MICROSTRUCTURAL DESIGN OF NANOMULTILAYERS
(FROM STEEL TO MAGNETICS)
Greg Jan Kusinski and Gareth Thomas
MMFX Technologies Corporation
Irvine, CA 92612, USA
ABSTRACT
The development of high-tech materials requires functional, multicomponent microstructures including
multilayers, designed, processed and controlled at the micron and nanometer levels. It is now quite well
recognized that to optimize and design materials for specified properties, materials are best utilized as
composites. The nature of the components, their structures, morphologies and interfacial characteristics are
most important. In particular, multilayered nano-structures are attractive for mechanical and many
functional properties (e.g. magnetic). In order to understand the properties of any such multilayered system,
and hence to be able to design pre-determined sets of properties, it is necessary to know their structure. For
this reason, characterization of physical, chemical, and magnetic structures at relevant length scales is of
particular importance.
In this paper, attention is drawn to multilayered retained austenite/martensite (γret/MS) steels and
multilayered Co/Pt films for high density magnetic recording.
INTRODUCTION
Materials Science and engineering is concerned with understanding the relationships between processing,
microstructure, and properties. With this understanding it is now possible to manipulate the microstructure
down to the nano size, so as to achieve a set of defined properties. Since it is normally impossible to obtain
such sets of properties in a monotonic monolayer, it is necessary to design the needed properties via
composites within which the multilayer becomes an attractive morphology, and one in which control of the
interfacial structure and bonding can be achieved.
Clearly, the ability to characterize such nano materials requires advanced imaging and probe methods,
within which high resolution (now at the atomic level), is essential. The results given in this paper, have
involved detailed applications of electron microscopy, diffraction and microanalyses [1,2].
There are size limits that must be considered. Often no recognizable structures can be resolved when layer
thicknesses are about <10 Å [3] and nano sized components do not necessarily have the same crystal
structure as in bulk. Crystal structure determination is non-trivial and atomic imaging for real space analysis
is becoming a viable approach [4]. Also as the dimensions of the components in the multilayers, approach
nano scale and below, there is a marked increase in surface to volume ratios. Thus unique properties would
be logically expected in such small scale composites
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In this paper we present results to show the microstructural similarity of nano composite structures in
“traditional materials” such as steels (to improve the mechanical and corrosion properties) and in the
currently developing vertical magnetic recording materials for improving the information storage capacity of
magnetic recording materials.
The following abbreviations will be used through out this paper: Transmission Electron Microscopy (TEM);
Bright Field imaging in TEM (BF); Dark Field imaging in TEM (DF); Selected Area Diffraction (SAD );
High Resolution Transmission Electron Microscopy (HRTEM); Multilayer (ML); Perpendicular Magnetic
Anisotropy (PMA); martensite (MS), austenite (γ) martensite start transformation temperature (TMS), growth
temperature (TG), Co layer thickness (tCo), coercivity HC.
NANO-LAYER MICROSTRUCTURES
Steels-microcomposite martensite
The austenite → martensite reaction in steels produces two main types of transformation products:- twinned
plates when the transformation occurs at low temperatures (below about 250C), and packets of dislocated
laths at higher temperatures (above about 320C). Figure 1 shows a schematic of the multilayer structure that
is desired for improved properties. Crystallographic analysis, see Figure 2, shows that the orientation
relationship is that of the Kurdjumov-Sachs (K/S) with {111}γ habits. Thus there is a maximum of four
martensitic packets of laths per prior austenite grains. Notice also that the microstructure is designed to be
free from precipitates such as carbides, carbonitrides etc. The basis for this is to greatly reduce the formation
of microgalvanic cells between particles and ferritic matrix, which is necessary to improve corrosion
resistance, especially in saline solutions. The current steels used in construction are generally ferrite/pearlite
mixtures such as in ASTM A615 steels, which are unsatisfactory for good corrosion resistance (see [5]).
Figure 1. Schematic of Microcomposite steels.
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Figure 2 shows an actual example of the multilayered austenite/martensite compostie -note the scale of the
multilayer structure in the packet. Many investigations have been carried out on martensitic steels, but it
was not until 1973 that the microcomposite nature of packet martensite was elucidated [6-8]. The
microstructure can be rationalized in terms of the temperature dependence of the resolved shear stresses for
slip vs. twinning—the latter being preferred at low temperatures [8]. Twinned martensites are relatively
brittle, and are generally to be avoided in designing for toughness at high strength levels [9,10]. However
since the relationship between composition and martensite start transformation temperature (TMS) is quite
well known, it is possible to design a packet multilayer martensitic steel by choosing compositions which
keep TMS above about 320ºC [8]. This restricts the %C to be usually below about 0.35 wt%. Typical alloys
include the Fe/Cr/Mn/C and Fe/Si/C systems [7-10]. A model to explain the multilayer martensite/austenite
composite has been published recently and is illustrated in Figure 3 [11].
The multilayer composite contains the tough work-hardened laths of martensite linked coherently (with the
Kurdjumov-Sachs K/S orientation relation) to the untransformed, ductile austenite, giving a packet of
alternating layers of austenite (~50 atoms wide) and dislocated martensite laths. Unlike pearlitic/bainitic
microstructures, there are no carbides or other particles, and the high strength, high toughness derives from
this microstructure. A variant would be dual/triple phase steels (DFM) in which the packets are mixed with
ferrite. In all these alloys, the key to attractive properties is in maintaining the multilayer austenitemartensite composites in the packets by keeping TMS above ~320ºC, and cooling fast enough to avoid nonmartensitic products. Such microstructures offer a range of attractive properties [12-14], and some examples
are summarized below and in Figure 4, which compares some properties of the multilayered steels with
commercial ASTM A615 steel (ferrite – pearlite).
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Figure 2: Transmission electron micrographs showing the interlath retained austenite. (a) Bright field
image – packet lath martensite. (b) Corresponding dark field image formed with 111 γ reflection – thin
films of retained austenite reverse contrast. (c) SAD pattern showing K-S relationship of martensite
and austenite and hence, coherent interfaces. DF image (b) was formed with 111 γ reflection.
4
Figure 3: Schematic to illustrate the transformation of austenite (γ) to martensite (Ms) and the
stabilization of untransformed austenite between martensite laths. (a) “entrapment” of stabilized
untransformed austenite between laths in the martensite packets, (b) “Final State”, Martensite and
austenite layers not to scale.
Figure 4: (a) The tensile properties of Microcomposite steels compared to ASTM A615-ferrite
pearlite steel, (b) the superior fatigue properties compared to A615, (c) superior DBTT and low
temperature Charpy toughness, (c) corrosion samples after several weeks in 3% chloride. Note the
A615 steel is almost totally destroyed by corrosion.
Some properties
(A) Variable strength may be designed by varying the %C in the laths since this is a linear relationship, or in
DFM steels by varying the % packets+ferrite. Note there is no strict yield point, Figure 4(a), in the stressstrain curves from such structures since the laths are already rich in dislocations, and so the steel is already
in the plastic condition. Typically the 0.2% offset in the stress-strain curve is taken as the “yield stress”, and
the high tensile to yield ratios (1.5 or more), indicate excellent cold formability, e.g. for sheets (autos), wire
and strands. This microstructure and its coherency allows plastic relaxation ahead of cracks, giving K1C
values > 100 ksi√ins [8-10].
(B) The multilayered structure contains no precipitates (carbides etc.) and hence there are few microgalvanic
corrosion pitting centers. Thus there is a considerable gain in corrosion resistance especially in saline
conditions. These benefits are now being utilized in new and repaired infrastructures in which corrosion
limits the lifetime of steel in concrete structures (See the web site: [5]). It is an astounding fact that
infrastructure repair costs are estimated to be approximately 4 trillions $ over the next decade.
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Co/Pt MAGNETIC MULTILAYERS
Basis
The realization in the past decade or so that many important magnetic properties are microstructure sensitive
has led to rapid developments in the field of magnetic devices, involving atomically engineered thin films,
nanostructures and multilayers. In these composites the structure of the interfaces between dissimilar
materials or/and at the grain boundaries governs the novel properties (e.g. perpendicular magnetic
anisotropy, giant magnetic resistance, etc.) [15]. In particular, magnetic multilayers (MLs) composed of
modulated ferromagnetic-nonmagnetic layers such as Co/Pt MLs [16,17] with large perpendicular magnetic
anisotropy (PMA) and high coercivity have recently been proposed as future magnetic media in Terabit/in2
magnetic recording systems [18]. In addition, recent interest in Co/Pt was sparked by the discovery that the
magnetic properties can be locally modified by ion-beam irradiation [19,20]. Characterization of the
microstructure and the magnetic domain structure in such films is important from a technological as well as
a fundamental perspective [21-23].
In this paper, the influence of growth temperature and multilayer thickness on the structure and magnetic
properties of Co/Pt MLs are discussed. Other parameters which affect the magnetic performance of the MLs
include magnetic patterning by ion irradiation which allows the separation of the vertical and in-plane
magnetization. Further details can be found elsewhere [21,24-26].
Microstructural characterization
For the specific experiments discussed here, Co/Pt multilayers with representative structure: / 20 nm Pt seed
/ N×(t nm Co/1 nm Pt) / 1 nm Pt cap layer /, were fabricated by electron beam evaporation using a 10-8 – 10-9
Torr base pressure deposition system [27,28].
Figure 5 shows a plan-view, bright field Transmission Electron Microscopy (TEM) image and Selected Area
Diffraction (SAD) patterns of a Co/Pt multilayer grown at 250°C. The fine structure visible in some of the
grains is attributed to Moiré fringes caused by the small lattice mismatch between Co and Pt. The plan-view
SAD pattern shows the typical ring spacing associated with polycrystalline face-centered cubic (fcc) Pt
except for ring splitting due to the presence of highly strained Co layers, which will be addressed later. The
intensity distribution of the rings indicates a strong <111> texture with only some grains oriented randomly
and contributing to the (111) and (002) ring intensities.
Figure 5: (a) Plan-view TEM micrograph and (b) SAD pattern of a 10×(3 Å Co / 10 Å Pt) multilayer
grown at TG = 250°C, (c) SAD with 30° sample tilt showing arcing of the rings.
One of the questions to be resolved by careful characterization is whether the multilayers are alloyed or not,
i.e. are the Co/Pt “pure” components. Careful examination of the SAD patterns revealed that all of the rings
associated with the fcc structure were split, with separation ∆g increasing with the diffraction vector g. For
each doublet, the inner ring corresponding to a larger real space parameter is always more intense and the
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outer ring is weaker. This implies two distinctive lattice parameters through the thickness of the Co/Pt
multilayer stack. Figure 6 shows (a) the SAD pattern for the discussed Co/Pt multilayer and simulated ring
diffraction patterns of (b) Ptfcc and Co Pt3, (c) Ptfcc and Cohcp, (d) Ptfcc and Cofcc. On all patterns, triangles
label Pt rings: (111), (200), (220), (113) and (222). By comparing the simulated SADs to the experimental
data shown in Figure 6, it is clear that the layers must be “pure” Co-Pt both fcc. No evidence for alloying or
ordered compound formation was detected. However, the ratio of the equilibrium Cofcc to Ptfcc spacing in
the simulated pattern is not correct. As shown, any two Co and Pt rings on the recorded diffraction pattern,
for example Co(220) and Pt(220), are much closer to each other than on the simulated pattern (d). This
implies that the Co layers are strained. Figure 6 (e) and (f) shows simulated diffraction rings as a function of
increasing tensile strain of the Co layers, with image (f) showing a good match. Hence, Starting from the
seed Pt layer, the Pt(111) layers are stacked according to the equilibrium fcc stacking sequence observed in
bulk Pt. The Co layers, which, in bulk are hcp, follow the fcc stacking sequence of underlaying Pt.
Figure 6: SAD pattern for the Co/Pt multilayer and simulated plan-view diffraction patterns for (b)
Ptfcc and CoPt3, (c) Ptfcc and Cohcp, and (d) Ptfcc and Cofcc, triangles label Pt rings. (e) and (f)
Simulated plan-view diffraction patterns of Ptfcc and Cofcc at various strains: (e) aCo = 3.65 Å, σ=
+2.85%, aCo/aPt = 0.92 and (f) aCo = 3.85 Å, σ= +8.45%, aCo/aPt = 0.98. Numbers label rings: 1)
Pt(111), 2) Co(111), 3) Pt(002), 4) Co(002), 5) Pt(022), 6) Co(022), 7) Pt(113), 8) Pt(222), 9)
Co(113), 10) Co(222), (see text).
Figure 7 shows a lattice image with fcc multilayers stacking and good coherency across the Co/Pt interfaces.
This (111) stacking further implies a small volume anisotropy of Co. Thus the perpendicular easy axis of
magnetization is attributed to the interface anisotropy and in particular to strain contributions [29].
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Figure 7: Cross-sectional micrograph of a 10×(3 Å Co and 10 Å Pt) multilayer grown at TG = 250°C.
Analysis of the multilayers as a function of growth temperature
The microstructural evolution of multilayers as a function of growth temperature (TG) is presented in Figure
8, which shows plan-view BF images, SAD patterns and DF-hollow cone images obtained using (111) rings,
(111)DF, of ML’s grown at TG = 190°C, TG = 250°C and TG = 390°C. As depicted in the BF images, Figure
8(a), (e) and (i), all of the investigated samples were polycrystalline. The average grain size was found to
increase with increasing TG. As discussed above the SAD patterns, Figure 8(d), (h) and (l) showed typical
ring spacing associated with the polycrystalline fcc Pt structure. For all samples, a strong out-of-plane
<111> texture was measured, as seen from very weak (111) and (002) rings and an intense (022) ring. Only
some grains oriented randomly contributed to (111) and (002) rings. The SAD patterns from larger areas
showed a uniform intensity distribution around all rings, indicating random in-plane orientation and only
out-of-plane <111> texture.
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Figure 8: Plan-view BF TEM images, SAD patterns and (111)DF images of the Co/Pt ML’s grown at
TG = 190°C, TG = 250°C and TG = 390oC, displayed in rows. All SAD’s were collected with a 5 µm
aperture [24].
The (111)DF images, Figure 8(b), (f) and (d), show the same areas as the respective BF images and display
only the grains without the <111> texture. Figure 8(c), (g) and (k) are lower magnification images showing
the distribution of such grains. As presented, samples grown at the low temperature of 190°C, (b) and (c),
had a large number of grains lacking the <111> texture. The distribution of such grains was very uniform.
With increasing TG, the number of such misoriented grains decreased. For the TG = 250°C sample, (f) and
(g), the majority of misoriented grains were smaller than the average grain size, although a few, large
misoriented grains were also found. For samples grown at 390°C, a good <111> texture was observed, with
only very few grains lacking the <111> texture, as shown in (j) and (k). Moreover, such misoriented grains
were much smaller than the <111> textured grains. Hence, as clearly shown by the set of (111)DF images
the <111> texture was found to improve with an increase in sample growth temperature. For samples grown
at TG up to 300°C, all diffraction rings were split, indicating two distinctive lattice parameters for the Co and
Pt layers, as discussed above. For samples grown at TG = 390°C above a critical transition temperature, Tcrit,
the ring splitting attributed to the two separate Pt and Co parameters was not detectable, indicating a
continuous gradient in lattice parameter between that of strained Co and that of Pt.
Structure as a function of Co thickness
The structure of Co/Pt multilayers was also investigated as a function of Co layer thickness (tCo). In
agreement with previously published data [30], it was found that the stacking sequence in the Co layers
changes with increasing Co thickness. At 0.2 nm – 0.6 nm Co thickness, the (Co/Pt) multilayer structure
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showed fcc stacking with 2.2 Å layer spacing. When the Co thickness increased beyond ~0.6 nm, hcp
stacking faults (SF) were observed and their density increased with increasing Co layer thickness. Above
approximately tCo = 15 Å, the Co layer stacking is mainly hcp. Co/Pt multilayers with Co layer thickness
below approximately 6 Å were uniform and relatively defect free. The only indication of the Co vs. Pt layer
in the HRTEM image was a slightly higher intensity in the Co layer due to preferential thinning of the Co
layer during specimen preparation by ion-milling and a lower atomic number contrast. This intensity
variation had a periodicity of the analyzed multilayers. With increasing tCo, roughening of the MLs was
observed, with increased number of defects (Figure 9). As shown, the measured intensity is much higher
within the Co layer; however, the precise location of the Co/Pt interface is not known. The (111) stacking
sequence in some of the locations is fcc, (same vertical atom position for every third layer, ABCABC), but
some hcp stacking (same vertical atom position for every second layer, ABAB) is also seen.
Figure 9: Cross-section HRTEM image of (111) 15 Å, Co layer in the Co/Pt multilayer.
Magnetic properties
In addition to affecting the grain size and the <111> texture, TG is also known to influence the magnetic
properties, such as coercivity (HC). Magnetic hysteresis loops of the Co/Pt multilayers grown at different
growth temperatures are shown in Figure 10. As shown, the perpendicular HC of the Co/Pt ML’s was found
to increase with increasing TG [27,24]. The coercivity increases by almost a factor of two on increasing the
TG from 200°C to 300°C. However, when TG is increased beyond a critical transition temperature, Tcrit, a
decrease in perpendicular coercivity is observed, and the HC for samples grown at 390°C is reduced to 5.8
kOe. A further increase in multilayer growth temperature results in loss of perpendicular anisotropy.
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Figure 10: Mgnetic hysteresis vs. multilayer growth temperature. (a) TG = 190°C, (b) TG = 250°C, (c)
TG = 300°C and (d) TG = 390°C.
These results show that the magnetic properties of such multilayer films depend on microstructure. With an
increase in TG up to a certain critical temperature, Tcrit, an increase in grain size and an improved <111>
texture were found, both contributing to an increase in HC. For ML’s grown at TG = 390°C > Tcrit, a
decrease in HC was measured, and small magnetically decoupled domains were observed. The size of these
domains was similar to the grain size. This correlates well with Co depletion at the column grain boundaries
and with the diffuse Co/Pt interfaces, measured as one set of Co-Pt rings. Such diffuse interfaces are known
to reduce the PMA and, hence, reduce HC. Also with increasing Co layer thickness, a decrease in HC was
observed. The coercivity is therefore a function of both the microstructure (grain size and texture) and
interface quality, which is strongly influenced by TG. Thus the magnetic properties can be engineered by
appropriate choice of growth parameters and multilayer repetition.
CONCLUSIONS
In this brief review we have shown that controlled multilayer structures provide superior properties such as
mechanical and corrosion in austenitic/lath martensitic steels, and magnetic properties in Co/Pt multilayers.
ACKNOWLEDGEMENTS
Work at NCEM/LBNL was supported by the U.S. Department of Energy under Contract No. DE-AC0376SF00098. The authors acknowledge contributions from K.M. Krishnan, E.C. Nelson and G. Denbeaux of
Lawrence Berkeley National Laboratory and B.D. Terris, C.T. Rettner, A. Kellock and J.E.E. Bagglin of
IBM Almaden Research Center.
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