Processing and properties of particulate reinforced steel matrix composites

Materials Science and Engineering A246 (1998) 221 – 234

Processing and properties of particulate reinforced steel matrix composites E. Pagounis *, V.K. Lindroos Laboratory of Physical Metallurgy and Materials Science, Helsinki Uni6ersity of Technology, 02150 Espoo, Finland Received 15 May 1997; received in revised form 27 August 1997

Abstract Metal matrix composites are an attractive choice for aerospace and automotive applications because of their high stiffness-toweight ratio. Composites with aluminum and magnesium matrices have been investigated extensively, while less work has been carried out on steel matrix composites. In the present study the hot isostatic pressing (HIP) process of steel matrix composites is described, and the factors influencing the reinforcement distribution, interface processes, as well as the mechanical and corrosion properties, are revealed. Both stainless steels and tool steels were used as the matrix material, and the particulate reinforcements were Al2O3, TiC, Cr3C2, or TiN. The results are compared with those of the corresponding unreinforced alloys and also with those of aluminum and magnesium matrix composites. It was found that the incorporation of a relatively low volume fraction of ceramic particulate reinforcements significantly increases the wear resistance of the steel matrices, without deteriorating the corrosion properties. On the other hand, reductions in the tensile strength, ductility and toughness were observed. The superaustenitic stainless steel–TiN and hot work tool steel – Cr3C2 composites may offer the best combination of properties. © 1998 Elsevier Science S.A. All rights reserved. Keywords: Steel matrix composites; Stainless steel; Tool steel; HIP; Mechanical properties; Corrosion properties

1. Introduction The incorporation of ceramic reinforcements into metal matrices to produce composite materials with improved properties has been the subject of intensive investigation during the past three decades. Metal matrix composite (MMC) technology improves the elastic modulus, wear resistance and strength of unreinforced metals and alloys. The first effort has been directed towards the development of high performance composites with very high strengths and moduli, for use in specialized aerospace applications. These materials were composed mainly of a light metal matrix (e.g. Al-, Mg-, or Ti- alloy) reinforced with aligned continuous fibers (e.g. Al2O3, SiC, B). However, the cost-effective production of these composites has restricted their commercial applications significantly. Therefore, an attempt * Corresponding author. Current address: Nokia-Maillefer Oy, P.O. Box 44, 01511 Vantaa, Finland. E-mail: emmanouel.pagounis@ nokia-maillefer.com 0921-5093/98/$19.00 © 1998 Elsevier Science S.A. All rights reserved. PII S0921-5093(97)00710-7

to assess the potential to develop low cost aluminum or magnesium matrix composites reinforced with inexpensive ceramic particulates was initiated. These composites have been investigated extensively [1–4] and are presently of commercial interest to the automotive and aerospace industries. The ability to use iron and its alloys as the matrix material in composite systems is of great importance because it is the most widely used metallic material with a variety of commercially available steel grades. Additionally, steel has higher stiffness, strength and toughness compared with aluminum, it shows good machinability (in the annealed condition) and weldability, and some grades offer improved corrosion resistance. However, until recently, little attention has been given to the development of steel matrix composites. The most important of them are tool steels and high speed steels (HSS) reinforced with various carbides (TiC, WC, VC, NbC, Cr3C2, etc.) [5–7], but stainless steel matrix composites containing Al2O3 or Y2O3 have

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Table 1 Chemical composition of steel matrices used for composite production Grade

Chemical composition (wt%)

Powder size (mm)

Stainless steels (SS) Ferritic – austenitic (duplex) Austenitic (316L) Superaustenitic (654 SMO)

26 Cr–6.6 Ni–3.3 Mo–1.1 Mn – 2.5 Cu – 0.36 Si – 0.02 C – 0.22 N – Fe balance 17 Cr–12 Ni–2.5 Mo–1.6 Mn – 1.0 Si – 0.1 C – Fe balance 24 Cr–22 Ni–7.3 Mo–3.5 Mn – 0.5 Cu – 0.3 Si – 0.01 C – 0.5 N – Fe balance

B75 (fine) 75 – 250 (coarse) B75 B75

Tool steels High speed steel Hot work steel White iron

4.2 Cr–5.0 Mo–6.3 W–3.1 V – 1.3 C – Fe balance 4.9 Cr–1.6 Mo–0.6 V–0.35 C – Fe balance 26 Cr–2 C–Fe balance

B106 B106 B150

654 SMOTM is a trademark of Avesta AB, Sweden.

also been reported [8 – 10]. The unique properties of these composites have recently attracted great attention [11–15]. Steel matrix composites have been proposed for use as wear and corrosion resistant parts in the chemical and process industry, or as substitutes for the more expensive cemented carbides (e.g. WC – Co) [16]. The powder metallurgy is an attractive processing route for steel matrix composites because a wide range of reinforcement volume fraction and size can be used and the distribution of the embedded particles is more uniform. In addition, extensive metal – ceramic interface reactions can be avoided by using a solid state process. Compared with their unreinforced alloys, steel matrix composites offer higher hardness and wear resistance and their elastic modulus is the highest among all machinable materials and materials which can be hardened [16]. On the other hand, as in all metal matrix composites, reductions in the fracture properties (ductility, toughness) are to be expected. Finally, the corrosion and oxidation properties of these composites are also attractive [17,18]. The purpose of the present paper is to present the powder metallurgical production of steel matrix composites and to reveal the factors influencing their mechanical and corrosion properties. A variety of steel matrices and ceramic reinforcements have been used for composite production in order to provide more insight on the limited information available today for this class of materials. The process-structure-property relationship is also discussed and potential applications of steel matrix composites are proposed.

2. Experimental procedures

2.1. Materials Two groups of steel grades were selected for the matrix material of the composites examined in the present study (Table 1). The first was composed of three highly alloyed powder metallurgical stainless steels and the second of three wear resistant tool steels.

All grades were gas-atomized powders suitable for a hot isostatic pressing process. Various ceramic particulates were appraised for incorporation in the steel matrices; these included oxides, carbides and nitrides (Table 2). The particulate reinforcements were an irregular shape with an aspect ratio close to unit. In order to examine the influence of the reinforcement particle size on the microstructure and properties of the composites, fine and coarse reinforcements were used.

2.2. Processing The materials (unreinforced alloys and composites) examined in the present study were processed by hot isostatic pressing (HIP). This technique offers an advantage: in a commercial steel-processing route the only additional step for producing steel matrix composites is the mixing with the ceramic powder. Additionally, nearly fully densified materials and near net shapes can be achieved. Reference unreinforced alloys and composites containing up to 30 vol.% reinforcements were produced. The mixing of the steel and ceramic powders was carried out in a Turbula mixer, with 4 wt% ethanol to prevent segregation of the lower density ceramic particles (composites containing Cr3C2 reinforcements were dry mixed because the two powders have similar densities). After mixing, the powders were filled in mild steel capsules and dried for 18 h. The capsules were then sealed by welding, evacuated at 500°C for 3 h, and inserted into the HIP equipment. The HIPing parameters for all the materials were 1180°C, 100 MPa pressure and 3 h holding time. After HIPing the materials were heat treated inside the capsule to avoid decarburization. The heat treating cycles for the composites were identical to those of their corresponding unreinforced alloys, which are summarized in Table 3. Stainless steel-based materials were heat treated in order to avoid the formation of the embrittling s-phase, which may precipitate during the very slow cooling after the HIPing process.

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Table 2 Ceramic particulates used as composites’ reinforcement Reinforcement Particle size (mm)

Fine Coarse

Al2O3

TiC

Cr3C2

TiN

44 – 74 105 – 149

5.6 – 22.5 44 – 105

5 – 25 45 – 125

5 – 36 20 – 36

2.3. Materials characterization and testing After processing the materials were sectioned, polished and examined by means of optical microscopy (OM), scanning electron microscopy (SEM) and transmission electron microscopy (TEM). The most important microstructural features for investigation were the reinforcement distribution, metal – ceramic interfacial region, and the subsequent microstructural development (e.g. dislocations, precipitates). Tensile tests were carried out to evaluate the strength and the ductility of the stainless steel matrix composites. The tests were implemented according to the European standard EN 10002 using a Fiskars 40 Mp machine equipped with a computer data acquisition system (Matertest Oy, Finland). Cylindrical specimens with a cross-sectional diameter of 7 mm, a gage length of 40 mm, and threaded ends were used. The rate of straining did not exceed 0.008 s − 1. The reported ultimate tensile strengths and elongations to failure were based on the average of three tests. The three-body abrasion wear tests were carried out according to the ASTM G 65-91 standard procedure with a rubber wheel abrasion equipment. Quartz sand was used as the abrasive and the total wear distance was 1432 m. Details on the procedure are given elsewhere [12]. The impact toughness of the materials was measured with a pendulum impact testing machine (standard Losenhausenwerk test equipment), using unnotched samples with the size of 5×10 ×55 mm3. Impact test measurements were chosen instead of the more widely used fracture toughness tests for comTable 3 Heat treating cycles for the various steel grades Grade

Austenitizing treatment

Subsequent treatment

Ferritic – austenitic (duplex) Austenitic (316L) Superaustenitic (654 SMO) High speed steel Hot work steel

1160°C, 1 h, PAC

—

1100°C, 1 h, PAC 1180°C, 1 h, PAC

— —

1180°C, 10 min, PAC 1030°C, 20 min, air cooling 1160°C, 1 h, air cooling

3×550°C, 2 h 540°C, 1 h+2× 500°C, 1 h 510°C, 2 h

White iron

PAC, pressure air cooling.

posites, because they are a better representation of the deformation rates in real applications where toughness is an important material parameter [13,19]. The average of four tests was used for the abrasion wear loss and impact energy of the materials. Because stainless steel matrix composites are mainly proposed as wear and corrosion resistant parts, their corrosion properties were also examined. Both immersion and electrochemical polarization measurements were carried out in order to reveal the corrosion susceptibility and the pitting resistance of the composites. Before testing, the samples were polished to a 1 mm diamond paste finish. At least two samples were run for each test and the average values determined the results. The pitting resistance of the composites was measured according to the ASTM G 48-76 standard procedure. Specimens with known weight and dimensions were immersed in an aqueous ferric chloride solution (6 wt% FeCl3 in distilled water) at 50°C and kept for 72 h. After the test the specimens were rinsed with distilled water and scrubbed with a nylon bristle bush under running water to remove the corrosion products. They were then ultrasonically cleaned with acetone, dried and finally the weight loss was measured. The susceptibility of the stainless steel matrix composites on localized corrosion was electrochemically measured according to the standard procedure ASTM G 61-86. The measurements were carried out using an Avesta cell [20] with an aqueous 3.5 wt% NaCl solution (pH 6.7–7.2) at room temperature. The anodic polarization curves of the designed composites were obtained potentiodynamically using an AutoTafel equipment (ACM Instruments, UK). The potential in these measurements was increased continuously at the standardized scanning rate of 0.167 mV s − 1. The anodic polarization curves of the potential versus log current were then recorded automatically by means of a logcurrent converter and a X–Y recorder. All potentials were measured versus a saturated calomel reference electrode (SCE).

3. Results and discussion

3.1. Structures The distribution of the added ceramic particles is of great importance to the mechanical properties of the

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Fig. 1. Microstructures of the composites showing the reinforcement distribution after using different matrix – reinforcement powder size ratio, at a reinforcement volume fraction of 20%: (a) coarse (75–250 mm) – fine (44 – 74 mm), (b) fine (B75 mm) – fine (44 – 74 mm), (c) coarse (75–250 mm) – fine (5.6 – 22.5 mm), (d) fine (B 75 mm)–fine (5.6–22.5). Magnification marker in all photographs the same as in (a).

composite. Clusters of reinforcing particles, resulting from insufficient mixing or electrostatic phenomena, deteriorate the properties of the material by serving as sites for microporosity and crack initiation. On the other hand, a homogeneous distribution of the reinforcements ensures isotropic properties and uniform

distribution of stresses into the material. Typical microstructures of the HIPed composites showing the reinforcement distribution are presented in Fig. 1. As can be seen, the distribution varies with different combinations of matrix and reinforcement powder size. In composites processed with relatively coarse matrix pow-

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der size the reinforcement distribution is inherently uneven. The ceramic reinforcements surround the coarser matrix powder and local agglomeration takes place at the junctions of three or four matrix particles (Fig. 1(a), (c)). On the other hand, the use of fine matrix powder size results in a more uniform reinforcement distribution (Fig. 1(b), (d)). A detailed investigation on the consolidation behavior of composite powders during HIPing has been carried out by the authors elsewhere [21]. A very important factor influencing the structure and mechanical properties of composite materials is the metal–ceramic interface [14,22]. In the steel matrix composites examined in the present study, reinforcements with different chemical behavior in contact with iron alloys have been used, therefore, different interfaces should be expected. Alumina is an ionic ceramic and difficult to wet by metals, since its electrons are tightly bound and its surface represents large discontinuities in charge [23]. In the HIPed steel matrix composites reinforced with

Fig. 2. (a) Optical micrograph showing good wettability of irregular jagged Al2O3 particles by the steel matrix. (b) High magnification TEM micrograph revealing the clean steel–Al2O3 interface.

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Fig. 3. TEM micrograph of a TiC containing stainless steel matrix composite revealing spherical precipitates at the interface (arrows) and an increased dislocation density in the austenitic matrix.

Al2O3 particles good wettability was observed, regardless of the iron alloy matrix, but there were no interface reactions (Fig. 2). As can be seen in Fig. 2(a) although the alumina particles have an angular shape they have been wet quite well by the steel matrix. This should be attributed to the alloying elements present in the steel matrix. Fischmeister et al. [24] examined the effects of alloying elements on the structural stability and cohesion between phases in the Al2O3 –steel system. They reported that elements such as Cr, Mo, Mn and Ni increase the wettability and work of adhesion between the two phases. Howe [25] demonstrated that alloying elements with greater affinity to oxygen than iron segregate at the Al2O3 –steel interface, and decrease the interfacial energy according to Gibbs adsorption isotherm. Moreover, the HIPing process eliminates problems such as wettability, evaporation of components or sintering activity, encountered frequently in conventional liquid phase sintering [26]. However, no interface processes (e.g. diffusion alloying, reactions) took place in the Al2O3 reinforced composites, as TEM observations confirm (Fig. 2(b)). Titanium carbide is thermodynamically stable when in contact with iron alloys during sintering or HIPing and therefore, is suitable to reinforce steel matrix composites [12,14,16]. Fig. 3 presents a TEM micrograph of the metal–ceramic interface in a TiC containing stainless steel matrix composite. The iron alloy matrix has wet the TiC particle and small spherical precipitates have formed in the TiC side of the interface. These precipitates were found by high-resolution energy dispersive spectroscopy to be Fe- and Cr-rich.

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Naidich [27] and Ku¨barsepp [28] have studied the wettability of TiC by iron alloys. Naidich studied the contact angle u at various atmospheres and Ku¨barsepp in vacuum. The latter also examined the influence of alloying elements such as Ni, Cr, and Si on the wetting behavior. According to his results, the increase of Ni content up to 12 wt% in Fe – Ni alloys leaves the wetting almost unaffected. Provided the Ni content exceeds 15 wt%, it will show interfacial activity and the wettability increases (work of adhesion increases). Chromium hardly influences the wetting, while the increase of Si content (which is a well-known u reducer) results in the diminishing of u. The rise of C content in the iron alloy results in the worsening of wetting, while an increased Al content reduces u. As it has been mentioned, TiC particles are usually thermodynamically stable in an iron alloy matrix at the sintering or HIPing temperatures, and no extensive solubility takes place. Ku¨barsepp reported [28] that the solubility of TiC in Fe during sintering is practically negligible (0.5–1%), increasing as the Ni content in the iron alloy increases or the Ti and C contents decrease. This has also been reported in the experimental results of Bolton and Youseffi [29]. They only found some diffusion of iron from the steel matrix into the TiC particles; Ku¨barsepp reported that the solubility of Fe in TiC could be up to 2%. This limited solubility may have caused the formation of the (Fe, Cr)-rich precipitates at the interface, shown in Fig. 3. The thin (B 200 nm) precipitate-rich interface layer can be beneficial for the strength of the interface, as results in Al–SiC composites confirm [30]. The Cr3C2 shows a dual chemical behavior in contact with iron alloys. When it is incorporated within a highly alloyed stainless steel matrix it reacts heavily, forming a thick interface reaction layer (Fig. 4). On the other hand, within a lower alloyed tool steel matrix no

Fig. 4. Optical micrograph showing extensive interface reaction in a stainless steel – Cr3C2 composite.

Fig. 5. Influence of reinforcement volume fraction on the tensile strength of stainless steel matrix composites: (a) duplex–Al2O3, (b) duplex – TiC, (c) duplex – Cr3C2, (d) 654 SMO – TiN.

interface reactions take place as reported in the authors’ previous work [15]. It seems, therefore, that the presence of a substantial amount of alloying elements in the steel matrix enhances its reactivity in contact with ceramic reinforcements.

3.2. Properties of stainless steel matrix composites Fig. 5 shows the influence of the reinforcement volume fraction on the tensile strength of stainless steel matrix composites. As can be seen, the incorporation of ceramic particulates decreases the strength of the stainless steel. The limited data currently available on the strength of steel matrix composites indicate the same trend [7,31]. However, this behavior is in contrast to the results reported for aluminum and magnesium matrix composites, where an increased strength was reported compared with that of the unreinforced alloy [1–4]. Therefore, the strengthening mechanisms, which have been proposed to take place in those composites, may not be applicable in composites with a steel matrix. During the past decade a large number of investigations has been carried out to reveal the strengthening mechanisms of metal matrix composites, and both continuum and micromechanical models have been developed [32]. As a result of these investigations the mechanisms, which may contribute to the strengthening of a composite, have been deduced: (a) Load transfer from the matrix to the reinforcement via shear stresses at the interface between the components (shear-lag theory) [33]. (b) Increased density of dislocations generated during cooling because of the differential thermal contraction between matrix and reinforcement [34]. (c) Artificially reduced interparticle spacing due to the creation of Orowan loops (Orowan strengthening) [35].

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(d) Increased initial work hardening rate of the composite due to the presence of the ceramic particles [32]. All the preceding strengthening mechanisms, from a theoretical point of view, also seem suitable to take place in steel matrix composites. For example, loadbearing capacity of the ceramic reinforcement undoubtedly occurs mainly during the period of elastic deformation. Additionally, an increased dislocation density arising from the thermal mismatch straining of the matrix may also be plausible, as TEM observations confirm (Fig. 3). On the other hand, only the composites reinforced with 10 vol.% reinforcements possessed a work hardening coefficient, while for higher volume fractions rupture of the specimen occurred, usually before yielding. Humphreys et al. [32] and Manoharan and Lewandowski [36] reported that the addition of a brittle ceramic in a high strength alloy might decrease its tensile strength. Indeed, experimental results have proven that strength increase compared with the unreinforced alloy takes place for composites with low strength (100–400 MPa) 1050, 6061 or 2024-Al matrices [32,37–39]. However, in an Al – Zn – Mg alloy (7075-series) with a tensile strength of 638 MPa great strength reductions were observed after the incorporation of ceramic reinforcements [36]. The same behavior may also be applied with the high strength stainless steel matrices examined in the present study. One plausible explanation is that as composites with a high strength matrix are strained, the stresses on the reinforcement become large. Fracture of the reinforcement can then occur, particularly in the presence of a pre-existing flaw in the reinforcement formed probably during earlier processing. In fact, it has been pointed out by a number of authors that high matrix strength results in more particle-related damages, while composites with lower matrix strength exhibit a smaller percentage of reinforcement damage at the fracture strain [40,41]. Once the reinforcement fractures, the net load-carrying capacity of the composite decreases, therefore, strength might decrease. It is also possible that the thermoelastic mismatch between reinforcement and matrix leads to a large stress concentration near the reinforcement. Provided no sufficient plastic relaxation takes place, the matrix in that region fails prematurely during straining. This may be the case for steel matrix composites where large temperature differences are encountered during processing or heat treatment. Furthermore, this interpretation is supported by the experimental results of Bendixen and Mortensen in HIPed steel–Al2O3 composites [42]. In lower strength matrices, the stresses reached locally might not be large enough for either of the preceding effects to occur, leading to strength improvements through the processes (a)–(d) outlined earlier.

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The influence of the reinforcement particle size on the strength is presented in Fig. 6. The composites reinforced with coarse ceramic particles show higher tensile strength and this effect is more evident in the low reinforcement volume fraction (Fig. 6(a)). It appears that fine particles tend to agglomerate and this affects the tensile behavior of the composite. In regions of local reinforcement clusters the matrix probably does not have sufficient internal ductility to redistribute the very high-localized stresses, therefore, the composite fails before being able to reach stable plastic flow and normal ultimate strength [36]. In composites with a higher reinforcement volume fraction the difference in the ultimate tensile strengths is not significant because the extent of local agglomeration of particles may be similar, while porosity effects should also be considered [21]. Typical stress-strain curves for the stainless steel matrix composites are plotted in Fig. 6(b). These curves clearly show the great influence of the reinforcement

Fig. 6. (a) Influence of the reinforcement particle size on the tensile strength of duplex stainless steel matrix composites reinforced with 10 and 20 vol.% Al2O3. (b) True stress versus strain curves for 654 SMO – 10 vol.% TiN composites reinforced with fine and with coarse particles.

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particle size, with the corresponding effect on the distribution, on the tensile ductility of the composite. In the absence of fine particles the reinforcement distribution became more uniform, which increased the elongation to failure of the composite significantly. The beneficial influence of coarse reinforcement particle sizes on the tensile properties of steel matrix composites is in contrast to the results reported so far for aluminum matrix composites [39,43]. In those materials, fine particle sizes resulted in improvements of the strength and ductility of the composite due to the more favorable conditions for the mechanisms (a) through (d) to take place. Scanning electron microscope observations on the fracture surface of tensile specimens revealed that fracture of the steel matrix composites occurred either through particle cracking or interfacial decohesion (Fig. 7), while matrix cracking was not visible on the micrographs. As mentioned earlier, particle-related damages is the fracture mode observed in composites with a high strength matrix [40,41]. Particle cracking was mainly observed in composites reinforced with coarse particles, while interfacial decohesion in composites reinforced with fine or with Al2O3 particles. The latter is an indication of poor bonding strength of the alumina particles. It is interesting to note a cracked Al2O3 particle revealed on the fracture surface of the composite in Fig. 7(a). Given the weak particle–matrix interface, it seems most plausible that this particle was damaged during fabrication of the composite and subsequently fractured during straining. This observation supports the interpretation given in the previous paragraph for the decreased strength of steel matrix composites. Further experiments and observations are planned (e.g. on sectioned samples below the fracture surface) to reveal the role of the fractured particles in the decreasing strength of steel matrix composites. In composites with a low reinforcement content a ductile fracture mode was observed, with a fine lacy dimple network in the matrix (Fig. 7(a), (b)). The specimens exhibited a smooth chisel-point shear fracture, but without the reduction in the cross-sectional area observed in the unreinforced alloys. At reinforcement contents of 20 – 30 vol.%, the fracture became flat and granular across the entire area, with no evidence of a chisel-point shear lip (Fig. 7(c)). Composites exhibiting this type of fracture mode failed in a brittle manner, with failure strains even less than 2%. Fig. 8 presents the elongation to failure of stainless steel matrix composites as a function of the reinforcement volume fraction. The decrease in the ductility of the matrix alloy after the incorporation of ceramic particles is consistent with previous experience on aluminum and magnesium matrix composites [1 – 4]. The main aim of incorporating ceramic reinforcements in stainless steel matrices is to increase their wear resistance. Wear problems are frequently encountered

Fig. 7. SEM fractographs of stainless steel matrix composites revealing different fracture modes: (a) particle – matrix interface decohesion (arrows) in a 10 vol.% Al2O3 containing composite, (b) particle cracking in a 10 vol.% TiC containing composite, and (c) cleavage fracture in a 30 vol.% Al2O3 containing composite.

in the chemical, pulp and offshore industry, where highly alloyed stainless steels are widely used. As Fig. 9 demonstrates, the addition of ceramic particulate reinforcements significantly decreases the abrasive wear loss

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Fig. 8. Elongation to failure versus reinforcement volume fraction for (a) 654 SMO – TiN and (b) duplex–Al2O3 composites.

of the stainless steels. Over 30% reduction in the wear loss takes place in most steel grades after incorporating only 5 vol.% reinforcements. The wear loss decreases with the increasing reinforcement volume fraction because extensive ploughing and cutting of the matrix by the abrasive particles is reduced by the presence of the harder ceramic particles. It is interesting to observe that composites reinforced with Al2O3 particles show the lowest wear resistance, particularly at the higher volume fractions. This may be attributed to the weak interface bonding of Al2O3 with the steel matrix, which causes spalling of the particles during wearing. Spalled particles increase the wear loss because they can contribute abrasively in addition to the quartz particles. At any volume fraction of reinforcements the austenitic 316L stainless steel matrix composite, which has the softest matrix (85 Rockwell B), exhibits the highest wear loss. On the other hand, the duplex – 30 vol.% Cr3C2 and the 654 SMO – 30 vol.% TiN composites

Fig. 9. Abrasive wear loss versus reinforcement volume fraction for various stainless steel matrix composites: (a) 316L–Al2O3, (b) duplex – Al2O3, (c) duplex–Cr3C2, and (d) 654 SMO–TiN.

Fig. 10. (a) Influence of nature and particle size of the reinforcement on the corrosion weight loss of duplex stainless steel matrix composites containing 10 vol.% ceramic particles. (b) Corrosion weight loss versus reinforcement volume fraction for (i) duplex –Al2O3, and (ii) 654 SMO – TiN composites.

have wear resistance comparable with or exceeding that of HIPed high speed steels. Stainless steel matrices used in the present study comprise some of the most corrosion resistant materials available. Particularly the superaustenitic steel 654 SMO has been reported to be the most advanced stainless steel grade [44]. On the other hand, the duplex stainless steel has a pitting resistance equivalent (PRE) number of 41–42 and is very resistant against pitting corrosion [45]. It is, therefore, of great interest that the addition of the ceramic particles will not significantly deteriorate the excellent corrosion properties of the stainless steels. The influence of the nature and particle size of the reinforcement on the corrosion weight loss of duplex stainless steel matrix composites is presented in Fig. 10(a). As can be seen, composites reinforced with Al2O3 particles present the lowest weight loss in the ASTM G 48-76 immersion test. Taking into account that the metal–ceramic interface may be a main source of pitting initiation sites, due to galvanic coupling between dissimilar materials, the clean and featureless Al2O3 – steel interface (e.g. Fig. 2) provides an important ad-

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vantage in achieving better corrosion resistance. In addition, composites reinforced with coarse particles have a slightly better corrosion resistance, which may be attributed to the less interface area exposed to corrosion attack, and to the more uniform reinforcement distribution. The results presented in Fig. 10(b) demonstrate that at the lower reinforcement volume fractions the corrosion resistance of steel matrix composites is similar to that of their base alloy. After a threshold of reinforcement volume fraction, however, accelerated corrosion takes place with a corresponding sharp increase in the material weight loss. In the 654 SMO – TiN composites this threshold is at 20 vol.%, while it is at 10 vol.% for duplex–Al2O3. Examination of the corrosion specimens in the SEM revealed that at lower reinforcement volume fractions the corroded surface of the composites was similar to that of the unreinforced alloy, with little or no breakdown of the passive layer. However, the presence of an increased amount of reinforcements might have locally weakened the passive layer of the composites, which led to a local attack of the corrosive environment. The weakening of the passive layer of the composite due to the reinforcements can be explained in three ways. Firstly, the existence of thermal residual stresses in the composite may cause the formation of a passive film with structural imperfections and lower adhesion to the underlying bulk material. The residual stresses arising from the differential thermal contraction of the constituents are higher in TiC containing steel matrix composites, as recent works of the authors confirm [14,46]. Secondly, the passive film in the vicinity of the reinforcements may be imperfect compared with that in other areas, as the composition of the matrix may have been shifted by mutual diffusion or chemical reaction. This may be the case mainly of composites containing Cr3C2 particles. Finally, the shape, size and distribution of the reinforcements can cause sharp changes on the surface contour, e.g. during mechanical polishing, and the passive film formed may be uneven and therefore, more vulnerable to breakdown. The breakdown of the passive film in composites with a reinforcement volume fraction above the threshold caused severe attack of the FeCl3 corrosive environment. The passive film on attacked samples broke down in some specific areas by the action of the Cl- ions, and the bulk material beneath dissolved. The SEM micrograph of Fig. 11(a) shows the breakdown of the passive film with corresponding material dissolution and large pit formation, in a duplex – 30 vol.% Cr3C2 composite. The large pits observed were created in areas of local reinforcement agglomerates, because the material in those areas was loosely connected due to insufficient densification [21]. Fig. 11(b) shows the breakdown of the passive layer in a 316L matrix com-

posite after incorporating only 5 vol.% TiC particles. On the other hand, the 654 SMO grade was totally immune in the aggressive corrosion environment of the immersion test, and this behavior could be maintained even in composites containing 20 vol.% TiN reinforcements (weight loss B 0.4 mg cm − 2). Polarization tests were carried out in the 654 SMO– TiN composite to confirm the excellent corrosion behavior of this material in the immersion test. In these tests the difference between the pitting and the protection potential (DE =Epit − Eprot) of the materials was determined from the anodic potentiodynamic polarization curves. A smaller difference in the potentials indicates lower susceptibility to localized corrosion [47]. The results of the tests are summarized in Table 4, where results for the conventional 316L grade are also reported. The measurements clearly demonstrate that all the composites, regardless of the reinforcement volume

Fig. 11. (a) SEM micrograph of the corroded surface of a duplex–30 vol.% Cr3C2 composite showing breakdown of the passive film (A), areas of material dissolution (B), and large pit formation (C). (b) SEM micrograph showing breakdown of the passive layer in a 316L – 5 vol.% TiC composite.

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Table 4 Results of anodic polarization measurements for the corrosion potential (Ecorr)a, pitting potential (Epit), protection potential (Eprot), and passivation current density (ipass) Material

Ecorr (mV)

Epit (mV)

Eprot (mV)

DE b (mV)

ipass (mA cm−2)

654 – 10 vol.% TiN fine 654 – 20 vol.% TiN fine 654 – 30 vol.% TiN fine 654 – 10 vol.% TiN coarse 654 – 20 vol.% TiN coarse 654 – 30 vol.% TiN coarse Unreinforced 654 SMO Unreinforced 316L

−72 −60 −9 −52 −54 −42 −96 −115

929 887 879 931 874 875 968 716

922 873 861 924 868 864 954 86

7 14 18 7 6 11 14 630

0.0058 0.0064 0.0061 0.0061 0.0074 0.0074 0.0027 0.0024

a b

The reported potentials are versus SCE. DE= Epit−Eprot.

fraction and size, have similar or slightly lower corrosion susceptibility compared with the unreinforced alloy. The corrosion resistance (as measured by the DE difference) of both the unreinforced alloy and the composites in the polarization test was excellent and significantly higher than that of the well-known 316L austenitic steel. The results of the polarization tests indicate that variations in the reinforcement volume fraction or size have a marginal effect on the susceptibility of the composites to localized corrosion, because the DE differences between the materials were practically negligible. It is interesting to observe that even the 654 SMO–30 vol.% TiN composites were practically immune in the corrosive environment. This is in contrast to the results of the immersion tests, which clearly demonstrated that changes in content, size and distribution of the reinforcing particles can dramatically shift the corrosion behavior of the composite (e.g. Fig. 10). Despite the fact that different corrosive environments were used in the two tests, this discrepancy indicates that additional work is needed to fully understand the corrosion behavior of stainless steel matrix composites.

was achieved by adding only 5 vol.% Cr3C2 particles. The percentage of the reduction decreased when the matrix material became more wear resistant. In composites with high speed steel and with white iron matrix the addition of ceramic particles decreased the wear loss until a certain volume fraction of reinforcements, after which the wear loss increased. On the other hand, in composites with a hot work steel matrix, the wear loss decreased continuously with the increasing reinforcement volume fraction, as demonstrated also in stainless steel matrix composites (Fig. 9). The different wear behavior should be attributed to the different microstructures of the matrix materials. Powder metallurgical high speed steels and white irons contain a high amount of primary and eutectic carbides with the size of B 10 mm in a predominantly martensitic matrix [12,49]. The carbides in the high speed steel of the present study were found to be mainly Mo2C, WC, VC and Cr3C2, while in the white iron they were (Cr4Fe3)C3 [12]. Their content was up to 26 vol.%, as measured by the point-count method. On the other hand, hot work steels and stainless steels consist of a

3.3. Properties of tool steel matrix composites Composites with a tool steel matrix and ceramic particulate reinforcements bring new possibilities in the production of inexpensive materials for wear resistance applications [12,16]. They can offer high hardness, which reduces the depth of penetration of abrasive grits, and sufficient toughness, which avoids brittle fracture [48]. Because these materials are potential candidates for wear resistant parts their abrasion wear resistance and impact toughness were examined. The results presented in Fig. 12 demonstrate that great reductions in the abrasion wear loss of tool steels take place after the incorporation of ceramic particle reinforcements. Particularly in composites with a hot work tool steel matrix a 75% reduction in the wear loss

Fig. 12. Abrasive wear loss versus reinforcement volume fraction for various tool steel matrix composites: (a) hot work steel–Cr3C2, (b) white iron – TiC, and (c) high speed steel – TiC.

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Fig. 13. (a) Impact energy versus reinforcement volume fraction for hot work steel matrix composites reinforced with fine and with coarse Cr3C2 particles. (b) SEM fractograph of a hot work steel – 10 vol.% Cr3C2 composite, revealing fractured Cr3C2 particles (arrows) and a micron sized dimple morphology in the matrix.

martensitic and a ferritic or austenitic matrix respectively, in which a low fraction of second phases (e.g. carbides, intermetallics, Cu-precipitates) with the size of B 1 mm is present [15,50]. In the composites with the high amount of in situ carbides in the matrix, the additional incorporation of ceramic particles increases the abrasion wear loss after a certain reinforcement volume fraction. This happens because the matrix is not able to continuously support a high fraction of second phase particles during the wear test, and spalling of the reinforcements takes place. A detailed SEM study of wearing surfaces, carried out by the authors in their previous work, supports this interpretation [12,13]. The influence of the reinforcement volume fraction on the impact toughness of hot work tool steel matrix composites is shown in Fig. 13(a), where the effect of the reinforcement particle size is also presented. A significant drop in the fracture resistance of the steel matrix occurs after the addition of ceramic particulates, and this effect appears to be strongest at the lower volume fractions. Similar results have been reported for

aluminum matrix composites [37,51]. The presence of the brittle ceramic reinforcements increases the fracture initiation sites and the fracture path propagates either through the ceramic particle or the metal–ceramic interface. A SEM fractograph of the hot work steel matrix composite containing 10 vol.% coarse Cr3C2 particles is presented in Fig. 13(b). As can be seen the fracture path proceeded through the particles, the fractured parts of which are still present on the surface. A micron sized dimple morphology is visible, indicating some ductility in the matrix. The volume fraction of Cr3C2 on the fracture surface was measured to be 14%, i.e. higher than that incorporated during mixing, which means that the fracture path proceeded in areas of reinforcement clusters. This should be attributed to the internal tensile stress field within the composite, which attracts the crack front towards particle clusters [52]. In aluminum matrix composites Flom and Arsenault [53] have demonstrated that 1.33 times more reinforcements can be found on the fracture surface than on any random plane cut through the composite. The results presented in Fig. 13(a) indicate that the reinforcement particle size does not affect the impact toughness of steel matrix composites. This contrasts the void nucleation and growth theory (VNG) which is well established for ductile materials [54,55] and aluminum matrix composites [36]. It predicts that particle size and spacing between void nucleating particles are critical microstructural parameters controlling the toughness of a given material. In the steel matrix composites of the present work no void formation associated with the ceramic reinforcements was observed, therefore, the VNG theory does not appear to be applicable. On the other hand, in aluminum matrix composites the increase of the reinforcement particle size increases the toughness [56,57]. This was attributed to the larger sizes of voids nucleated by larger particles, and to the increased interparticle spacing. The positive influence of voids formed during fracture on the toughness of MMCs is well-established [58,59].

Fig. 14. Abrasive wear loss versus hardness for unreinforced steel matrices ( ) and composites (").

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In Fig. 14 results of abrasion wear loss and hardness tests for the tool steel matrices and their composites are summarized. In the unreinforced tool steels the abrasion wear loss decreases to a first approximation, with an increase in hardness. However, in steel matrix composites no mutual correlation between hardness and abrasion wear loss is observed. Composites with similar hardness values have significant differences in the abrasion wear loss and vice versa. Khruschov found an inverse relationship between abrasion rate and hardness for pure metals [60]. On the other hand, Zum Gahr demonstrated that hardness is not a suitable parameter to describe reliable wear resistance of inhomogeneous or multiphase materials [48]. In steel matrix composites, variations in the reinforcement volume fraction and size, amount of in situ matrix carbides, matrix microstructure and hardness, and internal residual stresses, may have a more critical effect on the wear behavior than the overall composite hardness. Similar results have been reported for a white iron matrix composite the microstructure of which (e.g. amount of in situ carbides, hardness of martensite) was changed by systematic heat treatments [13].

3.4. Applications of steel matrix composites Depending on the matrix material, a number of potential applications of the steel matrix composites examined in the present study can be deduced. Owing to the good wear and corrosion properties of the stainless steel matrix composites, these materials can be used in environments where corroding liquids and gases are present, or where eroding mass containing impurities is flowing. Under these conditions superaustenitic 654 SMO and duplex stainless steels reinforced with TiN and Al2O3 particles, respectively, may offer the most advantages. They can be used as components in the chemical and petrochemical industry (scoring dies, pumps, valves and mixers), pulp and paper industry (digesters, bleaching vessels), heat exchanger tubes, high-pressure plungers, and as cladding material on steel tubes transporting corroding mass. Tool steel matrix composites can be used in conditions where wear problems are encountered, even at elevated temperatures. Hot work tool steel reinforced with Cr3C2 particles may offer the best solutions. These materials can be used as cutting blades and grinders in wood and plastic handling, as wear resistant parts in mining and earth-moving operations, and as metal forming tools (extrusion dies, rolling mill rolls, tubeforming rolls). It should be here mentioned that in actual service conditions the wear performance of the composite can vary, depending on the material and wear situation.

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4. Conclusions Studies were undertaken to evaluate the mechanical and corrosion properties of HIPed steel matrix composites reinforced with inexpensive ceramic particulates. From the results of the present investigation the following conclusions can be drawn. 1. Careful selection of the matrix and reinforcement powder sizes can improve the distribution of the ceramic particulate into the steel matrix. The presence of a substantial amount of alloying elements in the steel matrix enhances its reactivity in contact with ceramic reinforcements. The TiC reinforcement was found to have the best bonding behavior with the steel matrix because of the formation of a thin (Fe, Cr)-rich layer at the interface. 2. The wear resistance of the unreinforced steel matrices increases to a great extent after the incorporation of ceramic particulate reinforcements. Depending on the matrix alloy, up to 75% reduction in the abrasion wear loss can be achieved by adding only 5 vol.% ceramic particulates. In composites having a high amount of in situ matrix carbides the wear resistance decreases after a certain reinforcement volume fraction. No mutual correlation between the abrasion wear loss and the hardness of the steel matrix composites was observed. 3. The presence of ceramic particulates has little or no effect on the corrosion resistance of the stainless steel matrices until a certain reinforcement volume fraction, which depends on the matrix alloy. The additional incorporation of ceramic reinforcements causes local breakdown of the passive layer with a corresponding material dissolution and pit formation. 4. Tensile strength, ductility and impact toughness of the unreinforced steel matrices decrease after the incorporation of ceramic particulates. While the decrease in the fracture properties of the composites is consistent with previous observations in aluminum and magnesium matrix composites, the decrease in the tensile strength is in contrast to most of the previous findings. The use of coarse reinforcing particles increases the tensile strength and the ductility of steel matrix composites, which again contradicts the results reported for light metal matrix composites. 5. Superaustenitic 654 SMO reinforced with TiN and hot work steel reinforced with Cr3C2 were found to offer the best combination of properties.

Acknowledgements This work is part of a national project on iron-based composites funded by The Technology Development

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Centre of Finland (TEKES) and Finnish industry. The authors would like to acknowledge their support, as well as the assistance of X. Liu, M. Rastas, P. Korpiala, J. Hellman, and O. Mattila.

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