectively. In metallic films the majority contribution to the transport behaviour comes from within the grains because electrons spend most of the time inside the grain rather than at the grain boundaries. Therefore, we have considered that number of phonon modes within the grains is the relevant parameter. Since N is decided by the crystal structure of these grains, which remains unchanged for Au after irradiation, we have considered the quantity N to be practically unchanged. Similarly, for a particular crystal structure the velocity of sound remains constant. The overall velocity of sound through the entire film of-course varies as a function of fluence. However, since the majority contribution to the electronic transport comes from individual grains one can assume it to be almost constant for practical purposes. The quantity V, on the
Table 2 Simulated parameters used in simulating r(T) by BG formula. Simulated parameters
Pristine 5 1014 ions/ cm2
3 1015 ions/ cm2
1 1016 ions/ cm2
rCo(107 U-m) rAu(107 U-m)
2.2 1.7 5 2.05 5 445 120
6.7 5.5 5 2.2 5 235 155
8.5 50 5 2.25 5 210 160
nCo1 nCo2 nAu QCo(K) QAu(K)
2.4 0.9 5 2.05 5 335 130
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V. Siva et al. / Journal of Alloys and Compounds 680 (2016) 722e728
from resistivity data also corroborated to the observed fcc to hcp phase transformation in Co. A major advantage of this method is the capability of selective area modification through patterned mask ion irradiation. The coercivity of the irradiated and nonirradiated regions will be significantly different which can be mapped through MFM or other local magnetic probes.
Acknowledgements
Fig. 8. Fluence dependent variation of Co Debye temperature (left side axis) and normalized intensity of hcp Co peak (right side axis). Black and blue dashed lines are guides to the eye. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
The authors acknowledge the funding from National Institute of Science Education and Research (NISER), Department of Atomic Energy (DAE), India. We thank G. Santosh babu for the useful discussions and help in XRR analysis. We also thank all members of Ion Beam Laboratory of Institute of Physics, Bhubaneswar for irradiation and RBS measurements. We thank A. K. Sinha and all members of beam line-12, Indus-2, RRCAT for their help during the XRD measurements. Thanks are due to Sudipta Mahana for her help during the transport measurements.
References other hand, changed significantly for Au as indicated by the RBS data. Therefore, we believe the change in Debye temperature is primarily a consequence of changing volume of the Au film. Unlike Au, Table 2 shows a decrease in Debye temperature for the Co layer as a function of ion fluence. However, the initial decrease saturates at the highest fluences, as evident from Fig. 8. On the right hand axis of Fig. 8 we have plotted the normalized intensity of the (1010) peak of hcp Co, from Fig. 1(a). Clearly, the appearance of the hcp phase of Co has a direct relationship with the decrease in the Debye temperature. Although both fcc and hcp structures are close packing structures with a packing fractions of 0.74 and a co-ordination number 12, the atomic stacking patterns are different. Along the c-axis of the conventional unit cell, hcp structure has two types of layers (atomic arrangements) stacked alternatively, whereas the fcc structure has three types of layers [1] along the c-axis. Therefore one would expect a reduced number of allowed phonon modes in hcp structures compared to the fcc structures. Referring to Eq. (4), reduced phonon modes results in reduced Debye temperature, as observed in our case. Debye temperatures can be altered for different phases of a material, when the lattice structures are different [27]. 4. Conclusions In conclusion we have shown that medium energy heavy ion irradiation is a suitable way to obtain fcc to hcp phase transition in Co thin films. Due to the competing electronic and nuclear energy loss mechanisms at 1.5 MeV, amorphization of the films were observed followed by re-crystallization at higher fluences. Above a fluence of 5 1014 ions/cm2 nucleation of hcp phase started which saturated beyond a dose of 7 1015 ion/cm2. The phase evolution was confirmed by various complementary characterization techniques. Debye temperatures of Co as a function of fluence, extracted
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Ion irradiation induced phase transition of Co in Co/Au multilayers Vantari Siva a, Siddharth S. Sahu a, D.P. Datta a, P.C. Pradhan b, M. Nayak b, V. Solanki c, D. Topwal c, Kartik Senapati a, Pratap K. Sahoo a, * a
School of Physical Sciences, National Institute of Science Education and Research, Bhubaneswar, Jatni 752050, India Indus Synchrotrons Utilization Division, Raja Ramanna Centre for Advanced Technology, Indore 452013, India c Institute of Physics, Sachivalaya Marg, Bhubaneswar 751005, India b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 5 January 2016 Received in revised form 14 April 2016 Accepted 16 April 2016 Available online 20 April 2016
We have studied the structural phase transformation in thin Co films using moderate energy Au ion irradiation. We found that local energy transfer along the track of the ion beam is an effective method for transforming fcc phase of Co into technologically appealing hcp phase. As a function of the irradiation fluence, a gradual crossover to the hcp phase was observed in X-ray diffraction data. Such phase transition was explained on the basis of evolution of grains under increasing ion fluence. Bloch-Grüneisen simulations to the four contact resistivity data showed a decrease in Debye temperature with evolution of hcp-phase Cobalt, consistent with the reduced modal freedom for phonons in the hcp structure compared to the fcc structure. © 2016 Elsevier B.V. All rights reserved.
Keywords: Phase transformation Ion irradiation RBS AFM Resistivity
1. Introduction In last few decades, Cobalt, one of the strongest ferromagnetic elements with very high Curie temperature (1388 K) [1], has received much attention [2e5] in the context of various magnetoactive devices. Magnetoresistive sensors, magneto-optic recording media [3], and thin film spintronic devices [6] are of major interest. Among the various possible phases of Co, hcp and fcc are the most stable ones in thin films. However, Bulk Co exists in hcp phase at room temperature and transforms into high temperature fcc phase above ~693 K [7]. These phase transformations can be induced by the post deposition treatments such as thermal annealing [7] and ion irradiation [8e11]. Ion irradiation has the advantage of being a low temperature process along with possibility of spatial selectivity, compared to usual thermal annealing processes. Structural phase transition in Co films has been studied by Zhang et al. [9] and Yan et al. [11] using 200 keV Xe ion irradiation in Co/Fe bilayers and Ag/Co multilayers, respectively. Similarly, Zhang et al. have studied this transition by irradiating 200 keV Xe ions on Co films [10]. Gupta et al. [8] have irradiated Co thin films using 320 KeV and 120 MeV
* Corresponding author. E-mail addresses: [email protected] (K. Senapati), [email protected] (P.K. Sahoo). http://dx.doi.org/10.1016/j.jallcom.2016.04.155 0925-8388/© 2016 Elsevier B.V. All rights reserved.
Au ions, to understand the effects of nuclear and electronic energy loss on this phase transition. Johansen et al. [12] has also reported the same transition in Co single crystals by 150 keV Pb ion implantation. All these reports have found that phase transformation occurs from hcp to fcc phase of Cobalt. In contrast, here we show phase transformation of Co from fcc to hcp phase using ion irradiation. We note here that among all the stable phases of Co, hcp phase has the highest coercive field, which is desirable for ultra-high density storage applications [13]. Unlike previous studies, we have used 1.5 MeV Au ion for irradiation of Co/Au multilayer films. It is well known [14] that irradiated ions deposit energy in material primarily through two channels i.e. electronic energy loss and nuclear energy loss. While electronic energy loss is dominant at higher energies, nuclear energy loss dominates at low energies. All previous experiments on ion irradiation induced phase transformation in Co have focused on either the low energy or swift heavy ions. Here we have used moderate energy (1.5 MeV) ions which have similar contribution towards nuclear energy loss (Sn) and electronic energy loss (Se). The moderate energy range also has the added advantage of relatively lower sputtering rate of material compared to the lower energy ions. However, one can not completely rule out irradiation induced sputtering at this moderate ion energy. Therefore, we have used Au layers on the surface to minimize damage to the Co layer. The reason behind the specific choice of Au as a damage protective layer
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for Co is the positive heat of mixing (DHmix ¼ þ11 kJ/mol) [16] in the Co/Au system. This implies that Au and Co are immiscible, thereby ensuring no ion-beam induced AueCo phase formation happens at the interface of Co and Au. The overall effects of the ion irradiation on the multilayer Co/Au films, including the phase transition of cobalt from fcc to hcp phase, has been studied using glancing angle X-ray diffraction, scanning electron microscopy, atomic force microscopy, and electrical transport measurements.
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A series of multilayers of Co and Au thin films were deposited on Si substrates by e-beam evaporation technique in a high vacuum chamber with a base pressure below 1 107 mbar. The thickness of each layer was 18 nm, shown by quartz crystal monitor (QCM). The rates of deposition, 0.6 nm/min and 0.4 nm/min for Co and Au respectively, were maintained throughout the deposition for all samples. The samples were irradiated at room temperature by Au ions of 1.5 MeV energy at different fluences in the range of 5 1014 to 1 1016 ions/cm2. For 1.5 MeV Au ions, the calculated value of Se (from SRIM [15]) is ~3.3 keV/nm and Sn is ~8 keV/nm in Co. The energy of these ions was chosen in such a way that the range of these ions (~220 nm) is more than the total thickness of the films, so that the phase transition of Co would be uniform throughout the layers. The structural properties of Co and Au were investigated using glancing angle X-ray diffraction (GAXRD) with the synchrotron radiation X-ray source at Indus-2, RRCAT, Indore. The X-ray reflectivity (XRR) measurements were also performed in the pristine sample in order to verify the thickness and interface roughness of the samples. The modifications in surface topography were observed by Field emission scanning electron microscopy (FESEM). To understand the morphological changes in detail, we have performed atomic force microscopy (AFM) on the samples. The information of interlayer diffusion, sputtering and thickness was obtained from Rutherford back scattering spectroscopy (RBS). Resistivity of all the samples in the temperature range of 2 Ke300 K was measured in the standard four-probe geometry in a liquid cryogen free cryostat.
along with the simulated curve. The fitting to the reflectivity data (indicated by the solid line in Fig. 1) was performed using Bruker software. The thickness and interface roughness of Au and Co layers obtained from XRR fitting are shown in Table 1. The simulated thickness was found to be very close to the quartz crystal monitored thickness. Maximum interface roughness was found to be less than 1 nm, indicating reasonably sharp interfaces, which can be attributed to the immiscibility of Co and Au. Fig. 2(a) shows X-ray diffraction patterns of the pristine sample and irradiated ones using synchrotron radiation at different fluences. From Fig. 2(a), we find that the native phase of as deposited Co is fcc as seen for the pristine sample with only the fcc (111) peak of Co. We note here that (200) peak of cubic Au also falls at the same angle as the (111) fcc peak of Co. With the lowest irradiation dose of 5 1014 ions/cm2 we observe a small hump near 2q ¼ 22.2 , corresponding to the (1010) peak of hcp Co. In the sample irradiated at an intermediate dose of 3 1015 ions/cm2, we observe a broad and prominent peak in between the positions of hcp and fcc phases of Co, which is separately shown in Fig. 2(b). This may correspond to an intermediate phase of Co. On further increasing the fluence, a clear hcp peak segregates out of the broad intermediate peak and transforms into hcp and fcc phases of Co. However, the possibility of Co fcc-phase presence can not completely be ruled out based on the presented data only. The remaining intensity at the position of fcc (111) Co is due to the Au (200) planes. In all the samples, we have observed the cubic phase of Au with space group Fm3m. No alloy phases or compounds of CoeAu has been found after ion irradiation, due to positive heat of mixing (as mentioned earlier). Moreover, we have noticed a shift in the peak position of all the peaks towards the higher angles due to the compressive stress developed in the irradiated samples. We have extracted FWHM values for the different Au peaks and plotted in Fig. 2(c). As shown in Fig. 2(c), the FWHM values increase upto a fluence of 7 1015 ions/cm2 followed by a sharp decrease at higher fluences. An initial increase in FWHM represents random amorphization while the sharp decrease in FWHM values at higher fluences indicates re-crystallization through ion induced local melting.
3. Results and discussions
3.2. Scanning electron microscopy (SEM)
3.1. X-ray diffraction (XRD)
Surface evolution of the films with increasing ion fluence, obtained from field emission SEM, is shown in Fig. 3. Typically bulk cobalt remains in hcp phase along with a small fraction of fcc phase down to particle sizes of about 40 nm. For grain sizes below ~ 20 nm, it has been reported that Co remains almost completely in the fcc phase [7]. In the range of 20e40 nm grains, a mixed phase of Co is observed with almost equal percentage of fcc and hcp phases [7,8]. The pristine sample consists of small grains distributed uniformly over the surface. The rate of evaporation of the Co layer in the present study compared to the other reports in the literature is much smaller (at least 5 times), which is a possible reason for formation of fcc phase. In fact, it is shown in Ref. [7] that the crystal structure of evaporated Co grains is dependent upon size
2. Experimental methods
Fig. 1 shows the X-ray reflectivity data of as-deposited sample,
Table 1 Thickness and interlayer roughness values in nm obtained from XRR and RBS. Layer
Thickness XRR
Fig. 1. X-Ray reflectivity curves of as deposited bilayers. The red solid line is fitting to the data. A schematic of the bilayer is shown in the inset. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
Au Co Au Co
16.1 18.4 16.0 17.3
Roughness RBS
± ± ± ±
0.09 0.10 0.79 0.40
16.3 17.8 17.1 17.8
± ± ± ±
0.5 0.6 0.7 0.6
XRR
RBS
0.86 ± 0.03 1.0 ± 0.2 0.79 ± 0.09 0.9 ± 0.31
0.7 0.6 0.7 0.5
± ± ± ±
0.02 0.03 0.04 0.02
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Fig. 2. (a) Glancing angle XRD patterns of pristine and irradiated samples, (b) Deconvolution of the peak in 2q range of 21.5 e23.4 from the sample irradiated at a fluence of 3 1015 ions/cm2 and (c) The FWHM of the different Au peaks as a function of the irradiation fluence.
distribution of the grains. We observe that the low rate of evaporation in the present study compared to Refs. [8e11], leads to a reduction in the particle size distribution, resulting in the formation of fcc phase in pristine sample. The size of distributed particles on the surface is seen to increase (up to ~ 50 nm) at higher fluences, possibly due to ion induced local melting. Local melting takes place in the material due to energy deposition by ions which gives rise to a local annealing effect, leading to the agglomeration of small grains into the bigger ones. Consistent with this general feature of
Co grains, we observe the evolution of hcp-phase Co for increased grain size as observed from electron microscopy. An additional feature of the SEM images is the black and grey colour contrast observed on the surface of all the irradiated samples (as shown in Fig. 3(b) and (c)). The darker contrast corresponds to regions where the top Au layer is sputtered out due to irradiation. A Cross-sectional SEM image of a sample irradiated to a fluence of 1 1016 ions/cm2 is shown in Fig. 3(d), which explains that the colour contrast is indeed due to the formation of hillocks and
Fig. 3. Panels (a), (b) and (c) show the planar SEM images of pristine, irradiated films at 5 1014 and 1 1016 ions/cm2. The panel (d) represents a cross sectional view of the sample irradiated at 1 1016 ions/cm2 (Inset is the magnified image of panel d).
V. Siva et al. / Journal of Alloys and Compounds 680 (2016) 722e728
craters. This can be due to the sputtering at lower fluences, while at higher fluences, it can be attributed to a combined effect of ion induced local melting and sputtering. 3.3. Rutherford back scattering (RBS) Fig. 4 shows RBS spectra of the pristine and irradiated samples. The fitting of the data was performed using RUMP software [17,18]. As shown in Table 1, the thickness and interface roughness values obtained from RBS analysis are in very good agreement with those found from XRR data fittings. The integral intensity (which is a measure of the amount of material) of a peak corresponding to top Au layer was observed to decrease with the increasing ion fluence systematically, which is a clear indication of the sputtering (consistent with the observation from SEM). Such sputtering, though much less in amount, can also be observed in the top Co layer beyond a fluence 3 1015 ions/cm2. To understand this phenomenon in more detail, we have estimated the sputtering yields from TRIDYN [19] simulations and results are shown in Fig. 4(b). The sputtering yield of Au drops sharply beyond a fluence of 2 1015 ions/cm2. It is well known that the TRIDYN allow dynamic simulations, which takes preferential sputtering as well as stoichiometry changes into account within the collision cascade [20]. Therefore, at the higher fluences, as most of the top Au layer is sputtered, the collision cascade is expected to be altered. This explains the
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reduction in sputtering yield with increasing fluence, which is due to decrease in amount of Au (as observed from SEM and RBS studies discussed above). The sputtering yield in case of Co is almost zero at lower fluences since the Co layer is buried under the top Au layer. On the other hand, it increases at higher fluences as the Au layer is systematically sputtered out and Co layer underneath gets exposed. Moreover, a tailing effect of the peaks can be noticed in Fig. 4(a), which is not due to the interfacial mixing but due to the increase in interface roughness with the ion fluence. The possibility of interfacial mixing of CoeAu system can be safely excluded as it is an immiscible system. This is the reason why we did not notice any trace of alloy phase formation in XRD data. In fact, a tailing effect at the high energy side of Si is also observed in the RBS data at higher fluences, but the presence of alloy phases of Co and Si peaks are not observed from XRD analysis. Therefore, we can conclude that the slope noticed in the RBS spectra [21] is not due to mixing, but it is related to increased rms roughness, as found from AFM studies (discussed below). 3.4. Atomic force microscopy (AFM) To further understand the ion induced modifications of the films, we performed atomic force microscopy (AFM) on all the samples (shown in Fig. 5). As-deposited films consist of plain and uniform surface throughout the scanned region as observed from SEM. As soon as ion fluence reaches 5 1014 ions/cm2, formation of hillocks and craters can be clearly seen. The RMS roughness of the as-deposited sample is 4.6 nm, which is in good agreement with the total roughness (3.55 nm) obtained from XRR data. The rms roughness of the films irradiated at 5 1014 ions/cm2 has reached 4 times that of the pristine sample and it increases gradually beyond this ion fluence. The main reason for the increase in the roughness is the sputtering of the material from the surface, which is confirmed from RBS results. A difference in the size of the objects has been observed in the AFM images compared to the ones obtained from SEM. This can be explained as the AFM images show larger 3D features and the SEM images show details of surface morphology within such features at respective fluences. Another important aspect of AFM study is that the underlying mechanism of the surface modification can be extracted from AFM images by using power spectral density (PSD). The PSD can be mathematically defined by Ref. [22].
2 1 d2 r iq$r CðqÞ ¼ e < hðrÞ > t ∬ area 2p
(1)
where q is the spatial frequency and h(r) is the height of the surface at a point r ¼ (x, y). The PSD functions, determined using WSxM software [23], is plotted against the spatial frequency for all samples in Fig. 6. This PSD function corresponds to a surface that can be divided into two different regimes depending upon surface corrugation. In the low frequency regime (q
Fig. 4. (a) RBS spectra of the pristine and irradiated samples at the fluences 5 1014 ions/cm2, 3 1015 ions/cm2 and 1 1016 ions/cm2. (b) Variation of sputtering yields of Co and Au with ion fluence.
CðqÞ ¼
4 X 1
!1 ag qg
(2)
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Fig. 5. AFM images for (a) pristine and after 1.5 MeV Au ion irradiation at fluences of (b) 5 1014 ions/cm2, (c) 3 1015 ions/cm2 and (d) 1 1016 ions/cm2. Creation of Hillocks and craters after irradiation can be seen in (b). The scales are shown at the right hand side of each image and length of the scales are provided in the top-right corner of the respective images.
where ag is a constant. The exponent (g) of spatial frequency in the second regime provides the useful information of the underlying mechanism of the surface evolution. Herring et al. [25] reported using a linear dimensional analysis that the values of g are 1, 2, 3 and 4 corresponds to bulk viscous flow, evaporationcondensation, volume diffusion, and surface diffusion respectively [24]. In the pristine sample, we found the value of g to be 2, which is expected because the films are formed through the process of evaporation and condensation. For the sample irradiated to a fluence of 5 1014 ions/cm2, the value of g is 3, which indicates volume diffusion from the surface. This can be understood as the removal of material from surface which is basically the ion induced sputtering as is also observed from RBS spectra. Beyond this ion fluence, the value of g is 4, which means that the surface diffusion was driven by ion induced local melting (as noticed from XRD and SEM analysis). This is consistent with the findings from SEM and RBS.
layers as the parallel resistors. The solid lines represent the simulations to the experimental data. The Bloch-Grüneisen (BG) formula for metals is given by Ref. [26].
3.5. Electrical transport Fig. 7 shows the measured resistivity (r) as a function of temperature in the pristine sample and irradiated ones at the fluences of 5 1014 ions/cm2, 3 1015 ions/cm2 and 1 1016 ions/cm2. The temperature dependent resistivity of all the samples were simulated using Bloch-Grüneisen formula by considering the Au and Co
Fig. 6. Power spectral density (PSD) of all samples as a function of spatial frequency. The vertical arrows indicate the position of qo at respective fluences. The q2 and q4 dependences are also shown for comparison with variation of PSD.
V. Siva et al. / Journal of Alloys and Compounds 680 (2016) 722e728
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Fig. 7. Temperature dependent resistivity (r) of pristine sample and irradiated ones at 5 1014 ions/cm2, 3 1015 ions/cm2 and 1 1016 ions/cm2.The red solid lines represent the simulations of the respective experimental data. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
rðTÞ ¼ r0 þ A
T
QD
n QZD =2T 0
xn
dx ðsinhxÞ2
(3)
where r0 is the residual resistivity, A is a pre-factor of the BlochGrüneisen formula and qD is the Debye temperature of the metal. In general, the constant n takes different integral values depending on the nature of electronic interaction. In magnetic metals, in addition to the usual electron-phonon interactions, one expects electronmagnon interactions as well. Therefore, a third term was introduced to the above formula to account for resistivity due to magnon scattering in Co. In the Bloch-Grueisen formula used for simulating the resistivity data, the three exponents i.e., nCo1, nCo2 and nAu were actually constrained very close to the theoretical values because they arise from well established models of electron-phonon interaction and electron-magnon interactions. Therefore, essentially there were 4 parameters which determined the simulated curves. The simulated parameters obtained from the data are tabulated in Table 2. The values of resistivities of both Co and Au layers are seen to be increased monotonically with the ion fluence, which is expected for ion induced defect formation. From Table 2, it can be seen that Co has two different values of n i.e., 2, and 5 corresponding to electronmagnon and electron-phonon interactions respectively. Since, Au is a non-magnetic metal, it has only one value of n (equal to 5), which is due to electron-phonon interaction. The Debye temperature of Co(QCo ) for pristine sample is found to be 445 K, which is same as that of bulk Co. The values of QCo reduces continuously with the increasing ion fluence. In contrast, an increment in the Debye temperatures of Au is found, upon increasing the fluence. The Debye temperature (QD) can be expressed as the following [1].
QD ¼
1=3 Zn 6p2 N V kB
(4)
where Z, n, kB and N are reduced Planck’s constant, velocity of sound in the sample, Boltzmann constant and total number of phonons in volume V respectively. In metallic films the majority contribution to the transport behaviour comes from within the grains because electrons spend most of the time inside the grain rather than at the grain boundaries. Therefore, we have considered that number of phonon modes within the grains is the relevant parameter. Since N is decided by the crystal structure of these grains, which remains unchanged for Au after irradiation, we have considered the quantity N to be practically unchanged. Similarly, for a particular crystal structure the velocity of sound remains constant. The overall velocity of sound through the entire film of-course varies as a function of fluence. However, since the majority contribution to the electronic transport comes from individual grains one can assume it to be almost constant for practical purposes. The quantity V, on the
Table 2 Simulated parameters used in simulating r(T) by BG formula. Simulated parameters
Pristine 5 1014 ions/ cm2
3 1015 ions/ cm2
1 1016 ions/ cm2
rCo(107 U-m) rAu(107 U-m)
2.2 1.7 5 2.05 5 445 120
6.7 5.5 5 2.2 5 235 155
8.5 50 5 2.25 5 210 160
nCo1 nCo2 nAu QCo(K) QAu(K)
2.4 0.9 5 2.05 5 335 130
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from resistivity data also corroborated to the observed fcc to hcp phase transformation in Co. A major advantage of this method is the capability of selective area modification through patterned mask ion irradiation. The coercivity of the irradiated and nonirradiated regions will be significantly different which can be mapped through MFM or other local magnetic probes.
Acknowledgements
Fig. 8. Fluence dependent variation of Co Debye temperature (left side axis) and normalized intensity of hcp Co peak (right side axis). Black and blue dashed lines are guides to the eye. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
The authors acknowledge the funding from National Institute of Science Education and Research (NISER), Department of Atomic Energy (DAE), India. We thank G. Santosh babu for the useful discussions and help in XRR analysis. We also thank all members of Ion Beam Laboratory of Institute of Physics, Bhubaneswar for irradiation and RBS measurements. We thank A. K. Sinha and all members of beam line-12, Indus-2, RRCAT for their help during the XRD measurements. Thanks are due to Sudipta Mahana for her help during the transport measurements.
References other hand, changed significantly for Au as indicated by the RBS data. Therefore, we believe the change in Debye temperature is primarily a consequence of changing volume of the Au film. Unlike Au, Table 2 shows a decrease in Debye temperature for the Co layer as a function of ion fluence. However, the initial decrease saturates at the highest fluences, as evident from Fig. 8. On the right hand axis of Fig. 8 we have plotted the normalized intensity of the (1010) peak of hcp Co, from Fig. 1(a). Clearly, the appearance of the hcp phase of Co has a direct relationship with the decrease in the Debye temperature. Although both fcc and hcp structures are close packing structures with a packing fractions of 0.74 and a co-ordination number 12, the atomic stacking patterns are different. Along the c-axis of the conventional unit cell, hcp structure has two types of layers (atomic arrangements) stacked alternatively, whereas the fcc structure has three types of layers [1] along the c-axis. Therefore one would expect a reduced number of allowed phonon modes in hcp structures compared to the fcc structures. Referring to Eq. (4), reduced phonon modes results in reduced Debye temperature, as observed in our case. Debye temperatures can be altered for different phases of a material, when the lattice structures are different [27]. 4. Conclusions In conclusion we have shown that medium energy heavy ion irradiation is a suitable way to obtain fcc to hcp phase transition in Co thin films. Due to the competing electronic and nuclear energy loss mechanisms at 1.5 MeV, amorphization of the films were observed followed by re-crystallization at higher fluences. Above a fluence of 5 1014 ions/cm2 nucleation of hcp phase started which saturated beyond a dose of 7 1015 ion/cm2. The phase evolution was confirmed by various complementary characterization techniques. Debye temperatures of Co as a function of fluence, extracted
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