DENSIFICATION AND GRAIN GROWTH
MECHANISMS DURING SPARK PLASMA
SINTERING OF IONIC SOLIDS
A thesis submitted in partial fulfillment
for the award of the degrees
Bachelor and Master of Technology
by
Karthik Akkiraju
MM10B021
Department of Metallurgical and Materials Engineering
Indian Institute of Technology Madras
May 2015
For the love of cup noodles, Coke and coffee.
THESIS CERTIFICATE
This is to certify that the thesis titled ’Densification and Grain Growth Mech-
anisms during Spark Plasma Sintering of Ionic Solids’ submitted by Karthik
Akkiraju (MM10B021), to the Indian Institute of Technology Madras, Chennai for
the award of the degree of B.Tech and M.Tech, is a bona fide record of the research work
done by him under my supervision. The contents of this thesis, in full or in parts, have
not been submitted to any other Institute or University for the award of any degree or
diploma.
Prof. B.S. Murty,
Research Guide,
Dept. of Metallurgical and Materials Engineering
Place: Chennai
IIT-Madras, 600 036
Date: 5th May 2015
Acknowledgments
I shall forever be indebted to Prof. B.S Murty for introducing me to research, having faith
in me for all these years and inspiring me to push the boundaries all along. Dr. Lukas
Bichler, from the University of British Columbia has played an equal part in stitching this
journey together and making Kelowna a place I yearn to be at. This work would have
been incomplete had it not been the engaging discussions over chai with Niraj Chawake,
Dr. Ajit Srivastava, Koundinya NTBN and Dr. Sanjay Kashyap. Thank you very much
for being around on the trips to Ramu and CC. Also, I cannot thank Vishank Kumar,
Audrey Siebert Timmer, Karen Robles and Dr. Apara Ranjan enough for sharing research
woes with Tim Hortons and frequent dinner invites.
I would like to acknowledge the contributions of David Arkinstall (UBCO) and Amit
Sharma (IISc) for helping out with the SEM work. My special appreciation to Anirudha
Karati for helping with the TEM characterization. For their constant advice, Dr. Anand
Kanjarla, Dr. Chinmoy Chattopadhyay, Karthiselva, Pravin from the MME department
at IIT Madras. It was never going to be an individual effort and my research groups at
IITM (Advanced Materials Research) and UBCO have always made me feel at home and
I thank them for bearing me with till I graduate. A huge cheer to my friends Shukla,
Eureka, Mama, Chuddu, Bobo, Belly, Blk, Hotseat, Patil, Chetta Always Wins, and to
you for staying till the end.
Karthik
I
Abstract
In the present work, the densification and grain growth mechanisms during SPS of three
model oxide ceramics: NiO, ZnO and MgO were studied investigating to gain fundamental
understanding of the operative atomisitics. The role of applied pressure, dwell time,
heating rate and electric field was investigated in various temperature regimes.
In the case of NiO nanopowder, the morphological changes and the grain growth mecha-
nism during spark plasma sintering and annealing at different temperatures were studied.
Cuboid shaped grain morphology was observed in the case of pressureless sintering at tem-
peratures above 1000C. The grain growth mechanism involved Brownian motion assisted
grain coalescence via oriented attachment of irregularly shaped nanocrystals, leading to
the formation of nanocubes. The nanocubes subsequently combined to form dense 3-
dimensional microstructure. Similar grain growth mechanism was observed when the
powders were annealed in ambient atmosphere at and above 700 C. Twinning and dis-
location arrays were observed at the resulting interfaces. However, sintering below 1000
C resulted in conventional polyhedral morphology, which was driven by diffusion based
processes. The absence of electric field shifted the density and grain size trajectories to
higher temperatures.
Similarly, SPS of ultra-fine ZnO was carried out in the 600-1000C temperature range
with dwell times of 0-15 min. Between 600-700C, a sintering window was obtained where
grain size and relative density can be controlled. Observation of bridging necks indicated
that localized melting could take place during SPS of ZnO. The presence of liquid film
was related to the role of electric fields via nucleation of Frenkel defects. Formation of
nano grain clusters further indicated the possibility of densification via liquid phase aided
grain rotation. Grain growth and densification mechanisms were evaluated based on the
stress and grain growth exponents. Pore growth and negative shrinkage behavior was
observed at 1000C.
II
MgO nanopowder was sintered in the temperatures range of 900-1400C. Applied pres-
sure was found to lower the sintering temperatures to achieve a given density. A higher
heating rate had a detrimental effect on the densification. Two distinct windows were
observed in the grain size and relative density plots when plotted against temperature.
The sintering behavior was found to be identical in the presence and absence of the elec-
tric field with identical shrinkage curves, grain sizes, relative density and microstructural
features. The grain growth was characterized by the presence of pyramidal-spiral struc-
tures and elongated structures growing from the grain surfaces. It is expected that a
evaporation-condensation mechanism is operative for their growth.
III
Contents
Acknowledgments I
Abstract II
List of Figures VI
List of Tables X
1 Introduction 1
1.1 Sintering . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1
1.2 Hot Isostatic Pressing . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2
1.3 Field Assisted Sintering Techniques . . . . . . . . . . . . . . . . . . . . . 2
1.3.1 Spark Plasma Sintering . . . . . . . . . . . . . . . . . . . . . . . . 2
1.3.2 Flash Sintering . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3
2 Literature Survey 5
2.1 Classical Sintering Theory . . . . . . . . . . . . . . . . . . . . . . . . . . 5
Mass Transport Mechanisms . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5
2.1.1 Surface Diffusion . . . . . . . . . . . . . . . . . . . . . . . . . . . 5
2.1.2 Volume Diffusion Lattice Diffusion . . . . . . . . . . . . . . . . . 6
2.1.3 Grain Boundary Diffusion . . . . . . . . . . . . . . . . . . . . . . 6
2.2 Grain Growth . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6
2.3 Recent Advances in Sintering Theory . . . . . . . . . . . . . . . . . . . . 7
2.3.1 Effect of Pressure . . . . . . . . . . . . . . . . . . . . . . . . . . . 7
2.4 Material Selection . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9
2.5 Objectives of the Work . . . . . . . . . . . . . . . . . . . . . . . . . . . . 10
2.6 Scope of the Thesis . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 11
IV
3 Experimental Details 12
3.1 Raw Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 12
3.2 Spark Plasma Sintering and Annealing . . . . . . . . . . . . . . . . . . . 12
3.3 Characterization Techniques . . . . . . . . . . . . . . . . . . . . . . . . . 14
3.3.1 X-ray Diffraction . . . . . . . . . . . . . . . . . . . . . . . . . . . 14
3.3.2 Scanning Electron Microscopy . . . . . . . . . . . . . . . . . . . . 14
3.3.3 Transmission Electron Microscopy . . . . . . . . . . . . . . . . . . 15
3.3.4 Hardness Measurement . . . . . . . . . . . . . . . . . . . . . . . . 15
4 Results and Discussion: Spark Plasma Sintering of nano NiO 16
4.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 16
4.2 Results . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 17
4.2.1 Reduction of NiO . . . . . . . . . . . . . . . . . . . . . . . . . . . 18
4.2.2 Coalescence Driven Grain Growth of Nanocrystals . . . . . . . . . 19
4.2.3 Effect of Boron Nitride layer . . . . . . . . . . . . . . . . . . . . . 23
4.3 Discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 25
4.3.1 Reduction Behavior . . . . . . . . . . . . . . . . . . . . . . . . . . 25
4.3.2 Grain Coalescence . . . . . . . . . . . . . . . . . . . . . . . . . . 25
4.3.3 Effect of Electric Field . . . . . . . . . . . . . . . . . . . . . . . . 30
4.4 Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 30
5 Results and Discussion: Spark Plasma Sintering of ZnO 31
5.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 31
5.2 Results and Discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . . 32
5.2.1 Densification and Grain growth . . . . . . . . . . . . . . . . . . . 33
5.2.2 Observation of sintering necks . . . . . . . . . . . . . . . . . . . . 35
5.2.3 Grain Coalescence . . . . . . . . . . . . . . . . . . . . . . . . . . 38
5.2.4 Effect of Dwell Time: Sintering Analysis . . . . . . . . . . . . . . 39
5.3 Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 42
6 Results and Discussion: Spark Plasma Sintering of MgO 43
6.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 43
6.2 Results and Discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . . 43
6.2.1 Decomposition of Mg(OH)2 . . . . . . . . . . . . . . . . . . . . . 43
V
6.2.2 Grain growth and Densification . . . . . . . . . . . . . . . . . . . 45
6.2.3 Microstructural Characteristics . . . . . . . . . . . . . . . . . . . 47
6.2.4 Terraced Oxide Growth . . . . . . . . . . . . . . . . . . . . . . . 50
6.3 Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 52
7 Conclusions 53
7.1 Recommendations for Future Work . . . . . . . . . . . . . . . . . . . . . 54
References 55
VI
List of Figures
1.1 Schematic of the SPS setup.[3] . . . . . . . . . . . . . . . . . . . . . . . . 3
1.2 (a)Schematic of the flash sintering setup while (b) shows the linear shrink-
age during the flash sintering process.[4] . . . . . . . . . . . . . . . . . . . 4
2.1 Schematic of the possible sintering pathways across the particle interface.[6] 5
3.1 Showing the current path in the case of BN cased NiO (left) and conven-
tional SPS (right). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 13
4.1 Dark field TEM image of agglomerate with an average crystallite size of
10 nm while (b) shows the indexed diffraction pattern. . . . . . . . . . . 17
4.2 XRD patterns of the top surface of the sintered pellet showing increased
phase fraction of the reduced Ni layer. . . . . . . . . . . . . . . . . . . . 18
4.3 (a) Line scan across the top surface showing a Nickel layer which gets
oxidized after the thermal etching in (b). The marked line shows the
interface. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 18
4.4 (a) Microstructure of the Ni/NiO interface, (b) the growth of the Ni layer
into the NiO particle, (c) reoxidized surface of the Ni/NiO interface, (d)
growth of NiO scales after thermal etching. . . . . . . . . . . . . . . . . . 19
4.5 Fractured surfaces of the samples sintered for 10 min at 50 MPa pressure
(a) at 900C with polyhedral grains and (b) with the faceted cuboidal
morphology developing at 1000C. . . . . . . . . . . . . . . . . . . . . . . 20
4.6 Fractured surfaces of the samples sintered for 10 min at 5 MPa pressure
with (a) cuboidal grains at 1000C, (b) a few larger faceted grains are seen
when sintered at 1050C and (c) cuboidal voids at 1200C with the inset
showing the cubic pores at a higher magnification. . . . . . . . . . . . . . 20
VII
4.7 Plot showing the change in grain size and relative density (inset) as a
function of sintering temperature for the samples sintered with a pressure
of 5 MPa and 50 MPa. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 21
4.8 Sequence of the images showing the coalescence mechanism (a) number of
smaller cubes attached to a larger surface, (b) larger cuboidal block and a
cube attachment, and (c) the chain of particles formed as a result of the
coalescence. All images are taken from the powders annealed at 900C for
1hr. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 22
4.9 TEM images of the powders annealed at 700C for 1h with the initial
irregular morphology still being retained, while (b) shows the HRSEM
image evolving cube with some of the retained 111 planes. . . . . . . . 22
4.10 HRTEM of the interface of an impinged particle, (b) showing the indexed
FFT pattern with‘t’ as the twinned spots, and (c) is reconstructed image
of the selected portion showing the twinned region. (d) HRTEM image of
the interface between two particles, (e) reconstructed image using the FFT
showing an array of dislocations and partial dislocations at the interface
with the inset showing a magnified viewed of the defect region. The sample
was annealed at 700 C for 5 min. . . . . . . . . . . . . . . . . . . . . . . 23
4.11 Comparison of the (a) grain size (b) relative density (c) hardness for the
various configuraions studied. . . . . . . . . . . . . . . . . . . . . . . . . 24
4.12 Microstructure of the sintered pellet with BN blocking at (a) 1050C and
(b) 1250C. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 24
4.13 (a) SEM image of the sample annealed at 1200C showing the initiation
of the attachment process and (b) TEM image of an annealed powder at
700C for 1h showing the cubic building blocks within a large cluster and
a smaller faceted particle is shown to attach to this cluster. . . . . . . . . 28
4.14 Graphical representation of the coalescence mechanism . . . . . . . . . . 30
5.1 SEM micrograph of the raw powder with the inset showing a magnified
image of the nanoparticles. . . . . . . . . . . . . . . . . . . . . . . . . . . 32
5.2 XRD micrographs of the received ZnO powder and pellets sintered at var-
ious temperatures. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 33
VIII
5.3 (a) Evolution of sintering trajectory and (b) grain size-relative density
trend, as a function of sintering temperature for a dwell time of 5 min. . 33
5.4 Fractured surfaces of ZnO sintered at (a) 600C (b) 800C (c) 1000C. . . 34
5.5 Presence of bridging necks between the particle 600C in (a), (b) and at
1000C in (c). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 35
5.6 Line scan across the grain interfaces showing a dip in Oxygen concentration
in (a) while the composition is uniform in (b) for the sample sintered at
1000C for 0 min. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 36
5.7 Line scan across the grain interfaces with the necks showing uniform Oxy-
gen concentration in (a) and (b) for the sample sintered at 1000C for 0
min. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 36
5.8 Growth of nano grain clusters from (a) 600C to (b) 800C and (c) 1000C. 38
5.9 (a) Change in relative density and (b) grain size as a function of dwell time
at 600C, 700C, 800C and 1000C. . . . . . . . . . . . . . . . . . . . . 39
5.10 (a) Calculation of stress exponent and (b) grain growth exponents. . . . . 40
5.11 (a) SEM images of fractured surfaces of samples sintered at (a) 700C with
15 min dwell and (b) 800C for a dwell time of 10 min. . . . . . . . . . . 41
5.12 Pore enlargement with increase in dwell time at 1000C. . . . . . . . . . 41
6.1 XRD patterns of the heat treated raw powders at various temperatures . 44
6.2 Bright field image of the heat treated powder with the inset showing the
ring pattern. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 44
6.3 Shrinkage curves of the samples sintered at 1200C with and without the
application of pressure. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 45
6.4 Comparative (a) grain size and (b) relative density plot of sintered MgO
with the clouds showing similar trends. . . . . . . . . . . . . . . . . . . . 46
6.5 Fracture surface of the sample sintered at 1200C with a heating rate of
200Cmin−1 and a pressure of 50 MPa. . . . . . . . . . . . . . . . . . . . 47
6.6 Microstrucutural feature of the sintered pellet at 1300C with BN (5 MPa)
showing the neck formation while (b) was sintered at 1300C (50MPa)
showing a stepped surface. . . . . . . . . . . . . . . . . . . . . . . . . . . 48
IX
6.7 Comparison of the microstructure of the sintered samples at 5 MPa with
(a),(b) sintered at 1100C and (c),(d) at 1300C with and without BN
respectively. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 48
6.8 Comparison of the microstructure of the sintered samples at 50 MPa with
(a),(b) sintered at 1100C and (c),(d) at 1300C with and without BN
respectively. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 49
6.9 Fractured surface of the sample sintered at 1100C with 5MPa showing
terraced grain structure and (b) showing a pyramid-like terraced structure. 50
6.10 (a) Micron sized rod-like structures while (b)shows the magnified image of
a smaller rod growing from the tip of the terrace structure. The sample
was sintered at 1100C with 5 MPa pressure. . . . . . . . . . . . . . . . . 51
6.11 (a) Triangular islands growing on the surface on the existing micron sized
grains,(b) magnified image of such a structures.The sample was sintered
at 1300C with 50 MPa pressure with BN. . . . . . . . . . . . . . . . . . 52
X
List of Tables
2.1 Comparative properties of the materials used in this study . . . . . . . . 10
3.1 SPS Parameters used in each of the case studies . . . . . . . . . . . . . . 14
XI
Chapter 1
Introduction
Sintering is one of the processing techniques used for ceramic production. In this chapter,
various sintering techniques shall be discussed leading to the debate on spark plasma
sintering and it’s mechanisms.
1.1 Sintering
Sintering is a well established process for producing bulk polycrystalline ceramic mate-
rials. Traditionally sintering has been the thermally driven process to convert the green
bodies into dense ceramic parts (below the melting point). As a result of atomic diffusion,
densification and grain growth occur.
Sintering is classified as solid state and liquid state sintering. Solid state sintering
is driven by the difference in the curvature of the particle surfaces,where the driving
force is given by the Gibbs Thomson equation. Recently, a chemical reaction or external
parameters such as external electric field, magnetic field or pressure are being used to
enhance the sintering kinetics. In addition to this, in liquid phase sintering, the presence
of a small fraction of a liquid phase at the grain boundaries enhances the diffusivities
by 6-8 orders of magnitudes [1]. This help in lowering the sintering temperatures which
can reduce grain growth apart from increasing the economic viability. Apart from the
conventional sintering which uses binders and long sintering times, several techniques
such as Hot Isostaic Pressing (HIP), Field Assisted Sintering Techniques (FAST) such as
Flash sintering and Spark Plasma Sintering, Microwave sintering have been introduced
to reduce the sintering temperatures and have a better control on grain growth and
1
densification. These techniques shall be briefly discussed.
1.2 Hot Isostatic Pressing
In hot isostatic pressing (HIP), a high pressure (100-200 MPa) is applied by means of an
external gas in a sintering furnace upto very high temperatures (2000C). The powder
is encapsulated in a metal container that is placed in vacuum and pressurized. Also,
samples sintered via conventional sintering can be used placed in this furnace for complete
densification. With HIPing near net shape components can be developed, thus eliminating
the need for welding and machining in many cases.
1.3 Field Assisted Sintering Techniques
Apart from applied external pressure, electromagnetic fields can enhance densification
depending on the compatibility of the material with the applied field.
1.3.1 Spark Plasma Sintering
In Spark Plasma Sintering (SPS), sintering is achieved by means of a pulsed electric
field that is applied across the die and the sample along with a pressure. SPS differs
significantly from conventional processes due to the rapid heating rates (>800C min−1)
possible. This is ascribed to the internal Joule heating of the dies and the powders due
to the passage of large currents (thousands of Amperes) rather than the external heating
in conventional techniques [2].
The schematic of the process is shown in Fig. 1.1 [3]. Pressure is applied by means
of stainless steel ramps through which also the current is passed. The temperature
is measured either by a thermocouple (<1000C) that is inserted into the surface of
the die or a pyrometer ( >1000C) focused on the surface of the die. Sintering can
be done either in a vacuum medium or using an inter gas such as argon or helium.
Another feature of the SPS unit is the rapid cooling involved because of the efficient
water circulation. A controlled cooling rate also enables additional heat treatment steps
after the sintering process. Using a computer interface, the temperature and pressure
profiles can be programmed. Linear shrinkage, voltage, current, applied pressure and
2
chamber pressure are constantly monitored during the process. However the scaling of
SPS for industrial applications is an active area of research, with aims to incorporate
larger samples and reduce the power consumption.
Figure 1.1: Schematic of the SPS setup.[3]
Many of the studies in SPS report achievement of near theoretical density at lower
sintering temperatures, shorter times while minimizing grain growth for a wide range of
materials. The underlying mechanisms for such a rapid kinetics shall be discussed in the
next chapter.
1.3.2 Flash Sintering
This recent technique by Cologna and Raj at the University of Colarado involves the
use of extremely large electric fields to sinter ceramics. The schematic of the setup by
Francis et al. [4] is shown in Fig. 1.2(a). A dog-bone shaped green sample is prepared
and hooked using platinum electrodes [4]. The setup is placed in a preheated furnace.
It has been observed that at a particular temperature, remarkable sinterability is seen
in a few seconds. This anomalous densification is invariably accompanied by increase in
the counductivity of the specimen while limiting grain growth. The shrinkage curves for
3
this process are seen in Fig. 1.2(b) where with increasing field strength the flash event is
brought down to lower temperatures. Flash sintering has shown tremendous success with
sintering of nanograined YSZ, Y2O3, Mg doped Alumina and more recently with SiC [5].
The applicability of this process to other non-ionic systems will be an interesting follow
up. Despite this, the low cost of operation has led to the establishment of a company,
Ceram that commercially produces such furnaces.
Figure 1.2: (a)Schematic of the flash sintering setup while (b) shows the linear shrinkage
during the flash sintering process.[4]
4
Chapter 2
Literature Survey
2.1 Classical Sintering Theory
Mass Transport Mechanisms
Sintering occurs due to atomic diffusion at elevated temperatures. To achieve this various
pathways such as surface diffusion, grain boundary diffusion, volume diffusion, lattice
diffusion exist [6]. A schematic for the same is shown in Fig. 2.1.
Figure 2.1: Schematic of the possible sintering pathways across the particle interface.[6]
2.1.1 Surface Diffusion
The surface is source of defects such as extra atoms, vacancies, absorbed species. Due
to this atomic motion occurs between these defects. Atoms diffuse from a convex to a
5
concave curvature to minimize the surface energy. Since this is a re-distribution process,
it is non-densifying.
2.1.2 Volume Diffusion Lattice Diffusion
It is the replacement of atoms with vacancies. Pores, which are large vacancy clusters
are filled with atoms from the surrounding. Such a diffusion from the neck region to
the grain boundary can lead to densification. Pore growth can also occur due to volume
diffusion due to transfer of vacancies across pores.
2.1.3 Grain Boundary Diffusion
This is the most common densifying mechanism for many metals. The activation energy
is lower than that for lattice diffusion but higher than surface diffusion. Due to surface
diffusion, grain coarsening occurs, the driving force for surface diffusion decreases, leading
to the take over by grain boundary diffusion. Here, there is mass transfer from the grain
boundary region to the neck.
Apart from these common mechanisms, plastic deformation, evaporation condensation
have also been reported.
Classical sintering theory divides the densification into three stages : initial, interme-
diate and final. In the initial stage, necks develop at the particle contact points, and the
densfication proceeds via surface diffusion and particle rearrangement. The next stage is
characterized by the change of porosity from the open channels to closed ones. The final
stage, commonly accepted as >92% density involves the elimination of pores at grain
boundaries and grain growth. Mostly grain boundary diffusion and volume diffusion are
active in this regime.
2.2 Grain Growth
Rapid grain growth in nanoparticles is an important consideration for the production of
bulk nanograined materials. The motivation of many sintering techniques is to reduce
sintering temperatures and duration to avoid grain growth. Strategies such as novel
sintering techniques, control of the grain growth mechanism by doping or step sintering
processes are being explored to tackle this problem.
6
The study of grain growth mechanisms can be again divided depending on the size
range. Once the nanosized particles grow into micron sizes, the classical generalized
parabolic equations become applicable which is based on the Ostwald ripening phe-
nomenon in two phase materials. Among various models, the calculation of activation
energies in different regimes and evaluation of the exponent n can give some insights into
the grain growth mechanism. However, given the number of factors that go into such
equations, it often becomes difficult to zero in on a single mechanism.
The driving force for any particle sintering is inversely proportional to the particle
size as given by the equation,
σ = γκ = γ(1
R1
+1
R2
),
where γ is the surface energy of the material, κ is the curvature of the surface, and R1
and R2 are the principle radii of curvature [6]. Hence for nanosized particle, the driving
forces are higher [7]. The initial grain growth mechanisms in the nano range is still
not well understood with surface diffusion, grain boundary migration, coalescence and
reprecipitation.
2.3 Recent Advances in Sintering Theory
With the introduction of new sintering techniques (discussed in the previous chapter)
it has become necessary to develop new theories to explain the anomalous densification
behavior during SPS and flash sintering of nanosized particles. SPS, in particular, has
attracted attention from the sintering community for its rapid heating rate and the con-
tentious role of plasma that is believed to enhance the densification. Flash sintering can
be identified as an extreme case of SPS where very large fields are supplied and the re-
sults from such studies can aid in the understanding of SPS. Similarly, the pressure effects
from HIP or hot pressing studies can be compared to the SPS. Of the several processing
parameters during SPS, the effect of pressure, the rapid heating rate and the electric field
shall be briefly discussed.
2.3.1 Effect of Pressure
With increase in pressure, densification increases invariably. The maximum applied pres-
sure is limited by the strength of the dies. The driving force at a particle interface is
7
given by the equation [8],
µ1 = µ0i ∗ −σnΩI + ezφ,
Where µ0i is the standard reference potential, σn is the normal stress, e is the electron
charge, z is the valence of the diffusing species, φ is the local electric potential and ΩI is
the atomic volume of the diffusing species.
Densification will occur via plastic deformation if the yield strength of the material is
less than the applied pressure. MgO has been densified by SPS via this mode and maps
have been developed for various regimes [9]. Studies have been carried out to use the
models developed for creep and hot pressing for SPS data. Hence, by eliminating the
role of pressure by conducting pressureless sintering, more fundamental insights in the
process can be obtained.
Heating Rate
Rapid heating rate is one of the crucial parameter in SPS. The importance of heating
rate is attributed to the bypassing of the lower temperature surface diffusion regime [10].
The effect of heating rate on densification is also unclear.
Effect of Electric Field
The argument about the presence and absence of plasma has been debated extensively.
While have [11]shown that the plasma was not found even after rigorous experiments.Interestingly,
recently it has been shown by direct visual examination and microstructural character-
ization that a discharge may occur [12]. Chawake et al. [3] have recently shown that in
the case of metal powders such as Fe, Ni and Cu, it is the Joule heating arising from the
contacts of graphite die and punches that is responsible for the densification rather than
the internal current. While other investigations have shown that for conducting powders,
the electric current can enhance defect mobility.
Chen et al. [13] have investigated the role of pulsing during SPS by observing the
growth on the Mo-Si intermetallic layer with varying DC pulsing time. In this indirect
study no effect on the growth rate of the reaction layer was found on changing the pulse
patterns from 2:8 to 8:2 ON/OFF ratio. Although there are several works on the role of
pulsed current during SPS of metals , the fundamental role of the pulsed current during
8
sintering of ceramics in not well established. Shen et al. reported that both densifica-
tion and grain growth were affected by the pulse sequences [14]. At lower temperature
densification was retarded as the ON/OFF ratio increased, at higher temperature grain
size decreased. However, in a recent study no effect on the pulse pattern during sintering
of Al2O3 was found [15]. Pulse sequence ranging from 2:2 to 48:2 had no effect on the
densification of spark plasma sintered yttria stabilized zirconia [16].
The case for non-conducting materials or ionic ceramics is inconclusive with several
conflicting results on the role of grain boundary mobility. While very large electric field
have shown to enhance densification, the role of weak fields in SPS is currently being
debated. Ghosh has shown that weak DC electric fields can retard grain growth which
is contrary to the normal expectation [17].While YSZ has been used as model ceramic
for most such experiments, the role of weak fields in other ceramics is not understood.
Narayan [18] [19] strengthened the cause for electric field effects by proposing a surface
melting phenomenon and a new grain growth model based on his observations on MgO.
Although the electric field and current cannot be explicitly controlled in SPS, their
role in the sintering processes needs to be revisited.
2.4 Material Selection
By understanding the role of the applied pressure and the electric field during spark
plasma sintering of oxides in enhancing the densification during SPS, a better under-
standing of the atomistics during SPS can be gained.
During their in-situ experiments of field assisted sintering of Ni nanoparticles which
had a thin NiO layer, Bonifacio et al. observed the time dependent breakdown of the NiO
layer to given an Ni layer. Metallic Ni necks were formed between the particles, which
was followed by surges in the current value [20]. This experiment using an STM tip in
a TEM machine gave a fundamental insights of the field effects during ECAS of NiO.
Interestingly, Chaim and Bar-Hama [21] densified nano NiO and associated the rapid
densification to the creep and grain boundary sliding. Hence, these two studies provided
information about the possible mechanisms for NiO, both of which are possible during
SPS.
9
Similarly, Narayan developed his ’Surface Melting’ [19] [18] theory based on his obser-
vation during field treatment of MgO single crystals. In such two electrode experiments,
the MgO single crystals were found to fail by a Joule heating along dislocations and va-
cancy clusters [22]. It has been proposed that this heating can cause rapid densification
in FAST. In addition to this, sintering maps developed for SPS of MgO attribute the
rapid densification to the plastic deformation [9]. Here again, the proven role of electric
field in MgO, provides an opportunity for investigation into the operative mechanisms
during SPS of MgO.
Langer et al. have concluded that there is no difference in the densification mecha-
nisms and grain growth behavior between ZnO processed by HP and SPS, suggesting that
the current plays no significant role during the sintering process [23]. However, Misawa
et al. clearly observed particle boundaries in the case of DC pulse-free SPS processing,
while the boundaries were absent when DC-pulses were active [24]. Further, during SPS
of nanocrystalline ZnO at 500 C, an amorphous intergranular ZnO layer was observed,
presence of which resulted in abnormal electrical conductivity of the as-sintered mate-
rial [25]. These contradictory results indicate that there is still not a clear understanding
on the role of electric field during sintering of ZnO.
A comparison of the properties of three oxides is shown in Table. 2.1.
Table 2.1: Comparative properties of the materials used in this study
Material Melt. Temp.(C) Dielectric Constant Crystal Structure Band Gap(eV)
NiO 1955 9-12 Rocksalt 3.4
ZnO 1975 10 Wurztite 3.3
MgO 2852 10 Rocksalt 7.8
2.5 Objectives of the Work
The objectives of the present work are:
• To densify three model oxide systems: NiO, ZnO and MgO nanopowders using
Spark plasma sintering.
• Study of the role of process parameters by systematic deconstruction of each factor
by means of pressureless sintering and electric field blocking experiments using
10
Boron Nitride (BN).
• To evaluate the grain growth and densification mechanisms using microscopic tech-
niques and analytic models in each of such cases shall be probed.
• To identify the processing parameters among the pressure, heating rate, electric
field and dwell time that are significant for the densification of ceramic oxides.
2.6 Scope of the Thesis
Limited experimentation of the field effects during sintering of NiO, MgO and ZnO,
and the availability of classical densification mechanisms during SPS, provides a case
for the investigation of these oxide materials with emphasis on the role of pressure and
electric fields. This approach has been attempted by conducting conventional pressure
assisted sintering along with pressure less experiments and field blocking experiment
using hexagonal BN. The Results and Discussion of the thesis are organized separately
for the oxide systems with an underlying theme to gain a common understanding of the
process. With similar experimental design across the three materials, the conclusions
chapter compares the results and presents a common understanding.
11
Chapter 3
Experimental Details
3.1 Raw Materials
Raw powders of NiO (Inframat Advanced Materials, USA) , ZnO (Fischer Scientific,
USA), MgO (Inframat Advanced Materials, USA) were purchased and used directly for
the experiments except in the case of MgO where prior heat treatment was necessary to
eliminate the formed Magnesium Hydroxide.
3.2 Spark Plasma Sintering and Annealing
Different SPS parameters had to be employed for different materials, owing to the differ-
ences in the melting point. ZnO and NiO were sintered in a 20 mm diameter graphite
die lined with grafoil using the Model 10-3 (10 ton/3000 A) Thermal Technology SPS
unit while MgO was sintered using the Dr. Sinter SPS-5000 machine. For all sintering
experiments, a 5 MPa pre-load was applied to achieve a green pellet density of 45%.
Pressureless sintering was carried out with the 5 MPa load maintained throughout the
entire SPS run.
For the field blocking experiments, the ceramic samples were sandwiched between lay-
ers of hexagonal Boron Nitride (BN) as shown in Fig. 3.1. For pressure assisted sintering
experiments, 50 MPa load was used for all the three oxides. The optimized temperature
and dwell times investigated are reported in the table below. For temperatures inves-
tigated below 1000C, thermocouple inserted into the bottom die was used to monitor
the temperature. For temperatures above 1000C, a pyrometer focused on the surface of
12
the die was used. In this case, temperatures below 600C cannot be measured and the
heating rate in this range is uncontrollable. Real time sample shrinkage information was
obtained from the sample dimensions and the SPS ram displacement data collected by
the data acquisition unit.
Figure 3.1: Showing the current path in the case of BN cased NiO (left) and conventional
SPS (right).
The density of the sintered pellet was measured using the Archimedes principle.
ρ =ma
ma −mw
∗ ρH2O
Where ma and mw are the weight of the sample in air and in de-ionized water respectively.
The instantaneous relative density of the pellet was estimated using the relation,
ρi = ρf ∗hfhi
Where ρi is the instantaneous density, ρf is the final density and hf , hi are the corre-
sponding heights of the specimens.
The densification strain rate behavior was studied via Eq. 1,
1
ρ
dρ
dt= HD
φPa
RTGn
The equation governing grain growth can be written as Eq. 3, simplified from the
pore drag model by Kang [1],
Gn −Gn0 = kt
13
Where G0 is the average grain size at time t=0, k is the rate constant and n is the grain
size exponent. If G0 is small when compared to G, the slope of the lnG vs lnt plot yields
the grain size exponent n.
Table 3.1: SPS Parameters used in each of the case studies
Material Temperature( C) Dwell (min) Heating Rate (C/min)
NiO 900-1300 10 200
ZnO 600-1000 0-15 200
MgO 900-1400 3 200,100
In order to compare the grain growth occurring during sintering process and annealing
process, annealing experiments were conducted by placing the raw powders in a muffle
heat treatment furnace by directly loading them at the set temperature.
3.3 Characterization Techniques
3.3.1 X-ray Diffraction
X-ray diffraction (XRD) experiments were conducted using the X’Pert Pro (PANalytical)
X-ray Diffractometer in Bragg-Brentano geometry equipped with X’Celerator detector
with Cu-Kα radiation (45 kV, 30 mA) in the 30-90 2θ range.
3.3.2 Scanning Electron Microscopy
Powders (i.e. as-received and annealed nanopowders) along with fractured and polished
surfaces of the as-sintered ceramic samples were observed using a Tescan MIRA3 XMU
equipped with and FEI QUANTA 400 Field Emission Gun equipped scanning electron
microscopes (SEM). For grain size measurements, the sintered samples were cut, polished
to 1 µm and then thermally etched at 100C below the sintering temperature to reveal
the grain boundaries and the largest axis of the grain polygon was measured.
14
3.3.3 Transmission Electron Microscopy
Transmission electron microscopy (TEM) was carried out using a Tecnai T-20 unit op-
erating at 200 kV, while high resolution TEM (HR-TEM) imaging was carried out with
Tecnai F-30 field Emission TEM operating at 300 kV to observe grain interfaces in as-
sintered samples. Fast Fourier Transform (FFT) and reconstructed lattice images using
the Digital MicrographTM software (Gatan Microscopy Suite) were used to analyze the
HRTEM images. For the powders, the samples were deposited on a copper grid after dis-
persion in acetone. For the sintered samples, ion milling was carried out after dimpling.
Additional cross-sectional method was usd to prepared TEM samples as the ceramic
samples were too brittle for conventional techniques.
3.3.4 Hardness Measurement
Hardness was evaluated using the Vickers micro hardness indenter with a load of 1N for a
dwell time of 15 sec. An average of 20 measurements was used to estimate the hardness.
15
Chapter 4
Results and Discussion: Spark
Plasma Sintering of nano NiO
4.1 Introduction
Processing of nanocrystalline materials is fundamentally different from sintering of con-
ventional materials, since the larger surface area of nanocrystalline powders enhances
the sintering kinetics [7]. Ostwald ripening (OR) and oriented attachment (OA) are
the mechanisms that are usually suggested to explain the sintering and grain growth in
nanocrystalline materials [26]. OR is a classical grain growth mechanism where larger
grains grow at the expense of smaller ones, and the driving force for grain growth comes
from the excess surface energy, with the growth rate given by the Lifshitz-Slyozov-Wagner
(LSW) equation [27]. In contrast, OA grain growth model proposed by Penn et al. [28]
for nanocrystalline systems suggests that particles rotate until an epitaxial or twin con-
figuration is reached. Though both grain growth mechanisms involve reduction in surface
energy, the ability of grains to rotate is inversely proportional to their grain size [29]. As
a result, the OA mechanism is expected to be prominent in nanocrystalline materials.
The ability of the SPS process to rapidly sinter nanopowders with minimal grain
coarsening can create new pathways for control of grain growth kinetics and grain mor-
phologies [30]. Owing to the high surface energy of the nanograins, the anisotropy of the
surface energy has to be considered during sintering. Few studies in the literature have
reported the impact of surface energies on the grain growth behavior of nanomaterials
during sintering. A number of published investigations attempting to describe the grain
16
growth mechanisms in nanograined materials have been carried out with colloidal systems
during hydrothermal growth. In such systems, along with the temperature, surfactants
which adsorb differently on planes with different atomic density may affect the relative
growth rate of various planes [31]. However, in solid state sintering processes such as
SPS, rapid heating rate can also be a useful parameter to control the grain morphology
and the growth mechanism. Hu et al. [30] have reported grain growth by a novel multi or-
dered coalescence mode of cubic nanocrystals of SrTiO3, where microstructure evolution
was governed by the rate of heating.
In the present work, nanocrystalline NiO was used as a model system to study the
grain growth behavior during SPS of nanomaterials. NiO is an important material with
a potential for thermoelectric material applications [32], anodes in Lithium-ion batteries
[33] and supercapacitors [34], where the interplay of porosity and grain size can be used
to tune the ceramic’s functional properties.
4.2 Results
TEM micrographs revealed that the as-received NiO powders formed nanocrystalline
agglomerates. The average crystallite size was estimated to be 10 nm from the dark
field TEM images(Fig. 4.1(a)), which correlated well with the crystallite size of 12 nm
obtained from XRD analysis (Fig. 4.2). From Fig. 4.1(b), the ring pattern was the
identified as a rock salt structure of NiO.
Figure 4.1: Dark field TEM image of agglomerate with an average crystallite size of 10
nm while (b) shows the indexed diffraction pattern.
17
Figure 4.2: XRD patterns of the top surface of the sintered pellet showing increased
phase fraction of the reduced Ni layer.
4.2.1 Reduction of NiO
The sintered samples had a thin metallic layer on the outer surface. Elemental analysis
of this layer showed that the layer consisted of metallic elemental layer. After thermal
etching the nickel oxidized to form Nickel Oxide againas (Fig. 4.3(a,b)). The difference
in the composition in both these cases is illustrated in the line scans in Fig. 4.3. The
oxidized micro-structure consisted of beautiful scales of NiO, as seen in Fig. 4.4
Figure 4.3: (a) Line scan across the top surface showing a Nickel layer which gets oxidized
after the thermal etching in (b). The marked line shows the interface.
18
Figure 4.4: (a) Microstructure of the Ni/NiO interface, (b) the growth of the Ni layer
into the NiO particle, (c) reoxidized surface of the Ni/NiO interface, (d) growth of NiO
scales after thermal etching.
4.2.2 Coalescence Driven Grain Growth of Nanocrystals
In the case of pressure assisted sintering, polyhedral grains with average size of 0.386 nm
and 0.473 nm were observed when the NiO was sintered at 900 and 1000C, respectively.
Pores were located either at the triple junctions or at the grain boundaries in accordance
to classical sintering theory (Fig. 4.5). As the sintering temperature increased to 1100C,
the polyhedral grains and spherical pores were no longer observed. Instead cuboidal pores
were seen (Fig. 4.5(b)). Fractured surface of the pressureless sintered samples at 1000C
showed cubic morphology of the grains with an average grain size of 68 nm and relative
density of 79% as seen in (Fig. 4.6(a)). At 1025C, the relative density increased to
88%, while the average grain size increased to 100 nm. At 1050C, the density increased
to 90%, and the microstructure was marked with the presence of larger micron sized
cuboidal grains along with nanosized cubes in their vicinity (Fig. 4.6(b)).
19
Figure 4.5: Fractured surfaces of the samples sintered for 10 min at 50 MPa pressure (a)
at 900C with polyhedral grains and (b) with the faceted cuboidal morphology developing
at 1000C.
Figure 4.6: Fractured surfaces of the samples sintered for 10 min at 5 MPa pressure with
(a) cuboidal grains at 1000C, (b) a few larger faceted grains are seen when sintered at
1050C and (c) cuboidal voids at 1200C with the inset showing the cubic pores at a
higher magnification.
A further increase in the sintering temperature showed that the average grain size
remained unaffected until 1075C, at which point an abrupt increase in grain size was
observed (from 110 to 1400 nm) at temperatures between 1075 and 1100C) as shown in
Fig. 4.7. After 1100C, the grain size did not increase rapidly and reached 2.25 µm at
1200C. Transgranular fracture of the sample sintered at 1200C revealed cuboidal pores,
as observed in Fig. 4.6(c).
20
Figure 4.7: Plot showing the change in grain size and relative density (inset) as a function
of sintering temperature for the samples sintered with a pressure of 5 MPa and 50 MPa.
The observation of the cuboidal morphology of the grains and the pores in the pres-
sureless sintering experiments at and above 1000C and the change in the pore mor-
phology from spherical to cuboidal between 1000 and 1100C for the pressure assisted
case indicates the presence of an activation barrier for this morphological transformation.
This transformation appears to be independent of the grain size or the relative density
as the pressure assisted samples had much higher densities and grain sizes when com-
pared to the pressure-less case (Fig. 4.7). The presence of polyhedral shape at lower
temperature and the formation of nanocubes at higher temperature could be the result
of different grain growth mechanisms active at different temperature ranges. In order to
verify whether the grain growth mechanism is a feature of any spark/plasma event of the
SPS unit or an inherent nature of the nanocrystalline system, the microstructure of the
annealed as-received powders was examined.
To further investigate the facet development and coalescence behavior, the powders
were observed in a TEM. The powders annealed at 900C showed nanocubes with a
size range between 3-200 nm. The attachment of the smaller nanocubes to the existing
grains provides evidence that the stable larger grains can act as base for subsequent
growth (Fig. 4.8(a)). At some locations, a few cuboids and nanopores were observed
in the TEM images, indicating that the nanocubes were formed through assembling of a
number of nano grains. This could lead to the formation of a continuous chain of particles
(Fig. 4.6(a) and Fig. 4.8(b,c)).
21
Figure 4.8: Sequence of the images showing the coalescence mechanism (a) number of
smaller cubes attached to a larger surface, (b) larger cuboidal block and a cube attach-
ment, and (c) the chain of particles formed as a result of the coalescence. All images are
taken from the powders annealed at 900C for 1hr.
Akin to the sintered samples, the annealed powders also had the nanocube grains and
the microstructure appeared to consist of nanocube shaped building blocks. The cubic
grains first appeared after annealing at 700C for 1hr, but a few nanocrystals remained
untransformed, as seen in Fig. 4.9(a). Temperature driven nanograin coalescence appears
to be an inherent character of the nanocrystals. Ultra high resolution SEM imaging of
the cube structures revealed that in some cases the edges had triangular facets, while in
some locations edges were perfectly formed, as observed in Fig. 4.9(b).
Figure 4.9: TEM images of the powders annealed at 700C for 1h with the initial irregular
morphology still being retained, while (b) shows the HRSEM image evolving cube with
some of the retained 111 planes.
To investigate the interfaces formed during the attachment process, HRTEM imaging
was carried out on both the annealed powder and as-sintered NiO. FFT and reconstruc-
tion of the square region marked in Fig. 4.10 of the pellet sintered at 1075C showed
lattice fringes with an interplanar spacing of 0.21 nm. Indexing using standard database
22
(JCPDS PDF: 47-1049) identified these interfaces as the 200 family of planes of NiO.
Hence, the nanocubes can be assumed to be made up of 100 planes.
Figure 4.10: HRTEM of the interface of an impinged particle, (b) showing the indexed
FFT pattern with‘t’ as the twinned spots, and (c) is reconstructed image of the selected
portion showing the twinned region. (d) HRTEM image of the interface between two
particles, (e) reconstructed image using the FFT showing an array of dislocations and
partial dislocations at the interface with the inset showing a magnified viewed of the
defect region. The sample was annealed at 700 C for 5 min.
The coalescence of the 100 planes has been investigated by intentionally arresting
the annealing process and then examining the microstructure. For powders annealed at
700C for 5 min and the sample sintered at 1075C for 10 min, twin configuration was
observed at some locations (Fig. 4.10(b)). Extra spots in FFT pattern (inset of Fig.
4.10(c)) confirmed the presence of the twin. Coalescence along 111 contact plane led
to twin formation in the both the annealed and the sintered specimen with [011] twin
axis, as indexed in Fig. 4.10(b). At other interfaces, reconstructed images using the FFT
patterns revealed an array of dislocations along the interface of two coalescing particles
(Fig. 4.10(d,e)) while at other locations the interfaces were defect free.
4.2.3 Effect of Boron Nitride layer
A clear distinction in the relative density, grain size and hardness was observed in the
case of electric field comparison studies. For a given temperature, the BN samples always
23
had a lower grain size, relative density and hardness, as inferred from Fig. 4.11.
Figure 4.11: Comparison of the (a) grain size (b) relative density (c) hardness for the
various configuraions studied.
The microstructural difference between the samples sintered with and without the BN
casing as shown in Fig. 4.12(a,b). In the case of the BN samples, the microstructure at
1050C resembles that of the conventionally sintered samples, with no trace of the cubic
morphology. The grains begin to develop polyhedral facets instead of the 100 seen in
the pressureless case. As the sintering temperature increased to 1250C, a few cuboidal
pores were observed along with the spherical ones, while the grain morphology remained
polyhedral.
Figure 4.12: Microstructure of the sintered pellet with BN blocking at (a) 1050C and
(b) 1250C.
24
4.3 Discussion
4.3.1 Reduction Behavior
The presence of the Ni rich layer indicates that NiO has undergone reduction. There are
several reasons for this,
• In-situ reduction due to the electric field. Holland et al. [20]have reported that Ni
can form during in-situ TEM experiments when electric field is applied via an STM
tip.
• Reduction due to reaction with graphite.
• Decomposition to Ni due to presence of vacuum in the chamber, following the
Ellingham’s diagram.
As the Nickel layer was present in the case of the BN-NiO case as well, the loss of oxygen
can be ascribed to the low oxygen pressure in the chamber. However, the contribution
from other factors have still to be evaluated.
4.3.2 Grain Coalescence
The results suggest that the microstructure of the SPS and annealed samples is similar.
Thus, it is likely that formation of grain interfaces and hence the grain growth mecha-
nism followed a similar path in both the cases. It has to be noted that grain growth in
nanograins cannot be characterized by a single growth mechanism, since the attachment
and coarsening of the nanograins may be occurring simultaneously. Extensive in-situ
TEM experiments are necessary to unambiguously determine the grain growth mecha-
nism [35]. In this work, OA was assumed to be the dominant mechanism based on the
microstructural observations.
If the nanocrystals are treated as individual entities, understanding of their growth
mechanism must consider the driving force for an oriented attachment growth, the shape
of the final assembly and the mechanism of attachment [36]. The mode of growth by
crystal attachment has been well established during hydrothermal growth of nanoparticles
in colloidal systems and has been termed as grain rotation induced by grain coalescence
(GRIGC) [37].
25
Gong et al. have suggested that OA occurs in a stepwise process where the initial
nanocrystal aggregates rotate to facilitate the collision of high energy surfaces followed
by the removal of surfactants and formation of coherent interfaces [38]. By adjusting the
reaction temperature and annealing time of FeS2 nanoparticles they have obtained cubic
and sheet morphologies. In the present work, particle clustering in the as-received powder
and the annealed powders, along with the appearance of cuboidal grains at 1000C in
the sintered specimens, suggests that a critical temperature needs to be achieved before
sufficient kinetic energy is available for the nanocrystals to overcome attractive forces
within the raw powder agglomerate. For this to occur, sufficient Brownian motion of the
crystals has to be ensured, especially for solid state growth where surfactants or a liquid
medium which can aid attachment are not available. At high temperatures (above a
critical temperature, Tc), the rapidly moving nanoparticles may collide with a favorable
crystallographic plane and then coalesce [38]. The observed initial particle clustering
(where common crystallographic orientation may not be achieved) could aid in the OA
process as the cluster in Fig. 1 may act as the precursor for subsequent formation the
cubic grains [16]. The formed clusters then transform into larger single crystals, as seen
in the sintered sample in Fig. 4.5 and in the sample annealed at 900C (Fig. 4.9). The
formation of single crystal has been confirmed from the electron diffraction patterns.
The starting nanocrystals form clusters and assume a shape depending on the surface
energies of the planes. The equilibrium shape is given by Wulff plot [39], where the higher
surface energy faces grow faster and finally are not present in the equilibrium morphology.
Lee et al. [31] have proposed that intrinsic surface energy of crystallographic faces, amount
of adsorbing species, sufficient time for growth of thermodynamically stable structures
and molecular precursors are the four factors that govern the growth behavior in the case
of hydrothermal growth. In the case of solid state sintering without any sintering aids,
the surface energies of the crystallographic faces of the starting particles and sufficient
time at the growth temperatures are the applicable factors. Hence, the competition
between these thermodynamic and kinetic factors has to be studied to understand the
grain growth behavior.
The temperature dependence of the growth of various planes has to be considered,
taking into account that NiO crystallizes in the rock salt configuration with O2− at the
(0,0,0) position and Ni2+ at the (12
,12
,12) positions in the investigated temperature range.
26
Typically for grain growth of rock-salt crystals, a polyhedron is assumed to be formed
initially at low temperature and is bound by six 100 and eight 111 faces. The final
crystal morphology is dependent on the rate of evolution of these faces [40]. Theoretical
calculations have shown that 100 facet in NiO has the lowest surface energy (0.958
J m−2) while the 110 is the second lowest (1.26 J m−2) [41]. HRTEM images of the
NiO cubes showed that the interplanar spacing was 0.21 nm, corresponding to the 200
planes. Hence the surface of the NiO cube was formed by six 100 faces. If sufficient
thermal energy (kT) is provided, rapid growth occurs in the <111 >direction and cubic
morphology evolves. In case sufficient energy or time are not available, truncated octa-
hedron can result due to the slower growth kinetics of the 111 planes when compared
to the 100 faces. While for the pressureless sintered NiO specimens there was sufficient
thermal energy for the 111 planes to grow, in the pressure-assisted case, the polyhedra
remained. Hence the triangular planes observed in Fig. 5(b) are the 111 planes that
have not yet fully grown out.
Chaim and Bar-Hama [21] have previously densified nano-NiO using SPS with the
powders from the same manufacturer that were used in this study. Interestingly, they
have not reported the formation of any nanocubes and their TEM micrographs reveal only
conventional polyhedral structures at 900C. In the present study, when a pressure of 50
MPa was applied at 900C, NiO grains with polyhedral morphology were obtained in Fig.
4.5. Sahoo et al. also sintered nanocrystalline NiO using SPS in the temperature range of
400-800C and have reported clean grain boundaries [42]. Diffusion-driven mechanism can
be responsible for such grain growth at lower temperatures. However, there was a shift
in the temperature associated with the evolution of nanocubes, with the highest being
for pressure–assisted, then pressure-less sintered and finally for annealed samples. This
can be attributed to the steric hindrance (function of relative density) for the nanocrystal
movement [43].
The subsequent microstructure evolution is dependent on the attachment of the
nanocubes. For particle coalescence, attachment may begin at the edges and corners
(as seen in Fig. 4.9(a)), where the electric double layer that prevents coalescence is
absent [44]. The initial attachment is followed by the formation of the quasi sintering
neck, as seen in Fig. 4.13(b)). Formation of such necks was observed in the case of tin
oxide [37]. The kinetics of subsequent coalescence were described with the A1+A1 model,
27
where starting nanocrystals (A1 blocks) combined with each other by sharing a face and
decreasing the surface energy during coalescence. The blocks formed in Fig. 4.13(b)
appear to be the result of such a process. However, as seen in Fig. 4.9(c), a continuous
coalescence of different sizes of crystals was also observed. Therefore, the overall growth
process is likely described by the Ai+Ai model, where the attachment of primary crystals
to an existing higher-order crystal substrate takes place. Hence, the abrupt jump in the
grain size at 1100C may be associated with the primary cubes coming together to form
larger crystals. Similarly variation in grain size was observed in the case of SrTiO3 [30],
CeO2 [45] and it was suggested that the larger grains were formed by aggregation with
smaller ones. Interestingly, the assemblage of the smaller nanocubes could not be cap-
tured in the present work likely due to the narrow window of the exaggerated growth.
In-situ TEM studies by Theissmann et al. [43] revealed that the coalescence process oc-
curs in less than 1/25 s. Thus, the initial grain clustering and the sudden surge in grain
size indicate that microstructure evolution progressed via a continuous buildup of the
nanocubes.
Figure 4.13: (a) SEM image of the sample annealed at 1200C showing the initiation
of the attachment process and (b) TEM image of an annealed powder at 700C for 1h
showing the cubic building blocks within a large cluster and a smaller faceted particle is
shown to attach to this cluster.
Rapid heating in the SPS enables the initial nanocrystals to reach higher temperatures
without any coarsening or necking in order to promote the formation of nanocubes. Hu
et al. observed that multi ordered coalescence of cubic nanocrystals was promoted by
a high heating rate. In contrast, at lower heating rate the ordered coalescence was
significantly reduced due to the formation of necks at lower homologous temperatures.
28
These studies suggest that a critical temperature should be reached rapidly in order
to facilitate grain motion. In this study, a heating rate of 200C min−1 was applied
during SPS processing, which could have enabled sufficient Brownian motion for random
collision and coalescence of the nanoparticles. The pores observed in Fig. 4.5 were likely
the result of improper building of the microstructure due to excessive kinetic energy of
the nano sized units. Similarly, annealing of the nanopowders at a high temperature
enabled the activation of such grain motions and morphological transformations, further
establishing the significance of the heating rate in SPS. Such a pathway is active only when
sufficient Brownian motion is available for the nanocrystal motion without significant
grain coarsening.
Rotational activity continues until a common low energy configuration is obtained,
which is governed by either a parallel epitaxial relationship or a twin configuration (Fig.
8(b)). Twinning was observed in the case of other oxide particles, such as CeO2 and
TiO2 [46] during coalescence of 111 and 112 planes, respectively. In the present work,
attachment involved the contact plane 111. However, incomplete attachment process
can lead to the formation of line/planar defects which are local minima in the energy
configuration [47]. The presence of large number of line defects, planar defects and pores
cannot be explained by OR. In the model developed by Huang et al., it has been shown
that during coalescence dislocations arise as a result of the large strain mismatch and the
critical grain size for dislocation generation is inversely proportional to the strain [48].
Dislocations have been found during SPS of Al2O3, which were attributed to the strain
created at the grain boundaries during the sintering process. However, these dislocations
disappeared after annealing [49]. In the present work, the presences of such defects in the
annealed powders as well as the sintered samples indicate that the coalescence process
might be a source of these defects.
Bulk material with tailored interfacial defects, morphologies and porosity are impor-
tant for energy, sensing and catalytic applications. Along with the conventional use of
SPS for densifying ceramics, the results of this study indicate that it can also be used as
important tool to control the morphology as well the interface defects when the pressure
and temperature regimes are carefully chosen.
29
4.3.3 Effect of Electric Field
The role of the BN layer can be attributed to two factors, namely as an electric field barrier
and a thermal barrier. While previous studies and our own results on MgO have shown
that BN is a reliable material for such experiments without inducing large temperature
gradients, the cause for the shift in the morphologies and the data obtained in Fig. 4.11
and Fig. 4.12 have to be attributed to the field itself. Tang et al. [50] have shown that
there is difference in the hardness when a electric field was applied. The lowering of grain
size when a field electric field was applied has been investigated previously. Although,
we have not completely ruled out the presence of the temperature gradient, the results
here are an indication that the electric fields may enhance the densification in SPS.
4.4 Conclusions
A novel cubic growth mechanism of nanograins was reported for the first time during
sintering of NiO. For the 3-D microstructure construction, cubic nanostructures act as
building blocks. As evident from TEM analysis, larger cubes, due to their relative stability
can act as a substrate to which smaller cubes attach. A schematic of the process is
illustrated in Fig. 4.14. Crystal growth via the OA mechanism resulted in the formation
of building blocks, along with defects at the interfaces. The observation of a similar
growth behavior in the annealed powders and SPS samples suggests that the coalescence
behavior may not be a feature of SPS, but a characteristic grain growth mechanism of
nano sized particles at high temperature. The electric field has been shown to influence
the morphology and the properties, however the factor of temperature gradients needs to
verified.
Figure 4.14: Graphical representation of the coalescence mechanism
30
Chapter 5
Results and Discussion: Spark
Plasma Sintering of ZnO
5.1 Introduction
Zinc oxide (ZnO) is a wide band semiconductor with applications in the electronic, optical
and bio medical industry. Grain growth and densification studies of pure and doped ZnO
processed by conventional sintering, hot pressing (HP)and SPS, have been previously
reported [23]. A fundamental study by Holland et al. to evaluate the thermal contri-
butions due to rapid heating rates and athermal contributions arising from the applied
field during sintering of ZnO led to the conclusion that electric fields coupled with the
high heating rate enabled enhanced sinterability of ZnO [51]. However, Langet et al. [23]
showed that electric field did not influence the sinterability of ZnO. Such contradictory
conclusions on the role of electric field during sintering of ZnO in literature could be due
to the differences in the range of the strength of the applied electric field, particle size of
the starting ZnO powders and investigated temperature ranges.
The synergetic effect of electric fields, rapid heating rates and applied pressure make
SPS a difficult process to analyze from the existing framework and much uncertainty
remains on the operative mechanisms [52]. Therefore it is of interest to reinvestigate the
microstructure of ZnO ceramics processed via SPS, with an aim to gather evidence of
athermal contributions (if any), in order to yield a better understanding of the atomistic
phenomena during consolidation via SPS. As traditional sintering theories have been
insufficient to explain the rapid sintering associated with SPS, alternate mechanisms
31
should be explored along with careful evaluation of the microstructure sintering analysis
to understand the operative mechanism. In the current study, densification behavior and
microstructure evolution of ultra-fine ZnO powder processed via SPS was studied and
the grain growth mechanisms were evaluated critically with consideration of liquid-like
neck formation and nano grain rotation.
5.2 Results and Discussion
The as-received ZnO powder had a bimodal particle size distribution with faceted particles
in the range of 150-200 nm, along with fine spherical particles of 20-30 nm in diameter
(Fig. 5.1).
Figure 5.1: SEM micrograph of the raw powder with the inset showing a magnified image
of the nanoparticles.
XRD analysis confirmed that the starting powder and all the sintered samples had
peaks corresponding to hexagonal Wurtzite ZnO phase (Fig. 5.2).
32
Figure 5.2: XRD micrographs of the received ZnO powder and pellets sintered at various
temperatures.
5.2.1 Densification and Grain growth
Effect of Temperature
As the sintering pressure increased to 50 MPa from the pre load of 5 MPa, the rela-
tive density increased from 45% to 62% and remained constant until shrinkage began at
350C. The development of the relative density in Fig. 5.3(a) with respect to tempera-
ture indicates that the densification rate reached a maximum in the 600-700C sintering
temperature range.
Figure 5.3: (a) Evolution of sintering trajectory and (b) grain size-relative density trend,
as a function of sintering temperature for a dwell time of 5 min.
33
The difference in the kinetics of densification and grain growth were reflected in the
nature of porosity and grain morphology. At 600C, the microstructure contained open
interconnected porosity, while the grain surfaces were curved, with an average grain size
of 290 nm (Fig. 5.4(a)). With further increase in the sintering temperature, a gradual
change in the grain morphology towards a faceted polygonal morphology was observed,
with equilibrium isotropic tetrakaidodecahedron shape being obtained at 800C (Fig.
5.4(b)). At this temperature the porosity was predominately seen at grain vertexes. At
1000C, pores were seen on the facets of grain boundaries and at grain intersections (Fig.
5.4(c)).
Figure 5.4: Fractured surfaces of ZnO sintered at (a) 600C (b) 800C (c) 1000C.
As the sintering temperature increased, the density increased from 80.5% at 600C to
92.1% at 700 C, while the grain size increased from 166 nm to 443 nm with the increase
in sintering temperature. The relative density increased to 99.4% at 800C with a grain
size of 1.52 µm. Additional rapid grain growth occurred in the 900-1000 C range, with
the grain size increasing to 4.3 µm. This grain growth occurred after maximum relative
density had been achieved (Fig. 5.3b) shows that as the sintering temperature increased
between 600-800C, there was a rapid increase in density with minimal grain growth.
Further increase in sintering temperature resulted in a minor change in the density.
The change in pore morphology may be related to the relative mobility between
grain boundaries and pore as a function of temperature. Typically, at relatively lower
temperatures isolated pores at grain boundaries and grain intersections restrict grain
growth. As the sintering temperature increases, grain boundary mobility exceeds pore
mobility, and the pores are absorbed into the grains. With the grain boundary pinning
no longer in operation, rapid anisotropic grain growth occurs [53]. Formation of intra-
34
granular surface porosity could be a result of closed pores rising to the surface due to
curvature differences between the leading and trailing ends of the pore [54].
5.2.2 Observation of sintering necks
Several locations of the as-sintered disc often contained regions with a liquid-like necks
(Fig. 5.5) suggesting the possibility of a local melting phenomenon. Similar liquid-
like morphology has been previously observed at low homologous temperatures during
SPS processing of Cu [55], Inconel 718 [56] and Mn-Zn ferrites [57]. In the case
of Copper, the necking and liquid-like morphology was attributed to a spark discharge
effect. For Mn-Zn ferrites the liquid like film was attributed to electro-magnetic effects,
while internal current was cited as the reason for the melting of Inconel 718. However,
recent rigorous SPS experiments (using Al2O3 and Cu) remain inconclusive in confirming
the absence/presence of a discharge effect [11] , and the role of electric field and current
has to be further investigated in the case of ionic ceramics.
Figure 5.5: Presence of bridging necks between the particle 600C in (a), (b) and at
1000C in (c).
Further, elemental analysis was carried on the raw powders, sintered samples and in
the neck regions in order to confirm the formation of a new phase or the presence of an
impurity. Elemental maps on the raw powders showed the presence of only Zn and O in
the raw powders and sintered specimens. In the grain boundary regions of the sintered
samples, line scan profiles were taken to look for any compositional gradients. In Fig. 5.6
along the clean boundaries, a uniform composition of Zn and O atoms was observed.
Similar line scans across the sintered necks showed that the composition was ZnO as
seen in Fig. 5.7, further driving the case of high local temperatures.
35
Figure 5.6: Line scan across the grain interfaces showing a dip in Oxygen concentration
in (a) while the composition is uniform in (b) for the sample sintered at 1000C for 0
min.
Figure 5.7: Line scan across the grain interfaces with the necks showing uniform Oxygen
concentration in (a) and (b) for the sample sintered at 1000C for 0 min.
36
Hence the further part of the section shall deal with the theoretical arguments for the
formation of such sintering necks. Thermal gradients can be expected during SPS, both
at the macro scale and at particle interfaces. On the macroscopic scale, the gradients
can arise due to thermal conductivity of the material, radiation effects and the sample
geometry. Finite element models (FEM) have confirmed the possibility of macroscopic
temperature gradients between the edge and the center of the disc in SPS [58]. Melting
of ZnO particles, however, would require temperatures of 2000C, which is unlikely to
occur on the macroscopic scale. Meanwhile, local gradients at the grain boundaries and
particle interfaces can also occur due to preferential Joule heating of the grain boundaries,
dielectric breakdown or the development of a space charge layer [51]. The possibilities
of such field effects in ZnO leading to the observed necks are examined further.
In the present work, the applied field strength was 6.5V/cm−1, whereas techniques like
flash sintering use much larger field strengths. The FEM model developed by Holland
et al. [59]however showed that the local field strength at the particle interfaces can
magnify several times in the case of dielectric materials (due to charge polarization). For
ZnO (dielectric constant =10), a magnification of 30x has been predicted. Narayan [60]
proposed a new mechanism for rapid sintering of ceramics, where accumulation of point
defects along preferred grain boundaries and dislocations increases conductivity. At high
electric field strengths localized melting is possible. Similar localized melting can also
occur in varistor ZnO during electric runaway. Also ZnO treated with pulsed DC current
may cause breakdown by puncture mode and cracking mode, depending on the nature of
the current. In the puncture mode, formation of a through hole from the anode to the
cathode takes place due to vaporization of the ceramic.
Raj et al. [52] have postulated nucleation of Frenkel pair under the influence an exter-
nal field as a fundamental densification mechanism in Field Assisted Sintering Techniques
(FAST). Frenkel pair generation and migration in non-stoichiometric ZnO under electrical
and thermal fields has been reported in the grain boundary defect model for instability in
ZnO varistors [61], where the presence of Zinc interstitials is related to dielectric break-
down. Under thermal and electrical driving forces, intrinsic Zinc interstitials (Znix) are
formed by reaction of defects at the grain boundaries. Accumulation of the neutral Zinc
keeps taking place as long as the electric field is applied, thus increasing conductivity and
ultimately leading to breakdown.
37
From the current observation of necks with a liquid like morphology between sintered
particles (Fig. 5.5), it can be concluded that this phase could in fact aid the densification
and grain growth process. The presence of the liquid layer can enhance the diffusivities
by 6-8 orders of magnitude when compared to solid state sintering, leading to the rapid
densification often expected in SPS experiments. It can be postulated that defect driven
localized currents lead to the observed selective melting of the particle surfaces. Non-
uniform local field strength arising from grain morphology, particle size distribution, stage
of sintering and contamination could be some of the factors leading to the evolution of
the liquid film.
5.2.3 Grain Coalescence
Presence of liquid film possibly enabled densification by nano grain rotation in the samples
sintered at 600C, 800C and 1000C (Fig. 5.5). At 600C, the clusters appeared to be
particle agglomerations (20-30 nm), while at 800C and 1000C clusters were a part of a
well-defined grain (Fig. 5.8). Similar observations of nano grain clusters were observed
during SPS of Al2O3 [62] and Yttrium Aluminum Garnet (YAG) [63].
Figure 5.8: Growth of nano grain clusters from (a) 600C to (b) 800C and (c) 1000C.
Cluster formation behavior can be explained by the grain rotation mechanism in nano
crystalline ceramics [29]. Dense nano clusters (Fig. 5.8(b,c)) can form via a coalescence
route when individual nano grains with high angle grain boundaries rotate until they form
low angle grain boundaries with adjacent nano grains. This mechanism is in contrast
to the traditional grain growth by grain boundary migration, where larger grains grow
at the expense of smaller grains. The relative probability for grain rotation is inversely
38
proportional to the fourth degree of the grain size. Also the extent of subsequent rotation
decreases with each rotation step. Hence, only a fraction of the starting powder particles
are expected to undergo such a densification mechanism, while the rest of the larger
particles undergo curvature driven grain boundary migration [64]. At all temperatures,
there is competition between grain boundary rotation and grain boundary migration.
As the sintering temperature increased the nano grains grow and the tendency for grain
boundary migration increases while that of grain rotation decreases, ultimately leading
to the formation of a single larger grain.
Densification assisted by the grain rotation mechanism is believed to be accelerated
by the presence of a liquid layer. The observation of the liquid like-morphology and
the presence of nano-grains, suggests the possibility of liquid phase driven grain rotation
during SPS of ZnO.
5.2.4 Effect of Dwell Time: Sintering Analysis
For the samples sintered at 600C, increase in holding time from 0 to 15 min at 600C
caused the relative density to increase from 78.2% to 84.3%, while the change in grain
size was statistically insignificant (Fig. 5.9). The microstructure at the end of 15 min
also showed a reduced porostiy, possibly due to better particle rearrangement and break
down of any agglomerates.
Figure 5.9: (a) Change in relative density and (b) grain size as a function of dwell time
at 600C, 700C, 800C and 1000C.
A stress exponent l of 1.24 was calculated from the slope of the plot (Fig. 5.10(a)),
suggesting that the densification was aided by diffusional process, either lattice diffusion,
39
grain boundary diffusion or by viscous flow. From the presence of the liquid morphology
and the formation of nano grain clusters, it is expected that the ZnO particles which are
coated with the viscous layer underwent densification via diffusion through this medium.
Sintering at this temperature was seen to be driven by the formation of nanograin clusters,
the development of necks and minimal grain growth.
The plot of lnG vs lnt (Fig. 5.10(b)) enabled the calculation of the grain growth
exponent for densification. It is generally accepted that the value of the grain growth
exponent characterizes the nature of the grain growth behavior during sintering. At
700C, n was 2.13, while at 1000C and 800C, n was 2.63 and 2.71 respectively. These
values of suggest that the grain growth primarily occurred via volume diffusion through
a liquid layer (n= 3). Previous investigators have also reported a value of n=3 in the
case of undoped ZnO [65]. Such a value can be justified by the formation of the liquid
layer at the particle interfaces, across which migration of Zn2+ ions takes place.
Figure 5.10: (a) Calculation of stress exponent and (b) grain growth exponents.
At 700C sintering temperature, as the holding time increased to 15 min, the relative
density increased from 85.4% to 97%. At 0 min, rounded grains were present. At the end
of the 5 min sintering cycle, the porosity had changed to closed and well defined grains
started to appear (Fig. 5.11(a)). During this time there was a rapid increase in density,
but a slow increase in grain size (Fig. 5.9). Pores at grain quadrature also decreased in
size with increase in holding time. For a sample sintered for 15 min at 700C only a few
pores remained, as seen in Fig. 5.11(a).
At 800C sintering temperature, evolution of pores and grain morphology reached a
steady state. At a dwell time of 10 min, a nearly pore free microstructure was obtained
(Fig. 5.11(b)). At 1000 C fractured surfaces of samples revealed intra-granular surface
40
porosity and pores at grain intersections (Fig. 5.4(c)). The size of these pores increased
with dwell time (Fig. 5.12).
Figure 5.11: (a) SEM images of fractured surfaces of samples sintered at (a) 700C with
15 min dwell and (b) 800C for a dwell time of 10 min.
Figure 5.12: Pore enlargement with increase in dwell time at 1000C.
An increasing negative shrinkage was observed from the ram displacement data for
all the samples sintered at 1000C (Fig. 5.3(b)). A decrease in density after achieving
maximum density has been previously linked to the entrapment of inert gases in the
pores, where the increased gas pressure with an increase in temperature leads to pore
enlargement. However, no appreciable decrease in relative density was observed for any
of the samples sintered at 1000C (Fig. 5.9(b)). An additional sample sintered at 1200C
exhibited further negative shrinkage (from ram displacement data), while its density
41
remained approximately constant. The cause for the observed trend however needs further
investigation. Recently, Olevsky suggested that local temperature gradients due to pulsed
frequencies during SPS can cause thermal diffusion, which drives vacancies and atom
separation. He speculated that the early stages of this mechanism, necking would be
enhanced, but at the final stages, pores act as vacancy sinks under a diffusion gradient,
leading to pore coarsening and hindrance of the densification process. In the present
work, 700 and 800C, shrinkage of pores due to vacancy annihilation could have taken
place while the predicted trend reversal possibly occurred at 1000C.
5.3 Conclusions
Sintering behavior of ZnO prepared through the SPS route was characterized by studying
the microstructural changes and correlating them with possible densification mechanisms.
An increase in relative density was observed up to 86%, without significant grain growth.
Thereafter densification was associated with slow grain growth thereafter. Closure of
inter-granular grains took place, resulting in the achievement of near pore free microstruc-
ture at 800C. Further increase in sintering temperature led to increase in intragranular
surface pore size and negative shrinkage behavior.
Holding time was found to be more significant at lower SPS temperatures in aiding
densification. A window between 600-700C was identified, where the time was most
influential. In this temperature range interplay of sintering parameters controlled the
final porosity and grain size.
The evidence from microstructural analysis supported by literature suggests the pos-
sibility of local melting. Formation of necks with the liquid morphology indicated that
the densification and grain growth was assisted by diffusion through this medium. This
film may also have enabled dense nano grain clusters which contributed to densification
by grain rotation. These observations were supported by sintering analysis exponents (l,
n) which suggested a liquid phase aided densification process.
The significance of electric fields cannot be over emphasized as the effect of pressure;
pulsed current other possible mechanism have not been accounted for. However, the
microstructural evidence indicates that a melting event could have taken place.
42
Chapter 6
Results and Discussion: Spark
Plasma Sintering of MgO
6.1 Introduction
MgO is a well-investigated oxide because of its use as sintering aid and as a refractory.
Recently, transparent MgO ceramics were produced using hot pressing with LiF as an ad-
ditive [66]. Such transparent material are potential replacements for the sapphire windows
used in shock wave experiments. By sintering in vacuum, Misawa et al. have produced
transparent MgO at 1600C. For such applications, a near pore free microstructure is
required and SPS has become a proven technique for producing such pore free ceramics.
Chaim and co-workers have developed transparent MgO using SPS previously and also
have developed detailed densification maps for the sintering at high pressures [9]. Using
the grain growth kinetics, various mechanism were deducted in 900-1420C range [67].
But in light of the recent work by J. Narayan [19] [18], it becomes crucial to investigate
the role of each of the parameters during SPS of nano Mgo.
6.2 Results and Discussion
6.2.1 Decomposition of Mg(OH)2
The as-received powder had a significant amount of Mg(OH)2, as indexed in the XRD
pattern. While it is possible to carry out SPS experiments with this powder by allowing
43
the in-situ decomposition during SPS, the volume change during the phase transfor-
mation is significant to cause sudden displacements and large changes in the estimated
dimensions. Hence the powders were loaded into a muffle furnace and the annealing time
was optimized as 375C for 1h, as the decomposition reaction to form MgO occurs at
332C,as shown in Fig. 6.1.
Figure 6.1: XRD patterns of the heat treated raw powders at various temperatures
After the heat treatment, the powders were deposited on a carbon grid and loaded
into a TEM. The powders were agglomerated with an irregular morphology as shown in
Fig. 6.2. An average crystallite size of 12 nm was calculated. All further SPS experiments
were conducted on this heat treated powder.
Figure 6.2: Bright field image of the heat treated powder with the inset showing the ring
pattern.
44
6.2.2 Grain growth and Densification
The shrinkage curves obtained from the SPS machine for all the samples have been
compared at 1200C. The shrinkage curves for the samples sintered at 50 MPa, showed 3
distinct trends after the initial pressure application at the end of 4 minutes. The shrinkage
increased linearly with temperature till until a slight change of slope was observed in 1050-
1100C temperature range. This was followed by a negative shrinkage as marked in the
Fig. 6.3. However the samples did not show a significant decrease in the density value
that would have corresponded to the observed shrinkage as indicated. A change in the
slope of the shrinkage curve indicates a change in the sintering mechanism.
Figure 6.3: Shrinkage curves of the samples sintered at 1200C with and without the
application of pressure.
The shrinkage curves for both the pressureless samples, with and without the BN
casing showed a similar trend for all the samples. Here also, after the initial shrink-
age corresponding to particle rearrangement and the breakage of soft agglomerates, the
shrinkage increased linearly till the 1150C window discussed above. Here, a prominent
second sintering step was observed that was not observed clearly in the case when pressure
was applied.
The grain sizes estimated from the SEM images and the relative density (theoretical
density of MgO as 3.58 gcm−3) is plotted in Fig. 6.4. The areas marked show that within
experimental error the grain sizes and relative density of the sintered samples in that
window, whether BN layer was used or otherwise, follow a similar trend.
45
Figure 6.4: Comparative (a) grain size and (b) relative density plot of sintered MgO with
the clouds showing similar trends.
At 1100C, for the 50 MPa case, the grain size in the conventional case was 3.01
µm, while in the field blocking experiment it was 3.77 µm. The corresponding relative
densities were 96.32% and 97.51%. For the pressure case, the grain size in the both the
cases was comparatively lower, with 1.04 µm in the conventional case and 1.13 µm in the
BN case. The relative densities of both these samples were similar, with 89.1% for the
BN case and 88.1% for the other.
At 1200C, in the pressure assisted sintering case, the average grain size was 7.14 µm
while for the BN case it was 7.32 µm. The relative densities were 96.92% and 96.53%
respectively. In the pressureless case, for the BN case, the grain size was 3.579 µm while
in the other case it was 3.017µm . The relative density in the case of the BN sample was
92.81% while for the conventional sample it was 91.67%.
As marked in the Fig. 6.4, there are distinct regions in each of the plots. Pressure
application has improved the density of the sample to the window of 95-98% from the 89-
92% window in the pressure-less case. This is an expected behavior as the external driving
force increases, the sinterabilty of the material increases. Also the samples sintered at
higher loads always had a higher grain size. This may be misleading as the relative
densities are not comparable, and 92$ is the commonly suggested density value for the
initiation of the third stage of sintering that involves rapid grain growth. More crucially,
the samples sintered at a given pressure, whether a BN layer is used or not fall in the
same window as seen in Fig. 6.4. This implies that the role of BN seems to be limited in
the set of sintering parameters investigated in this study.
46
Further, the rapid heated sample at 200Cmin−1 showed a lower density at 84.4%
when a 50 MPa load was applied. All the rapidly heated samples crumbled after sintering,
although the microstrucuture shown in Fig. 6.5 had no open porosity that would have lead
to any cracking. The cracking of the sample probably occurred due to the in-homogeneous
thermal gradients generated as a result of the rapid heating.
Figure 6.5: Fracture surface of the sample sintered at 1200C with a heating rate of
200Cmin−1 and a pressure of 50 MPa.
6.2.3 Microstructural Characteristics
Microstructural evolution was tracked for all sample by investigating the fracture surfaces.
The microstrucutres of all the sintered samples was characterized by the presence of
sintering necks in the early stages and terrace-like features on the exposed grain surfaces
in the later stages, as seen in Fig. 6.6(a,b). The necking behavior falls in line with the
classical theory of diffusion driven sintering in the initial stages of sintering. Infact, most
of the microstructural evolution follows the changes that were observed in the case of
ZnO (Refer previous chapter), albeit with a shift in the temperature.
For the 5 MPa samples, at 1100C, the grains were still not faceted with the open
porosity closing down, as observed in Fig. 6.7(a,b). With an increase in sintering tem-
perature to 1300C, the tetrakaidecahodron grains develop with the porosity present at
grain boundaries(Fig. 6.7(c,d). When sintered at 1100C with 50 MPa, the particles
had already become faceted in both the regions as seen in Fig.6.8(a,b). For the sample
47
sintered at 1300C with a pressure of 50 MPa, the pores could be observed entrapped
within the matrix indicating that the grain boundary mobility had become more than
the pore mobility and the final stage of sintering is in process. (Fig. 6.8(c,d)).
Figure 6.6: Microstrucutural feature of the sintered pellet at 1300C with BN (5 MPa)
showing the neck formation while (b) was sintered at 1300C (50MPa) showing a stepped
surface.
Figure 6.7: Comparison of the microstructure of the sintered samples at 5 MPa with
(a),(b) sintered at 1100C and (c),(d) at 1300C with and without BN respectively.
48
Figure 6.8: Comparison of the microstructure of the sintered samples at 50 MPa with
(a),(b) sintered at 1100C and (c),(d) at 1300C with and without BN respectively.
In addition to the similarity of the shrinkage curves, the grain sizes and the relative
density, the microstuctural feature are remarkably similar between the samples sintered at
the same pressure. As seen in Fig. 6.7(a,b), the morphology of the grains is still spherical
in both the cases and evolves to the equilibrium morphology at the same temperatures.
All these observations question the role of electric fields during SPS. With no signif-
icant differences in the microstructure, grain size or the relative density, the utility of
SPS maybe be limited to the application of high loads and rapid heating rates. While
the results of the field experiments by Narayan and co- workers demonstrated the role of
electric field during field treatment of MgO, the fields applied are significantly large than
those commonly used in SPS. By testing the role of the electric field at high pressure,
where plastic deformation is expected to occur and at low temperature, the results indi-
cate that the electric field does not couple with pressure to enhance the sintering rate in
the investigated temperature and pressure regimes.
49
6.2.4 Terraced Oxide Growth
While the presence of steps has been previously attributed to plastic deformation, in the
temperature regime 773-1173K, the yield strength of MgO follows the equation given
below [9],
σy = −0.215T + 314.26
At 1173K the yield strength calculated is 64 MPa. Hence for the temperature range of
1373-1673K investigated in this study it is assumed that the yield strength shall drop
below 50 MPa. Hence, if the presence of terrace features (Fig. 6.9)are attributed to the
plastic deformation, the presence of such structures in the pressure-less experiments at 5
MPa does not fit into this picture.
Figure 6.9: Fractured surface of the sample sintered at 1100C with 5MPa showing
terraced grain structure and (b) showing a pyramid-like terraced structure.
The pyramidal-like shaped terraces observed on the grain surface can also be due to a
vapor-solid process. Using Atomic Force Microscopy Maestre et al. [68] investigated the
growth of such planes in SnO2 samples. The growth such pyramidal features depends on
the crystallographic orientation, surface defects.
Rod Growth During SPS
As seen in Fig. 6.10, large micron sized elongated structures were observed to grow from
the existing grain surfaces and grain boundaries. The nucleation point of such structures
was observed at the tips, at higher magnifications it was observed that the rods also grew
50
from the terrace structures. Maestre et al. [69] had observed similar elongated structures
during sintering of TiO2 along with the observation of the terraced structure. They have
observed that the rods and terraces grew when the sample was subjected to different
sintering cycles. The mechanism for such a growth is a vapor-solid process in which the
source and the substrate for the growth is the compacted grain itself. The nucleation
site for these elongated structures are the single nucleation sites present on the tip of the
terraces.
Figure 6.10: (a) Micron sized rod-like structures while (b)shows the magnified image of
a smaller rod growing from the tip of the terrace structure. The sample was sintered at
1100C with 5 MPa pressure.
Another feature that was observed on the surface of the grains were triangular island
shaped structures. On the fractured surface of the grains, prismatic structures were
decorated on the grain boundaries as seen in Fig. 6.11(a). These structures also have
a terrace-like structure (Fig. 6.11(b)). Whether these are the initial stages of the rod
growth needs to be still investigated. Previously it has been reported that oxidation of a
Cu-5%Pt alloy and the subsequent growth of the Cu2O phase occurred via such an island
nucleation [70]. These triangular protrusions are expected to grow by consuming the local
terrace atoms. Using first principle density functional slabs Jennison and Bogicevic [71]
have shown that even species like (OH)−1 can destabilize the bonds on the surface and
cause nucleation of islands even at elevated temperatures. In the case of SPS of MgO, the
presence of residual Mg(OH)2 can release the hydroxide ions that can aid in the growth
of such structures.
51
Figure 6.11: (a) Triangular islands growing on the surface on the existing micron sized
grains,(b) magnified image of such a structures.The sample was sintered at 1300C with
50 MPa pressure with BN.
6.3 Conclusions
The effect of electric field, pressure and heating rate on the densfication and grain growth
of nano MgO was investigated. It has been observed the pressure helped in lowering the
sintering temperature, while the electric field did not have any significant effect on either
the densification or the grain growth. The trends observed at a pressure of 5 MPa and
50 MPa were same, regardless of the presence of an electric field. Heating of 200Cmin−1
led to a decrease in the density.
Elongated structures were seen to have grown during the sintering process. The pres-
ence of pyramidal terrace structures and triangular structures can aid the nucleation of
such micro and nano rods. The final stage of sintering is characterized by an evaporation-
condensation type mechanism. The role of defect structures, especially oxygen vacancies
needs to be investigated using spectroscopic techniques to fully understand the growth
mechanism.
52
Chapter 7
Conclusions
While the conclusions for each of the materials have been summarized independently,
here a comparative approach is undertaken. Based on the results obtained across the
three oxide sintering experiments, the following conclusions can be drawn:
• In all the three cases, the sintering and grain growth behavior agreed with the
classical sintering theory in the conventional case where pressure was applied. It
was only the use of nanopowders, pressureless sintering experiments and the BN
casing that led to unexpected behavior.
• Pressure assisted sintering always leads to enhanced sintering rates. A lowering of
temperature by atleast 200C was observed to achieve the same density, with all
other parameters remaining the same. Dwell time did not lead to any significant
shrinkage in the case of MgO and NiO. Even in the case of ZnO, only at lower
temperatures, time improved the relative density. Hence shorter dwell times, <3min
can be used to minimize the grain growth.
• The grain growth in nanoparticles is different from micron sized particles as seen
in the case of ZnO and MgO. Coalescence was observed be the preferred pathway
rather than Ostwald Ripening. However, the rapid grain growth to the micron
sized particles needs to be avoided. As seen in the case of NiO, thermally activated
attachment can take place during SPS.
• However, such a behavior was not observed in the case of MgO. MgO and NiO have
the same crystal structure, comparable starting particle size, but the morphologies
observed are different in the case of pressureless sintering.
53
• From the results of MgO and NiO, it appears that the role of electric fields in limited.
In the case of NiO, annealing and SPS had shown similar mechanisms. The BN
experiments also revealed that for MgO, similar trends occur with and without the
fields, while for NiO, a shift in the temperatures was observed in the densification
and grain growth plots. This is likely due to thermal gradients rather than the
electric fields. To conclude this a FEM model with the temperature distributions
needs to be developed. A more quantitative understanding can be developed by
measuring defect driven properties like electrical conductivity.
7.1 Recommendations for Future Work
• A wider range of materials, including metal powders can be investigated for the
role of electric fields. However for a systematic comparative study, all the other
parameters such as particles sizes, plastic deformation need to taken out of the
equation.
• As observed in the case of NiO and MgO, the surface energies of various crystal-
lographic planes seems to play an important role in determining the grain growth
mechanism. Hence, a quantitative thermodynamic description of the surface ener-
gies during sintering needs to be undertaken to predict the growth morphologies.
• Many of the sintering processes lead to the formation of atomistic defects. A quali-
tative description of the defect chemistry using spectroscopic tools will help under-
stand the process better.
• As most of the sintering process occurs in a short period, in-situ microscopy will a
useful tool to monitor the sintering process.
54
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