Crystallization of Ni-P Fabricated by Electroless

Crystallization of Ni-P Fabricated by Electroless Deposition:
Microscopic Mechanism
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2014 MRS Fall Meeting
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Zhan, Xun; Case Western Reserve University, Materials Science and
Engineering
Ernst, Frank; Case Western Reserve University, Materials Science and
Engineering
nucleation & growth, transmission electron microscopy (TEM), amorphous
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Crystallization of Ni–P Fabricated by Electroless Deposition:
Microscopic Mechanism
Xun Zhan and Frank Ernst
Department of Materials Science and Engineering
Case Western Reserve University
10900 Euclid Avenue, Cleveland, Ohio 4106-7204, USA
ABSTRACT
We investigate the crystallization of amorphous Ni–P with near-eutectic composition,
fabricated by electroless plating as a 10 µm thick continuous layer. Aiming to
understand phase transformations that occur upon heating and, in particular, the
microscopic mechanism of crystallization, we combine a variety of complimentary
characterization techniques. DSC (differential scanning calorimetry) during isothermal heating reveals the crystallization kinetics. Conventional-, high-resolution-,
and analytical TEM (transmission electron microscopy) and TEM-based electron
diffraction provide high-spatial-resolution information on phase nucleation and spatial distribution of atom species, particularly the crystallography of the nucleating
crystalline phases (Ni3P and Ni) and the spatial distribution of phosphorus in the
partially and completely crystallized alloy. Our results indicate that crystallization
proceeds by homogeneous nucleation of Ni3P grains. Internally, these exhibit a
microstructure of radially oriented subgrains containing Ni nano-platelets in a
specific crystallographic OR (orientation relationship) with Ni3P. However, the preferred Ni–Ni3P OR differs from those reported in the literature for similar material.
Combining our observation on the structure and microstructure of partially and
completely crystallized Ni–P with the observed crystallization kinetics provides a
deeper understanding of the microscopic mechanism of crystallization.
INTRODUCTION
Owing to unique properties, amorphous Ni–P alloys fabricated by ELP (electroless plating) are
widely used as buffer layers for rigid-memory disc platters. In particular, amorphous Ni–P can be
overgrown with a homogenously nucleating Cr alloy layer – as required for the subsequent
deposition of the functional (memory storing) ferromagnetic layer. Since amorphous Ni–P is
paramagnetic, it does not interfere with the ferromagnetic layer. However, the amorphous
structure is unstable and tends to crystallize into a phase mixture with ferromagnetic properties
that interfere with those of the functional layer. Therefore, understanding the micromechanism of
crystallization and the phases it produces is of scientific as well as technological interest [1-6].
For near-eutectic compositions of Ni–P, it was found that crystallization proceeds via formation
of barrel-shaped spherulites of continuous BCT (body-centered tetragonal) Ni3P containing
subgrains with a dispersion of rod-like FCC (face-centered cubic) Ni nanoparticles. This
microstructure has been observed regardless of the method by which the amorphous phase was
fabricated, e.g. electrodeposition, electroless plating, or melt spinning. Preferred ORs between Ni
and Ni3P have been reported by Watanabe et al. [2] and Lu et al. [5]. However, controversies
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Fig. 1(a) Isothermal DSC curves at different annealing temperatures. Inset: crystallization volume fraction as a
function of time. (b) Avrami plot for annealing at 598 K (325 °C).
exist owing to overlapping spots of Ni3P (BCT) and Ni (FCC) in electron diffraction pattens.
Watanabe et al. [2] proposed the preferred OR to be
<110>BCT║<110>FCC ,
<001>BCT║<1–12>FCC .
(1)
Lu et al. [5], in contrast, proposed
<111>BCT║<110>FCC ,
(2)
which is incompatible with (1).
Some authors discuss the crystallization process as classical nucleation and growth [2-4,7].
Lu et al. [6], in contrast, proposed (i) formation and growth of BCT Ni3P precursor clusters and
(ii) crystal nucleation and growth by shearing and deposition of precursor clusters at the
crystalline front. Until now, however, the reason for the observed arrangement of crystallized
phases is unclear and the question about the exact micromechanism of crystallization remains
open. Details of the short- and medium-range order in the amorphous structure may play an
important role in the crystallization process, and different synthesis methods may actually
produce different amorphous structures. Therefore, the mechanisms proposed before may not
apply to the ELP Ni–P of our study.
EXPERIMENTAL METHODS
The as-received material was a 10 µm thick layer of amorphous Ni–P with near-eutectic
composition, deposited by ELP onto an Al substrate by Atotech Deutschland GmbH. After
dissolving the Al substrate in a 10% NaOH solution at 313 K (40 °C), we punched out discs with
3 mm diameter for TEM specimen preparation. The composition was quantified by SEM-XEDS
(X-ray energy dispersive spectroscopy) with a Ni2P standard and NIST DTSA-II software.
Accordingly, the phosphorus concentration is (19.8 ± 0.1) at.%.
Isothermal crystallization kinetics was investigated using a Netzsch Pegasus 404 F1 DSC
under pure Al at a flow rate of 0.3 ml/s. Al pans with lids were used as sample holders.
Temperature and enthalpy were calibrated by measuring the melting points and the heat of fusion
of In, Sn, Bi, and Zn standard specimens. A layer of Ni0.802P0.198 was heated to four target
annealing temperatures, 583 K, 588 K, 593 K and 598 K, at a rate of 0.7 K/s. In order to study
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Fig. 2: Ni0.802P0.198. (a-c) SAED patterns. (a) As received. (b) After 0.9 ks at 598 K. (c) After 1.5 ks at 598 K.
(d) Profiles of diffracted intensity versus spatial frequency, obtained by rotation averaging of the patterns (a-c). (e)
is TEM bright-field image after 1.5 ks at 598 K (field of view represents SAED aperture) and corresponding SAED
pattern of the interface region between crystallites and amorphous matrix.
the microstructure evolution during isothermal annealing, a Ni0.802P0.198 film was annealed at
598 K for three different periods of time: 0.9 ks, 1.5 ks, and 2.4 ks.
TEM specimens were prepared from material tempered for 0.9 ks and 1.5 ks at 598 K. The
material was first thinned by Ar+ ion-beam milling using a PIPS (precision ion polishing system,
Gatan) with an ion energy of 4.5 keV and an incident angle of 4°, then polished with 2.5 keV Ar+
ions to minimize ion-beam induced structural damage at the specimen surface. The specimens
were investigated using a Tecnai F30 transmission electron microscope (300 kV, FEI).
RESULTS
Crystallization Kinetics
Figure 1a shows DSC curves obtained at different annealing temperatures. The curves exhibit a
single exothermic peak after a certain incubation time ts. With increasing annealing temperature,
ts decreases owing to increased diffusivity. Assuming a heat flow proportional to the total mass
of crystalline phases yields the volume fraction v of transformation as a function of time t,
shown in the inset. The transformation kinetics can be described by the JMA (Johnson–Mehl–
Avrami) model [8,9], derived from assuming a sequence of nucleation, growth, and impingement:
v[t] = 1 – exp[ –k (t – ts)n ] ,
(3)
where k is the reaction-rate constant and n the Avrami index, which reflects the type of
transformation. Usually, n ranges between 1 and 4 and can be obtained by plotting ln[ln[1-v[t]]-1]
versus ln[t-ts], Fig. 1b. Figure 1b reveals three distinct stages of crystallization in Ni0.802P0.198.
However, only the slope of the second stage can be safely used to analyze the crystallization
mechanism; the DSC data for first and the last stage suffer from too much uncertainty in the
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DSC signal. Within the second stage, the measured
Avrami index for isothermal annealing at 598 K is 4.
Microscopic Mechanism
To understand how crystallization proceeds,
Ni0.802P0.198 were studied by TEM in as-received
condition and after tempering at 598 K for 0.9 ks,
1.5 ks, and 2.4 ks. Figure 2 shows SAED (selectedarea electron diffraction) patterns of as-received
material and after tempering for 0.9 ks and 1.5 ks,
respectively. All patterns show the diffuse rings
characteristic of amorphous material. To obtain lownoise radial intensity profiles, we applied rotational
averaging [10]. However, the profiles do not reveal
significant differences.
Figure 3 presents TEM results for 2.4 ks at 598 K.
The conventional bright-field TEM imaging Fig. 3a
exhibits globular grains of BCT Ni3P with a typical
size of ≈ 5 µm, randomly oriented in some remaining
amorphous matrix, impinging on each other, and
containing radially arranged elongated subgrains. The
orientation of the subgrains varies by only a few
degrees. The HRTEM (high-resolution) image of a
subgrain in Fig. 3, viewed in <001>BCT || <110>FCC,
reveals a dispersion of elemental Ni (FCC) nanoplatelets. The platelets are extended in <110>BCT
directions. They have a typical length of 20 nm and
the thickness of 5 nm. The elemental map of
phosphorous (inset) obtained by ESI (electron-spectroscopic imaging) confirms this. Similar observations were reported earlier [2-5]. As seen in the
HRTEM image as well as in the diffraction pattern of
Fig. 3, the Ni lattice always makes and the following
OR with the Ni3P lattice:
<001>BCT║<110>FCC ,
–
{330}BCT║{111}FCC .
Fig. 3: Ni0.802P0.198 after tempering at 598 K for
2.4 ks. (a) Crystallites with radial subgrains.
(b) SAED pattern from crystallites and
amorphous matrix. (c) Corresponding HRTEM
image. The inset is a phosphorus map obtained
by ESI. Dark regions: elemental Ni.
(4)
Owing to spot diffuseness, Fig. 3b might also be inter–
preted as {510}BCT║{002}FCC. We discard this option as these lattice planes are less populated and
therefore of less physical significance. Defining the
lattices of Ni3P and Ni as “aligned” when the
corresponding fundamental translations of the lattices
are aligned, the OR (4) can be described as a rotation
of 46° about a <912> direction of the FCC Ni lattice.
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Fig. 4: Unit cells of Ni and Ni3P, drawn to scale. (b) Ni3P (BCT) with a {110} plane
highlighted. (a) Ni (FCC) in OR (5) with Ni3P (this work). A {111} plane of Ni is
highlighted, as well as a <211> direction in this plane. (c) Ni in OR (1) with Ni3P
(after [2]).
Fig. 5: Quantitative illustration of misfit at the two
different variants of the
Ni–Ni3P interface observed
in Fig. 3c.
In the BCT Ni3P lattice, this rotation vector corresponds to a <914> direction or a {911} plane
normal. Further, the HRTEM image reveals two preferred crystallographic orientations for the
Ni–Ni3P interface plane:
{110}BCT║{111}FCC,
(6)
{110}BCT║{211}FCC .
(7)
The first one occurs at the extended (“long”) interface sections in Fig. 3, the second one at the
small (“short”) sections.
DISCUSSION
The Avrami index of n = 4 we obtained by DSC and JMA analysis is characteristic of
homogeneous (uniform) nucleation with constant nucleation rate and 3D (three-dimensional)
growth. However, assuming that each stable nucleus eventually leads to the formation of a
barrel-shaped grain with subgrains and Ni nano-platelets, the number density of nuclei is rather
low (order of magnitude: 1016 m–3). On the other hand, the size of the grains and that they make
contact indicates that the growth rate is relatively large. (The early stages of these grains may
correspond to what has been described as “barrel-shaped” grains in the literature [2-5].)
Combining these observations leads to the conclusion the kinetics of crystallization are
controlled more by the energy barrier of nucleation than the mobility of the involved atom
species. On the other hand, from Fig. 3c the typical spacing of the Ni nanoparticles is only 5 nm.
Correspondingly, the crystallizing microstructure contains identity of Ni–Ni3P interfaces. This
indicates that the phase separation into Ni3P and Ni occurs under high driving force and limited
diffusivity of either Ni or phosphorus.
Figure 4 illustrates the OR (4) and compares it to the OR (1) by showing the relative
orientation of the unit cells of Ni and Ni3P drawn to scale. Actually, both ORs feature Ni {111}
planes parallel to Ni3P {110}. The difference between OR (4) and OR (1) can be described as a
90° <111> rotation of the Ni lattice (or equivalent under FCC point symmetry). The OR (4) is not
compatible with the result (2) (the closest directions to <111> in to which <110>FCC directions
are rotated are <335>BCT directions, about 11° away from <111>BCT.)
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It appears that minimization of misfit strain energy is the physical origin of the OR (4), the
interface orientations (6), (7) and the platelet-like morphology of the Ni particles observed in Fig.
3c. The misfit
a–b

 [a,b] := 2 a + b
(8)
between the spacings {111}FCC and {330}BCT is only –0.037 (the Ni spacing is smaller). This is
the misfit at the “short” interfaces (7). At the “long” interfaces (6), the 70.5° inclination of the
{111}FCC against the interface plane further reduces the misfit of corresponding interatomic
spacings along the interface to 0.021, as quantitatively illustrated in Fig. 5. With this small misfit,
the spacing of hypothetical b = ½<110> misfit dislocations in Ni would be as large ≈ 10 nm, about
half of the typical length of the “long” interfaces observed in the HRTEM image and the
phosphorus map of Fig. 3c.
CONCLUSIONS
Our observations of amorphous layers of Ni–P made by electroless deposition support that
crystallization begins by homogeneous nucleation. The nuclei grow radially at more or less
isotropic rates, generating microstructure of subgrain boundaries and elemental Ni nano-platelets.
The nucleation barrier is more rate-controlling than kinetic limitations of growth. Crystal growth
itself, on the other hand, is limited by small available atom transport distance for the phase
separation into Ni3P and elemental Ni. This leads to a nano-dispersion of Ni nano-platelets and a
correspondingly high density of Ni–Ni3P interfaces. Energetically, becomes possible by a special
OR (orientation relationship) between the Ni and the Ni3P lattice and an alignment of the
interface plane parallel to {110} planes of Ni3P and close-packed planes of Ni, enables highly
coherent interfaces will effectively minimize the interface energy.
ACKNOWLEDGEMENT
We thank Atotech Deutschland GmbH for financial support and for providing samples. We also
thank Z. Vashaei and Z. Li for help with the experimental work.
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