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Chapter 10
Powder Injection Moulding of Tool Materials and
Materials Containing One-Dimensional Nanostructural
Elements
Leszek A. Dobrzański and Grzegorz Matula
Leszek A. Dobrzański and Grzegorz Matula
Additional information is available at the end of the chapter
Additional information is available at the end of the chapter
http://dx.doi.org/10.5772/67353
Abstract
As modern manufacturing methods have been developing, the application methods of
powders have changed, and they do not always have to be moulded prior to sintering.
The powder injection moulding (PIM) method is suitable for large-lot and mass production; still, powder consumption is not too high. The metal injection moulding (MIM) is an
advanced technology and not as developed as classical pressing and sintering but constantly and dynamically developing. The technology is developing towards micro-MIM,
that is, production of very small parts for miniaturised devices. The chapter presents the
overview of powder injection moulding as specialist powder metallurgy method and its
application for fabrication of tool materials. Specially, the fabrication of high-speed steels
and carbide-steels on their matrix by powder injection moulding is descripted. In last
part of the chapter, the results of own investigations of the structure with nanostructural
elements of high-speed steels and carbide-steels on their matrix fabricated by powder
injection moulding are presented.
Keywords: powder injection moulding, debinding, sintering, tool materials
nanostructural elements
1. Introduction
As modern manufacturing methods have been developing, the application methods of
powders have changed, and they do not always have to be moulded prior to sintering. This
depends on the technological process applied, for example, hardfacing, thermal spraying, or
selective laser sintering [1, 2] exclude the need of powder forming. Powder forming itself does
© 2017 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons
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224
Powder Metallurgy - Fundamentals and Case Studies
not have to be connected with pressing or with shapening the item being fabricated, as seen
with the example of pressure-free moulding of coatings or foils using polymer-powder slips.
Classical powder metallurgy is however based on pressing and at the same time forming the
powder in such a way that the item has the dimensions and shape of a ready product or semiproduct after sintering, requiring only minor final treatment [3, 4].
The introduction of atomisers into mass production of metallic powders allowed to reduce
the production costs of sintered materials. For instance, the manufacturing cost of high speed
conventionally cast and sintered steel at the end of the last century was comparable and
dependent mainly on the price of alloy additives. The automotive industry is now estimated
to be the key user of elements moulded and sintered using metallic powders [5]. Powderpressed and sintered gears intended for gearboxes are characterised by very silent work,
which is associated with the suppression of vibrations in porous materials. The teeth of such
gears are burnished and heat treated, which ensure their high hardness and resistance to
abrasive wear. Uniaxial pressing in closed dies is still most widespread and is dedicated to
symmetric elements with relatively simple shapes. Products with a highly developed surface, small dimensions, and massively produced are manufactured by the injection moulding
of powders witnessing very dynamic growth in the recent years. The moulding of ignition
plugs’ insulators was one of the first applications of this method in 1937 [6].
The method was employed in the 1970s of the last century for producing metals and ceramics used in electronics. The earlier limitations of this method were connected with binder
removal, which is technologically difficult, time-consuming and is the source of defects in the
form of cracks, gas bubbles, and distortions. In the majority of cases, binder removal is based
on its thermal degradation, often without access of oxygen, that is, in the process of controlled
pyrolysis. The growth of the technology in the recent years has been chiefly linked to a search
for alternative solutions ensuring debinding, which is fast and safe for the products manufactured and for the environment. The methods of solvent degradation under heightened pressure in a vacuum and in supercritical pressure conditions were employed as a result of such
investigations. The catalytic degradation method, adopted and developed by BASF, offering
at that time ready granulates and binder removal devices, turned out to be a breakthrough
method [7].
Although the powder injection moulding (PIM) method is suitable for large-lot and mass
production; still, powder consumption is not too high. In general, the powders of metals
and powders of metal alloys produced are principally designed for moulding and sintering products, and only a small fraction of powders is used in welding processes and in
other processes, for example, in chemical industry. For instance, the United States have produced about 365 thousand tonnes of iron and steel powders in 2013, of which the powders
used for other purposes than powder metallurgy account for only 10%. The remaining 90%
include also powders used in the metal injection moulding (MIM) process. The MIM technology represents only 1% of all the metallic powders used in Europe in powder metallurgy. This signifies that this is an advanced technology and not as developed as classical
pressing and sintering, but constantly and dynamically developing, which is confirmed by
numerous Refs. [8–10]. It should be noted that the comparison of consumption in tonnes of
powder for the MIM technology with other manufacturing methods is not the best indicator
Powder Injection Moulding of Tool Materials and Materials Containing One-Dimensional Nanostructural Elements
http://dx.doi.org/10.5772/67353
of this method’s popularity, because it is employed for producing small elements. The sales
of parts manufactured by this method have grown very sharply in, for example, Europe and
have gone up from less than 150 to over 250 mEUR between 2003 and 2013. The technology
is developing towards micro-MIM, that is, production of very small parts for miniaturised
devices [8, 11, 12].
2. Overview of powder injection moulding as specialist powder
metallurgy method and its application for fabrication of tool materials
Powder forming and sintering technologies offer unlimited opportunities for selection of
the chemical composition of the tool materials produced. Classical powder metallurgy,
based on uniaxial pressing and sintering with potential isostatic pressing at a high sintering
temperature, is used for the fabrication of the most popular tool materials such as sintered
carbides produced as, for example, inserts for turning tools. Frequently, the simple shapes
of such inserts do not require any other powder forming technique, whereas uniaxial pressing, especially one-sided pressing, is the easiest powder forming method, with no need to
use a binder or a large inclusion of slipping agents, which ensures small shrinkage of the
material after sintering. An uncomplicated technological process of fabrication of sintered
carbides, introduced for the first time by a German company, Krupp, in 1926, is still enjoying
great popularity, and sintered carbides themselves are the most often used tool material. For
example, the annual growth rate for the market of sintered and superhard tool materials in
2007–2012 was 7.5% per annum, with the projected growth for the subsequent years of 11.3%
[8, 13]. It is estimated that the global market of sintered carbides and superhard materials
may reach the threshold of USD 20.2 billion in 2018 owing to the development of industry and production in developing countries and due to the resulting increased demand for
highly efficient machining tools [14]. This method imposes, however, substantial constraints
if carbide inserts with highly complex shapes are produced. An alternative, and actually
an unrivalled method in such case, is powder injection moulding. It is injection moulding’s
strength that we can produce a complicated geometry or products in the final or nearly
final form at once. We can come close to the final goal at once, and this is the crucial reason
for choosing MIM [15]. The technology is dedicated to mass production solely, due to high
investment costs necessitating the use of high class equipment. Metal injection moulding
(MIM) is especially significant in metallic powder forming. Table 1 lists the key advantages
of the PIM method as compared to general disadvantages and advantages of other manufacturing techniques of elements with their shapes and dimensions similar to the final shapes
and dimensions.
It is assumed that the powder injection moulding technology was introduced in 1849 by pressure casting of nonferrous metal alloys [8]. Further development was associated with proficiency in the polymer processing technology and with the first piston moulders constructed
in the USA and Germany. PIM’s history in Europe is relatively short and dates back 30 years
ago. The first elements manufactured by this method are orthodontic hooks produced in
Germany in the 1980s [17].
225
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Powder Metallurgy - Fundamentals and Case Studies
Selected characteristics
Technology
PIM
PM
Casting
Machining
Density
98%
90%
95–99%
100%
Strength versus solid
material
100%
70%
98%
100%
Complexity
High
Low
Medium
High
Part weight
Between 0.003 g and
17 kg
Between 0.1 g and
10 kg
From 1 g
From 0.1 g
Wall thickness
0.1–10 mm
2 mm
5 mm
2 mm
Surface manufacturing
accuracy
0.4–0.8 μm
2 μm
2 μm
0.4 do 2 μm
Production scale
Mass
Mass
Medium/low
Medium/low
Table 1. Selected characteristics of elements manufactured by various technologies [16].
The injection moulding of a polymer-powder slip allows to produce relatively small parts
with complicated shapes and developed area and requires no plastic working or machining, which is in line with the modern direction of producing ready parts, that is, “near–netshape”. Despite numerous advantages, the injection moulding process is not suitable for
fabrication of large parts. The largest dimension should not exceed 100 mm. This is because
subsequent debinding is required prior to sintering, as the polymer materials contained
in a binder are undergoing gassing during thermal degradation, and the pressure of the
gas closed in the pores is rising due to being abruptly heated to sintering temperature.
The wall thickness of the parts manufactured by this method is not more than 10 mm,
they have complicated shape and high manufacturing precision and low production costs
[9, 18]. This method is most often applied for manufacturing parts hard to produce by
other techniques, for example, elements with the smallest mass of not more than 0.5 g, in
particular orthodontic hooks [16]. The technology, considering a possibility of production
automation, high speed and dimensional repeatability and high costs of injection moulders and heating devices, is designed for large-lot or mass production [19, 20]. Owing to
the principal advantage of this technique—where ready parts are produced without additional treatment being necessary—the technique is used more and more extensively for
producing hard materials, including tool materials, which are exceptionally difficult and
costly to machine [6, 9, 13, 15, 21]. The formability of metallic and ceramic powders and
their mixtures enables to fabricate metal tools with relatively high ductility, ceramic tools
with high hardness or metal matrix composites (MMC) and ceramic matrix composites
(CMC) combining high properties distinct for metals and ceramics [22–25]. The forming of
metallic or ceramic powders in a matrix of polymer binders, especially injection moulding
or extrusion, has been applied for producing tool materials, including high-speed steels,
carbide-steels and sintered carbides, characterised by their structure similar to commercial
sintered tool materials [26–29]. It is not possible to achieve such high bending strength as
for high-speed steels fabricated by hot isostatic pressing or for sintered carbides pressed
Powder Injection Moulding of Tool Materials and Materials Containing One-Dimensional Nanostructural Elements
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isostatically during sintering. The main reason for lower bending strength of the manufactured tool materials is not, however, powder forming technologies, but the process of selfsintering, which does not allow to eliminate completely the locally occurring pores being
material discontinuity lowering the material’s mechanical properties. It is beyond doubt a
benefit of polymer-powder slip forming, mainly powder injection moulding, that plastic
working and machining operations are eliminated, hence reducing the related costs, as a
result of which application possibilities of the so manufactured tool materials are offered,
especially where they are not subject to strong dynamic loads. Moreover, the parts produced this way do not have to act only as tool material but act as an element working
in the conditions of tribological wear. An additional benefit of injection moulding of the
investigated tool materials is that protective and/or reactive atmospheres can be utilised
for sintering, permitting to use furnaces cheaper than vacuum furnaces and hence allowing
automation. The use of modern, polymer binder-based powder forming technologies, in
particular injection moulding, for preparing tool materials, offers promising perspectives
for the manufacture of tool materials and is consistent with the valid development trend of
this technology. The sales of products manufactured by PIM and MIM have been growing
over the last 25 years and signify the industry’s huge interest in such methods [30].
In powder injection moulding, an injected preform is produced, which should then undergo
debinding and sintering to achieve high functional properties [16, 18, 27, 30, 31]. Classical
injection moulding in moulders is the same as the moulding of thermoplastic polymers
(Figure 1).
Figure 1. General chart of powder injection moulding.
227
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Powder Metallurgy - Fundamentals and Case Studies
Metallic or ceramic powders, and even their mixtures, are used as feedstock in powder injection moulding. The use of thermoplastic polymers as a binder, which is binding metallic or
ceramic powder, enables to transport it and mould it in an injector socket. Two types of binders based on, respectively, paraffin and polymer materials and an aqueous methylcellulose
solution (Wiech’s or Rivers’ solution), are used most often [8]. A powder-to-binder ratio is
closely linked to the shape, size of powder particles, material type, powder wettability by a
binder, and the properties of the binder itself and is also linked to mixture production conditions [32, 33]. By the rule, it is more difficult to mix ceramic powders with a binder and to
inject them than is the case with metallic powders. A binder with larger fraction has to be used
for this. Figure 2a shows the influence of powder fraction on viscosity curve.
Debinding must be carried out in such a way to prevent cracks, blisters, shape deformation,
gas bubbles, or semi-product delamination, so it should be deformed early enough, maintaining the shape of the element produced. The most popular method of debinding used in PIM
(Figure 2b) is thermal degradation and solvent degradation. A mixed method is linked to
combined degradation techniques, for example, solvent and thermal technique.
Currently, for exceptionally small preforms formed, for example, in the μPIM process,
debinding is associated with heating to a sintering temperature [8, 12]. Regardless the semiproduct size, debinding is making a preform becoming very brittle. Reducing atmospheres
consisting of hydrogen with a high concentration of hydrogen (85%) ensure highest strength,
hence, increase carbon concentration in the preform after sintering. Preform density is about
60% of the theoretical density when the binder is completely removed. This density depends
on the binder fraction, and this fraction depends on the type of the powder formed. Most
often, metal powders are so selected that they are characterised by a circular shape with good
wettability, which allows to reduce the binder fraction as much as possible to only 30%. The
binder fraction in case of ceramic powders may reach even 55%.
Preform sintering takes place after debinding. Irrespective of the preform density, it is subjected to densification and shrinkage due to sintering. This is natural regardless the preform
Figure 2. (a) Influence of type and amount of powder on viscosity, (b) debinding methods in PM technologies.
Powder Injection Moulding of Tool Materials and Materials Containing One-Dimensional Nanostructural Elements
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forming method; however, as a result of high porosity, shrinkage is very strong and an uncontrolled change of the sintered elements’ shape may occur. Because a change in the volume of
the sintered material is inversely proportional to its density after forming, growth in preform
density reduces the risk of shape defects of the preform being formed. Injection pressure,
powder size, and binder fraction have influence on increase in preform density. Sintering is
usually the last operation of the technological process decisive for density and properties of
the ready product [18]. If the ready element should have high mechanical properties, final
heat treatment and often machining are used to ensure accurate dimensions of the produced
sinters. The sintering of injection moulded or pressure-free moulded powders does not differ
largely from the sintering of powders formed by other methods. The key properties of sinters,
especially tool materials, are improved by the formation of carbonitrides as a result of the
interacting atmosphere containing nitrogen during debinding and sintering [23]. Sintering
is initiated by growth in the concentration of carbon resulting from debinding, however, in
case of some materials such as stainless steels or high-speed steels, carbon concentration must
be closely monitored due to their properties or influence on heat treatment. The type of the
binder used influences the final carbon concentration [18]. Sintering is also influenced by
particle size. Injection-moulded fine-grained powders with a larger specific surface area are
filling the volume of the sintered preform more thoroughly and are subject to faster remelting.
The powder particle size is also decisive for surface roughness and for the value of the sintered
material’s edge radius. In case of classical pressing, fine-grained powder is not filling the die
socket so well due to low-powder liquidity. An atmosphere inside the furnace chamber is an
important factor conditioning sintering [26]. Although the vacuum is not related to direct gas
costs, vacuum sintering is a costly alternative considering that furnaces are equipped with
vacuum systems and require maintenance. Nonetheless, vacuum is often used for sintering
high-speed steels, especially those pressed in a die or pressed isostatically. The sintering of
injection-moulded high-speed steels in high vacuum is quite difficult due to the gas products
being released, coming from the thermal degradation of base polymer residues. Final heat
treatment is required to achieve high mechanical properties for high-speed steels or sintered
carbides manufactured by PIM [34–38].
3. Fabrication of high-speed steels and carbide-steels on their matrix
by powder injection moulding
Two types of tool materials, that is, high-speed steels and carbide steels with the matrix of the
same high-speed steels were fabricated by the powder injection moulding method. The powders of high-speed steels with the chemical composition given in Table 2 and with the particle size determined by the laser diffraction method in the Malvern Mastersizer 2000 device
shown in Figure 3, and with density and technological properties presented in Table 2, were
fabricated by atomising inert gas (Figures 3 and 4) by Sandvik Osprey Ltd. A commercially
available mixture of carbides, with the trade name of Tetra Carbides (in this chapter TetraC)
by Treibacher Industrie AG containing WC, TiC, TaC, and NbC carbide powders with the
volume fraction and technological properties presented in Table 3 and other properties given
in Figure 3 and morphology given in Figure 4 were used as hard carbide phases.
229
230
Powder Metallurgy - Fundamentals and Case Studies
Steel grade
Mass concentration of elements (%)
C
W
Mo
V
Co
Cr
Mn
Si
EN HS 6-5-2
0.84
6.54
4.81
1.95
–
3.97
0.36
0.35
EN HS 12-0-5-5
1.47
11.8
0.06
4.75
4.64
4.76
0.5
0.43
Table 2. Concentration of alloy elements in high-speed steels powders used.
Figure 3. Particle size distribution of (a) EN HS 6-5-2, (b) EN HS 12-0-5-5 high-speed steels powder, and (c) carbides WC,
TiC, TaC, NbC mixture (TetraC).
Figure 4. Morphology of (a) EN HS 6-5-2, (b) EN HS 12-0-5-5 high-speed steels, and (c) mixture of EN HS 12-0-5-5 highspeed steel and carbides WC, TiC, TaC, NbC mixture (TetraC).
Powder
Powder properties
Density (g/cm3)
Bulk density (g/cm3)
Tap density, (g/cm3)
EN HS 6-5-2
8.16
2.89
4.02
EN HS 12-0-5-5
8.19
3.05
4.25
Mixture 47% WC, 14% TiC,
33% TaC, 6% NbC (TetraC)
10.23
4.38
5.64
Table 3. Technological properties of used powders.
Powder Injection Moulding of Tool Materials and Materials Containing One-Dimensional Nanostructural Elements
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The applied high-speed steel powders meet the condition indicating that ball-shaped powder
with the average particle size of below 20 μm is most suitable for injection moulding of powders. The ball shape is most desired due to high wettability, low slipping agent viscosity, and
high-packing density of particles. The size distribution of particles is also an important factor.
If particle size distribution is relatively broad, then the pores forming between large grains
may be filled by small particles, as signified by the Sw curve inclination coefficient Eq. (1) [18].
2.56
Sw = ______
D .
(1)
90
log ___
D
10
The powder with the Sw coefficient of about two is most recommended for injection moulding.
It is not recommended to mould powder with the Sw coefficient of seven with its powder grain
size distribution characteristic being very narrow. The powder particle size analysis is shown
in Table 4. Pure steel powder, EN HS 6-5-2, has the highest Sw coefficient value. If a mixture
of WC, TiC, TaC, and NbC carbides is introduced into the powder of high-speed steels, this is
broadening the powder grain size distribution characteristic, thus advantageously reducing
the Sw coefficient whose lowest value for the mixture of high-speed steels EN HS 12-0-5-5 with
carbides is 2.69. The calculated Sw coefficients of the applied powders do not exceed the value
4, which signifies that they can be used for injection moulding [18]. Table 5 presents information about the components of the formed powder and polymer-powder mixtures.
The purpose of the stearic acid used was to cover the surface of carbide powders before mixing them with high-speed steel powders. Stearic acid covering the carbides with a thin layer
is increasing their wettability when mixing with other binder components and decreasing a
ready feedstock’s viscosity [8, 13, 31, 39]. Next, the covered carbides together with a binder
and high-speed steel powder were initially agitated in a chamber of a universal Rheomex
CTW100p stirrer by Haake enabling to measure torque and rotation speed of stirrers and
charge temperature. The results of rheological tests for the mixtures of high-speed steel powders, carbides and binder, carried out with a capillary rheometer, Rheoflixer by ThermoHaake,
at 170, 180 and 190°C, with the homogenisation speed of 10–10000 s−1 and with the torque
depending on the phase composition of the examined polymer-powder mixtures and their
homogenisation time, allowed to select a mixture with relatively low viscosity, high volume
fraction of powders, good powder wettability by applying a binder and thus with the ability
of fast homogenisation.
Powder
High-speed steels
EN HS 6-5-2
EN HS 12-0-5-5
Undersize D10 (μm)
3.5
4.68
Undersize D90 (μm)
16
Coefficient Sw
3.87
Carbide mixture
in TetraC
Mixture of appropriate high-speed
steels with mixture of TetraC
EN HS 6-5-2
EN HS 12-0-5-5
1.45
2.8
2.8
28.5
7.3
17
25
3.25
3.64
3.26
2.69
Table 4. The curve slope coefficients of the particle size distribution Sw calculated on the basis of undersize D10 and D90.
231
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Powder Metallurgy - Fundamentals and Case Studies
Metal matrix
Carbides
Binder
EN HS 6-5-2
–
PW, PP, HDPE, SA
EN HS 12-0-5-5
–
EN HS 6-5-2
Mixture 47% WC, 14% TiC, 33% TaC, 6% NbC (TetraC)
EN HS 12-0-5-5
Table 5. Components of the formed powder and polymer-powder mixtures.
The possibility of injection moulding of the applied polymer-powder mixture was initially
evaluated based on the results from measuring the torque of stirrers during skip homogenisation. The maximum fraction of solid particles was thus established while maintaining
relatively low viscosity of the polymer-powder mixture enabling injection moulding or extrusion. A high fraction of powder ensures minimum shrinkage in sintering, while an increase
in binder fraction ensures easy forming but extends degradation time and increases sinter
shrinkage which often causes the occurrence of distortions [9]. Irrespective of the polymerpowder mixture type, paraffin fraction was always equal to the fraction of the main binder
component, that is, polyethylene or polypropylene.
Figure 5a shows the influence of high-density polyethylene (HDPE) and polypropylene on
the torque curve during homogenisation of polymer-powder slip containing 70% of EN HS
12-0-5-5 steel powder. Regardless the homogenisation time and type of high-speed steel powder, a mixture containing a polypropylene-paraffin (PP/PW) binder is characterised by the
lowest torque of stirrers in relation to a mixture containing an HDPE/PW slip. The minimum
torque value for stirrers during the homogenisation of a mixture containing polyethylene for
3 h is about 2.2 Nm. In case of a feedstock containing polypropylene, torque after such long
homogenisation time is 1 Nm. A curve for a mixture with polyethylene applied is not stable,
which may signify inhomogeneous distribution of metallic powder in the binder matrix,
Figure 5. Torque measurements of feedstock with (a) 70% of EN HS 12-0-5-5 and two different binder compositions:
A – HDPE/PW, B – PP/PW, (b) different amount of EN HS 6-5-2 powder: C – 75%, D – 70%, E – 65%, F – 60%.
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despite long homogenisation time. Torque for a mixture with PP applied is going down over
the entire homogenisation range, that is, over 3 h.
A correctly selected binder should wet the powder and reach the homogenous state within
short time of about 30 min, and the torque characteristic should then stabilise. The longer time
of homogenisation may lead to partial degradation of the low melting binder component, paraffin (PW) in this case, or to breaking PP and HDPE chains and to re-netting. The consequence
of this would be higher torque, which is precluded by a falling tendency of the analysed curves.
Figure 5b shows torque variations depending on the fraction of high-speed steel EN HS 6-5-2
powder. If powder fraction is decreased by 5%, the torque characteristic is more uniform
and the torque value after 30-min homogenisation falls from 4.0 to about 3.1 Nm. If powder
fraction is again decreased by 5 and 10%, this causes the characteristic to be further equalised
and torque to be reduced to the minimum value of about 1.2 Nm. The torque of the stirrers
homogenising (for 1 h) a polymer and powder slip containing 60% of TetraC and a PP/PW
binder is about four times higher in relation to the torque of the similar mixture containing
EN HS 6-5-2 steel powder. The torque value is substantially decreased by using stearic acid
(SA) covering the surface of carbides. Stearic acid was not used for covering high-speed steel
powders due to its adverse impact on the high-speed steel structure after sintering, characterised by large precipitates of carbides on grain limits [31].
A low value of torque of mixtures containing polypropylene (PP) and paraffin (PW) corresponds
to low viscosity. Powders of the EN HS 12-0-5-5 steel mixture with PP and PW exhibit smaller
viscosity in relation to a mixture containing HDPE instead of PP, and this is independent of the
homogenisation rate (Figure 6a). A mixture of high-speed steel with TetraC, which is accounting
for 10% of the volume fraction, also has low viscosity similar to the mixture of EN HS 12-0-5-5
steel with PP and PW. Paraffin, apart from lessening the density, allows to use solvent degradation expediting the rate of thermal degradation and shortens the duration of the whole cycle.
Figure 6. Viscosity curve for selected feedstock with (a) EN HS 12-0-5-5 and two different compositions: A – HDPE/PW,
B – PP/PW, (b) different amount of EN HS 6-5-2 powder: D – 70%, E – 65%, F – 60%.
233
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Powder Metallurgy - Fundamentals and Case Studies
Metallic powder with a 60, 65 and 70% fraction containing a polypropylene-paraffin binder
generally shows low viscosity regardless the fraction of powder. A mixture with the lowest
content of powder exhibits the lowest viscosity (Figure 6b). Polypropylene and paraffin are
the main binder component regardless the content of stearic acid.
Stearic acid improves wettability of metallic and ceramic powders by covering their oxidised
(polar) surface that adsorbs the hydrophilic part of the chain as a result of the existing electrostatic forces between the powder and the wetting agent. The nonpolar part of the chain
should be mixed without limitations with other polymers present in the binder. Apart from
reducing viscosity, stearic acid is acting as a slipping agent in contact between the powder
and the die surface or the surface of another particle. It also prevents powders from migrating
during high-speed homogenisation. A process of migrating the powder inside the capillary or
the destruction of the binder structure occur most probably during the high-speed homogenisation of a mixture not containing stearic acid. This is manifested by strongly falling viscosity together with an increased homogenisation rate. The viscosity of the mixtures containing
stearic acid is not so much dependent on a homogenisation rate; hence, the growing rate of
homogenisation does not have such a strong effect on the structure of a homogenous mixture.
The torque is only negligibly decreased by increasing the content of stearic acid by another
4%; hence, its content did not exceed 4%.
Considering the feedstock viscosity, the maximum applicable fraction of carbides not coated
with stearic acid is 50%. If stearic acid is used for a mixture containing 50% of carbides, viscosity is greatly reduced and a higher volume fraction of carbides can be obtained. The maximum volume fraction of powders which can be applied in a mixture for injection moulding
was established by analysing the technological properties of polymer and powder mixtures
containing binder-carbides, and hence, a mixture with only a 10% volume fraction of carbides
was applied. Four different polymer-powder mixtures for injection moulding were selected
and prepared. The fraction of such mixtures was presented in Table 6. The letter F was used
in denomination to differentiate the injector’s feedstock from pure powder.
The injection moulded materials were then subjected to solvent degradation, thermal degradation, sintering, and heat treatment. The specimens were sintered in a pipe furnace in the
atmosphere of a flowing mixture of N2/10%H2 gases with the maximum sintering temperature
of 1450°C. The sintering time of the moulded parts, regardless the furnace type and sintering
Injector feedstock
Fraction volume of component (%)
EN HS 6-5-2 (O)
EN HS 12-0-5-5
TetraC
PP
HDPE
PW
SA
F EN HS 6-5-2
68
–
–
–
16
16
–
F EN HS 12-0-5-5
–
68
–
–
16
16
–
F EN HS 6-5-2/TetraC
58
–
10
14
–
14
4
F EN HS 12-0-5-5/ TetraC
–
58
10
14
–
14
4
Table 6. Types of injection-moulded polymer and powder mixtures.
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temperature, was 30 min each time. A sintering temperature was selected experimentally. The
sintering temperature range was determined based on preliminary tests of 1180–1300°C with
20°C steps. Heating speed to sintering and cooling temperature was 5°C per min. Further heat
treatment was performed with classical hardening from an austenitisation temperature of
1220, 1240 and 1260°C, and triple tempering in the same temperature of 540, 570, 600 or 630°C
for 1 h was carried out directly after hardening in oil.
The conditions of solvent degradation were selected experimentally taking account of the
element mass depending on the paraffin solving time and bath temperature. Correctly performed debinding should ensure a uniform concentration of carbon within the entire volume
of the specimen. EN HS 6-5-2 steel subjected to thermal degradation at 450°C, in the atmosphere of N2-10%H2 and sintered at 1240–1260°C, reaches density close to theoretical density,
that is, it does not exhibit any pores (Figure 7a). The density of EN HS 6-5-2/TetraC carbidesteel sintered within the same temperature range was shown in the Figure 7 to compare the
influence of carbon additives on the optimum sintering temperature ensuring highest density. EN HS 6-5-2/TetraC carbide-steel reaches the maximum density of 8.77 g/cm3 after sintering at 1260°C. The density values were deliberately not referenced to the theoretical density
of the material fabricated, which, based on the calculations, should be 8.69 g/cm3. A varying
chemical composition of the sinter as a result of an increased concentration of carbon, dependent upon debinding and upon the nitrogen coming from the atmosphere during sintering
and forming the carbonitrides, does not allow to determine accurately what should be its
maximum density. In addition, the WC, TiC, TaC and NbC carbides introduced are dissolving
in a high-speed matrix during sintering at a high temperature and create M6C and MC-type
carbonitrides or MX carbonitrides in it, in case of sintering in an atmosphere containing nitrogen, identified with the diffraction methods. The density of the newly created phases differs from the carbides introduced, which is impacting the sinter’s overall density. For this
reason, the sinter porosity presented in Figure 7b is a very interesting piece of information.
Regardless the sintering temperature, EN HS 6-5-2/TetraC carbide-steel is characterised by
low porosity not exceeding 1%, appropriate for tool materials. The share of pores in the most
Figure 7. Influence of sintering temperature on (a) density of EN HS 6-5-2 and EN HS 6-5-2 /TetraC, (b) porosity of EN
HS 6-5-2/TetraC, (c) density of EN HS 12-0-5-5 depend on debinding at 450 or 475°C and sintering temperature.
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interesting range of the sintering temperature of 1240–1260°C is about 0.6%. The pores are ball
shaped in the majority of cases and do not exceed the size of 2 μm; however, large pores may
exist locally, which are forming most probably as a result of gas bubbles being formed during injection moulding or thermal degradation. Changes in the EN HS 12-0-5-5 steel density
presented in Figure 7c points out that a lower degradation temperature is reducing a sintering temperature. The maximum density in the sintering temperature of 1240°C occurs after
degradation at 450°C. If the degradation temperature is increased by 25°C, it becomes necessary to increase the sintering temperature by 30°C to achieve the maximum sinter density.
Generally, EN HS 12-0-5-5 steel can be sintered within a broader temperature range versus
EN HS 6-5-2. It was found by observing the structure of the sintered carbide-steels in which
TetraC was applied covered with stearic acid that both, in EN HS 6-5-2/TetraC material and
in EN HS 12-0-5-5/TetraC, the carbides surrounding the high steel grains are inhibiting their
growth. The grain size of the sinter matrix is comparable to the particles of the high-speed
steel powder employed. It is feasible to manufacture tool materials with high functional properties by selecting an appropriate binder, binder fraction and by controlling the technological
process conditions, which enables to increase a concentration of carbon coming from thermal
degradation products of the binder surrounding the nitride particles and activating a sintering process. Unfortunately, the locally occurring bubbles of the gas created during injection
moulding or thermal degradation are not eliminated during free sintering, despite the presence of the liquid phase, which effectively compromises the bending strength of the sinter.
Due to large gas bubbles existing in the material, further heat treatment does not influence
bending strength, and the maximum value of about 1400 MPa is lower than the strength of
conventional EN HS 6-5-2 cast and heat treated steel. The binder used, apart from its key task,
that is, to enable injection moulding, plays an additional role as a source of carbon activating
sintering. A rising concentration of carbon is dependent here upon the type of the binder and
its debinding conditions.
4. Results of own investigations of the structure with nanostructural
elements of high-speed steels and carbide steels on their matrix
fabricated by powder injection moulding
In case of high-speed steels, dispersion carbides are released during tempering, and more
rarely as pre-eutectoid precipitates when cooling from austenitisation temperature. In
carbide-steels, small dispersion precipitates of the nanometric size occur extensively after
quenching and tempering and after sintering. Such precipitates are surrounding large M6C
carbides and primary austenite grains blocking their growth. Moreover, they dissolve more
quickly in austenitisation because of their small size, and alloy additives responsible for the
secondary hardness effect are passing more easily to the matrix. The structure of EN HS 6-5-2/
TetraC carbide-steel sintered at 1230°C shows small growth and coagulation of carbides in
relation to the high-speed steel structure, sintered at the same temperature. Despite this, the
structure of carbide-steels is still homogenous and fine-grained with carbides surrounding
circular grains of high-speed steel. When comparing EN HS 6-5-2 and EN HS 6-5-2/TetraC
Powder Injection Moulding of Tool Materials and Materials Containing One-Dimensional Nanostructural Elements
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materials sintered in the same conditions at the temperature of 1300°C, one may conclude that
a strong distortion occurred in the both cases, signifying an extensive liquid phase existing
during sintering. A structure of carbide-steel sintered at 1300°C (Figure 8c) does not exhibit,
however, an eutectic typical for a high-speed steel presented in Figure 8b. Carbides in highspeed steel with a characteristic “fishbone” shape are revealed as light phases in the image of
secondary electrons and contain mainly Fe, W, Mo and a small concentration of V. The precipitates are sized up to 1 mm. If a mixture of carbides is employed, in particular interstitial
phases with a regular network, stable at a high sintering and austenitisation temperature, the
growth of primary austenite grains is inhibited and the hardness of carbide-steels after heat
treatment is improved in relation to high-speed steel. Irrespective of the type of the binder
used, which is influencing the way a polymer-powder slip is formed, and also irrespective of
the sintering temperature and atmosphere, the precipitation of large, eutectic carbides typical
for high-speed steels sintered freely at temperature exceeding the solidus line was completely
eliminated in carbide-steels and also sometimes austenitised in such conditions, causing local
remelting. Owing to such property, it is not necessary in sintering to use heating devices
equipped with very accurate measuring and control systems ensuring temperature stability
within a very narrow range, that is, approx. 5°C for the EN HS 6-5-2 steel applied.
The average size of carbides in carbide-steel sintered at 1300°C is rising by approx. 0.1 μm in
relation to carbides existing after sintering at 1280°C. The average and maximum size of carbides and their fraction volume in EN HS 6-5-2/TetraC sintered within the entire temperature
range are given in Table 7. The so-selected chemical composition of carbide steel sintered in
the atmosphere of a flowing mixture of N2-10% H2 gases enables sintering within a wide temperature range ensuring a homogenous structure. If the protective atmosphere of N2-10% H2
is applied, surface oxidisation is prevented during thermal degradation and especially during
sintering, while introducing at the same time, the desired nitrogen into the sinter forming fine
nanometric carbonitride precipitates limiting the grain growth of other carbon precipitates
and matrix grains, in particular carbide steels. When sintering temperature is raised from
1200 to 1300°C, the average size of carbides is increased by about 70%.
WC carbides with light colour based on secondary electrons are presented in Figure 9. A phase
with grey colour rich in W, Fe, V, and Mo was formed from the alloy additives situated in
high-speed steel or from elements coming from the solved carbides introduced into the steel.
No existence of WC carbides was found at an elevated sintering temperature, which is con-
Figure 8. (a) Microstructure of metal matrix composite on the basis of EN HS 6-5-2/TetraC with large spherical pore, (b)
and (c) structure of sintered at temperature 1300°C, (b) EN HS 6-5-2, (c) EN HS 6-5-2/TetraC.
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Sintering temperature (°C)
1200
1220
1240
1260
1280
1300
Carbides amount (%)
36.2
35.9
33.8
33.2
34.2
34.2
Average size of carbides (μm)
1.65
2.06
2.2
2.66
2.71
2.82
Maximum size of carbides (μm)
21.57
25.68
28.54
31.58
38.17
40.58
Table 7. Average and maximum size of carbides precipitation and their amount versus sintering temperature.
Figure 9. (a) Structure of EN HS 6-5-2/TetraC carbide-steels sintered at 1200°C and a charts of scattered X-ray radiation
for the areas marked in (a) (b) A and (c) B, respectively.
firmed by the results of an X-ray phase analysis. M6C and MC carbides typical for high-speed
steel were identified, though. The growth of light carbides, shown in the image of secondary
electrons, rich in W, Mo, and Fe, is limited by the surrounding grey, spherical precipitate rich
in Ti and V, presented in Figure 9a, the size of which does not exceed mostly several dozens
of nanometres, and the largest ones do not exceed 1 μm. In order to better present such nanometric precipitates, Figure 10 shows the morphology of carbides in the form of sediment of
electrolytically isolated precipitates.
Figure 10. Structure of EN HS 6-5-2/TetraC carbide-steel sintered at 1280°C, (a) sediment of electrolytically isolated
precipitates, SEM, (b) X-ray diffraction pattern of electrolytically isolated precipitates from EN HS 6-5-2/TetraC carbidesteel tempered and quenched at 540°C.
Powder Injection Moulding of Tool Materials and Materials Containing One-Dimensional Nanostructural Elements
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The existence of M2C carbides typical for EN HS 6-5-2 high-speed steels, rich in molybdenum,
in the form of long needles, precipitated at a high sintering temperature and considerably
reducing mechanical properties, was not found [40]. M6C and MX precipitates exist only in
the sediment of electrolytically isolated carbides (Figure 10).
Nanometric precipitates in EN HS 6-5-2/TetraC carbide-steel are presented in Figures 11a
and 12a present results of examinations of precipitates coming from EN HS 6-5-2/TetraC
carbide-steel after heat treatment in a transmission electron microscope. The diagram of the
function of scattered X-ray radiation for such precipitates presented in Figures 11b and 12b
proves that carbides are rich in tungsten and molybdenum, which rather indicate the M6C
phase, not MC or M2C phase released in high-speed steels during tempering. Undoubtedly,
their size does not exceed 200 Nm. It should be added that typical dispersive MC or M2C
carbides, precipitating during tempering, have their size ten times smaller. A higher fraction
of retained austenite in carbide-steel in relation to classical high-speed steels is caused by a
higher concentration of carbon remaining after debinding. In addition, carbon is released
as a result of carbonitrides being formed and the introduced carbides being dissolved. If
carbon concentration is increased, the start temperature of martensite transformation is
decreased; hence, a fraction volume of retained austenite after quenching in such materials is up to 62%. A higher tempering temperature of approx. 600°C has to be applied due
to variations in retained austenite fraction in the matrix of the analysed carbide-steels, in
relation to high-speed steel with the analogous chemical composition of the matrix. EN HS
6-5-2 high-speed steel should be quenched at below 600°C. In case of carbide-steels and a
large fraction of retained austenite in such steels after hardening, it is necessary to temper at
a higher temperature ensuring the transformation of retained austenite into martensite and
high hardness after heat treatment. Unfortunately, M6C carbides may precipitate at such high
temperature; hence, most probably, the precipitates seen in Figures 11 and 12 are the result
of it. Precipitates with a nanometric size, formed during sintering, are rich in other elements
such as V, Ti, and N.
Figure 11. (a) HAADF image of precipitate in EN HS 6-5-2/TetraC carbide-steel after quenching and tempering twice at
630°C, SEM, (b) diagram of the function of scattered X-ray radiation for the carbide shown in (a).
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Powder Metallurgy - Fundamentals and Case Studies
Figure 12. (a) HAADF image of precipitate in EN HS 6-5-2/TetraC carbide-steel after quenching and tempering twice at
630°C, SEM, (b) diagram of the function of scattered X-ray radiation for the carbide shown in (a).
Additional information
The results of the research carried out partially in the research Project “NANOCOPOR—
Determining the importance of the effect of the one-dimensional nanostructural materials on
the structure and properties of newly developed functional nanocomposite and nanoporous
materials”, funded by the DEC-2012/07/B/ST8/04070 of the Polish National Science Centre in
the framework of the “OPUS” competitions, headed by Prof. Leszek A. Dobrzański were used
in this chapter.
Author details
Leszek A. Dobrzański* and Grzegorz Matula
*Address all correspondence to: leszek.adam@gmail.com
Faculty of Mechanical Engineering, Silesian University of Technology, Gliwice, Poland
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