Fe composites

Fe composites

Journal of Materials Processing Technology 127 (2002) 131–139 Self-propagating high-temperature synthesis and liquid-phase sintering of TiC/Fe compos...

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Journal of Materials Processing Technology 127 (2002) 131–139

Self-propagating high-temperature synthesis and liquid-phase sintering of TiC/Fe composites Patrik Perssona, Anders E.W. Jarforsb,*, Steven Savagec a

The Swedish Defence Materials Administration, Bane´rgatan 62, S-115 88 Stockholm, Sweden Department of Materials Processing, Royal Institute of Technology, S-100 44 Stockholm, Sweden c Swedish Defence Research Establishment, S-172 90 Stockholm, Sweden


Abstract The present paper addresses a possible route for the manufacturing of iron-based metal–matrix composites (MMCs) with a high level of reinforcement. The ceramic reinforcement studied was titanium carbide (TiC). TiC is one of the hardest materials to be found, which is why a TiC/Fe composite has the potential to be used as armour. Two manufacturing routes have been evaluated experimentally, i.e., liquid-phase sintering (LPS) and self-propagating high-temperature synthesis (SHS). LPS was found to be more effective due to easier process control and due to the process yielding a more homogeneous material. Both LPS and SHS produced a material with a relatively high degree of porosity. The porosity in the LPS experiments could be decreased by varying the process parameters, but this was not possible in the SHS process. Metallographic investigation shows that the TiC/Fe system is feasible and that applications utilising TiC/Fe composites are possible in the future. This is due to the fact that despite the porosity, an improvement of the MMC properties is detected, compared to those of the matrix material. # 2002 Elsevier Science B.V. All rights reserved. Keywords: Self-propagating high-temperature synthesis; TiC/Fe composites; Liquid-phase sintering

1. Background The purpose of the present paper is to investigate the possibility of producing an iron-based metal–matrix composite (MMC) effectively as a component of light armour. This is due to the continuing development of new weapons and improved projectiles with increased penetration ability. This development leads to the need of more effective protection of both military personnel and military vehicles. Effective light armour should be able to protect against a wide range of small calibre projectiles, including projectiles from direct fire to splinters. Traditional armour consists of materials such as steel, aluminium and titanium alloys, whilst new armours include the use of ceramics. One of the problems with ceramic materials is that they crack after being hit by a projectile, deteriorating their protective capability, which why ceramic materials have to be backed by a more ductile material. The ideal armour material should therefore be hard enough to disintegrate the projectile and ductile enough to absorb the impact energy. One promising type of materials is MMCs, since they are made from a softer matrix and a harder reinforcing phase, often a ceramic. * Corresponding author. Tel.: þ46-8-790-7989; fax: þ46-8-216-557. E-mail address: [email protected] (A.E.W. Jarfors).

In the present study, the hard phase is titanium carbide (TiC) and the matrix phase is an iron-based alloy. The aim and scope of the present paper is to investigate liquid-phase sintering (LPS) and self-propagating high-temperature synthesis (SHS), for the manufacturing of TiC/Fe composites with a high level of reinforcement. The most critical feature for this manufacturing is pore closure, to reach a high relative density, thereby to enhance the ductility of the resulting MMC.

2. LPS LPS is today a well known and established manufacturing route in powder metallurgy [1]. In LPS a powder mixture consisting of the alloy powder and sometimes a binder/ lubricant is pressed into a green body. This green body is then heated to a temperature where a liquid phase is formed. The consolidation of the green body into a dense body consists of three stages [1]: (i) particle rearrangement due to surface forces within the compact; (ii) solution reprecipitation type of rearrangement of the particles; and (iii) finial pore closure and grain growth. In the first stage, capillary forces dominated the densification. Agglomerates are disintegrated, the liquid penetrating

0924-0136/02/$ – see front matter # 2002 Elsevier Science B.V. All rights reserved. PII: S 0 9 2 4 - 0 1 3 6 ( 0 2 ) 0 0 1 1 3 - 9


P. Persson et al. / Journal of Materials Processing Technology 127 (2002) 131–139

the grain boundaries and covering the particle surfaces. This leads to a rapid densification and the elimination of the porosity. This process is strongly dependent on the amount of liquid present and on the capillary forces. The capillary forces are also dependent on the solubility of the solid phase in the liquid, indicating that there may be a strong temperature dependence on this initial stage. In general, an increased temperature improves the densification rate, especially in systems with a low mutual solubility of the solid and liquid phases [1]. Wetting also depends on the particle size and the particle size distribution of the solid phase. The mobility of the solid phase is also dependent on the degree of compaction. When a hard phase such as TiC is pressed to a high density, the particles will tend to lock each other, making rearrangement more difficult. It is thus not completely evident that a high green density will yield an improved sintered density [2]. The initial stage of densification is often finished within a few minutes [1]. Ostwald ripening and coalescence of the solid phase dominate the second stage. Both of these processes lead to an increase of the grain size, especially at high sintering temperatures. To obtain a ductile material it is necessary to maintain a small grain size, as well as close to the porosity. The dominating mechanism for pore closure is diffusion and mass transfer of the different species in the liquid phase. The main part of the densification is normally finished after around 20 min, including both the first and the second stages [1]. The third and final stage occurs when a generally solid skeleton has been formed by coalescence. At this point, densification is slow and the dominating event is grain growth. After roughly 1 h of sintering, normally no further increase in density is found [1]. In general, it can be said that if the green density, the density after compaction, is high, then the shrinkage will be lower. Typically the volume changes during LPS are around 15–20% [3]. 2.1. Experimental procedure and results for the LPS experiments The experimental procedure follows the typical manufacturing route for LPS, i.e. (i) powder characterisation; (ii) milling/sieving and mixing; (iii) powder pressing; (iv) sintering; and (v) metallographic and mechanical properties investigation. The parameters investigated were the following: (i) sintering temperatures of between 1473 and 1733 K; (ii) the compaction pressures of between 90 and 1300 MPa; (iii) volume percentages of TiC between 50 and 85; (iv) the use of a binder (ethyl alcohol); (v) the use of a protective sintering atmosphere; and (vi) sintering time. The TiC-hard phase, supplied by London and Scandinavian Metallurgical, was in the form of their XTiCTM-powder. This powder contains 20% Fe, by weight, whilst the remainder is TiC, which is present in the form of micron-sized

Fig. 1. Die geometry for the compaction of the samples.

particles. The as-received powder consists of agglomerates up to a size of 2 mm. This powder was milled, using a planetary ball mill, to reduce the sizes of the agglomerates. After milling, the agglomerates ranged from 45 to 250 mm. Sieving limited the maximum agglomerate size to 80 mm. The milled and sieved powder was then mixed with carbonyl iron of an average size of 5 mm and graphite powder. Two mixtures were made. The first mixture contained 70% TiC by volume with 1.5% C, by weight, to control the melting point of the matrix phase, whilst the second mixture contained 50% TiC, by volume and 2% C, by weight. Mixing was performed in a V-shaped blender for 120 min. After mixing, the powder was pressed into a compact with a diameter of 25 mm and a typical thickness of 4 mm. The die geometry is shown in Fig. 1. The experimental conditions can be found in Table 1, together with the experimental results. To investigate the temperature dependence on the relative density, two experimental sub-series were made, one series being without a protective atmosphere (samples 8–11), and one series with a protective atmosphere (samples 18–22). All of these samples were pressed at 148 MPa, keeping the initial conditions constant. The results are shown in Fig. 2. The samples sintered under a protective atmosphere show the highest relative density. The temperature dependence of the relative density is weak but a tendency towards an increase in relative density is seen. It can be concluded that the use of a protective atmosphere increases the relative density by approximately 10–20%. The amount of TiC particles present also has a marked effect. Not surprisingly, it can be seen that a lesser amount of reinforcement produces a greater relative density. Changing the amount of TiC from 70 to 50% increases the relative density by around 5%. The use of a protective atmosphere thus has the strongest influence. This suggests that the surface conditions of the solid and liquid phases are important for a successful initial stage consolidation. The influence of the compaction pressure was also studied. The relative density as a function of the compaction pressure, after sintering, was found to obey R ¼ 4:999 lnðP  106 Þ þ 51:015


where R is the relative density after sintering (–), and P the pressure (Pa).

P. Persson et al. / Journal of Materials Processing Technology 127 (2002) 131–139


Table 1 Experimental conditions and results for the LPS experiments Number

Sintering temperature (8C)

Sintering time (min)

Compaction pressure (MPa)

TiC (vol.%)


Protective atmosphere

Density (kg/m3)

Relative density (%)

4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25

1460 1200 1200 1200 1200 1250 1300 1350 1200 1250 1250 1200 1225 1225 1200 1250 1300 1350 1375 1300 1300 1300

30 30 30 30 30 30 30 30 30 30 30 30 30 60 60 60 60 60 60 60 60 60

148 148 148 148 148 148 148 148 140 140 90 148 148 148 148 148 148 148 148 1300 462 1231

75 75 85 80 70 70 70 70 50 50 50 50 50 50 50 50 50 50 50 50 50 50

– Y Y Y Y Y Y Y Y Y Y Y Y Y Y Y Y Y Y – – –

– – – – – – – – – – – – – – Ar Ar Ar Ar Ar Ar Ar Ar

– – – – 3460 3640 3860 2830 4000 4140 4040 4250 4400 4100 4930 4800 4840 5040 4620 5700 5020 5490

– – – – 55.9 58.1 62.3 45.7 62.0 64.2 62.6 66.0 68.3 63.6 76.4 74.4 75.1 78.1 71.6 88.2 77.8 85.2

This also implies that it will be difficult to reach above 90% relative density, under the present conditions, as illustrated in Fig. 3. Another feature was also found by measuring the size of the compact and the sintered part. Sintering did not alter the size of the part significantly. The typical linear contraction is in the order of 2–3%. This corresponds to a volume change of 4–9%, which in turn corresponds to a change in relative density of 4–10%. The possible densification under the present conditions is thus most likely limited by particle inter-locking. The parameters determining the final porosity, under the present conditions, are thus the compaction pressure and the compact green density. Metallographic analysis of the specimens after sintering reveals porosity, as indicated by the density measurements. It is also found that the wettability of the iron matrix on the

Fig. 2. Relative density after LPS as a function of the sintering temperature.

TiC particles is reasonably good, since the agglomerates are penetrates by the melt. It is also clear that the agglomerates still appear as clusters, creating an inhomogeneous microstructure on the microscopic scale, Fig. 4. This is again evidence of a limited mobility of the TiC particles by particle inter-locking. However, the microstructure on the macroscopic scale is homogeneous, giving the specimen relatively homogeneous and isotropic properties. The Vickers macro-hardness (load 10 kp, pyramid angle 1368) was also measured on the most successful samples. These results are shown in Table 2, together with one result from the SHS experiments.

Fig. 3. Relative density after LPS as a function of the compaction pressure.


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initial porosity and the liquid solubility after pressing to obtain a fully dense body flmin ¼ e0

1  XiL flmin XiL ) e ¼ 0 XiL þ e0 ð1  XiL Þ ð1  XiL Þð1  flmin Þ

0:5XiL XiL ; ¼ L ð1  Xi Þ  0:5 1  XiL 0:3XiL 0:43XiL e0 ð0:3Þ ¼ ¼ L ð1  Xi Þ  0:7 1  XiL ) e0 ð0:5Þ ¼

Fig. 4. Microstructure after LPS, showing the penetration of a cluster.

2.2. Discussion of the LPS experiments In the present study the porosity was not eliminated, there being several reasons for this. The microscopic distribution of the hard phase TiC was not significantly altered during sintering, even though the agglomerates were penetrated. This suggests that the particles within the agglomerates are locking each other as a result of the initially present TiC agglomerates and of pressing. This can be said due to the fact that no evidence of coalescence between the TiC particles was found. Particle inter-locking and particle mobility is determined by the residual porosity after compaction and by the amount of liquid present during sintering. In order to densify a compact completely, a sufficient amount of liquid phase must be present. There exists a limit for the fraction necessary for this to occur. German [1] has given this as flmin ¼ e0

1  XiL XiL þ e0 ð1  XiL Þ


where e0 is the initial porosity (–), and XiL the solubility of element i in the liquid (–). In the present case there are two distinct features. First the amount of liquid present is given by the fraction of TiC in the sample. This fraction ðflmin ¼ 0:30:5Þ is generally independent of temperature in the present temperature ranges. This gives the following relationship between the necessary Table 2 Vickers hardness of the samples made by LPS and SHS Number

Sintering temperature (8C)

Compaction pressure (MPa)

Relative density (%)

Hardness (HV)

20 23 24 25

1300 1300 1300 1300

148 1300 462 1231

75 88 78 85

103 285 189 194




Since the packing density after pressing is in the order of 0.7, the initial porosity will be of the order of 0.3, so the required solubility of TiC in the melt is 0.23–0.41. This is far from the actual value, that is of the order of 0.015 [4]. This on the other hand demands that the compact must be pressed to a relative density of 0.985–0.993 to be able to reach full density. Regarding this, it does not seem possible to reach full density, using traditional pressing and LPS. To reach the highest density possible in the system under the given conditions it is necessary to optimise the porefilling process, in order to provide conditions to create a homogeneously distributed small-sized porosity. Since the liquid is wetting the solid, any menisci between the particles will have a small radii of curvature. Park et al. [5] have shown that in the case of LPS, pore-filling will occur if and only if the radii of curvature of the menisci is equal to that of the pore to be filled. According to Park et al. [5], the meniscus radius is directly proportional to the solid grain radius. It is from this point of view extremely important to control both the size and the size distribution of the initial powder mixture to avoid combinations giving small particles and large voids. From the pore-filling behaviour, it is advantageous to have large grains and a high fraction of liquid, since these factors decrease the curvature of the liquid meniscus [5]. This is, however, not an option for the present application. The best solution for the present powder types is to reduce the agglomeration of the TiC particles further, by extensive milling, since the agglomerates are the main cause of the large-size porosity in the present case. Another solution would be to use TiC powder of a wide size distribution to reduce the size of the porosity after pressing. 2.3. SHS Reactive processing as a concept may be considered as a generic name for a large number of different processes. The first effort in reactive processing was made by Merzhanov and Borovinskaya [6], studying the combustion of Ti–B compacts. The factor that all methods have in common is that a reaction is started and the material is converted partly or fully into a compound from pure elements or compounds. The processes and reactions are conveniently classified into two main groups, convergent and divergent, which in turn may be divided into two subgroups, fuelled and self-heated

P. Persson et al. / Journal of Materials Processing Technology 127 (2002) 131–139


[7]. In a convergent self-heated reaction the entering components produce a compound without the formation of products that need separation, i.e., 3Al þ Ti , Al3 Ti


3Ni þ Al , Ni3 Al


Ti2 Fe þ 2C , 2TiC þ Fe


or similar. The first two reactions are commonly used in the XDTM-process and the last in the XTiCTM-process [7]. The starting material may be pure elements or a compound. The amount of heat is often substantial, which is why the introduction of reaction products or intermediary compounds can be made before the reaction to reduce the amount of heat released. If this is made in large portion, the heat evolved is decreased so that the heat loss from the body is greater than the rate of heat generation and the reaction will no longer be self-heated. Clearly, a wide range of different types of reactions is possible occurring as solid/solid, solid/liquid, liquid/liquid, solid/gas, liquid/gas reactions [8]. The present work is focused on the so-called solution precipitation reaction synthesis, using a self-heated convergent reaction, i.e. when a solid phase is precipitated from a solid/liquid or liquid/liquid system Ti þ C , TiC


The reaction occurring is exothermal and due to this the temperature of the compact is increased. The maximum temperature attained, Tad, is the so-called adiabatic temperature defined as [8] Z Tm Z Tad Cp ðtÞ dt þ fl DHm þ Cp ðtÞ dt (4) DHf ¼ T0


where Tm is the melting point (K), T0 the reference temperature (K), DHm the latent heat of fusion (J/kg), DHf the heat of reaction (J/kg), fl the fraction of melt present (–), and Cp the specific heat (J/kg K). In a steady combustion the reaction front moves continuously through the compact, disregarding the initial and final transients. In a non-steady combustion, the non-linear coupling between mass transport and the reaction rate causes the reaction to become oscillatory or spiral. The first is a temporal instability related to reaction kinetics whilst the latter is a spatial instability related to mass transport to the reaction front. Munir and Lai [9] suggested, schematically, regions in the adiabatic temperature and dilution plane for thermal explosion and extinction and reaction modes, Fig. 5. Reaction processing from powder materials normally generates a dispersion of extremely fine particulate typically well below 1 mm. From a geometrical point of view the reaction may occur in the bulk or start at one point and propagate though the body. In the case where the reaction is started at a single point, propagating with a planar and

Fig. 5. SHS stability diagram, after Ref. [9].

narrow reaction front, the reaction front velocity, u, is given by [8,10] sffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi      Cp Tad RTad Q u ¼ sn K0 (5) exp  RTad r DHf Q where K0 is the frequency factor, from the reaction rate expression K ¼ K0 expðQ=RTÞ (s1), R the general gas constant (J/mol K), T the temperature (K), Q the activation energy (J/mol), sn the geometrical constant (–), and r the density (kg/m3). Independent of the point of initiation of the reactions, there are some essential features. The first important step is that a substantial reaction will in most cases not occur without the formation of a liquid, due to the lack of reaction area. This is a generic feature for solution precipitation reactive synthesis. This liquid spreads and increases the liquid/solid contact area for the reaction so that the heat evolved is equal to the heat loss to the surroundings and steady state is obtained. As the liquid spreads, entrapped gas will expand while the wetting forces will contract the body, causing densification. Reaction occurs until all species are consumed or the balance between heat evolution and heat loss is such that the reaction is no longer sustained. The initial stage of spreading and reaction is similar to that of reactive brazing, which has been treated by Wang and Lannutti [11] from a phenomenological point of view. The specific area wetted is given by ! X sj X Law X @ew ¼ Lwj þ nas ms (6) @t T T s a j where ew is the specific area wetted (m2/m3), sj the surface tension (J/m2), m the chemical potential (J), n the stoichiometry coefficient (–), and L the phenomenological constant (m K/J s, K/m J s). For a binary system (M–X) with the reaction: M þ X , MX



P. Persson et al. / Journal of Materials Processing Technology 127 (2002) 131–139

the expression for the rate of wetting is given as segl

 sgs þ ssl cosðjÞ Llw  Lwgl GM @ew þ ¼ Lwj lnðaM Þ T @t R Llw  Lwgl GC lnðaC Þ þ (7) R where j is the wetting angle (rad), G the surface excess of the species (mol/m2), and s the surface tension (N/m). This is under the assumption that the MX compound is stoichiometric compound and can be treated as a pure substance, i.e. aMX ¼ 1. Clearly, the rate of wetting of a specific surface is constant as long as the composition of the wetting liquid is constant and the local specific surface of the powder is constant. If spreading essentially occurs in one direction, the spreading speed in this direction is constant. The spreading velocity was constant when one-dimensional, as described above. At the propagation front, heat generates fresh liquid that wets the material ahead of the reaction front. Provided that the mixing of the powder is homogeneous on the scale of the reaction front width, the liquid generated will be homogeneous and thus also relatively constant during reaction. The rate of spreading will also be constant. Since the front is similar to a zone-refining process, there will, however, exist an initial and a final transient, where the composition changes and the rate of penetration and of reaction will be different. Since the densification of the compact depends partly on the rearrangement of the particles and the time in the liquid state, it will be advantageous to use material combinations where spreading is faster than combustion wave propagation. This will on the other hand expand the reaction zone. In the case of a bulk reaction it would be advantageous to first create a nonreacting liquid dissolving the reactive species that subsequently reacts with another species that is still solid. In general, it is advantageous to have a matrix phase that is in liquid state of the largest possible temperature range to promote particle rearrangement after reaction. A limited porosity decreases the conductivity stabilising of the reaction due to a concentrating effect on the reaction front [12]. Porosity also decreases the reaction rate and narrows the reaction front. A large fraction of porosity will thus lead to an extinction of the reaction and thus a destabilisation of the reaction. The reinforcement conductivity relative to the matrix conductivity is a critical parameter. If the conductivity is lower than the matrix, the action is similar to that of porosity with an initial stabilisation and a final destabilisation due to the reduction of reaction heat. If the conductivity is higher than that of the matrix the only action is destabilisation due that the reaction front being cooled by an increased conductivity and a decreased amount of heat coming from the reaction. The greatest problem that these materials struggle with is undoubtedly the reduction of porosity. Unlike sintering, combustion synthesis normally exhibits expansion during processing. The origin of the porosity may be extrinsic or intrinsic having the following causes [8]: (i) extrinsic

porosity; (ii) the inheritance of the porosity of the green compact due to lack of external pressure on insufficient capillary action; (iii) heterogeneous phase formation leading to Kirkendall micro-porosity; (iv) a high reaction temperature causing ‘degassing’ of absorbed gases and evaporation of volatiles of impurities; (v) intrinsic porosity; and (vi) changes in molar volumes during reaction. LaSalvia et al. [10,13] made an experimental and theoretical analysis of the formation of TiC particles in a Ti–Ni– Mo melt by the addition of C in the form of graphite. The formation of TiC starts with the formation of a polycrystalline layer on the added C particles. This reaction, under which TiC was formed, does not start without the formation of a liquid and its capillary spreading. The reaction layer is, however, unstable and cusps are formed. Depending on the thickness of the initial carbon particle, the TiC particle may be spherodised and released from the graphite particle. The reason behind this behaviour is a minimisation of the total stress and strain energy. The origin of the stress is due to variation in the TiC phase composition. TiC is not a purely stoichiometric compound, but varies between TiC0:96 > TiCx > TiC0:46 . The cusp minimises the strain because at the cusp the strain is concentrated, unloading the rest of the material. The work of LaSalvia et al. [10,13] shows a possible route for the control of precipitate morphology and also the possibility to create small completely spherical TiC precipitates. Both the size and the morphology of the TiC particles have the potential to give a strong and positive influence on the properties. Small particles are advantageous from a ductility point of view. 2.4. Experimental procedure and results for the SHS experiments Two different types of experimental series were performed to study the consolidation of powders using SHS. In the first series the powder was vibrated in the mould to increase particle packing and the influence of mould heating and mould insulation was studied. This series will be referred to as pressure-less experiments. Furthermore, the effect of the ignition technique was investigated. In the second series a compact was produced which was reacted. The MMC to be produced in the SHS experiments was set to contain 60% TiC by volume. This means that the power mixture had the composition of Fe10Ti12C4 by weight. The powder used was carbonyl iron with an average diameter of 5 mm, titanium powder (99.5%) with a size below 200 mesh and graphite powder with an unknown size distribution. The powder was mixed using a V-shaped blender for 120 min, as in the case for the LPS experiments. In the pressure-less experiments a simple mould was constructed to yield a planar reaction front, Fig. 6. Mould heating was applied, with the powder inside, using an oxyacetylene torch. The first tests were made without insulation of the mould, resulting in specimen cracking, which seemed to arise from stresses caused by shrinkage and thermal

P. Persson et al. / Journal of Materials Processing Technology 127 (2002) 131–139


Fig. 7. A typical specimen from the pressure-less series showing cracking.

Fig. 6. Mould geometry for the pressure-less SHS experiments.

stresses. After the first trials all pressure-less experiments were made using mould insulation to prevent the samples from cracking. Leading an electrical current through the mould enabled mould heating when the mould was insulated. Ignition of the powder was made by heating the powder located in the circular part of the mould, the leftmost part as shown in Fig. 6. Two different ignition techniques were investigated, ignition using an oxy-acetylene torch and ignition by an electrical spark. In general, it was more difficult to ignite the loose powder using the gas torch due to the strong convection caused by the gas flow. In the case of electrical ignition, the powder was also lost if the voltage was increased too quickly. The proper way to ignite the loose powder was found to be to slowly increased the voltage, where at around 10 V the powder was ignited. Samples ignited with electricity showed higher densification than the samples ignited with gas. Mould preheating had little influence. More important was then the mould insulation. Samples produced in an insulated mould showed higher densification and a lower tendency for cracking. The samples were, however, still highly porous and cracking occurred to some extent in all of the samples, Fig. 7. The most important factor, however, was the powder packing before ignition. This was also the reason why the intention to continue the pressure-less experiments was abandoned and the use of a compact was chosen instead. The experimental set-up for the high-pressure experiments is shown in Fig. 8. The same die configuration as

Fig. 8. High-pressure experimental set-up.

for the LPS experiments was used for compacting the powder, Fig. 1. The experimental conditions used are shown in Table 3. Cracking of the high-pressure samples was prevented, resulting in samples of a relatively high structural integrity, Fig. 9. The green density of the powder after compaction was not measured in all samples. However, in the samples measured, the green density was around 70%. The relative density after reaction decreased, which is seemed to be more pronounced in the samples ignited with the gas torch, which indicates that residual gas might have an influence. To try to increase the density, some trials were made under vacuum. These samples showed a tendency toward a higher density, but in some cases density decrease occurred. A metallographic investigation revealed that the TiC particles formed were perfectly spherical, as in the work

Table 3 Experimental conditions and results from the SHS experiments Number



Compaction pressure (MPa)

Density (kg/m3)

Relative density (%)

1 2 3 4 5 6 7 8 9 10 11 12 13

Gas Electricity Gas Electricity Gas Electricity Gas Gas Gas Electricity Electricity Electricity Electricity

Air Air Air Air Air Air Air Air Air Vacuum Vacuum Vacuum Vacuum

20 20 40 40 60 60 70 150 180 61 61 61 61

2902 3129 2584 3165 2948 3231 2881 3003 3539 3328 2981 3312 2827

47.6 51.3 42.3 51.9 48.3 53.0 47.2 49.2 58.0 55.4 48.8 54.3 46.3


P. Persson et al. / Journal of Materials Processing Technology 127 (2002) 131–139

Fig. 9. Examples of the appearance of the high-pressure specimen after processing.

by LaSalvia et al. [10,13], Fig. 10. The particles were densely packed in the pore surfaces, indicating that there may be particle inter-locking in this case as well as in the LPS experiments. 2.5. Discussion of the SHS experiments The densification during reaction was not homogeneous in the pressure-less experiment. Due to the mould insulation it was not possible to visually observe the reaction front. However, due to the low particle packing and large porosity it is likely that the reaction may be unstable. This is further emphasised by the fact that the non-insulated experiments were highly unsuccessful in becoming fully reacted. In the high-pressure experiments the reaction front propagates more or less continuously at relatively constant velocity through the specimen, indicating a relatively stable reaction. On the surface of the specimen, marks from the

Fig. 10. An SEM micrograph showing a fracture surface. Clearly visible are spherical TiC particles on the surface of a pore.

reaction front are visible, Fig. 9, these marks indicating a slight oscillatory nature of the reaction front. This is quite plausibly due to the spreading velocity of the melt formed and the reaction front velocity not being completely matched. The adiabatic temperature for a pure mixture of Ti and C has been reported to be of the order of 3150 K [10,13]. The congruent melting point for TiC is 3349 K [14]. The TiC being precipitated can be concluded to be solid, since the adiabatic temperature will be decreased due to the dilution to 60% TiC by volume, which was the composition of the present SHS specimen. The temperature in the samples will, however, be high enough so that a significant amount of liquid will be present. This should provide a situation in which the particles formed during reaction can rearrange themselves significantly, keeping particle inter-locking at a minimum. Despite this, the density decreased in the specimen compared to the green density. A possible explanation of this is the effect of the changes in molar volume occurring on reaction. When TiC is formed from pure substances the volume decrease is approximately 24% [8]. If no particle rearrangement occurs, the porosity will increase with reduction in the density, the porosity increase being 0:6  24 ¼ 14:4%. This in turn implies that the relative density changes from approximately 70% to roughly 53%. This is also of the order of the densest specimen produced using SHS. It can thus be concluded that the intrinsic porosity formed due to volume changes in the reaction is the main contributor to the porosity in the SHS experiments. The reaction is fast and only requiring a short time at hightemperature. As the temperature is decreased the TiC will precipitate, mainly on the existing particles. This will again cause particle inter-locking, hindering rearrangement of the solid phase. The main cause of this is that the time at elevated temperature is too short to allow for the necessary motion of the particles, as well as insufficient capillary action in the absence of external pressure.

P. Persson et al. / Journal of Materials Processing Technology 127 (2002) 131–139

3. Conclusions For the manufacturing of Fe/TiC composites for armour application, the most viable route is LPS. Manufacturing of Fe/TiC composites using SHS will encounter large difficulties in handling the intrinsic porosity during the reaction. It will, however, be difficult to produce a fully dense material using LPS. In both cases the compaction has a strong influence on the resulting density after processing.

Acknowledgements Mr. Martin McGarry and Dr. Jonathan Ellis, formerly with the London and Scandinavian Metallurgical Co. Ltd., are acknowledged for their kind supply of raw material for the LPS experiments. References [1] R.M. German, Liquid Phase Sintering, Plenum Press, New York, 1985. [2] P. Persson, Processing and properties of TiC/Fe composites, Thesis, Department of Materials Processing, Royal Institute of Technology, Stockholm, Sweden, 1997.


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