Characterisation and Optimisation of Blowing Agent for Making
Improved Metal Foams
vorgelegt von Diplom-Ingenieurin Metallurgie
Biljana Matijaevi-Lux aus Belgrad, Serbien und Montenegro
von der Fakultt III - Prozesswissenschaften der Technischen Universitt Berlin
zur Erlangung des akademischen Grades
Doktor der Ingenieurwissenschaften - Dr. - Ing. -
genehmigte Dissertation
Promotionsausschuss: Vorsitzender: Prof. Dr. F. Behrendt Berichter: Prof. Dr. J. Banhart Berichter: Prof. Dr. H. Schubert Tag der wissenschaftliche Aussprache: 24. Februar 2006
Berlin 2006
D 83
CONTENTS
i
1 Introduction......................................................................................................................................... 1
2 Metal foams ......................................................................................................................................... 3
2.1 Definition of solid foam................................................................................................................ 3 2.2 Aluminium foams ......................................................................................................................... 4
2.2.1 Factors influencing aluminium foam quality ........................................................................ 6 2.3 Blowing agents (hydrides) .......................................................................................................... 10
2.3.1 Covalent hydrides ............................................................................................................... 10 2.3.2 Saline hydrides.................................................................................................................... 11 2.3.3 Metallic hydrides................................................................................................................. 11
2.3.3.1 Titanium-hydrogen phase system ................................................................................... 12
3 Experimental procedures ................................................................................................................. 15
3.1 Materials...................................................................................................................................... 15
3.1.1 Powder specification ........................................................................................................... 15 3.1.2 Preparation of blowing agent powders................................................................................ 16 3.1.3 Preparation of precursors .................................................................................................... 17 3.1.4 Preparation of metallic foams ............................................................................................. 19
3.2 Characterisation methods............................................................................................................ 20
3.2.1 Powder analysis................................................................................................................... 20 3.2.1.1 Oxygen content analysis ................................................................................................. 20 3.2.1.2 Particle size analysis ....................................................................................................... 20
3.2.2 Microscopy analysis............................................................................................................ 21 3.2.2.1 Light microscopy (LM) analysis..................................................................................... 21 3.2.2.2 Scanning electron microscopy (SEM) analysis .............................................................. 21 3.2.2.3 Transmission electron microscopy (TEM) analysis........................................................ 21
3.2.3 Thermal analysis ................................................................................................................. 22 3.2.3.1 Thermogravimetric (TG) measurements coupled with mass spectrometry (MS) and differential thermal analysis (DTA)................................................................................ 22
3.2.4 X-ray diffraction experiments............................................................................................. 23 3.2.5 In-situ X-ray diffraction experiments ................................................................................. 24 3.2.6 Expandometer experiments................................................................................................. 26
4 Results: Powders ............................................................................................................................... 27
4.1 Powder characterisation .............................................................................................................. 27 4.2 Characterisation of TiH2 powder ................................................................................................ 34
4.2.1 Decomposition behaviour of untreated TiH2 powder ......................................................... 35 4.2.2 Oxidation behaviour of untreated TiH2 powder.................................................................. 38 4.2.3 Isothermal pre-treatment of TiH2 powder........................................................................... 39
4.2.3.1 Thermal treatment under argon....................................................................................... 39 4.2.3.2 Thermal treatment in air.................................................................................................. 41
4.2.4 Decomposition of pre-treated TiH2 powders ..................................................................... 42 4.2.4.1 Decomposition of pre-treated TiH2 powders in argon .................................................... 42
CONTENTS
ii
4.2.4.2 Decomposition of pre-treated TiH2 powders in air .........................................................43 4.2.5 Isothermal experiments at 600C of TiH2 powders (after constant heating at 20 K/min) ..45 4.2.6 X-ray diffraction investigations...........................................................................................47
4.2.6.1 Ex-situ X-ray diffraction of TiH2 powder .......................................................................47 4.2.6.2 In-situ X-ray diffraction of TiH2 powder ........................................................................48
4.2.7 TEM investigations .............................................................................................................52 4.3 Characterisation of ZrH2 and HfH2 powder (as a blowing agent)...............................................55
4.3.1 Decomposition of pre-treated ZrH2 and HfH2 powders ......................................................55 4.3.1.1 Decomposition of pre-treated ZrH2 powders in argon ....................................................55 4.3.1.2 Decomposition of pre-treated ZrH2 powders in air .........................................................56 4.3.1.3 Decomposition behaviour of HfH2 powder.....................................................................57
5 Results: Foams...................................................................................................................................59
5.1 Isothermal experiments at 600C of AlSi6Cu4 precursor with TiH2 powders (after constant heating at 20 K/min)....................................................................................................................59 5.2 TEM investigations of Al precursor ............................................................................................60 5.3 Impact of pre-treatment on foaming behaviour...........................................................................60
5.3.1 Precursors with TiH2 powders.............................................................................................61 5.3.2 Precursors with ZrH2 and HfH2 powders ............................................................................70 5.3.3 Expandometry AlSi6Cu4 precursor with TiH2 powder ......................................................72
6 Discussion ...........................................................................................................................................79
7 Summary ............................................................................................................................................97
8 References ..........................................................................................................................................99
9 Acknowledgement ...........................................................................................................................108
1 Introduction
1
1 Introduction
Aluminium foams are porous metallic structures which combine typical properties of cellular materials
with those of metals [1]. The high stiffness - to-mass ratio and good crash energy dissipation ability has
led to a variety of applications [2] especially in automotive industry [3,4].
The properties of metal foams depend on many morphological features such as pore size distribution, cell
wall curvature, defects etc. [5]. Although the exact interrelationship between properties and structure is
not yet sufficiently known, one usually assumes that a uniform distribution of convex pores free of
defects is highly desirable. Therefore, the task for the experimentalist is to produce such structures. A
short look at existing foams shows that there is still much potential for development since these often tend
to be very irregular [6].
Metals can be foamed in various ways [7,8]. One very promising method is the powder compact foaming
process: a powder mixture of an alloy and a blowing agent is compacted to a dense precursor and then
heated to above the melting point of the alloy [9]. Above the decomposition threshold of the blowing
agent hydrogen is released which blows bubbles in the melting alloy. Titanium hydride (TiH2) was shown
to be a suitable blowing agent for aluminium alloys [10] although other hydrides can also be used [11].
Hydrogen release from TiH2 starts around 400C, which is markedly below the melting point of most
commercial aluminium alloys [9]. This difference between decomposition and melting temperature
causes the formation of irregular, crack-like pores in early expansion stages which then can lead to
irregularities in the final product [12]. To minimize this temperature mismatch the melting point of the
alloy can be lowered by further alloying [13] with possible unwanted side effects or the
decomposition threshold of the blowing agent can be raised by thermal pre-treatment [14]. If TiH2
powder is pre-heated in air an oxide layer is formed on the surface of the particles. This layer delays gas
release from the particles, so that ideally hydrogen is released during foaming only after the melting
temperature of the alloy has been reached [14]. On the other hand, pre-treatments lead to hydrogen losses
lowering foaming power. The task therefore is to find treatments which form a sufficiently thick oxide
layer while a minimum of hydrogen is lost. Some pre-treatments have been suggested in metal foam
literature [1421]. However, as the mechanisms leading to a retardation of gas evolution are not yet
known, these parameters were found in a mostly empirical way.
1 Introduction
2
It is well known that the mechanical properties of metallic foams show a large knockdown compared to
theoretical models of uniform, regular foams, because of the presence of imperfections in the cell
structure. The origin of the imperfections lies in the physics of the foaming process, which is poorly
understood compared to the foaming process of polymer foams. Polymer foams have properties that
follow the theoretical models of perfect foams [22]. Many commercially available types of foams like
aluminium foams are manufactured by the decomposition of TiH2 in the molten (melt-route) or semi-solid
powder metallurgical (PM) route.
Therefore, the aims of the present thesis are:
to investigate in detail the hydrogen release from untreated TiH2 powder;
to investigate the effect of pre-heat treatment of TiH2 on its hydrogen release characteristics;
to characterise gas release from TiH2 in temperature ranges relevant for Al foaming;
to clarify the influence of oxide layers on the structure of powders and gas release;
to clarify the difference between TiH2, ZrH2 and HfH2 powder;
to show how powder treatments influence foaming of various alloys.
Chapter 2 gives a general information about metal foams. An overview of the current state of the art with
regard to the PM production route of the porous metallic materials is given in this chapter as well.
Experimental procedures used in the present work for foaming agent preparation, foam production and
the used experimental methods are described in Chapter 3. The experimental results are presented in
Chapters 4 and 5. The results are discussed in Chapter 6 and summarised in Chapter 7.
2 Metal foams
3
2 Metal foams
2.1 Definition of solid foam
The term "foam" can be explained and defined as shown in Fig. 2.1. Fig. 2.1 lists the designations for all
possible dispersions of one phase in a second one and it shows that foams are uniform dispersions of a
gaseous phase in either a liquid or a solid [23]. The single gas inclusions are separated from each other by
portions of the liquid or solid, respectively. The term "foam" in its original sense is reserved for a
dispersion of gas bubbles in a liquid. The morphology of such foams, however, can be preserved during
the solidification, thus obtaining what is called "solid foam". Speaking about "metallic foams" generally
means solid foam, which is more commonly called a "cellular solid" [9].
Fig. 2.1. Dispersions of one phase into a second one. Each phase can be in one of the three states of
matter [23].
Currently there are two main methods to produce metallic foams, as shown in Table 2.1. The first one is
the direct foaming method where the gas is injected continuously from an external source into a specially
2 Metal foams
4
prepared molten metal containing uniformly dispersed non-metallic particles to which gas bubbles are
added to create foams [24] or where a gas releasing chemical agent is added. The indirect foaming
method start from a solid precursor which consists of metallic matrix containing uniformly dispersed
blowing agent particles [9]. For aluminium based alloys the hydrides of transition elements like titanium
or zirconium hydride powder are mostly used as blowing agent powders [11]. Above melting temperature
this precursor expands and forms a foam. In the present work the indirect foaming method is used.
Table 2.1. Basic foaming routes for metal foams and current manufacturers for aluminium-based foams
and their products [25].
direct foaming melt alloy indirect foaming prepare foamable precursor
make alloy foamable remelt precursor
create gas bubbles create foam
collect foam solidify foam
solidify foam
manufacturers Cymat, Canada (SAF) manufacturers alm, Germany (AFS)
(products) Foamtech, Korea (Lasom) Alulight, Austria (alulight)
Htte Kleinreichenbach (HKB),
Austria (Metcomb)
Gleich-IWE, Germany
Schunk, Germany
Shinko-Wire, Japan (Alporas)
(Distributor: Gleich, Germany)
2.2 Aluminium foams
Over the past 15 years metal foams have become one of the most challenging materials for scientific and
industrial investigations. Many metals can be foamed but foams from aluminium proved to be a serious
candidate for industrial applications especially in automotive industry [3,4].
In 1943 Benjamin Sosnick [26] tried to foam aluminium with the aid of Hg in a closed chamber under
high pressure. With the release of the pressure the vaporisation of the Hg appears at the melting
temperature of aluminium and leads to the formation of a foam. The fundamental ideas for the indirect
foaming process were developed in the late 1950s by Benjamin Allen [27]. In 1990, the method was
rediscovered at the Fraunhofer Laboratory in Bremen, Germany [28] and brought to a considerable level
of sophistication for manufacturing foams and foamed components.
2 Metal foams
5
Using the powder metallurgical route aluminium, zinc [9] and lead [29] foams can be prepared in a cost
effective manner. This process begins with the mixing of a metal powder (pure metal or suitable alloy)
with a blowing agent (TiH2, ZrH2, CaCO3, etc.) [30], after which the mixed powders are compacted e.g.
by hot pressing, extrusion or powder rolling to a dense foamable precursor [9]. In principle, the
compaction can be done by any technique that ensures that the blowing agent is embedded into the metal
matrix without any notable open porosity [31]. Heating of the precursor above the melting temperature of
the metal or alloy is the so-called foaming process. During this process gas evolves from the foaming
agent blowing bubbles in the melting alloy. The released gas forces the semi - solid compacted precursor
material to expand. The result is a product with a highly porous structure called metal foam [1], a material
that offers manufacturers a significant potential for lightweight structures, for energy absorption and for
thermal management [7]. Therefore, it can be concluded that the temperature profile during foaming
consists of a heating stage, isothermal holding stage and a cooling stage as it is shown in Fig. 2.2. The
pores continuously grow during foaming due to the gas supplied by the blowing agent and thermal
expansion. Expansion is controlled by a continuous decomposition of the blowing agent. Therefore,
foaming already starts during heating of the precursor and the gas supply depends on the decomposition
kinetics of the blowing agent. As it is mentioned in Ref. [32] complex relationships exist between the
decomposition kinetics of the blowing agent and the melting of the alloy, which has a beneficial influence
on metal foam quality. It the next chapter the factors influencing aluminium foam quality are explained in
more detail.
2 Metal foams
6
Fig. 2.2. Schematic representation of foam evolution by foaming precursors containing blowing agents
[32].
2.2.1 Factors influencing aluminium foam quality
Metal foams made by first compressing mixtures of aluminium alloy powders and TiH2 and then melting
and foaming the resulting densified compact [9,33] are now being exploited commercially under trade
names such as "Alulight" or "Foaminal" [34]. Production volumes are still low owing to both the still
relatively high costs and to the properties of foamed aluminium alloys which are not always sufficient for
a given application. For this reason, many research activities are directed towards improving the
properties of such foams and to make their manufacture more reliable, reproducible and, as a
consequence, less expensive.
Intuitively one tends to assume that metal foam properties are improved when all the individual cells of a
foam have a similar size and a spherical shape. This, however, has not really been verified
experimentally. There is no doubt that both the density of a metal foam and the properties of the matrix
alloy influence, e.g. modulus and strength of the foam [1,35,36]. A clear influence of cell size distribution
and morphological parameters on foam properties, however, has not yet been established. The reason for
2 Metal foams
7
this is that it has not yet been possible to control these parameters during foaming. Therefore, studies of
the interdependence of morphology and properties of foamed metals have in the past either been purely
theoretical or have been concentrated on more simple systems such as open cell structures or honeycombs
which can be manufactured in a more controlled way. Still, there are various opinions on what the ideal
morphology of a closed - cell foam yielding the highest stiffness or strength at a given weight should be.
Some authors find an insensitivity of strength on cell size distribution [37,38], others claim that a bimodal
distribution of cell sizes is especially favourable [39]. Irregular foams were found to have a higher tangent
modulus at low strains, whereas regular foams with a more unique cell size were stronger at higher strains
[40]. Size effects were found to be an important issue because a certain number of cells across a sample is
necessary to ensure improved mechanical properties (see e.g. Ref. [41]) which makes foams with small
pores look more favourable. Importance has also been ascribed to the shape of cell walls. Especially
corrugated cell walls seem to have a strong detrimental influence on foam properties [42]. Foams with
missing cell walls or cell walls containing holes or cracks were also studied and found to be mechanically
weaker than perfect foams (see Ref. [43] for an overview). The variability of mechanical properties in a
group of specimens of identical nominal dimensions and densities has been analysed statistically and was
found to be very large even for the most regular foams available at present [44]. It is likely that more
regular foams, even if they do not show better mechanical properties, lead to a lower variability of these
properties, which would also be a very valuable feature.
Despite the unclear situation researchers are striving to make more uniform foams, assuming that their
production method is the most valuable one (see e.g. Ref. [45]). As the foaming process comprises
various steps [32] and each step can be influenced by many parameters the properties of the final foamed
part can vary very much if production conditions are changed. Especially in the early days of metal
foaming the search for appropriate foaming parameters was difficult because the mechanisms of metal
foaming were unknown and the complex interdependence between parameters was not understood.
Another problem is that there is no simple measure for the quality of a foam. Usually one assesses the
uniformity of cells in a qualitative way by looking at sections through foamed specimens. Whenever a
foam possesses many cells of a similar size it is often considered "good", if there are many obvious
defects it is called "inferior". Fig. 2.3 shows four samples as taken from different batches of aluminium
foam sandwich (AFS) panels made by the company Applied Lightweight Materials (Saarbrcken,
Germany) which clearly show a difference in foam quality.
2 Metal foams
8
Fig. 2.3. Sections of aluminium foam sandwich materials as obtained in industrial production [46]. From
bottom to top: improving quality (Courtesy: H.-W. Seeliger).
The problem with such qualitative measures is that important morphological features such as the
asphericity of cells or crack alignments in three dimensions might not be detectable this way. Therefore,
tomography and 3D image analysis were proposed [12] and successfully applied [47]. However, for the
purposes of parameter optimisation this method is often too elaborate. Another problem encountered is
the very restricted significance of individual foaming trials. Systematic variations in foam quality
associated with the change of a processing parameter are not easy to detect and require a large number of
experiments. Finally, there also seems to be a size effect in foaming. The smaller the samples the more
uniform the pore structure. The size of the sample makes a comparison of different results difficult.
The current state of knowledge suggests that the following rules have to be obeyed to obtain good
aluminium alloy foams via the powder compact foaming technique:
The melting behaviour of the alloy system and the decomposition characteristics of the blowing agent
have to be coordinated. Ideally, gas evolution is suppressed below the solidus temperature of the alloy to
2 Metal foams
9
avoid formation of cracks before melting. After partial melting the gas released by the blowing agent then
leads to the formation of spherical pores provided that the liquid fraction is sufficiently high. As TiH2 has
a very low decomposition temperature starting at about 400C for untreated hydride and commercial
aluminium alloys have solidus temperatures above 525C there is an obvious gap. It is clear that untreated
TiH2 does not fit within the melting range of any of commercial aluminium alloy. The first foaming
experiments [48] were carried out with pure aluminium or AlCu4 and often yielded inferior results. A
first improvement was the replacement of these wrought alloys by casting alloys such as AlSi7 and
AlSi12. Later the alloy AlSi6Cu4 was proposed and successfully foamed [35]. Therefore, in this work the
detail investigation was carried out on AlSi6Cu4 alloys and how these pre-treated TiH2 powders influence
the foaming process. In the next step the variation of the copper content was applied in order to produce
relatively high fraction of liquid which is necessary for good foaming.
Furthermore, heating conditions during foaming are important. When foaming parts, especially large and
flat foam panels, differences in temperature over the entire extension of the part lead to non-uniform
foaming. This does not only make it difficult to fully expand the entire component without any collapse
taking place in the region of the highest temperature, but also leads to an inferior pore structure due to the
formation of cracks in the foam [49]. The heating profile during foaming is also important. In pre-heated
furnaces the temperature usually exponentially approaches the final temperature. In regulated furnaces
almost arbitrary heating profiles can be realised. It was found that the temperature profile around the
melting point of the precursor significantly influences foam evolution [50]. Overheating can have a
detrimental effect on foaming and should be avoided.
A further factor which influences foam uniformity is the cleanliness of powder processing. The use of
impure powders or the presence of adsorbed water, dirt or gases entrapped in the precursor during
compaction seem to have an adverse effect on foaming, in the sense that these impurities can act as nuclei
for big voids in early stages of gas evolution from the blowing agent. The voids are then thought to grow
to large pores [51]. These effects are rather spurious and have not been quantified but are based on the
experience of a number of researchers including the present author.
Other conditions during powder processing are also important. Overheating of the powder during
extrusion was found to be detrimental to foam expansion [52]. The compaction temperature during hot
pressing has also to be selected carefully since even slightly elevated temperatures can lead to a reduced
foam expansion [53, 54].
2 Metal foams
10
In conclusion, the best recipe for a good foam, i.e. a foam with a high expansion factor developing a
uniform cell size distribution and fairly smooth convex cells seems to be:
use a low - melting aluminium alloy,
use TiH2 as blowing agent which has been adapted to the alloy by pre - treatment,
process powders under clean and reproducible conditions,
take care to avoid overheating during powder compaction,
carefully select the temperature profile during foaming: high heating rates without overheating
and uniform temperature distributions are preferable.
Tailoring the decomposition characteristics of TiH2 powder is the main topic of this work. Therefore, in
the next section some general information about the hydrides is given.
2.3 Blowing agents (hydrides)
The term "hydride" can be described as a binary combination of hydrogen and a metal or metalloid [55].
By the nature of the hydrogen bond, hydrides are classified into three principal categories [56]: covalent
or volatile, saline or ionic and metallic.
2.3.1 Covalent hydrides
Covalent hydrides may be solid (usually polymeric), liquid, or gaseous. They have a very big similarity in
their properties. The bond between hydrogen and the element is of the nonpolar electron sharing type
where valence electrons are shared on a fairly equal basis between the elements held by bond. Molecules
of covalent hydrides are not strongly attracted to each other, and this absence of strong intermolecular
forces results in the high degree of volatility and low melting points of the covalent hydrides. They are
thermally unstable with increasing atomic weight of the parent element. They burn readily in air or in
oxygen with liberation of considerable quantities of heat. Typical covalent hydrides are: aluminium
hydride, tin hydride, boron hydrides and germanium hydrides.
2 Metal foams
11
2.3.2 Saline hydrides
Saline hydrides are formed by the reaction of the strongly electro-positive alkali metals with hydrogen
like MgH2, CaH2 and SrH2. The bond of saline hydrides results from the strong electrostatic forces
existing between the dissimilar electric charges of the two ions. Therefore, they are highly polar. These
hydrides are crystalline, exhibit high heats of formation and high melting points and are electrical
conductors in the molten state.
2.3.3 Metallic hydrides
Metallic hydrides have metallic properties in the accepted sense: these include high thermal conductivity
and electrical resistivity, hardness and in some instances useful mechanical properties. The difference
between metals and hydrides is that they are usually quite brittle. There is a wide range of possible
applications. Hydrides of transition elements were shown to be used as blowing agents for the foaming of
metals [11]. TiH2 powders are also interesting for hydrogen storage applications or as getter material in
vacuum systems [57-60]. Metal hydrides possess properties that make them desirable for nuclear
application, for the preparation of powders and very pure metals, for chemical reducing agents, for
deoxidation and desulfurization of molten ferrous alloys and for use as high energy fuels. Some of the
hydrides can be used as portable sources of hydrogen. It is interesting to mention that zirconium hydride
contains about twice as much hydrogen atoms per unit volume then liquid hydrogen [61]. Metal hydrides
can also be used for preparing a coating on metals. Titanium, thorium and zirconium hydrides are used
effectively for metal-nonmetal bonds.
2 Metal foams
12
0
2.3.3.1 Titanium-hydrogen phase system
A relatively large amount of information is available about the Ti-H system. Despite such a large amount
of information this system is still not well investigated. The phase diagram used for explanation of the Ti-
H system is published by San-Martin and Manchester [62]. Therefore, investigations for this work mostly
used this source for the details presented below.
Fig. 2.4 shows an equilibrium titanium - hydrogen phase diagram. The diagram is one of the eutectoid
types and consists of the following phases: (1) cph ; (2) bcc ; (3) , a fcc hydride; (4) fct , a
tetragonally distorted hydride with axial ratio c / a
2 Metal foams
13
The hydrogen atoms in the hydride at room temperature occupy the tetrahedral interstitial site CaF2 type
structure [66] as shown in Fig. 2.5. There are different opinions within the literature about the occupation
of hydrogen atoms within TiH2 lattice. Some of the authors assume that hydrogen lies on the tetrahedral
interstices in an fcc metal lattice [64]. They claim that it is possible at the same time to have hydrogen
occupancy of the octahedral interstices at room temperature [64] as well.
Fig. 2.5. The tetrahedral lattice sites of H atoms in the TiHx lattice on the room temperature. Ti-red, H-
dark blue.
Some authors preformed several x-ray diffraction measurements [65-67] in order to determine the system
phase boundaries at lower temperatures. There are still some discrepancies on this in the literature, arising
mainly because of unreliable characterisation of the Ti-H system namely in the concentration and the
influence of the impurities.
It is an interesting observation that at the concentration of 66% of hydrogen at a certain temperature the
fcc hydride starts to be transformed to fct hydride. When the fcc hydride starts to be transformed to
fct hydride the temperature was observed around 25C [68].
2 Metal foams
14
Data which illustrates this effect is shown in Fig. 2.6. The transformation is manifested by splitting and
shifting of the diffraction peaks of the original cubic phase as the temperature is lowered.
Fig. 2.6. Cell dimensions of titanium hydride as a function of temperature [68].
This transformation is quite rapid and reversible. By cycling in the temperature range from 20 to 30C
the repeated transformation from cubic to tetragonal structure and vice versa was observed by Crane [68].
Yakel [69] and Takasaki [70] reported a different critical transformation temperature namely at 37C.
Important facts like purity and the hydride preparation methods as reported from [65] and as pointed out
by San-Martin and Manchester [62], that a critical concentration specification is necessary miss a more
specific temperature reference than "room temperature".
3 Experimental procedures
15
3 Experimental procedures
3.1 Materials
3.1.1 Powder specification
TiH2 powder supplied by three different manufacturers was used in this study. As it is shown in Table 3.1
there is a slight variance in the purity of these powders and particle sizes as reported from the
manufacturer. Detail investigations were carried out on TiH2 powder supplied from the company
Chemetall GmbH [71]. ZrH2 and HfH2 powders supplied by Chempur GmbH were used in this study as
well. The powder specifications given by the manufacturer are shown in the Table 3.1. Metallic powders
used for making aluminium alloys are specified in Table 3.2.
Table 3.1. Specification of the used metal hydride powders (specifications by manufacturer).
Powder Producer Purity
[%]
Density
[g cm-3]
Particle size
[m]
Impurities [wt.%] H [wt.%]
TiH2 Chemetall GmbH, Frankfurt 98.8 3.76 < 63 Si
3 Experimental procedures
16
3.1.2 Preparation of blowing agent powders
The powders were characterised in the "untreated" state and after several pre-treatments. TiH2 known as
the hydride with the best foaming potential, was investigated in more detail. ZrH2 powder and HfH2
powder were shown for foaming experiments.
Pre-treatments of the powders were carried out isothermally at various temperatures and times both in an
argon atmosphere (purified by a zirconium getter to 10-12 ppm) inside a tube furnace and under air in a
chamber furnace. In the former case the powder was filled into flat ceramic crucibles (powder height 2
mm) which were then placed into a horizontal quartz tube. After repeated evacuation and filling with
argon a furnace was slid over the tube for the time required. In the latter case the ceramic crucibles were
simply placed into a large volume chamber furnace and were left there for the time specified. After heat
treatment the powders were gently homogenised by tumbling in a container. Two further TiH2 samples
were obtained by using initially untreated samples which had then been subjected to an isothermal heat
treatment for 180 minutes at 520C inside the mass spectrometer both under argon (Fig. 4.11) flow and
air flow (Fig. 4.12) to study the influence of continuous hydrogen removal during treatment. Pre-
treatment parameters for each powder were chosen as shown in Table 3.3. The powders could be treated
without individual particles sticking together (this became a problem at T 540C). It is worth
mentioning that the pre-treatment parameters are applicable in the serial production of the metal foam
precursors.
3 Experimental procedures
17
Table 3.3. Parameters of the pre-treatment of the blowing agent powders (furnace).
Air Argon Powder
T [C] t [min] T [C] t [min]
440 180 440 180
480 180 460 180
500 180 480 180
520 90 520 90
520 180 520 180
TiH2, Chemetall GmbH
520 360 - -
TiH2, SeJong Materials GmbH 480 180 - -
TiH2, GFE GmbH 480 180 - -
400 180 400 180
480 180 480 180
520 90 520 90
ZrH2, Chempur GmbH
520 180 520 180
HfH2, Chempur GmbH 480 180 - -
To obtain bulk samples suitable for TEM investigations the TiH2 powders were cold pressed to small
pellets of 6 mm diameter and 3 mm height prior to pre-treatment in order to avoid deformation after heat
treatments which are known to modify the decomposition characteristics [11]. For the in-situ diffraction
experiments the samples were prepared the same way.
3.1.3 Preparation of precursors
Foamable precursors were prepared by mixing aluminium powder, silicon powder and copper powder in
fractions that led to different commercial aluminium alloys. First pure Al powder was used, then AlSi6,
AlSi6Cu4 and AlSi6Cu10. In these alloys 0.5 wt.% TiH2 was added either in the untreated state or after
one of the various pre-treatments under air. 0.5 wt.% ZrH2 and 0.5 wt.% HfH2 powder were added in
AlSi6Cu4 alloy as shown in Fig. 3.1. The current PM route used for metallic foam production in this
work is shown in Fig. 3.2.
3 Experimental procedures
18
Fig. 3.1. Overview of precursor preparation.
blowing agent
mixing
metal powder(s)
uni-axial pressing
foamableprecursor
foaming
blowing agentblowing agent
mixingmixing
metal powder(s)metal powder(s)
uni-axial pressing
foamableprecursor
foaming
Fig. 3.2. Production steps of PM foams [72].
Homogenization of powder mixture is done for a certain period of time in the appropriate powder
blenders. Depending on the application there is a different kind of blenders for dry mixing of powders, as
shown by Hausner [73]. The powders for metal foam production were mixed for 90 minutes in a tumbling
mixer to produce a homogeneous distribution of all components. The powder mixture prepared for uni-
axial pressing is shaped in special tools on a hydraulic press.
3 Experimental procedures
19
The requirements put to these pressing tools are very important as the pressures applied in pressing are
around 200 MPa and the tool is simultaneously heated to the desired temperature. Pressing die and the
punches must be of hardened steel or hard metal. The pressing tool was made of steel 1.2344. The powder
mixtures were first pre-pressed at room temperature in a cylindrical die of 36 mm diameter, after which
the die containing the powder mixtures was heated up to 450C and held there for temperature
equalisation. Subsequent hot pressing at about 200 MPa for 30 min yielded tablets with more than 99%
density. The density achieved during pressing depends on the compressibility of the powders and the
pressing force. To prevent the sticking of the compressed sample, the pressing tool was spray coated
before each pressing with a molykote spray, which is a molybdenum disulfide aerosol.
Hot pressing is a kind of heat treatment known as sintering in powder metallurgy. It is very important that
during simultaneous heating and pressing the particles are brought close to one another, as the atom
mobility is increased forming a new atomic network. This occurs on the contact particle spots and enables
forming of contact spots on which the powder particles are firmly interconnected. With this the sintering
is started. It continues with further atom exchanges in the interior of a particle. The pressing procedure is
explained in more detail in Refs. [74, 75].
As AlSi6Cu4 with TiH2 powder as a blowing agent is one of the common aluminium alloys for metal
foams production it was chosen for further investigations on the hydrogen release from the precursor. To
monitor hydrogen release from the precursor during foaming small cubes (l 3 mm) of precursor material
were formed. For the foaming experiments thin slices of 36 diameters and 4 mm thickness were cut off
the pressed tablets by electric discharge machining (EDM).
For the TEM observations of TiH2 particles embedded in an aluminium matrix the procedure described
above was applied with the only difference that no Si and Cu was added.
3.1.4 Preparation of metallic foams
The foamable precursor was converted into foam by heating it above its melting temperature. The
different types of precursors used are listed in Fig. 3.1. For foaming a sample was placed on a steel block
which was pre-heated to the foaming temperature inside a furnace and held at the same temperature. A
thermocouple was in contact with the samples at all times so that it was possible to monitor the
temperature inside the foam during the entire process. Each foaming experiment was stopped as soon as a
sample reached a given porosity, and the foaming time tf and final foaming temperature Tf was recorded.
3 Experimental procedures
20
All samples were tried to be foamed to the same final volume corresponding to about 60% porosity. As it
was not possible to have visual control of the foaming experiment several samples of each composition
with slightly varying foaming times were prepared and selected. The ones which had porosity in the
desired range were chosen for further examination. To foam Al precursors the foaming temperature inside
a furnace was 750C. All alloys were foamed at 650C. The preparation technique is explained in more
detail in Refs. [19, 74, 76]. Samples were retrieved after cooling, sectioned (using a microtome or EDM)
and photographed with a Digicam.
3.2 Characterisation methods
3.2.1 Powder analysis
3.2.1.1 Oxygen content analysis
The oxygen content was determined by carrier gas hot extraction (CGHE) [77] using the nitrogen/oxygen
analyser Horiba EMGA 620 WC. The procedure of this method comprises the quantitative reduction of
oxygen in the sample, the release of gaseous reaction products and their selective and quantitative
detection. About 5 mg of powder were first filled into a tin, then into a nickel capsule. This package was
dropped into a graphite crucible which then was electrically heated up to about 2500C for 40 s. The
capsules act as metal additives to obtain a homogenous reduction-assisting melt. At these high
temperatures the oxygen in the sample reacts with carbon forming CO which was detected by infrared
absorption analysis.
3.2.1.2 Particle size analysis
Particle size distributions (PSD) of all powders were measured using a Sympatec Helos Vectra particle
analyser. Powders were dispersed in water assisted by ultrasonic treatment to destroy agglomerates. In the
applied laser diffraction spectrometry technique (LDS), the angular distribution of light is measured after
passing through an optically dilute dispersion of the suspended particles. The Frauenhofer approximation
was the optical model for the analysis of diffraction spectra. So the particle size distribution was then
derived from the patterns of a diffracted laser beam.
3 Experimental procedures
21
3.2.2 Microscopy analysis
3.2.2.1 Light microscopy (LM) analysis
Samples for LM were prepared by using standard metallographic preparation techniques: embedding the
cut samples into embedding resin (Kulzer 4071 or 5071), grinding with a silicon carbide paper 1200-
4000, rough polishing with the 3m diamond grinding medium at MD-Dac (Struers). Final polishing was
done with 1m diamond at MD-Dac (Struers). End polishing was done using OPU-Enpolitur (SiO2) at
MD-Chem (Struers). Microstructure was observed on light microscope Zeiss Axioplan 2, supported with
a software Zeiss Axiovison 4.
3.2.2.2 Scanning electron microscopy (SEM) analysis
The morphology of all powders was characterised by SEM using a Philips XL 20 operated at 10 to 20 kV.
Structure analysis of the powder particle is very important to evaluate and optimise the process of the
powder production which affects the properties of the final product. This technique can visualize the
morphology of the surface, the presence of satellite, impurities and agglomerates as well as the
deformation or fracture of powder particles [78,79]. EDX analysis was used in order to get detail
information about the content of oxygen around individual TiH2 powder particles after the pre-treatment
in air at 480 and 520C for 180 min.
3.2.2.3 Transmission electron microscopy (TEM) analysis
TiH2 powders were characterized by TEM employing a Philips CM30 operated at 300 kV. Cold pressed
pellets were cut into sizes suitable for TEM preparation and were then mechanically polished and
subsequently ion-thinned in an ion beam milling machine. Cold pressing does not lead to a merger of
individual particles so that these could still be separated in the TEM images. For studying TiH2 particles
embedded in an aluminium matrix after compaction, i.e. in the foamable precursor, focused ion beam
(FIB) milling was used to cut a thin foil out of a bulk sample including an interface between the Al matrix
and a TiH2 particle. After this a micromanipulator was used to transfer the foil from the original sample to
a TEM grid coated with a porous carbon support film.
3 Experimental procedures
22
3.2.3 Thermal analysis
3.2.3.1 Thermogravimetric (TG) measurements coupled with mass spectrometry (MS) and differential thermal analysis (DTA)
To study the decomposition of blowing agent powders as a function of temperature or time a Netzsch
simultaneous thermal analyser STA 409C was used which allowed simultaneous differential thermal
analysis (DTA), thermo-gravimetric analysis (TGA) and mass spectrometry (MS) of both hydrogen and
water. The analyser was equipped with a silicon carbide furnace connected to a quadrupole mass
spectrometer via a skimmer coupling system. The distance between skimmer and sample holder
amounted to 12 mm. In all the measurements a constant amount of 150 mg of either the untreated or pre-
treated blowing agent powders or the foamable precursor were filled into the alumina crucible which was
placed on top of a ceramic sample carrier, as shown in Fig. 3.3. The first pressure reduction step located
in front of the skimmer and between the sample holder, realized by a ceramic shield with an orifice,
reduces the pressure to 10-1
mbar. The skimmer orifice further reduces the pressure down to 10-5 mbar
where the gases arrive at the analyzer ion source at the head of the quadrupole mass spectrometer,
positioned in cross beam geometry. A vertical gas flow of an inert gas (argon) through the furnace core
and the short distance between sample and skimmer guarantee the transport of even heavy molecules to
the analyzer ion source. To obtain an optimum sensitivity the mass spectrometer was operated with a
secondary electron multiplier.
After the positioning of the crucible the thermo-balance was evacuated three times and back-filled with
argon prior to heating the sample under a continuous flow of argon (flow rate 150 cm3 min-1).
3 Experimental procedures
23
Fig. 3.3. Simultaneous thermal analyser - Mass spectrometry with skimmer coupling system (Courtesy:
Netzsch Company).
3.2.4 X-ray diffraction experiments
A Bruker - AXS Advance diffractometer model D8 equipped with a graphite monochromator and a
scintillation counter was used to determine the phases present in the TiH2 powders before and after pre-
treatments in air at 520C. The diffraction patterns were measured for 2 ranging from 15 to 90 at steps
of 0.02. An acquisition time of 5 s per step using CuK radiation was sufficient for a successful phase
analysis. The XRD patterns were evaluated using evaluation program (EVA) from Bruker - AXS and
compared to a JCPC standard diffraction database.
3 Experimental procedures
24
3.2.5 In-situ X-ray diffraction experiments
The in-situ diffraction experiment was performed at the KMC2 beamline at the BESSY synchrotron light
source using a photon energy of 8 keV. The use of an area sensitive detector (HiStar, Bruker AXS) to
simultaneously acquire the photons scattered in 2 angle between 30 and 60 degrees combined with the
intense synchrotron beam allowed to acquire a complete scattering pattern every 10 seconds.
(a)
(b)
3 Experimental procedures
25
(c)
Fig. 3.4. Experimental set up for the in-situ diffraction measurements; a) sample holder, b) high
temperature chamber, c) diffraction geometry.
The sample was mounted in a high temperature chamber on a Huber goniometer as shown in Fig. 3.4. The
chamber was equipped with a beryllium dome, which is transparent for the photon energy used. This
allows heating the sample up to 1000C in vacuum or in controlled atmosphere and simultaneously to
perform scattering experiments in a wide range of incident and diffracted beam directions. For current
experiment the chamber was flushed with He to accelerate the H2 emission from the sample.
Since the area detector can span only 30 degrees in 2 (Fig.3.4c), the range between 30 and 60 degrees,
was chosen. In this range all major peaks of all expected phases appear. Samples were heated from room
temperature up to 780C with a heating rate of 10 K/min. Starting from 200C an XRD pattern was
recorded every 10 sec. Temperature was controlled using a Eurotherm process controller and a
thermocouple was placed in the vicinity of the sample. It was not possible to mount the thermocouple
directly on the sample because of a possible interference with the scattering experiment.
3 Experimental procedures
26
3.2.6 Expandometer experiments
The precursor tablets of AlSi6Cu4 prepared with untreated and differently pre-treated TiH2 powders were
foamed in a so-called "expandometer". This dilatometer was specially constructed for metal foam
characterisation and allows to heat up a piece of precursor in a quartz glass tube using lamps with a given
heating profile while recording sample temperature and vertical expansion. Expansion can be interrupted
by turning off the lamps and starting to cool with pressurised air [80]. The used experimental procedure is
explained in more detail in Refs. [54, 81]. The cooling rates achieved are about 2 K s-1.
Two types of experiments were carried out:
1. Experiments in which the precursor sample was foamed without interruption under the influence
of a given temperature profile, and
2. Experiments in which the foaming process was interrupted to reach a pre-selected expansion.
In the former case the temperature profile chosen included a constant heating rate of 1 K s-1 up to 650C
and a constant temperature with a given dwell time. In the latter case the final temperature was kept
slightly lower at 630C to slow down the foaming process. Samples were retrieved after cooling,
sectioned (using a microtome or EDM), photographed with a Digicam and further investigated by SEM.
In order to check the reproducibility of the expandometer measurements, about five experiments were
carried out for each set of parameters. As the resulting curves showed some variation a measurement was
accepted whenever three curves were identical within an error limit given by the precision required for
the analysis.
4 Results: Powders
27
4 Results: Powders
4.1 Powder characterisation
In Fig. 4.1 the morphology of both untreated TiH2 powder and TiH2 pre-treated under air is made visible
by SEM. The powders are angular shaped, the particle size is variable and the powder surface is irregular.
Obviously, pre-treatment in air does not significantly influence the morphology of the powders. Fig. 4.2
shows the EDX analysis of pre-treated TiH2 in air at 480 and 520C for 180 min obtained with an energy
of 10 kV. As oxygen is a light element it is difficult to measure quantitatively. Qualitatively, EDX
analysis shows that the pre-treatment in air at 520C for 180 min produces a higher oxygen content on the
surface of the powder particle.
(a)
4 Results: Powders
28
(b)
Fig. 4.1. SEM images of TiH2 powder particles; a) untreated, b) pre-treated in air at 480C for 180 min.
(a)
4 Results: Powders
29
(b)
Fig. 4.2. EDX analysis of TiH2 powder particles pre-treated in air; a) 480C for 180 min, b) 520C for
180 min.
The oxygen content of pre-treated TiH2 powders under conditions given in Table 3.3 was determined by
hot gas extraction as shown in Fig. 4.3. Untreated TiH2 powder contains about 1 wt.% oxygen. Treatment
in air leads to a continuous increase up to 7 wt.% for a treatment at 520C for 360 minutes. Treatment
under argon leads to a slight decrease of the oxygen content down to 0.65 wt.%. Changes in the colour of
the TiH2 powder were observed during heat treatment in air. The initially dark grey powder turned blue at
480C, green at 500C and light brownish to dark brown at 520C, depending on annealing time, as
shown in Fig. 4.4. Apparently the heat treatment in air does not change the morphology of the powder but
due to the light interference the optical properties changed as shown in Fig. 4.4.
4 Results: Powders
30
0 100 200 300 400 500
0
1
2
3
4
5
6
7
8ox
ygen
con
tent
[wt.%
]
annealing temperature [C]
Ar 180 min 90 min 360 min
450 475 500 5250,6
0,7
0,8
0,9
1,0Air
180 min 90 min 360 min
Fig. 4.3. Oxygen content obtained by CGHE of untreated TiH2 powder and pre-treated at various
temperatures and times under air and argon.
a) b) c)
4 Results: Powders
31
d) e) f)
Fig. 4.4. TiH2 oxidised at different conditions. The colour of the powder depends on the thickness of the
oxygen layer a) unterated b) 480C for 180 min c) 500C for 180 min d) 520C for 90 min e) 520C for
180 min f) 520C for 360 min.
Fig. 4.5 shows the morphologies of untreated ZrH2 and HfH2 powder obtained by SEM. Evidently the
powder morphologies are comparable with TiH2 powder, as well. They are angular shaped and the
particle size is variable. The particle size distribution and oxygen content analysis of used metallic
hydrides are given in Table 4.1. Analysis of the various metal powders used for metal foam precursor
preparation yielded to the values as shown in Table 4.1. The morphologies of the used powders as
obtained by SEM are shown in Fig. 4.6.
(a)
4 Results: Powders
32
(b) Fig. 4.5. SEM images of untreated powder particles; a) ZrH2, b) HfH2.
Table 4.1. Physical and chemical properties of the ZrH2, HfH2, TiH2, Al, Cu and Si powders used.
Particle size distribution [m] Material
Oxygen analysis [wt.%]
D(v.90) D(v.50) D(v.10)
ZrH2 Chempur GmbH 0.58 21.14 8.66 2.62
HfH2 Chempur GmbH 0.19 70.91 21.42 2.39
TiH2 Chemetall GmbH 1.00 30.79 14.14 3.18
TiH2 SeJong Material 0.51 25.32 13.46 2.82
Blowing
agents
TiH2 GFE GmbH 0.88 23.90 9.01 1.92
Aluminium Eckart 0.57 150.00 66.50 19.70
Copper Chempur GmbH 0.46 121.00 84.00 51.00
Metal
powders Silicon lschlger 0.53 76.10 27.80 4.70
4 Results: Powders
33
(a)
(b)
(c)
Fig. 4.6. SEM images of metal powders with obtained particle shapes a) aluminium, irregular, b) copper,
spherical c) silicon, angular.
4 Results: Powders
34
4.2 Characterisation of TiH2 powder
The onset and peak temperatures of mass spectrometric (MS) or thermoanalytic data are specified in this
and in the next chapter. The peak is defined as the location of the maximum value of the (MS) curve
under consideration, while the onset is defined as the point where the MS curve has risen from the
baseline to 1% of the peak value. Only in Fig. 4.16 this criterion had to be changed to 2.5%, while the
sharp edge in Fig. 5.1 allowed to use this point to determine the onset point. These definitions are adapted
to the quality of the data available. The onset temperatures obtained this way are a bit arbitrary and do not
reflect any physical transition in the material but they allow to accurately identify changes due to pre-
treatments.
Most measurements were repeated two or three times to be able to assess statistical errors. An excellent
reproducibility was found and therefore only one curve is shown with very small error bars (look Fig.
6.7). In Fig. 6.7 most of the results are summarized and error bars are given to demonstrate accuracy.
As TiH2 leads to the best foaming characteristics this powder was investigated in more detail. Therefore,
an extensive thermoanalytical programme is carried out. The experiments are organised as outlined in
Fig. 4.7. In Chapter 4.2.1 and 4.2.2 untreated TiH2 powder is characterised by thermal analysis in air and
in an argon atmosphere to obtain a first overview of the concurring processes "hydrogen release" and
"oxidation". Isothermal experiments at 520C which is a typical pre-treatment temperature are carried
out in Chapter 4.2.3. In a second step in Chapter 4.2.4 characterisation of TiH2 powders pre-treated in air
or argon prior to thermal analysis are shown. Again in Chapter 4.2.5 isothermal experiments of untreated
and pre-treated TiH2 powders in air are performed this time at 600C which is a typical foaming
temperature for aluminium alloys. Finally, in Chapter 5.1 hydrogen release from the foamable aluminium
precursor is characterised under the same isothermal conditions as shown in Chapter 4.2.5.
The phases present in TiH2 before and after pre-treatment in air are determined by XRD. The oxide layers
around individual TiH2 particles are investigated with TEM. Finally, in Chapter 5.3 it is shown that pre-
treatment actually improves the structure of aluminium foams.
4 Results: Powders
35
Fig. 4.7. Overview of all the thermo-analytical experiments performed.
4.2.1 Decomposition behaviour of untreated TiH2 powder
Fig. 4.8a shows the release of hydrogen and the decomposition behaviour of untreated TiH2 powder as a
function of time and temperature. An argon gas flow and a heating rate of 10 K/min were applied. The
temperature profile was linear throughout the entire range so that T(t) is not shown explicitly but can be
read from the upper axis. As no oxygen is present the decrease in mass has to be ascribed to the release of
hydrogen. The mass curve (TG) shows distinct changes in slope which become visible when showing the
derivative (DTG). Two peaks and one shoulder can be observed in the DTG curve as well as in the mass
spectroscopic (MS) curve. The first step of hydrogen release as identified by mass spectrometry applying
the 1% criterion starts at 415C (marked as "1"), the peak position is located at 524C. A second
dehydrogenation step starts at about 540C after deconvolution (Fig.4.8b) and reaches the maximum at
626C. The shoulder around 800C indicated a third step which, however, is not further investigated since
such high temperatures are never reached during foaming of aluminium. Mass spectrometry of hydrogen
(MS) and the derivative of the mass (-DTG) agree well even though they are slightly shifted against each
other. The DSC tends to a very similar result. Because of the higher sensitivity of the DSC measurement
4 Results: Powders
36
the starting point of decomposition becomes even visible at a lower temperature of around 390C ("2").
The negative maximum of the DSC curve indicates an endothermal reaction. Mass reduction continues up
to temperatures of around 920C ("3"). The total mass loss up to this temperature is 3.78%, which is close
to the calculated value of 4% (TiH2Ti corresponds to 5048 = -4%). The areas of the first and the
other peaks are consistently related as 1:3.21, 1:3.22 and 1:3.01 for the MS, DTG and DSC curve,
respectively. Hence, about 25% of the hydrogen gas is released in the first reaction. Some water is
detected during heating of the powder over a wide temperature range, the water ion signal is about 30
times smaller than that of hydrogen. The MS curve of hydrogen in Fig. 4.8b is shown after deconvolution
and shows three Gaussian fits of hydrogen release.
0 20 40 60 80 100 1200.0
0.2
0.4
0.6
0.8
1.0
1.2
97
98
99
100
101200 400 600 800 1000 1200
0.00
0.05
0.10
0.15
0.20
0.25
13
TG
MS hydrogen
- DTG
temperature [C]
MS
ion
curr
ent [
10-9A]
time [min]
MS (H2)
TG
DTG
[%/m
in]
TG [%
]
200 400 600 800-2,0
-1,5
-1,0
-0,5
0,0
2DS
C [
W/m
g]
DSC
temperature [C]
DTG
(a)
4 Results: Powders
37
200 400 600 800 1000 12000.0
0.2
0.4
0.6
0.8
1.0
1.2 MS (H2) Gauss fit MS (H2) Gauss fit peak 1 Gauss fit peak 2 Gauss fit peak 3
MS
(H2)
ion
curr
ent [
10-9A
]
temperature [C]
(b)
Fig. 4.8. Analysis of untreated TiH2 powder heated from 30C to 1200C at 10 K/min under argon. (a)
Calorimetric (DSC), mass (TG), mass change (-DTG=-d/dt(TG)) and mass spectrometric (MS) signals for
H2 are shown. The numbers in circles mark specific temperatures discussed in the text. At (b) MS (H2)
curve after deconvolution and respective Gaussian fits are shown.
Fig. 4.9 shows the decomposition behaviour of untreated TiH2 powders produced from various
manufacturers. Powder specifications from the manufacturer, pre-treatment parameters and powder
characterisation are given in Tables 3.1, 3.3 and 4.1, respectively. Applying the 1% criterion, it can be
seen that hydrogen first starts to be released from GFE powder at 360C. Powder from SeJong Materials
releases hydrogen at 365C applying the same criterion. 415C is the onset temperature of gas release
from TiH2 powder supplied by Chemetall. Therefore, the variance in the onset temperatures is 55 K. Gas
release ends at 900C for GFE powder and 950C for Sejong Materials as well as Chemetall TiH2
powder. The peak maximum shifts to +18 K for SeJong Material powder and to -17 K for GFE powder
compared to TiH2 powder from Chemetall.
4 Results: Powders
38
400 600 800 10000.0
0.2
0.4
0.6
0.8
1.0
1.2peak positions: 609
644
626
3
2
1
Chemetall SeJong
MS
ion
curr
ent [
10-9A
]
temperature [C]
GFE
ig. 4.9. Mass spectrometric analysis of untreated TiH2 powders supplied from different manufacturer
4.2.2 Oxidation behaviour of untreated TiH2 powder
ig. 4.10 shows the behaviour of untreated TiH2 powder heated at a rate of 10 K/min in air. Oxidation
F
during heating from 30C to 1200C at 10 K/min in an argon atmosphere (low temperature range
omitted). Curve for untreated powder from Chemetall is already given in Fig. 4.8 and is shown for
comparison. Peak positions are specified in the diagram (in C). The numbers in the circles mark specific
temperatures discussed in the text.
F
starts at 250C ("1") and leads to a continuous mass increase. Due to strong exothermal effects the
temperature curve departs from the linearity at 600C and leads to a peak at 680C ("2") associated with a
sudden increase of the slope of the TG curve. Both, hydrogen and water evolution was observed.
Hydrogen is released above 380C ("3"), while water evolution sets in at 510C ("4"). The total mass gain
from room temperature to T=1200C is 59.7%, which fits very well to the calculated increase of 60%
(TiH2TiO2 corresponds to 5080 = +60%).
4 Results: Powders
39
0 20 40 60 80 100
200
400
600
800
1000
1200
4
3
2
1MS hydrogen
MS water
tempe
rature
MS
ion
curr
ent [
10-9A
]
TG [%
]
time [min]
temperature
tem
pera
ture
[C
]
time [min]
0 20 40 60 80 100
100
110
120
130
140
150
160
TG
0.0
0.5
1.0
1.5
2.0
2.5
MS (H2) MS (H2O)
ig. 4.10. Analysis of untreated TiH2 powder heated from 30C to 1200C at 10 K/min in air. Mass (TG)
4.2.3 Isothermal pre-treatment of TiH2 powder
4.2.3.1 Thermal treatment under argon
ig. 4.11 shows experiments on untreated TiH2 performed under argon applying a heating rate of 10
F
and mass spectrometric (MS) signals for H2 and H2O are shown. The numbers in circles mark specific
temperatures discussed in the text.
F
K/min from room temperature up to 520C followed by an isothermal treatment at this temperature for
180 min. Different regimes of reaction can clearly be recognized in the MS, TG, -DTG and DTA curves.
The first regime ("A") is the temperature range from 400 to 520C (40 to 60 min) and is associated with a
loss of weight and strong evolution of hydrogen. The maximum of gas release is observed after 50
4 Results: Powders
40
minutes at 500C shortly before isothermal conditions have been reached. The negative maximum of the
DTA curve indicates an endothermal reaction. The second regime ("B") of hydrogen release becomes
dominant under isothermal conditions in the time range from 60-100 min and constant temperature. The
MS, TG, DTG and DTA curves notably change their slopes and develop a shoulder. In the third regime
("C") the release of hydrogen levels off quickly and tends to zero. Obviously, the TG, MS and DTA
curves are coupled. The transition from the second to the third regime can best be detected from the
changing slope of the MS-curve. The shoulder of the MS curve matches with the shoulder of the DTA
curve (see inset in Fig. 4.11).
50 100 150 200
97.5
98.0
98.5
99.0
99.5
100.0
100.550 100 150 200
100
200
300
400
500
600
-0.20
-0.15
-0.10
-0.05
0.00
0.05
0.00
0.05
0.10
0.15
0.20
0.25
0.30
0.35A B C
MS
ion
curr
ent [
10-9
A]
tem
pera
ture
[C
]
DTG
[%/m
in]
TG [%
]
time [min]
time [min]
10 K
/min
MS hydrogen TG
- DTGtemperature
50 100 150 200
-0.4
-0.3
-0.2
-0.1
0.0
t
TDTA
CBA
DTAT
100200300400500600
ig. 4.11. Analysis of untreated TiH2 powder heated up to 520C and held there under Ar flow. F
Calorimetric (DTA), mass (TG), mass change (DTG) and mass spectrometric signals (MS) for hydrogen
are given. Letters (numbers) in circles denote reaction ranges (temperatures) as discussed in the text.
4 Results: Powders
41
4.2.3.2 Thermal treatment in air
Fig. 4.12 shows the results of an analogous isothermal experiment under air. The increasing TG curve
indicates significant oxidation above 250C (TG curve) ("1") after a small initial mass change and reveals
that the highest oxidation rate is found in the range between 460 and 520C. Hydrogen release as given
by the MS curve gets notable at approximately 350C ("2") and reaches a maximum at 507C after 49
min. This is slightly earlier than in Fig. 4.10 reflecting inaccuracies in the starting conditions of each
experiment. After some time at 520C hydrogen evolution rapidly drops and reaches a low level after
about 100 min ("3"). Associated with H2 evolution, the formation of water was observed from the
corresponding mass spectrometric signal. If the isothermal holding temperature is 480C instead of 520C
oxidation proceeds much slower which can be seen from the corresponding TG curve included in Fig.
4.12.
0 50 100 150 200100
101
102
103
104
105
106
time [min]
32
1
TG (480C)
calc. TG
TG
MS water
MS hydrogen
10 K
/min
temperature
MS
ion
curr
ent [
10-9A
]
tem
pera
ture
[C
]
TG [%
]
time [min]
0 50 100 150 200
100
200
300
400
500
600
0.00
0.05
0.10
0.15
0.20
Fig. 4.12. Analysis of untreated TiH2 powder heated up to 520C and held there under air. Mass (TG) and
mass spectrometric signals (MS) for hydrogen and water are given. Letters (numbers) in circles denote
reaction ranges (temperatures) as discussed in the text. The dash-dotted line is a TG curve obtained at
480C, dotted lines are fit curves calculated using Eqs. 2 and 3 (p. 85).
4 Results: Powders
42
4.2.4 Decomposition of pre-treated TiH2 powders
The following experiments are based on powders pre-treated in air or argon prior to characterisation. At
this stage no further oxidation occurs since thermal analysis is carried out under argon. TiH2 powder
solely releases hydrogen which is also the case when metals are foamed.
4.2.4.1 Decomposition of pre-treated TiH2 powders in argon
Fig. 4.13 shows the decomposition behaviour of TiH2 powders previously pre-treated under argon.
Starting from the curve for untreated powder which was already described in Fig. 4.8, it can be seen that
heat treatment at 440C for 180 min eliminates the first decomposition stage and leaves only a single
stage. The single peak structure remains the same even for treatments at higher temperatures or for longer
durations. The position of the remaining peak is slightly shifted to lower temperatures compared to that of
the second peak of untreated TiH2. The maximum peak shift is -12 K. Another observation is that the
onset of gas evolution at 415C is not changed significantly by pre-treatments under argon, with all onset
temperatures being between 413C and 415C. Gas release levels off for temperatures beyond the peak
and ends between 950C for untreated powders or powders pre-treated at 440 and 460C, and 1100C
for powders pre-treated at 480 and 520C.
The decomposition curve for the powder pre-treated under the conditions described by Fig. 4.11 (curve
marked with an asterisk in Fig. 4.13) looks very different to the corresponding curve of the powder pre-
treated in a closed argon-filled container for almost the same time and temperature. While onset and peak
position are found at the same temperature, much less hydrogen is detected.
4 Results: Powders
43
400 600 800 10000.0
0.2
0.4
0.6
0.8
1.0
1.2
614*
untreated
pre-treated in argon at:
440C,180 min 460C,180 min 480C,180 min 520C,90 min 520C,180 min
*pre-treated in argon (in TG): 520C,180 min
peak positions:
614
621623
624
625
626
524
MS
ion
curr
ent [
10-9A]
temperature [C]
Fig. 4.13. Mass spectrometric analysis of pre-treated TiH2 powders in Ar during heating from 30C to
1200C at 10 K/min in an argon atmosphere (low temperature range omitted). Curve for untreated powder
as already given in Fig. 4.8, shown for comparison. Peak positions are specified in the diagram (in C).
4.2.4.2 Decomposition of pre-treated TiH2 powders in air
The powders pre-treated in air prior to thermal analysis show a different decomposition behaviour as
demonstrated in Fig. 4.14. As it was already observed for pre-treatment under argon, heat treatment for
180 min at 440C in air eliminates the first decomposition stage. Unlike treatment under argon, however,
treatment in air significantly changes the temperature of both the onset of gas evolution and the
temperature at which maximum gas expansion occurs. Treatments up to 360 min or treatments at higher
temperatures up to 520C shift the onset up to +170 K and peak positions up to +52 K. For all treatments
gas release is smaller than that for untreated powders at T750C and ceases earlier, namely at 900C
as compared to 950C for untreated powder.
4 Results: Powders
44
400 600 800 10000.0
0.2
0.4
0.6
0.8
1.0
1.2
677*
peak positions:
untreated
pre-treated in air: 440C, 180 min 480C, 180 min 500C, 180 min 520C, 90 min 520C, 180 min 520C, 360 min
*pre-treated in air (in TG): 520C, 180 min
657643
677678
673663
626
524
MS
ion
curr
ent [
10-9
A]
temperature [C]
Fig. 4.14. Mass spectrometric analysis of pre-treated TiH2 powders in air during heating from 30C to
1200C at 10 K/min in an argon atmosphere (low temperature range omitted). Curve for untreated powder
as already given in Fig. 4.8, is shown for comparison. Peak positions are specified in the diagram (in C).
The decomposition curve for the powder exposed to the temperature profile given in Fig. 4.12 prior to
characterisation (180 minutes at 520C after a heating ramp, curve marked with an asterisk in Fig. 4.14)
shows the same onset and peak temperature but less hydrogen. Compared to the analogous effect under
argon the reduction in hydrogen evolution is much smaller.
Mass spectrometric curves of hydrogen release from pre-treated TiH2 powder in air supplied from various
manufacturers are shown in Fig. 4.15. It can be seen that in all three cases the same pre-treatment
parameter (480C for 180 min) as shown in Table 3.3 was used. The maximum peak shift is -9 K and -14
K for GFE and SeJong Materials, respectively in relation to TiH2 powder supplied from Chemetall
Company. Applying the criterion of 1% the onset temperature of H2 release is 495C, 510C and 528C
for the GFE, SeJong Materials and Chemetall producer, respectively. It shows that the maximum
difference in the onset temperature is 33 K after pre-treatment in air under the same conditions.
4 Results: Powders
45
400 600 800 10000.0
0.2
0.4
0.6
0.8
1.0
1.2
Chemetall
MS
ion
curr
ent [
10-9A]
temperature [C]
peak positions:
648
643657
GFE SeJong
Fig. 4.15. Mass spectrometric analysis of TiH2 powders pre-treated in air at 480C for 180 min supplied
from different manufacture during heating from 30C to 1200C at 10 K/min in an argon atmosphere (low
temperature range omitted). Curve of TiH2 powder produced from Chemetall is already given in Fig. 4.14
and is shown for comparison. Peak positions are specified in the diagram (in C).
4.2.5 Isothermal experiments at 600C of TiH2 powders (after constant heating at 20 K/min)
Isothermal experiments at 600C with a heating rate increased to 20 K/min intended to simulate the
temperature profile which occurs during foaming of common aluminium alloys. First experiments were
carried out with TiH2 powder both in the untreated state and after heat treatments in air at various
temperatures (440, 460, 480, 520C) and times (90 min, 180 min). Fig. 4.16 shows a first release of
hydrogen gas from the untreated powder after 21 min at 420C ("1") applying the 2.5% criterion (see
Chapter 4.2). A gas release peak appears after 34 minutes. At this time the maximum temperature has
been reached.
4 Results: Powders
46
The curve resembles the one shown in Fig. 4.12 although there the heating rate and isothermal
temperature were lower. Employing oxidised powders, the onset of hydrogen evolution is shifted to
between 510C ("2") and 595C ("3") depending on the pre-treatment in analogy to the constant heating
rate experiments in Fig. 4.14. The maximum of gas release is also shifted to later times for pre-treated
powders.
0 10 20 30 40 50 60 70 130 140 1500.0
0.2
0.4
0.6
0.8
1.0
0 10 20 30 40 50 60 70 130 140 150time [min]
peak positions:
tem
pera
ture
[C
]
20 K
/min
temperature
43
41
3836
3534
untreated 440C, 180 min 460C, 180 min 480C, 180 min 520C, 90 min 520C, 180 minM
S io
n cu
rrent
[10-
9 A]
time [min]
100
200
300
400
500
6003
2
1
Fig. 4.16. Mass spectrometric analysis of loose TiH2 powders pre-treated under air (analysis is carried out
under Ar). Isothermal conditions at 600C apply after heating at 20 K/min. Numbers denote times at
which H2 evolution peaks (minutes) occur.
4 Results: Powders
47
4.2.6 X-ray diffraction investigations
4.2.6.1 Ex-situ X-ray diffraction of TiH2 powder
XRD patterns for TiH2 powder pre-treated in air at 520C for 90, 180 and 360 min are shown in Fig. 4.17
and are compared with those for the untreated powder. The unmarked peaks correspond to TiH2 or sub-
stoichiometric TiHx compounds. The untreated powder was found to be single-phased. Its composition
was identified as TiH1.924 (3.89 wt.% hydrogen) with a fcc crystal structure (a=0.4448 nm) using powder
diffraction files. No signs of a tetragonal phase were found. After annealing in air the peaks
corresponding to TiHx are shifted to higher angles (see Chapter 3.2.4) corresponding to a shift in lattice
constant to 0.4392, 0.4385 and 0.4378 nm for 90, 180 and 360 min at 520C, respectively. In addition,
peaks corresponding to Ti3O and TiO2 (rutile) are detected, corresponding to a hexagonal and a tetragonal
crystal structure, respectively.
d )
c )
b )
a )
O
O
OO
O
OO
O
O
O
O
O
O
XXXX
XX
XX
X
X
XX
X
X
2 th e ta [d e g re e s ]
O
2 0 3 0 4 0 5 0 6 0 7 0 8 0 9 0
O
XX
X X
Fig. 4.17. XRD patterns for TiH2 powder; a) untreated, b-d) pre-treated in air at 520C for 90, 180 and
360 min; x: TiO2, o: Ti3O.
4 Results: Powders
48
4.2.6.2 In-situ X-ray diffraction of TiH2 powder
Fig. 4.18 shows selected XRD patterns during an in-situ experiment at the current temperatures during
heating of untreated TiH2 powder. There are three different kinds of diffraction peaks which occur during
the measurement. The present phases are marked on the diagram. As it is shown in Fig. 4.18, at room
temperature the -TiH2 phase is present. With the heating to 548C a high temperature -Ti phase is
present and some presence of -Ti is observed, as well. At the temperature of 740C the presence of -Ti
is detected.
30 35 40 45 50 55
Ti Ti
Ti
Ti
Ti
Ti Ti
TiHx
TiHx
Ti
T=740C
T=548C
T=200C
inte
nsity
[a.u
.]
2 theta
Fig. 4.18. Scattering curves corresponding to the phases present during heating of untreated TiH2 powder
Fig. 4.19 shows the phase change during heating of untreated TiH2 powder. Three phases are recognised
in this diagram. -TiH2 phase is present from room temperature up to 540C, -Ti phase increase at
520C and after 550C it slowly decreases. Already at around 520C -Ti appears and the presence of this
phase increases until the end of the experiment. Around 420C the lattice constant of -TiH2 phase
4 Results: Powders
49
decreases. Obviously, this change is caused by the emission of hydrogen due to the decreasing hydrogen
amount in the sample.
200 300 400 500 600 7000
20
40
60
80
100
Ti Ti TiH2ph
ase
com
posi
tion
[vol
. %]
temperature [C]
Fig. 4.19. Changes of phases during heating of untreated TiH2 powder from 200C to 780C at 10 K/min
in He atmosphere.
Fig. 4.20 shows the scattered intensities during in-situ heating of untreated TiH2. Three phases are
recognised in this diagram: -TiH2 from room temperature to 540C, -Ti phase appears at around 520C,
and -Ti phase at around 520C as well. From 750C the whole sample transforms to -Ti phase. Before
first hydrogen emission takes place the peak sifting to lower angles is observed. This happens due to the
thermal expansion of the sample during the heating from room temperature. Important detail is the bend
(marked with an arrow on Fig. 4.20) to larger scattering angles before sample reaches the first transition
temperature. This bend denotes a large emission of hydrogen. It is assumed that the change in the lattice
constant in pure -phase is proportional to the emission of hydrogen.
4 Results: Powders
50
Fig. 4.20. Scattered intensities are plotted in grey scale. Dark regions are high, and light regions are low
intensities of untreated TiH2 powder during heating from 200C to 780C at 10 K/min.
Fig. 4.21 shows the change of the lattice constant of -TiH2 with temperature. The thermal expansion (red
curve) can be estimated from the slope below 400C. Since it is known that at 520C the phase boundary
is reached it is possible to calculate the hydrogen concentration in the sample.
4 Results: Powders
51
200 300 400 500 600 7000.980
0.985
0.990
0.995
1.000
1.005
1.010
1.015
1.020a/
a 300
temperature ]C]
Fig. 4.21. Change of lattice constant (a) relative to the lattice constant at T=300C (a300) of untreated TiH2
powder during heating from 200C to 780C at 10 K/min.
Fig. 4.22 shows the phase change during heating of TiH2 powder which was before pre-treated in air at
480C for 180 min. There is an obvious similarity to the experiments shown in Fig. 4.19. Three phases
are recognised in this diagram as well. -TiH2 phase is present from room temperature up to 550C. -Ti
phase increases suddenly at 550C and after 565C it slowly decreases. At room temperature the presence
of -Ti is found. At approximately 520C there is a sudden increase in the amount of -Ti and it is
increasing until 100% and holds it to the end of the experiment.
4 Results: Powders
52
200 300 400 500 600 7000
20
40
60
80
100
Ti Ti TiH2ph
ase
com
posi
tion
[vol
. %]
temperature [C]
Fig. 4.22. Changes of phases during heating of pre-treated TiH2 powder in air at 480C for 180 min from
200C to 780C at 10 K/min in He atmosphere.
4.2.7 TEM investigations
Fig. 4.23 (left) shows a bright field (BF) TEM micrograph of an untreated TiH2 particle. The surface of
the particle does not show any signs of oxide layers in this image. In contrast, TiH2 particles pre-treated in
air at 480C for 180 min exhibit a clearly visible surface layer with a rather constant thickness of about
100 nm (Fig. 4.23, right). In the example given the TiH2 fragment was so thin that oxide layers can be
identified on both sides. The oxide layers are polycrystalline as verified by the selected area electron
diffraction pattern (SAED) which is given in Fig. 4.23 together with the pattern corresponding to the non-
oxidised TiH2 core. From the position of the diffraction peaks a number of interplanar spacings dexp can be
derived. These are given in the first column of Table 4.2.
4 Results: Powders
53
A qualitative analysis of the oxygen content inside the oxide layer is also carried out. The EDX
measurements were performed along a line diagonally across the oxide layer. As a quantitative
determination of light elements is difficult without precise stand
Recommended
