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2007 Atlas of Cast Metal-Matrix Composites Structures

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Andrzej Wojciechowski Krystyna Pietrzak Dariusz Rudnik Jerzy Sobczak Natalia Sobczak Rajiv Asthana Motor Transport Institute Motor Transport Institute Foundry Research Institute Foundry Research Institute 2007 ® ATLAS METAL-MATRIX OF CAST OF CAST COMPOSITES STRUCTURES Commision 8.1 Cast Composites WFO W F O
Transcript

Andrzej Wojciechowski Krystyna Pietrzak Dariusz RudnikJerzy Sobczak Natalia Sobczak Rajiv Asthana

MotorTransportInstitute

MotorTransportInstitute

FoundryResearchInstitute

FoundryResearchInstitute

2007

®

ATLASMETAL-MATRIXOF CASTOF CAST

COMPOSITESSTRUCTURES

Commision 8.1Cast Composites

W F OW F O

Jerzy Sobczak Natalia Sobczak

Rajiv Asthana Andrzej Wojciechowski

Krystyna Pietrzak Dariusz Rudnik

ATLAS OF CAST

METAL-MATRIX COMPOSITE

STRUCTURES

Part I Qualitative analysis

Motor Transport Institute - Warsaw Foundry Research Institute - Cracow 2007

Translation assistance provided by Tomasz Drecki

With the sponsorship of:

- World Foundrymen Organization - Ministry of Science and Higher Education of Poland - Motor Transport Institute - Foundry Research Institute

ISBN 978-83-60965-00-9

Copyright © 2007 by Motor Transport Institute - Warsaw Foundry Research Institute - Cracow All rights reserved

No part of this book may be reproduced, stored in a retrieval system, or transmitted, in any form or by means, electronic, mechanical, photocopying, recording, or otherwise, without the prior written permission of the publisher

Published by Motor Transport Institute - Warsaw

Foundry Research Institute - Cracow

First edition

PRINTED IN POLAND

Contents

Page

Summary 5 Foreword 7 From the Authors 9 I. Particulate dispersion-reinforced composites 11 I.1. Aluminum composites containing fly ash (ALFA composites) 11 I.2. Aluminum composites containing graphite 31 I.3. Aluminum composites containing silicon carbide 47 I.4. Aluminum composites containing aluminum oxide 61 I.5. Aluminum composites containing graphite and silicon carbide (hybrid) 64 I.6. Copper composites containing graphite 70 I.7. Lead composites containing fly ash 80 I.8. Lead composites containing graphite 81 II. Short fiber-reinforced composites 83 II.1. Aluminum oxide-reinforced aluminum composites 83 II.2. Carbon felt-reinforced aluminum composites 85 II.3. Carbon felt-reinforced zinc composites 92 II.4. Carbon felt-reinforced lead composites 95 III. Long fiber-reinforced composites 96 III.1. Carbon fiber-reinforced aluminum composites 96 III.2. Carbon fiber-reinforced copper composites 99 III.3. Sapphire fiber-reinforced nickel composites 101 IV. Gas “reinforced” composites (gasars) 106 V. Nano-composites 109 VI. Typical structural defects in metal-matrix composites 114 VI.1. Particulate dispersion-reinforced composites 114 VI.1.1. Macro-scale nonuniform arrangement of the reinforcing phase 114 VI.1.2. Micro-scale nonuniform arrangement of the reinforcing phase 115 VI.1.3. Porosity - gas bubbles and shrinkage 119 VI.1.4. Excessive reactivity between the metal matrix and reinforcement 121 VI.1.5. Contaminations and non-metallic inclusions 125 VI.2. Porous ceramic preforms reinforced composites made of short, long fibers or particulates 126 VI.2.1. Structural discontinuities as a result of insufficient preform infiltration 126 VI.2.2. Structural heterogeneity and discontinuities resulting from the improper preform preparation 129 VI.2.3. Excessive reactivity between composite components 131 VI.2.4. Non-metallic inclusions 134 VI.3. Hybrid composites 135 VI.4. High volume fraction particulate-reinforced composites 137 VI.5. Long fiber-reinforced composites 140 Acknowledgements 143

5

Summary The ATLAS displays microstructures of metal-matrix composite materials based on aluminum, copper, zinc, lead, nickel (nickel-aluminide) as well as magnesium metal matrices, manufactured by the authors. The composites were reinforced with different types of particles (fly ash, graphite, silicon carbide and aluminum oxide), short fibers (aluminum oxide and carbon felt) and long fibers (carbon and sapphire). Microstructures of gasars (a new group of composites with gas as ”reinforcing phase”) are also presented together with microstructures of aluminum matrix nanocomposites. The analysis of structural characteristics of each kind of material is supplemented by a description of the composite type, method of manufacture as well as experimentally proven and potential service properties. Typical structural defects of the cast metal-matrix composites are covered in detail.

7

Foreword

Over the last fifty years, metal-matrix composites have seen wide-ranging applications beyond the aerospace sector for which they were originally intended. Consumer industry, electronic packaging industry, and most importantly, transportation industry have been the most prominent sectors to adopt metal-matrix composites in niche applications such as the brake disc and parts of the suspension, power train and the engine including pistons and cylinder blocks' sleeves. The remarkable design flexibility that these composites, and their higher-temperature counterparts, ceramic-matrix composites, exhibit enables the composites practitioner to design and adapt the material to suit the needs of the user industry.

In the ATLAS, the authors have put together an impressive collection of metal-

matrix composite microstructures synthesized by them in Poland and USA using the solidification and casting techniques. The microstructures of cast composites based on low-melting, medium-melting, and high-melting point metal matrices ranging from lead, zinc, magnesium, aluminum and copper to nickel-based intermetallics, are presented. The diversity of these composites is captured in hundreds of scientifically valuable and aesthetically beautiful photomicrographs, each of which is accompanied with a description of the prominent structural features, including structural defects, if any, processing conditions as well as proven and potential properties of the manufactured composites. The photomicrographs display the composites containing a wide variety of reinforcement including silicon carbide, aluminum oxide, graphite and fly ash. In addition, the ATLAS presents the microstructure and processing of new and emerging metal-based nanocomposites and gasar materials that are sure to gain wider acceptance in industry in years to come.

I am unaware of any previous effort of this magnitude and scope at

systematizing and assembling the unique structural attributes of cast composites in a single resource. The principal authors are leaders in the metal-matrix area and deserve to be congratulated for an excellent contribution to the technical literature that will undoubtedly be of immense value and use to the researcher and practitioner alike. July 30, 2007 Mrityunjay Singh, Ph.D.

Chief ScientistOhio Aerospace Institute

NASA Glenn Research CenterCleveland, OH, USA

9

From the Authors Following successful practical applications in the aerospace and aviation industry, increasing use of metal-matrix composites in car design (particularly as brake discs material, parts of the suspension, power train and parts of the engines) makes the slogan “get the composites down to earth” ever more up to date and real. It has generally been accepted that this new group of construction materials still refers to the niche materials in civil engineering (mainly electronics and land transport). The question is: will metal-matrix composites still remain the niche materials? They will probably stay as such at least for the next decade, though this is not an opinion shared by all the experts. One of the reasons is the fact that metal-matrix composites are, because of their nature, a kind of creative materials that can potentially be designed and adapted to the customer’s individual needs. Composites are also more expensive than metals and monolithic alloys; hence the prospect of their wider use demands offering something “extra” with respect to properties and serviceability. As a consequence, introduction of these materials in practice will always have a precisely defined reason. Probably, it is just those unique characteristics that carve out niches for composite materials. Some of those niches could, however, be quite big and significant. For example, vital parts of the combustion engines, such as pistons and cylinder blocks’ sleeves are quite a big niche, just like electronic systems enclosures. Together, those niches make a substantial assortment of engineering applications, and their amount and size is ever growing. Moreover, each new application brings new engineering experiences, which significantly reduce barriers for the next applications, even if each time they are exceedingly specific and unique. The main obstacle in common use of composite materials may prove to be accessibility of not only information about the level of their properties together with the methods of forming a particular products out of them, but also comprehensive knowledge regarding specific structural features, particularly, with respect to the interpretation. The structure of the material may be regarded as a subjectively variable area of interpretation, from high level of abstraction, bringing in “high imprecision” of system interconnections of the mutually linked parts, to relatively simple quality and quantity evaluation of selected structural components. From the point of view of the methodology of testing materials, characterization of structure, although an extremely important means to obtain the desired state of the material, is not an end in itself. It is how the structure manifests itself as a concrete, desired set of properties, mainly physical, chemical and mechanical properties, that ultimately determines the aim ! In comparison with monolithic materials (metals and alloy), the structure of heterogenic materials, particularly composite materials, gets enriched and complicated by the presence of the second component - the reinforcing phase and its interaction with the matrix itself. The degree of complexity in evaluating such heterogeneous structures increases, which may bring about a loss of required degree of adequacy in relating the chemical composition of the matrix, the presence of the reinforcement and its type, the technological parameters of the composite synthesis process, and the composite’s structure and properties. Hence the importance of credible and ample supply of information from the area of tested

10

materials structures in connection with initial attributes of their creation as well as detailed definition of the condition of their forming.

We present you with the first part of the atlas of metal-matrix composite structures covering those composites that are primarily based on nonferrous metal matrices, and which, in the vast majority of cases, were created and tested during the last decade, by the authors, both in Poland (at the Foundry Research Institute in Cracow and the Motor Transport Institute in Warsaw), as well as in the USA, particularly at the Composites Centre of the University of Wisconsin - Milwaukee and at the NASA Glenn Research Center in Cleveland, Ohio. From the onset, this part of the planned trilogy is devoted to the quality structure description, mainly at the micro-scale, generated under the optical microscope, in connection with manufacturing parameters of various kinds of cast composites. The second part of the atlas will refer to quantitative structural analysis, while the third and the last part, will cover the structural attributes of metal-matrix composites as revealed by the scanning electron microscope. The microstructure characterization was conducted mainly under the optical OLYMPUS PMG3 microscope using polarized light technique and phase contrast as well as using stereoscopic OLYMPUS SZ40 microscope and appropriately selected method of preparing samples allowing for the specifics of the ground (polished) surface to be analyzed. A few structural tests were also carried out on electron scanning STEREOSCAN 420 microscope. The structure of each kind of composite, apart from structural characteristics, is also accompanied by abbreviated description of the type of its constituents, manufacturing methods, and expected service properties, usually verified in the course of self-conducted tests. The authors would like to express modest hope that this publication will somewhat contribute to increasing the knowledge about structure of the family of advanced materials such as metal-matrix composites, their chemical contents, qualitative properties, and the manufacturing techniques developed on the basis of know-how used. The know-how is often either protected and undisclosed in the current open literature or it may be unknown and rarely used in creating advanced innovative materials due to the difficulties experienced at the experimental level. We would like to express our gratitude to professor Pradeep Kumar Rohatgi, Director of the Composites Center at the University of Wisconsin-Milwaukee, for his extensive help, particularly at the beginning of our activities in the field of metal-matrix composites as well as numerous ever inspiring discussions. Our thanks also go to professor Ludmila Boyko from the State Metallurgical Academy of Ukraine in Dnepropetrovsk, for directing our attention towards highly porous metallic media with controlled contents and gas bubble morphology, known as gasars.

Particulate dispersion-reinforced composites 11

I. Particulate dispersion-reinforced composites I.1. Aluminum composites containing fly ash (ALFA composites) 52K/x fly ash composite (x = 5, 10 and 15 vol. %) Matrix: 52K aluminum alloy (USA) - 9.18 Si, 2.55 Zn, 2.50 Cu, 1.01 Fe, 0.23 Mn, 0.13 Mg, 0.08 Ni, 0.07 Cr, 0.07 Pb, 0.05 Ti (wt. %). Reinforcing phase: fly ash from PSI Energy electrical power plant (USA) contains mainly cenospheres (microspheres) and about 10 % microgranules of the following chemical composition (converted into oxides, wt. %):

Oxide SiO2 Al2O3 Fe2O3 K2O MgO TiO2 CaO Na2O SO3 P2O3 Microspheres 58.0 32.2 5.05 3.5 1.6 1.2 0.8 0.7 0.3 0.1 Microgranules 38.5 17.1 25.1 2.5 1.0 1.5 4.0 0.5 1.9 0.5

An example appearance of the external cenospheres collected from another electrical power plant (Dayton Power and Light Co., USA), used in the subsequent tests, is shown in Fig. I.1.1. Manufacturing method: mechanical stirring method (vortex) with gravity casting into grey iron die. Following melting, the composite slurry was subjected to squeeze casting under a pressure of 220 MPa for 45s (die temperature - bottom – cavity/top – squeeze plunger: 190-230/170-200°C). The casting temperature of the composite suspension was maintained within 670-690°C range. Squeeze cast cylinder of 80 mm diameter and 25-30 mm height was subjected to the following heat treatment according: solutionizing at 540°C/12 h, cooling in water at 70°C; artificial aging at 160°C/24 h. Properties: lower (microgranules reinforcement) and very low (cenospheres reinforcement) density, increased and satisfactory mechanical properties, significant resistance to grinding (wear off) and thermal shock, decreased thermal expansion and thermal conductivity, good machinability, advantageous economic and ecological manufacturing parameters (mainly due to waste material character of the reinforcing phase). Structural characteristics (Figs. I.1.2 - I.1.7): dominating effect of the external pressure during solidification is noticeable in the fly ash-reinforced aluminum composites, characterized by low content of cenospheres and microgranules relatively uniformly dispersed in the metal matrix. Microstructural investigation proved that the hydrostatic pressure, applied to the solidifying composites prevents the growth of eutectic phase crystals in the favored directions after the eutectic temperature has been achieved. External pressure also leads to the increase of the α-phase amount in the structure and solidification of the cooled eutectic in the direct vicinity of the dendritic solidification front. As a result of the pressure application, liquid aluminum alloy fills up the microspheres, solidifying inside them in a much pulverized structure. Modifying influence of the external pressure is also observed on the structure of the whole composite matrix.

12 Particulate dispersion-reinforced composites

Fig. I.1.1. As-received fly ash reinforcing phase used for manufacturing ALFA composites. This is fly ash from the Dayton Power and Light Co electrical power plant, (USA), in the shape of glassy-like cenospheres with the following chemical composition (converted into oxides, wt. %): 61.0 SiO2, 25.8 Al2O3, 5.0 Fe2O3, 3.6 K2O, 1.6 MgO, 1.0 TiO2, 0.8 CaO, 0.7 Na2O, 0.3 SO3, 0.1 P2O3. Cenospheres, representing mainly amorphous material, are of varying size as well as non-uniform color due to diversified chemical composition (e.g. ginger brown color revels high content of the iron oxides)

a

b

c

Particulate dispersion-reinforced composites 13

d e

f g

Fig. I.1.2. Structure of 52K/5 vol. % fly ash aluminum alloy composite, obtained by vortex method. Following melting, the composite was subjected to squeeze casting and heat treatment (precipitation hardening). Non-etched sample. Polarized light

a b

c d

Fig. I.1.3. Structure of 52K/5 vol. % fly ash aluminum alloy composite, obtained by vortex method. Following melting, the composite was subjected to squeeze casting and heat treatment (precipitation hardening). Sample etched with Keller reagent. Polarized light

14 Particulate dispersion-reinforced composites

a b

c d

e f

Fig. I.1.4. Structure of 52K/10 vol. % fly ash aluminum alloy composite, obtained by vortex method. Following melting, the

composite was subjected to squeeze casting and heat treatment (precipitation hardening). Non-etched sample. Polarized light

a b

Particulate dispersion-reinforced composites 15

c

Fig. I.1.5. Structure of 52K/10 vol. % fly ash aluminum alloy composite, obtained by vortex method. Following melting, the

composite was subjected to squeeze casting and heat treatment (precipitation hardening). Sample etched with Keller reagent. Polarized light

a b

c

Fig. I.1.6. Structure of 52K/15 vol. % fly ash aluminum alloy composite, obtained by vortex method. Following melting, the

composite was subjected to squeeze casting and heat treatment (precipitation hardening). Non-etched sample. Polarized light

16 Particulate dispersion-reinforced composites

a

b

c

d

e

f

g h

Fig. I.1.7. Structure of 52K/15 vol. % fly ash aluminum alloy composite, obtained by vortex method. Following melting, the

composite was subjected to squeeze casting and heat treatment (precipitation hardening). Sample etched with Keller reagent. Polarized light

Particulate dispersion-reinforced composites 17

AK12/9 wt. % fly ash composite Matrix: AK12 aluminum alloy (AlSi12CuNiMg) – 12.20 Si, 1.15 Cu, 0.95 Ni, 0.90 Mg, 0.60 Fe, 0.20 Zn, 0.19 Mn, 0.09 Ti (wt. %). Reinforcing phase: fly ash from Dayton Power and Light Co. electrical power plant (USA) contains mainly microgranules of the following chemical composition: 55.9 SiO2, 30.2 Al2O3, 5.4 Fe2O3, 1.3 CaO, 2.7 K2O, 1.0 MgO, 1.6 TiO2, 0.2 Na2O, 0.3 SO3, 0.4 P2O3. Manufacturing method: introduction of initially enriched reinforcing phase by mechanical stirring (vortex) under argon and squeeze casting under the pressure of 170 MPa for 45 s, (die temperature – bottom/top - 180/220°C). The slurry was poured into the die at the temperature of 750°C and squeeze cast in order to produce combustion engine pistons of 90 mm diameter. Structural characteristics (Fig. I.1.8): good acceptability of the reinforcing phase by the metal matrix is noticeable due to intensive chemical character of the interaction between the composite constituents. This statement reflects the presence within the structure of the large primary silicon crystals, formed as a result of the reaction between the oxides (mainly SiO2 and FexOy) with liquid aluminum according to the oxy-redox reaction 3MeO+2Al=Al2O3+3Me. Also significant amount of cracked fly ash microgranules is an indirect proof of intensive interactions taking place in a given composite.

a

b c

18 Particulate dispersion-reinforced composites

d

e

f g Fig. I.1.8. Structure of squeeze cast AK12/9 wt. % fly ash composite. As-cast. Non-etched sample. Polarized light (a, b, d,

f) and phase contrast (c, e, g)

Particulate dispersion-reinforced composites 19

AG10/5 wt. % fly ash composite Matrix: AG10 aluminum alloy (AlMg10) - 9.50 Mg, 0.10 Mn, 0.20 Fe (wt. %). Reinforcing phase: fly ash from Skawina electrical power plant (Poland) of the following composition: 38.5 SiO2, 17.1 Al2O3, 25.1 Fe2O3, 4.0 CaO, 2.5 K2O, 1.0 MgO, 1.5 TiO2, 0.5 Na2O, 1.9 SO3, 0.5 P2O3 (5 wt. %). Manufacturing method: mechanical stirring method (vortex) using the device for refining and modifying Al alloys (URM-1 type, designed by Foundry Research Institute) and gravity casting into permanent mould (made from grey iron). Following melting, the composite was subjected to squeeze casting under a pressure of 230 MPa for 45 s (die temperature – bottom/top: 190-230/170-200°C). The casting temperature of the composite slurry was maintained within 670-690°C range. Cylindrical castings (70 mm diameter and 90 mm height) were subjected to the following heat treatment: solutionizing at 430°C/16 h, cooling in water at 70°C; natural aging for 96 h. Properties: lower density, satisfactory mechanical properties, significant resistance to grinding (wear off) and thermal shock, decreased thermal expansion and thermal conductivity, good machinability properties, advantageous economic and ecological manufacturing parameters (mainly due to waste material character of the reinforcing phase). Structural characteristics (Fig. I.1.9): The characteristic feature of the composites manufactured using fly ash from Skawina electrical power plant is the less uniform arrangement of the reinforcing phase in the metal-matrix when compared with composites obtained using fly ash from the PSI Energy electrical power plant (USA) (Figs. I.1.2 - I.1.7). It is, however, a deceptive effect because it is linked with the reinforcing phase properties in the as-received state and not with an effect of agglomeration during the composite manufacture. In the case of coal burning in the Skawina electrical power plant, resultant fly ash contains a large amount of primary agglomerates, consisting of a few to a few dozens of different size particles, permanently bridged with one another. The squeeze pressure facilitates filling with the liquid metal the cavities between the particulates and even micro-cracks and open porosity within individual particles. While visible in the photomicrographs, the structural discontinuities are within large agglomerates and form enclosed porosity.

a b

c d

20 Particulate dispersion-reinforced composites

e

f

g

Fig. I.1.9. Structure of AG10/5 vol. % fly ash composite: (a, b) gravity cast and (c-g) squeeze cast. Non-etched sample.

Polarized light (a, b, d-g) and phase contrast (c)

Particulate dispersion-reinforced composites 21

Al/25 wt. % fly ash composite Matrix: pure aluminum (99.5 wt. % Al). Reinforcing phase: fly ash from the Dayton Power and Light Co electrical power plants (USA) mainly containing microgranules of the chemical composition (wt. %.): 55.9 SiO2, 30.2 Al2O3, 5.4 Fe2O3, 1.3 CaO, 2.7 K2O, 1.0 MgO, 1.6 TiO2, 0.2 Na2O, 0.3 SO3, 0.4 P2O3. Manufacturing method: mechanical alloying method involving mixing of matrix powder with fly ash particles in a rotating and vibrating mill, under argon atmosphere, and then compacting the mixture under pressure. Properties: lower density, increased mechanical properties (particularly ultimate tensile strength), significant resistance to grinding (wear off) and thermal shock, fundamentally decreased thermal expansion and thermal conductivity, good machinability, advantageous economic and ecological manufacturing parameters (mainly due to the waste material character of the reinforcing phase). Structural characteristics (Figs. I.1.10 - I.1.11): The characteristic features of the initial constituent in the form of composite powder are (Fig. I.1.10):

1) presence of a substantial amount of the reinforcing phase, 2) a high degree of reinforcing phase break-up, stemming from the milling effect of the fly ash

particulates, while a significant part of microspheres of diameter below 10 µm does not get destroyed,

3) a high degree of homogeneity in the arrangement of the reinforcing phase within metal matrix, 4) relatively low presence of structural discontinuities in the regions of large composite particles,

which are essentially agglomerates of composite particulates (from several dozens to several hundred times smaller).

The characteristic features of the composite obtained from composite powders using powder compaction under pressure are (Figs. I.1.11 - I.1.12):

1) complete lack of structural discontinuities, 2) complete disappearance of microspheres as well as larger fractions of microspheres due to

their interaction with metal matrix, 3) more than ten times smaller size of the reinforcing phase precipitates in comparison with the

primary composite powder, 4) very minute reinforcing phase precipitates are created mainly by aluminum oxide particulates

produced due to oxy-redox reaction between aluminum and fly ash constituents (2Al+3MeO=Al2O3+3Me), mainly silicon oxide and iron oxide since the reactivity of titanium oxide is negligible, while calcium oxide and magnesium oxide do not react with the metal matrix,

5) banding of the fine aluminum oxide precipitates, resulting from Al/metal oxide interactions.

a b

22 Particulate dispersion-reinforced composites

c d

e f

g

h i

Particulate dispersion-reinforced composites 23

j k

l m

Fig. I.1.10. Structure of Al/25 wt. % fly ash composite (initial composite powders) obtained by mechanical alloying. Non-etched sample. Polarized light (a-f, g, h, j, l) and phase contrast (i, k, m)

a b

c d

24 Particulate dispersion-reinforced composites

e

f

g

h

Fig. I.1.11. Structure of Al/25 wt. % fly ash composite obtained by mechanical alloying followed by compaction of composite powders by know-how technology. Transverse section. Non-etched sample. Polarized light (b, c, e, g) and phase contrast (d, f, h). Figure (a) shows macrostructure of investigated section

a b

Particulate dispersion-reinforced composites 25

c

d

e

f

g h Fig. I.1.12. Structure of Al/25 wt. % fly ash composite obtained by mechanical alloying followed by compaction of composite

powders by know-how technology. Longitudinal section. Non-etched sample. Polarized light (b, c, e, g) and phase contrast (d, f, h). Figure (a) shows macrostructure of investigated section

26 Particulate dispersion-reinforced composites

a

b

Fig. I.1.13. Structure of Al/30 vol. % fly ash composite obtained by mechanical alloying followed by compaction of composite powders under a pressure of 3000 MPa. Non-etched sample

20 µm

10 µm

Particulate dispersion-reinforced composites 27

AK12/60 vol. % fly ash composite Matrix: AK12 aluminum alloy (AlSi12CuNiMg) – 12.2 Si, 1.15 Cu, 0.95 Ni, 0.90 Mg, 0.60 Fe, 0.20 Zn, 0.19 Mn, 0.09 Ti (wt. %). Reinforcing phase: fly ash from Dayton Power and Light Co. electrical power station (USA) cenospheres of chemical content (wt. %): 59.8 SiO2, 32.2 Al2O3, 5.1 Fe2O3, 0.4 CaO, 1.4 K2O, 0.8 MgO, 0.8 TiO2, 0.2 Na2O, 0.2 SO3, 0.3 P2O3. Manufacturing method: squeeze infiltration of the porous preform (porosity 40 vol. %) of cenospheres under a pressure of 170 MPa for 45 s (die temperature – top/bottom - 180/220°C). The matrix alloy was poured in the die at a temperature 760°C. A cylinder of 70 mm diameter and 30 mm height (in as-cast state, without heat treatment) was squeeze cast. Properties: significantly lower density (in the extreme case bordering on 1 g/cm3), decreased thermal expansion, electrical expansion and thermal conductivity, increased energy and sound absorption characteristics, good machinability, advantageous economic and ecological manufacturing parameters. Structural characteristics (Fig. I.1.13): in spite of the high content of the reinforcing phase, the preform is relatively well infiltrated by liquid matrix. The arrangement of the ash particulates in the matrix corresponds to their initial distribution in porous preform, and the rigid packing of particulates does not change in spite of a high pressure being exerted on the whole set. Parts of damaged particulates are completely filled up due to liquid metal infiltration through micro-cracks in the walls of microspheres. The apparent lack of metal penetration between the particulates, visible at high magnifications, is a consequence not so much of the imperfections in the technological process of squeeze infiltration method to manufacture the composites as it is of the process of preparing the preform itself (a binding compound was used to manufacture the preform). On one hand, binding compound causes bridges to appear, rigidly connecting the microspheres. But on the other hand, relatively good wetting of the preform particulates by the binder contributes to complete filling up and closure of the areas between some microspheres, making it impossible for the liquid metal to infiltrate such areas. In microspheres that get “coated” with bonding compound, structural discontinuities present in the micro sphere walls appear to get sealed, which increases the amount of microspheres that remain unfilled with the metal-matrix.

a

28 Particulate dispersion-reinforced composites

b c

d e

f g

h i

Fig. I.1.13. Structure of Al/60 vol. % fly ash composite, obtained by squeeze infiltration of the porous preforms made of

cenospheres. Non-etched sample. Phase contrast (c, e, g, i)

Particulate dispersion-reinforced composites 29

A356/5.0 wt. % fly ash composite Matrix: A356 aluminum alloy (AlSi7Mg) – 7.2 Si, 0.40 Mg, 0.60 Fe, 0.30 Mn (wt. % ). Reinforcing phase: aluminum composite powders containing 25 wt. % of fly ash from Skawina power station (Poland) and of the following content: 81.66 SiO2, 6.69 Al2O3, 4.37 Fe2O3, 4.29 CaO, 1.58 MgO, 0.40 K2O, 0.21 Na2O, 0.29 SO3, (wt. %), manufactured by mechanical alloying method, using mixing in the rotating-vibrating mill under argon atmosphere. Manufacturing method: mechanical stirring (vortex) of reinforcement with liquid metal of 720°C temperature for 24 h and gravity casting into the metal mould (cast iron). Properties: lower density, increased and satisfactory mechanical properties, significant resistance to grinding (wear off) and thermal shocks, decreased thermal expansion and thermal conductivity, good machinability, advantageous economic and ecological manufacturing parameters (mainly due to waste material character of reinforcing phase). Structural characteristics (Fig. I.1.14): during manufacturing of the initial reinforcing phase by mechanical alloying method, using aluminum powder and fly ash particulates, a pulverisation of main part of the ash particulates takes place together accompanied by their homogenous arrangement in the metal matrix. At the same time, mutual interaction between metal and ceramics takes place, leading to one of a kind final activation of reinforcing phase, aimed at facilitating its introduction into the metal bath (with the subsequent distribution of it within the entire volume). This way it is possible to obtain not only a high degree of metal-ceramic slurry homogeneity, but also its satisfactory stability. This is confirmed by relatively homogenous arrangement of fly ash particulates dispersed in the A356 alloy matrix, even in case of gravity casting. Reinforcement phase particulates, in their majority, are located in the areas between dendrites of the aluminum solid solution, where also eutectic (Al+Si), gets collected. In the given case, as in majority of ceramic reinforcing phases (SiC, Al2O3, graphite), fly ash particulates compose inactive nucleants for the stable aluminum solution and its growth causes reinforcement particulates to be pushed out into the inter-dendral areas. On the other hand, the same particulates form active nucleants for the eutectic silicon precipitations.

a

30 Particulate dispersion-reinforced composites

b

c

d

e

f

g

h i

Fig. I.1.14. Structure of A356/15 wt. % fly ash composite, manufactured by vortex method using initial preparation of the

reinforcing phase by mechanical alloying method. Conventional light (a, b, d, f, h) and phase contrast (c, e, g, i)

Particulate dispersion-reinforced composites 31

I.2. Aluminum composites containing graphite Al/GrNi composite (5 vol. % nickel-coated graphite) Matrix: pure aluminum (99.9 wt. % Al). Reinforcing phase: nickel-coated platelet graphite particulates (about 20 wt. % Ni) of 150 µm average dimension. Manufacturing method: mechanical stirring method (vortex), gravity cast into grey iron die and squeeze cast under the pressure of 110 MPa for 30 s (die temperature - bottom/top - 190-230/170-200°C). The casting temperature of the composite slurry was maintained at 720°C. A die-cast cylinder casting (80 mm diameter and 40 mm height) was annealed at 300°C for 4 h. Structural characteristics (Fig. I.2.1): nickel-coated graphite platelets are well wetted by liquid aluminum. This is a result of nickel coating dissolving in aluminum, and this process is accompanied by reactive metal penetration inside the individual platelets via micro-cracks, which in turn causes the platelet itself to divide into several smaller pieces.

Fig. I.2.1. Structure of squeeze cast Al/5 vol. % GrNi composite. Non-etched sample. Polarized light

10 µm

32 Particulate dispersion-reinforced composites

AK11/xGrNi composite (x = 2.5, 5.0 vol. % nickel-coated graphite) Matrix: AK11 aluminum alloy – 12.10 Si, 0.11 Mn, 0.42 Fe (wt. %). Reinforcing phase: nickel-coated graphite particulates (NOVAMET, 50 wt. % Ni) of 75-120 µm average dimensions. Manufacturing method: mechanical stirring method (vortex), gravity cast into grey iron die or squeeze cast under a pressure of 110 MPa for 30 s (die temperature - bottom/top - 190-230/170-200°C). The casting temperature of the composite suspension was maintained at 720°C. A die-cast cylinder (80 mm diameter and 40 mm height) was annealed at 300°C for 4 h. Properties: slightly increased density, satisfactory mechanical properties, improved tribological characteristics (decreased wear and friction coefficient in experiments with steel counter sample), especially after squeeze casting; a reduction in the thermal expansion coefficient is expected. Structural characteristics (Figs. I.2.2 - I.2.5): nickel-coated graphite particulates are well wetted by the liquid metal matrix which is why there are no structural discontinuities at the interface between matrix and graphite particulates as well as no particulate agglomerates form during composite manufacture. The susceptibility to wetting in the AK11/xGrNi system is linked with the high degree of chemical interaction due to intensive dissolution of nickel coating in the liquid matrix during composite manufacture. Excessive dissolution of the nickel coating, however, causes an undesired effect because of a change in the chemical composition of the metal matrix, its superfluous enrichment with nickel, and the formation of the brittle Al3Ni intermetallic phase precipitates as well as primary Si precipitates. An increase in the volumetric fraction of nickel-coated graphite being introduced in the matrix causes an increase in the amount of these precipitates. The significant fact is that in the case of gravity cast composites, apart from Al(Ni)+Al3Nieut eutectic, Al3Niprim primary precipitates also form. These precipitates have a shape of elongated platelets of differing dimensions. Irrespective of the composite manufacturing method, primary silicon precipitates have a tendency to grow on dispersed graphite particulates.

a

Graphite

Graphite

Siprimary

Al3Ni

Sieut

20 µm

Particulate dispersion-reinforced composites 33

b

c

Fig. I.2.2. Structure of AK11/2.5 vol. % GrNi composite: gravity cast (a), squeeze cast (b, c). Sample etched with Keller’s reagent. Polarized light

a

Si prim.

Al3Ni

Graphite20 µm

10 µm

20 µm

34 Particulate dispersion-reinforced composites

b

Fig. I.2.3. Structure of AK11/5 vol. % GrNi composite: gravity cast (a), squeeze cast (b). Sample etched with Keller’s

reagent. Conventional light (a) and polarized light (b)

20 µm

Particulate dispersion-reinforced composites 35

AK20/5.7 vol. % nickel-coated graphite composite Matrix: AK20 aluminum alloy - 22.34 Si, 1.50 Cu, 1.10 Ni, 0.60 Mg, 0.25 Mn, 0.45 Fe (wt. %). Reinforcing phase: nickel-coated graphite particulates (NOVAMET, 50 wt. % Ni) of 75-120 µm average dimensions. Manufacturing method: mechanical stirring method (vortex), squeeze cast under a pressure of 110 MPa for 30 s (die temperature - bottom/top - 190-230/170-200°C). The casting temperature of the composite slurry was maintained at 720°C. Cylindrical castings (70 mm diameter and 90 mm high, were subjected to heat treatment according to the following parameters: solutioning at 510°C/3.5 h, cooling at hot water, aging at 240°C/6.5 h. Properties: satisfactory mechanical properties, significantly improved tribological characteristics (decreased wear and friction coefficient in experiments with steel counter sample). A reduction in the thermal expansion coefficient is essential expected. Structural characteristics (Figs. I.2.4 - I.2.6): there are no structural discontinuities at the interface between the matrix and graphite particulates as well as no particulate agglomerates form during composite manufacture because nickel-coated graphite particulates are well wetted by the liquid metal matrix. These benefits are linked with the high degree of chemical interaction due to intensive dissolution of nickel coating in the liquid aluminum alloy during composite manufacture. Dissolution of the nickel coating results in changes of its thickness and structure accompanied with the formation of discontinuities, especially in its matrix-side layer. Despite the fact that nickel dissolution causes a change in the chemical composition of the metal matrix, application of squeeze casting suppresses its superfluous enrichment with nickel, resulting in the undesired formation of high volume fraction of brittle Al3Ni intermetallic phase precipitates, as it takes place in the case of gravity cast composites of the same system. Irrespective of the composite manufacturing method, primary silicon precipitates have a tendency to nucleate at the surface of dispersed graphite particulates.

a

36 Particulate dispersion-reinforced composites

b c

d e

f g

h i Fig. I.2.4. Structure of squeeze cast AK20/5.7 vol. % GrNi composite. Top of the casting. Non-etched sample. Polarized

light (a, b, d, f, h) and phase contrast (c, e, g, i)

Particulate dispersion-reinforced composites 37

a

b

c

d e

38 Particulate dispersion-reinforced composites

f

g

h

i

Fig. I.2.5. Structure of squeeze cast AK20/5.7 vol. % GrNi composite. Middle of the casting. Non-etched sample. Polarized light (a, b, d, f, h) and phase contrast (c, e, g, i)

a

Particulate dispersion-reinforced composites 39

b c

d e

f g

h i Fig. I.2.6. Structure of squeeze cast AK20/5.7 vol. % GrNi composite. Bottom of the casting. Non-etched sample. Polarized

light (a, b, d, f, h) and phase contrast (c, e, g, i)

40 Particulate dispersion-reinforced composites

A356/5.7 vol. % nickel-coated graphite composite Metal matrix: A356 aluminum alloy (AlSi7Mg) - 7.11Si, 0.30Mg, 0.48Fe, 0.31Mn (wt.%). Reinforcing phase: nickel-coated graphite particulates (NOVAMET, 50 wt. % Ni) of average dimension 75-150 µm. Manufacturing method: centrifugal casting. Nickel-coated graphite particulates (50Gr/50Ni) of average dimension 75-150 µm were heated to a temperature of 200°C for 2 h, before being introduced into the metal, following a newly developed procedure. Graphite mould temperature - 200°C, temperature of liquid metal 740°C (directly before the introduction of nickel-coated graphite), time of stirring composite slurry - 2 min, temperature of pouring - 740°C, rotational speed (graphite inlay placed tightly in a cast iron tube) - 1200 rpm, casting dimensions: diameter - 55 mm, length 100 mm. Properties: slightly increased density, acceptable mechanical properties, improved tribological characteristics (decreased wear and friction coefficient in experiments with steel counter sample), especially on internal layers of the centrifugal castings; a decrease in the thermal expansion coefficient could be expected. Structural characteristics (Fig. I.2.7): The centrifugal force causes loss of structural homogeneity of the gravity cast A356/5.7 GrNi composite (Fig. I.2.7) and generates four zones (Figs. I.2.8 and I.2.9). Zone I (a, c, e) forms an external layer on the casting, does not contain graphite particulates and is characterized by dispersive structure with relatively uniform distribution of a large quantity of Al3Ni phase, which forms as a result of the dissolution of nickel coating (from the graphite surface) in the liquid metal matrix. Zone II (g, i, k) bordering zone I also does not contain graphite particulates but, contrary to zone I, the Al3Ni phase precipitates in zone II have a shape of elongated plates with an arrangement clearly determined by the direction of solidification. Isolated graphite particulates appear in this zone and eutectic participates are much larger than in all other zones. Zone III (m) is a transition zone, between internal layer of the casting surface (zone IV) and zone II. In zone IV (o, r, t) there is accumulation of particulates of graphite which are driven out of other zones by the centrifugal force. In order to explain the specifics of structure of centrifugally cast A356/GrNi composite castings, some aspects linked with crystallization of the metal matrix as a ternary Al-Si-Ni alloy should be taken into account. This is because dissolution of nickel coating, initially spread over graphite, takes place during creation of metal suspension as well as in the course of its spinning and crystallization. Zone I, initially takes shape as a result of solidification of the suspension layer in contact with the mould, which is at a significantly lower temperature. Then, during spinning with simultaneous cooling of the suspension, its de-lamination takes place because centrifugal force segregates the heavier Al3Ni phase precipitates outside (i.e., farthest from the centre of rotation), and lighter graphite particulates on the inside periphery of the casting (i.e., nearest the centre of rotation). Internal concentrations of graphite particulates create a unique three-dimensional network (beginning of zone IV) inside of which grow the primary blocky Al3Ni phase precipitates during further decrease of the temperature. The movement of those precipitates, in comparison with other areas of the centrifugal casting, is severely restricted. As a result, these precipitates lodge themselves between relatively large graphite particulates. Further rotation and cooling of the suspension is accompanied by flowing out of the liquid metal into zones II and III, in which the pseudo-binary eutectic Al-Al3Ni crystallizes first at a temperature of 577°C with a preferred orientation. In the last stages of solidification at a temperature of 547°C, the crystallization of ternary eutectics Al-Al3Ni-Si finally occurs.

Fig. I.2.7. Structure of A356/5.7 vol. % Ni-coated graphite (50:50) composite made by vortex method and gravity cast into a grey iron die. As-cast. Etched with Keller reagent

100 µm

Particulate dispersion-reinforced composites 41

Fig. I.2.8. Structure of centrifugally cast A356/5.7 vol. % Ni-coated graphite (50:50) composite. The sequence of structure

evolution is shown, beginning with the external layer, through the middle zone, the transition zone, and ending with the internal (near centre) zone. Remarks on the picture refer to Figs. I.2.9 and I.2.10

a

b

c

d

a-b c-d e-f g-h i-j k-l m-n o-p r-s t-u

I II III IV

42 Particulate dispersion-reinforced composites

e

f

g

h

i

j

k l

Particulate dispersion-reinforced composites 43

m n

o p

r s

t u

Fig. I.2.9. Structure of centrifugally cast A356/5.7 vol. % Ni-coated graphite (50:50) composite. The sequence of structure evolution is shown, beginning with the external layer I (a - h), through the middle zone II (i - l), the transition zone III (m, n), and ending with the internal zone IV (o - u). Non-etched sample. Conventional light (on the left) and phase contrast (on the right). 100x

44 Particulate dispersion-reinforced composites

a

b

c

d

e

f

g h

Particulate dispersion-reinforced composites 45

i

j

k

l

m

n

o p

46 Particulate dispersion-reinforced composites

r

s

t u

Fig. I.2.10. Structure of centrifugally cast A356/5.7 vol.% Ni-coated graphite (50:50) composite. The sequence of structure

evolution is shown, beginning with the external layer I (a - h), through the middle zone II (i - l), the transition zone III (m, n), and ending with the internal zone IV (o - u). Non-etched sample. Conventional light (on the left) and phase contrast (on the right). 500x

Particulate dispersion-reinforced composites 47

I.3. Aluminum composites containing silicon carbide 2014/20 vol. % SiC composite Matrix: 2014 aluminum alloy (AlCu4SiMg) – 4.40Cu, 0.80Si, 0.8Mn, 0.60Mg, 0.30 Fe (wt. %). Reinforcing phase: single crystal silicon carbide platelets (hexagonal, α-SiC) (from Third Millennium Technologies, Inc., USA); of average platelet sizes 20-70 µm diameter x 0.50-5 µm thickness; as-received (i.e., uncoated). Manufacturing method: mechanical agitation using ultrasonic vibrations and gravity - die-casting. Pre-heated SiC platelets were spread over the melt poured in a graphite mould at 800°C, and the surface of the melt was touched with a titanium horn connected to an ultrasound generator. Structural characteristics: Only the finest variety of SiC platelets (20-70 µm x 0.50-5 µm) was dispersed in the melt using ultrasound as it was thought to present the greatest challenge to immersion and dispersion. Fig. I.3.1 shows a uniform distribution of the SiC in the matrix; no excessive agglomeration of SiC platelets was observed. The matrix structure showed finer dendrites as compared to a stir-cast composite to which Mg had been added for improved wettability. There was virtually no gas porosity evident at optical magnifications.

Fig. I.3.1. Structure of 2014l/xSiC graphite composite obtained by mechanical stirring coupled with ultrasonic vibration and

then gravity cast to mould. Non-etched sample. Conventional light

50 µm

48 Particulate dispersion-reinforced composites

Al/20 vol. % SiC composite Matrix: pure aluminum (99.9 wt. % Al). Reinforcing phase: nickel-coated silicon carbide particulates of 20-80 µm (SiCNi) diameter Manufacturing method: mechanical stirring method (vortex), squeeze casting. Structural characteristics (Fig. I.3.2): homogeneous arrangement of nickel-coated reinforcement phase particulates in the aluminum matrix, resulting from good susceptibility of nickel-coated areas to be wetted by liquid aluminum. This is linked to a process of dissolution of nickel coating in metal bath.

Fig. I.3.2. Structure of Al/20 vol. % Ni-coated SiC squeeze cast composite. Non-etched sample. Polarized light

20 µm

20 µm

Particulate dispersion-reinforced composites 49

F3S.15S (A359/15 vol. % SiC), F3S.20S (A359/20 vol. % SiC) as well as F3S.30S (A359/30 vol. % SiC) composite (F3S.xxS composites family) Matrix: A359 (AlSi9Mg) aluminum alloy – 8.50-9.50 Si, 0.45-0.65 Mg, 0.20 Fe, 0.20 Cu, 0.20 Ti (wt. %). Reinforcing phase: silicon carbide particulates of average diameter 20 µm. Manufacturing method: mechanical stirring method (vortex) similar to the method patented by DURALCAN Co., involving remelting and gravity casting into grey iron die and/or squeeze casting under the pressure of 230 MPa for 45 s (die temperature – bottom/top 190-230°C/170-200°C). The casting temperature of the composite slurry was maintained at 670-690°C. A cylinder (80 mm diameter and 40 mm height) was squeeze cast. A precipitation hardening heat treatment was applied: solutionizing at 538°C/8 h, cooling in water to 70°C, followed by aging at 154°C/5 h. Properties: increased mechanical properties, especially after squeeze casting (including fatigue properties), considerable improvement of tribological characteristics (significant reduction of wear, to the highest cast iron grades), a reduction in the coefficient of thermal expansion, an increase in the thermal conductivity. Structural characteristics (Figs. I.3.3 - I.3.8): SiC particulates are wetted by the liquid Al matrix and bond well with the matrix upon solidification. No structural discontinuities are observed at the interface between SiC particulates and metal matrix in either as-cast composite or fractured composite (Fig. I.3.3). The macro-scale structure of the F3S.20S gravity cast composite containing 20 vol. % SiC is characterized by a homogeneous distribution of the reinforcing phase particulates. However, at the micro-scale, SiC particulate agglomerates are observed in, the inter-dendritic regions, and within the surrounding (Al-Si)eut eutectic. This agglomeration is due to the particle pushing phenomenon created by growing (Si) phase, and taking place during composite solidification (Fig. I.3.4). The squeeze cast composite shows much higher degree of structure homogeneity resulting from the formation of Al(Si) solid solution. Another characteristic feature of such composites is the significant refinement of the primary (blocky) silicon crystals and modification of the eutectic silicon (Sieut) precipitates. Irrespective of the composite manufacturing method, there is a tendency for both Sieut phase and iron-rich phases to grow on the surface of the SiC particulates, which in turn leads to creation of bridges between reinforcing phase particulates. This is particularly clearly revealed under polarized light (Figs. I.3.4 e, I.3.5 e), when parts of SiC particulates become transparent, with an appropriate setting of the polarizer. Heat treatment of such composites causes an increase in the volume fraction of the Al(Si) phase precipitates as well as spheroidization of eutectic silicon precipitates. The effect of heat treatment on composite structure is shown in Fig. I.3.6. Observations of composite microstructure under polarized light (Fig. I.3.7) at optical magnifications demonstrate that the undesirable aluminum carbide (Al4C3) phase does not form. The second phase precipitates growing on SiC particulates (clearly visible under different polarizer settings, because SiC particle becomes transparent) represent eutectic silicon. An increase in SiC content to 30 vol. % is an additional factor which increases the homogeneity of the reinforcing phase distribution in the metal matrix (Fig. I.3.8).

50 Particulate dispersion-reinforced composites

Fig. I.3.3. SEM view of the fracture surface of squeeze cast F3S.15S (A359/15 vol. % SiC) composite. As-cast

a

b

c d

SiC

SiC

Particulate dispersion-reinforced composites 51

e

Fig. I.3.4. The structure of gravity cast F3S.20S composite (A359/20 vol. % SiC). As-cast. Non-etched sample. Conventional light (a, c, e) and phase contrast (b, d)

a

b

Si eut

52 Particulate dispersion-reinforced composites

c

d

e

Fig. I.3.5. The structure of squeeze cast F3S.20S composite (A359/20 vol. % SiC). As-cast. Non-etched sample. Conventional light (a, c, e) and phase contrast (b, d)

Particulate dispersion-reinforced composites 53

a

d

b

e

c f Fig. I.3.6. The influence of heat treatment on the structure of squeeze cast F3S.15S composite (A359/15 vol. % SiC):

(a-c) as-cast, (d-f) after heat treatment. Non-etched sample. Polarized light

a b

54 Particulate dispersion-reinforced composites

c

Fig. I.3.7. The structure of squeeze cast F3S.15S composite (A359/15 vol. % SiC) after heat treatment, shown under

polarized light (various images of the same place taken under different settings of microscope polarizer), illustrating growth of eutectic silicon precipitates on the SiC particulates. Non-etched sample

Fig. I.3.8. The structure of F3S.30S composite (A359/30 vol. % SiC). Non-etched sample. Phase contrast

Particulate dispersion-reinforced composites 55

2014/52 vol. % SiC composite Matrix: 2014 aluminum alloy (AlCu4SiMg) – 4.40Cu, 0.80Si, 0.8Mn, 0.60Mg, 0.30 Fe (wt. %). Reinforcing phase: single crystal silicon carbide platelets (hexagonal, α-SiC) (from Third Millennium Technologies, Inc., USA); of three average platelet sizes - 250-500 µm diameter and 25-50 µm thickness; 50-250 µm diameter and 5-25 µm thickness; and 20-70 µm diameter and 0.50-5 µm thickness; both as-received and copper-coated SiC platelets were utilized. Manufacturing method: counter-gravitational infiltration of packed binder-less silicon carbide platelets (packing density ~52%) by the molten 2014 aluminum alloy under nitrogen gas pressure; SiC platelets were packed in quartz tubes via multistage compaction of loose powders; infiltration was done in a stainless steel pressure vessel with a melting furnace, consisting of two semi-cylindrical heating elements, housed inside the pressure vessel; quartz tubes containing packed SiC were attached using leak-tight fittings to the lid of the pressure vessel which, upon closure, led to partial immersion of the tube in the molten alloy held in a graphite crucible inside the pressure vessel; an argon blanket was maintained over the SiC inside the quartz tube to reduce the oxidation of the liquid front; preheating time of SiC beds: 3-5 min; melt temperature: 800°C; nitrogen gas pressure: variable, 525-850 kPa; infiltration time: 3-5 min; slow cooling at the conclusion of infiltration in ambient air. Both as received and copper-coated SiC platelets were infiltrated. Cu-coating was done to promote the wettability and facilitate the infiltration. Copper coating was deposited on SiC using an electroless plating technique. For this, SiC platelets were stirred for 3 h in a solution of 10 g/l of SnCl2 and 40 ml/l of concentrated HCl. The platelets were then stirred for 15 min in a solution 0.5 g/liter PdCl2 and 1 ml/l concentrated HCl. After stirring, the solution was decanted, and platelets washed twice with distilled water. Finally, the platelets were stirred for 15 min in the plating bath of following composition: 10 g/liter CuSO4

.5H2O, 10 g/liter NaOH, 50 g/l sodium potassium tartrate, and 10 ml/l formaldehyde. The coating deposition was confirmed with the SEM examination. Properties: considerable improvement of tribological characteristics expected, significant reduction in thermal expansion coefficient and increase in the thermal conductivity, increased hardness and strength, and decreased ductility. Structural characteristics: A high volume fraction and a uniform distribution of SiC platelets are achieved in the SiC/2014Al composites via infiltration of fine and coarse SiC (Fig. I.3.9); however, some metal-starved cavities persist in the cast microstructure because of the difficulty of filling small gaps between closely-spaced or nearly-touching platelets. Color-etching of polished samples reveals the presence of micro-segregation in the alloy matrix between coarse (250-500 µm x 25-50 µm) SiC platelets (Fig. I.3.10); micro-segregation is nearly completely erased when inter-particle spacing becomes very small as in the case of fine (20-70 µm x 0.50-5) SiC platelets (Fig. I.3.11). The growth of primary Al dendrites is physically constrained by the SiC platelets as shown in Fig. I.3.12, and the inter-dendritic eutectic deposits on the SiC. A number of secondary phases in the alloy matrix of the infiltrated composite bars (Fig. I.3.13) are revealed by color etching, which include Al15(CuFeMn)3Si2, CuAl2, Al20Cu2Mn3, Mg2Si and CuAl2Mg.

a b

Fig. I.3.9. Structure of 2014/xSiC composites transverse to the infiltration direction, pressure: 710 kPa, melt temperature:

800°C, infiltration time: 3 min; as-received SiC platelet size: 20-70 µm diameter and 0.50 - 5.0 µm thickness (a); 250-500 µm diameter and 25-50 µm thickness (b)

20 µm

20 µm

56 Particulate dispersion-reinforced composites

Fig. I.3.10. Structure of 2014/xSiC composite transverse to the infiltration direction showing SiC distribution and micro-segregation; pressure: 525 kPa, melt temperature: 800°C, infiltration time: 5 min, SiC size: 250-500 µm diameter and 25-50 µm thickness, as-received SiC. The sample was color etched with a solution of 4 g KMnO4 and 2 g NaOH in 1 l distilled water

Fig. I.3.11. Structure of 2014/xSiC composite showing suppression of coring in the solidified matrix due to diffusional

constraints imposed by small inter-particle spacing in compacts made from fine SiC platelets; pressure: 750 kPa, melt temperature: 800ºC, infiltration time: 5 min, SiC size: 20-70 µm diameter and 0.50 - 5.0 µm thickness, as-received SiC. The sample was color etched with a solution of 4 g KMnO4 and 2 g NaOH in 1 l distilled water

200 µm

10 µm

Particulate dispersion-reinforced composites 57

a b Fig. I.3.12. Structure of 2014/xSiC composite showing two regions. Section transverse to the infiltration direction showing

physically constrained growth of Al dendrites, the inter-dendritic eutectic, and micro-segregation in cored dendrites growing within the network of packed SiC platelets; pressure: 825 kPa, melt temperature: 800°C, infiltration time: 5 min, SiC size: 50-250 µm diameter and 5-25 µm thickness, Cu-coated SiC; the sample was color etched with a solution of 4 g KMnO4 and 2 g NaOH in 1 l distilled water

a b Fig. I.3.13. Structure of two regions of a color-etched 2014/SiC composite transverse to the infiltration direction showing

secondary phase precipitation in the inter-particle regions. The composite was produced by gas pressure infiltration using nitrogen gas at a pressure of 525 kPa (melt temperature: 800°C, infiltration time: 5 min). The sample was color etched with a solution of 4 g KMnO4 and 2 g NaOH in 1 l distilled water. The phases that form during normal solidification of 2014 Al alloy include: Al15(CuFeMn)3Si2, CuAl2, Al20Cu2Mn3, Mg2Si and CuAl2Mg

10 µm

10 µm

10 µm

10 µm

58 Particulate dispersion-reinforced composites

AK12/60 vol. % SiC composite Matrix: AK12 aluminum alloy (AlSi12CuNiMg) – 12.20 Si, 1.15 Cu, 0.95 Ni, 0.90 Mg, 0.60 Fe, 0.20 Zn, 0.19 Mn, 0.09 Ti (wt. %). Reinforcing phase: silicon carbide particulates in the shape of a preform of 40 vol. % porosity. Manufacturing method: squeeze infiltration method under a pressure of 230 MPa for 45 s (die temperature - bottom: 200-220°C, pressure core: 190-200°C). The pouring temperature of AK12 alloy: 720°C. Properties: considerable improvement of tribological characteristics expected, significant reduction in thermal expansion coefficient and increase in the thermal conductivity. Structural characteristics (Fig. I.3.14 - I.3.15): further increase of reinforcing phase content up to 60 vol. % is possible because the preform of 40 vol. % porosity is fully infiltrated with liquid matrix under pressure. Structural analysis of such composites shows incomplete local filling of the areas between particulates with metal that are particularly clearly revealed under polarized light (Fig. I.3.9). However, this is only an apparent effect that results from the use of binder to fabricate the rigid porous preform. This is because liquid binder wets and links individual reinforcing phase particulates via bridges, which resemble structural discontinuities in appearance. Such a composite, with the reinforcing phase content over 50 vol. %, should in fact be called a ceramic-matrix composite material.

a b

c

Fig. I.3.14. Structure of squeeze cast AK12/60 vol. % SiC composite (silicon carbide preform, infiltrated with AlSi12CuNiMg

alloy under pressure. Non-etched sample. Phase contrast

Particulate dispersion-reinforced composites 59

a

b

c

Fig. I.3.15. Structure of squeeze cast AK12/60 vol. % SiC composite (silicon carbide preform infiltrated with AlSi12CuNiMg alloy under pressure). Non-etched sample. Polarized light

60 Particulate dispersion-reinforced composites

Al/70 vol. % SiC composite Matrix: pure aluminum (99.9 wt. % Al). Reinforcing phase: preform made of SiC particulates bonded using liquid glass binder. Preform porosity about 30 vol. %. Manufacturing method: gas pressure infiltration in an autoclave under a pressure of 7 MPa at 800°C for 30 min. Structural characteristics (Fig. I.3.16): complete preform infiltration with liquid matrix. Examination under the optical microscope magnifications does not disclose the existence of aluminum carbide precipitates, despite relatively high temperature as well as extended reaction time. This is related with the use of colloidal silicon dioxide as a binder (thin layer of SiO2 protects the SiC particulates from their direct contact with liquid aluminum).

Fig. I.3.16. Structure of Al/70% vol. % SiC composite, manufactured by gas pressure infiltration in the autoclave.

Non-etched sample. Polarized light

15 µm

Particulate dispersion-reinforced composites 61

I.4. Aluminum composites containing aluminum oxide A356/20 vol. % aluminum oxide composite (COMRALTM) Matrix: A356 aluminum alloy (AlSi7Mg) - 7.2 Si, 0.40 Mg, 0.60 Fe, 0.30 Mn (wt. %). Reinforcing phase: microgranules (incorrectly known in the literature as microspheres) of MICRAL-20 type, based on aluminum oxide (α−Al2O3 + mullite). Manufacturing method: mechanical stirring method (vortex). Following melting, the composite was subjected to squeeze casting under a pressure of 230 MPa for 45 s (die temperature - bottom/top: 190-230/170-200°C). The casting temperature of the composite suspension was maintained at 670-690°C. A cylindrical casting (70 mm diameter and 90 mm height) was produced but it was not heat treated. Properties: increased mechanical properties, considerable friction and wear resistance, and thermal shock resistance, decreased thermal expansion and conductivity, good machinability. Structural characteristics (Fig. 1.4.1-1.4.3): at the macro scale in the gravity cast composites, microgranules of aluminum oxide are homogeneously distributed. But at a micro-scale, there is a clear tendency for the reinforcing phase to gather in the interdendritic areas, as a result of the particle pushing phenomenon. The castings manufactured using squeeze casting method display significantly smaller segregation and more homogenous arrangement of microgranules in the metal matrix. This is linked with an essential influence of pressure on crystallization, and, consequently, with the refinement of metal matrix structure (including, both Al solid solution as well as Al-Si eutectic).

Fig. 1.4.1. SEM image of the reinforcing MICRAL-20 microgranules based on aluminum oxide with fine spinel particulates on their surface

10 µm

62 Particulate dispersion-reinforced composites

a

b

c

d

e

f g

Particulate dispersion-reinforced composites 63

h i

Fig. I.4.2. Structure of A356/20 vol. %. aluminum oxide (COMRALTM) composite, gravity cast into a metal die. Non-etched

sample. Conventional light (a, b, d, f) and phase contrast (c, e, g, i)

a

b

Fig. I.4.3. Structure of squeeze cast A356/20 vol. % aluminum oxide (COMRALTM) composite. Non-etched sample. Polarized light

20 µm

20 µm

64 Particulate dispersion-reinforced composites

I.5. Aluminum composites containing graphite and silicon carbide (hybrid) A356/15 vol. % SiC/5 vol. % graphite composite Matrix: A356 aluminum alloy (AlSi7Mg) - 7.2 Si, 0.40 Mg, 0.60 Fe, 0.30 Mn (wt. % ). Reinforcing phase: silicon carbide particulates of average 20 µm diameter and graphite particulates of about 100 µm diameter. Manufacturing method: mechanical stirring method (vortex) based on displacement neutral buoyancy principle. Properties: unique set of tribological properties, improved machinability (in comparison with silicon carbide-reinforced aluminum composites, e.g. DURALCAN composites). Structural characteristics (Fig. I.5.1) the gravity cast hybrid composites are characterized by very uniform arrangement of both reinforcing phases in the metal matrix. A high degree of metal-ceramic slurry stability, resulting from the displacement neutral buoyancy properties, has great significance in creating such structure during the slurry formation and its subsequent casting.

a

b

Fig. I.5.1. Structure of hybrid composite produced from A356 alloy reinforced with 15 vol. % SiC and 5 vol % graphite, gravity cast into die. Non-etched sample. Phase contrast (a) and conventional light (b)

Particulate dispersion-reinforced composites 65

A356/15 vol. % SiC/5 vol. % graphite composite Matrix: A356 aluminum alloy – 7.2 Si, 0.40 Mg, 0.60 Fe, 0.30 Mn (wt. %). Reinforcing phase: silicon carbide particulates of average 20 µm diameter and graphite particulates of about 100 µm diameter. Manufacturing method: mechanical stirring method (vortex) and centrifugal casting in graphite mould of 100 mm internal diameter, rotational speed of 1000 rpm. Properties: unique set of tribological properties, varying radially outward from the cylindrical cast axis, self-lubricating properties of the internal layer (graphite-enriched) and high resistance to wear of the external layer (silicon carbide-enriched). Structural characteristics (Fig. 1.5.2 - 1.5.3): Segregation of particulates of both reinforcing phases and shaping of the laminar structure indicate the important influence of the centrifugal force on the spatial distribution of the reinforcing phases according to their densities. SiC particulates being heavier than the metal-matrix are driven outward towards the external layer of the centrifugal cast. In contrast, lighter graphite particles tend to migrate towards internal layer, while the transition layer is devoid of the reinforcing phase. Besides SiC particulates, individual graphite particles of very large dimensions are observed in the external zone. In the internal zone, there are gas bubbles besides graphite. In comparison to the reinforced areas, the thickening of dendrites in the metal matrix is visible within this layer, which is devoid of SiC and graphite particulates. This suggests that the reinforcing particles restrict the growth of dendrites.

Fig. I.5.2. Structure of centrifugally cast A356/15vol.%SiC/5vol.% graphite composite. The sequence of the layers structure

is shown: graphite-rich internal layer (a-b, c-d, e-f), transition layer (g-h), and SiC-rich middle - external layers (i-j, k-l, m-n, o-p)

a-b c-d e-f g-h i-j k-l m-n o-p

66 Particulate dispersion-reinforced composites

a

b

c

d

e

f

g h

Particulate dispersion-reinforced composites 67

i j

k l

m n

o p

Fig. I.5.3. Structure of centrifugally cast A356/15vol.%SiC/5vol.%Gr composite. Internal layer (a-f), transition layer (g-h), and middle - external layer (i-p). Non-etched sample. Conventional light (on the left) and phase contrast (on the right). 100x

68 Particulate dispersion-reinforced composites

a

b

c

d

e

f

g h

Particulate dispersion-reinforced composites 69

i j

k l

m n

o p

Fig. I.5.4. Structure of centrifugally cast A356/15vol.%SiC/5vol.%Gr composite. Internal layer (a-f), transition layer (g-h),

and middle - external layer (i-p). Non-etched sample. Conventional light (on the left) and phase contrast (on the right). 500x

70 Particulate dispersion-reinforced composites

I.6. Copper composites containing graphite C90300/7.5 vol. % graphite composite Matrix: C90300 copper alloy (CuSn8Zn4) - 8 % Sn, 4 % Zn (wt. %). Reinforcing phase: graphite particulates of average 5 µm diameter. Manufacturing method: mechanical stirring method (vortex), gravity cast into graphite mould or grey iron die. The composite slurry casting temperature was maintained at about 1050°C. Cylindrical castings were made (40 mm diameter and 120 mm height). Properties: lower density, improved tribological characteristics (decreased wear and friction coefficients in tests with steel counter sample); a decrease in the thermal expansion coefficient can be expected. Structural characteristics (Figs. I.6.1 and I.6.2): despite a large difference in the densities of graphite and copper alloy, a homogeneous arrangement of the reinforcing phase in the metal matrix has been obtained both along the length and cross-section of the gravity castings. This occurred because a special operation of de-agglomerating graphite particulates was implemented prior to their introduction into the metal bath. Composite material gravity cast into grey iron die (Fig. I.6.2) has more uniform distribution of graphite particles, compared to that gravity cast into graphite mould (Fig. I.6.1).

I

II

III

Fig. I.6.1. Structure of the C90300/7.5 vol. % graphite composite, gravity cast into a graphite mould: I – top of the casting,

II – middle of the casting, III – bottom of the casting. Non-etched sample

200 µm

Particulate dispersion-reinforced composites 71

Fig. I.6.2. Structure of C90300/7.5 vol. % graphite composite, obtained by vortex method and gravity cast in a grey iron die. Non-etched sample. Conventional light

72 Particulate dispersion-reinforced composites

C90300/7.5 vol. % graphite composite Matrix: C90300 copper alloy (CuSn8Zn) - 8 % Sn, 4 % Zn, (wt. %). Reinforcing phase: graphite particulates of average 5 µm dimensions Manufacturing method: mechanical mixing method (vortex) with centrifugal casting in graphite. Properties: lower density, improvement of the tribological characteristics (decreased wear and friction coefficient in tests with steel counter sample), particularly of the centrifugal castings internal layer. Thermal expansion coefficient decrease can be expected. Structural characteristics: photographs shown on Figs. I.6.3 - I.6.5 depict structural details of individual areas visible on the cross-section of the C90300/7.5 vol. % graphite composite casting, under conventional light and using phase contrast. Centrifugally cast composites have layered structure, formed as a result of centrifugal force and density differences between graphite and copper alloy. With clear division into three zones. I – coarse grained external layer, practically not containing graphite phase, but with structural discontinuities in a shape of porosity. II - transition layer of relatively homogenous fine graphite particulates arrangement and pulverised metal matrix structure as a result of the introduction of technological additive (titanium acc. to know how method) into the metal matrix (up to the 1 wt. % level) in order to improve wettability of graphite particulates by liquid copper alloy. Both graphite particulates alone as well as very fine (below 1 µm) precipitations occurring as a result of an interaction between graphite and titanium dissolved in the alloy, form heterogeneous crystallisation nuclei. They cause clear brake up of the matrix grain in this zone of the composite casting. III – internal layer, graphite enriched and containing graphite bubbles and agglomerates.

Fig. I.6.3. Structure of centrifugally cast C90300/7.5 vol. % graphite composite. The sequence of the layers structure evolution is shown: I – external layer (graphite free), II – transition layer, III – internal layer (graphite enriched). Non-etched sample. Conventional light. Markings on the picture refer to Fig. I.6.4 and I.6.5

Structural pictures shown indicate various reasons for the porosity to occur depending on centrifugal casting spot being analyzed. In the external graphite free layer, the sporadic gas shrinkage porosity occurs (with a domination of shrinkage porosity) typical for the matrix alloy. In the internal layer however, enriched with graphite, the occurrence of large gas bubbles is linked with gassing up of the liquid metal-ceramic suspension (typical for the vortex method) and formation of large graphite particulates agglomerates, combined with bas bubbles, being pushed out towards the centre of the cylindrical casting under the centrifugal force.

a-b c-d e-f g-h i-j k-l m-n

I II III

Particulate dispersion-reinforced composites 73

a

b

c

d

e

f

g h

74 Particulate dispersion-reinforced composites

i j

k l

m n

Fig. I.6.4. Structure of centrifugally cast C90300/7.5 vol. % graphite composite. The sequence of the layers structure evolution is shown beginning with external layer I (a - h), through the middle one II (i - j), ending with III (k-n). Non-etched sample. Conventional light (on the left) and phase contrast (on the right). 100x

a b

Particulate dispersion-reinforced composites 75

c

d

e

f

g

h

i j

76 Particulate dispersion-reinforced composites

k

l

m n

Fig. I.6.5. Structure of centrifugally cast C90300/7.5 vol. % graphite composite. The sequence of the layers structure

evolution is shown beginning with external layer I (a - h), through the middle one II (i - j), ending with III (k-n). Non-etched sample. Conventional light (on the left) and phase contrast (on the right). 500x

Particulate dispersion-reinforced composites 77

Cu-Ni/Gr composite Matrix: pure copper (initial), Cu-Ni alloy. Reinforcing phase: nickel-coated graphite (50 : 50). Manufacturing method: gravity-assisted infiltration of loose nickel-coated graphite deposit (50 : 50) conducted in a graphite mould. Properties: improved tribological characteristics, especially at increased temperatures. Structural characteristics (Fig. I.6.5): very homogeneous arrangement of graphite particulates in the metal matrix. Good infiltration of the inter-particle cavities between the graphite particulates because of the dissolution of nickel coating in pure copper. Nickel segregation in the metal matrix takes place.

Fig. I.6.5. Structure of Cu-Ni/Gr composite obtained by gravity-assisted infiltration of loose nickel-coated graphite deposit

(50 wt. % Ni). Non-etched sample. Conventional light

78 Particulate dispersion-reinforced composites

CuFe10Sn7Zn3/2.0 vol. % graphite composite Matrix: CuFe10Sn7Zn3 copper alloy Reinforcing phase: flaky graphite. Manufacturing method: mechanical stirring method (vortex) and centrifugal casting in a graphite mould of 100 mm internal diameter, rotational speed: 1000 rpm. Properties: improved tribological characteristics expected (decreased wear and friction coefficients, particularly of the centrifugally cast internal layer). A decrease in the coefficient of thermal expansion can be expected. Structural characteristics: (Fig. I.6.6): during stirring of the composite slurry together with ground cast iron pieces containing flaky graphite, the dissolution of the iron in the liquid metal matrix takes place, resulting in in-situ matrix reinforcement with graphite particulates. During centrifugal casting, segregation of the metal suspension takes place because of the reinforcing phase segregation under centrifugal force. This is because "freed" graphite particulates gather in the internal layer of the centrifugal casing, while external layer becomes enriched with iron phase and contains pieces of partially melted cast iron.

a

b

c d

Cast iron

Cast iron

Cu

Cu

Particulate dispersion-reinforced composites 79

e f

g h

i j

k l

Fig. I.6.6. Structure of centrifugally cast CuFe10Sn7Zn3/2.0 vol. % graphite composite: (a-d) external layer, (e-h) transition

layer, (i-l) internal layer. Non-etched sample. Polarized light

80 Particulate dispersion-reinforced composites

I.7. Lead composites containing fly ash Pb/60 vol. % fly ash composite Matrix: lead containing 98.7 wt. % Pb. Reinforcing phase: fly ash from Dayton Power and Light Co. power stations (USA), cenospheres of chemical content (wt. %): 59.8 SiO2, 32.2 Al2O3, 5.1 Fe2O3, 0.4 CaO, 1.4 K2O, 0.8 MgO, 0.8 TiO2, 0.2 Na2O, 0.2 SO3, 0.3 P2O3. Manufacturing method: squeeze infiltration of the porous preform (porosity 40 vol. %) made of cenospheres, under a pressure of 10 MPa in 90 s (die temperature - bottom/top: 80/40°C). Lead was poured at 320°C. Squeeze cast cylinder of 50 mm diameter and 120 mm height was made. Properties: significantly lower density (approximately 5 g/cm3), decreased thermal expansion, increased energy and sound absorption characteristics, good machinability, and advantageous economic and ecological manufacturing parameters. Structural characteristics (Fig. I.7.1): in spite of a lack of wettability between liquid lead and fly ash, the application of external pressure makes it possible to produce a composite material of high reinforcing phase content, and devoid of structural discontinuities (e.g., porosity) between particulates. Some of the fly-ash particles, predominantly containing cracks or open porosity, are filled with liquid matrix during infiltration.

Fig. I.7.1. Structure of Pb/60 vol. % fly ash composite obtained by squeeze infiltration. Non-etched sample. Polarized light

20 µm

Particulate dispersion-reinforced composites 81

1.8. Lead composites containing graphite Pb/60 vol. % graphite composite Matrix: lead containing 98.7 wt. % Pb. Reinforcing phase: nickel-coated graphite particulates (NOVAMET, 50 wt. % Ni) of average dimension 75-150 µm. Manufacturing method: squeeze infiltration of the porous preform (porosity 40 vol. %) of nickel-coated graphite under a pressure of 60 MPa for 180 s (die temperature - top/bottom: 90/110°C). The matrix metal was poured in the die at a temperature of 320°C. Squeeze cast cylinder of 50 mm diameter and 120 mm height was made. Properties: essentially lowered density, increased energy and sound absorption characteristics, good machining properties, decreased thermal expansion. Structural characteristics (Fig. I.8.1): uniform distribution of reinforcing phase in Pb matrix and good filling of high volume fraction porous reinforcing preform with liquid matrix are obtained because of two effects, i.e. covering of graphite particles with metallic Ni layer and application of high pressure.

a

82 Particulate dispersion-reinforced composites

b

c

d

e

f g Fig. I.8.1. Structure of Pb/60 vol. % GrNi composite obtained by squeeze infiltration. Non-etched sample. Conventional light

(a, b, d, f) and phase contrast (c, e, g)

Short fiber-reinforced composites 83

II. Short fiber-reinforced composites II.1. Aluminum oxide-reinforced aluminum composites AK12/78 vol. % aluminum oxide composite Matrix: AK12 aluminum alloy (AlSi12CuNiMg) – 12.2 Si, 1.15 Cu, 0.95 Ni, 0.90 Mg, 0.60 Fe, 0.20 Zn 0.19 Mn, 0.09 Ti (wt. % ). Reinforcing phase: Short fiber preform based on aluminum oxide of initial porosity 78±2 %, made by MORGAN Ltd. (chemical composition of the fiber made of Al2O3 according to the manufacturer: 96 wt. % Al2O3, 4 % SiO2). Individual fibers bonded with colloidal silica in the amount of 5 wt. %, randomly spatially distributed and oriented. The preforms were of 95 mm external diameter and 30 mm height. Manufacturing method: squeeze infiltration of the preform (initially preheated up to 700°C at a maximum rate of 300°C/h) with a liquid aluminum alloy at 730°C, final pressure of 230 MPa, and infiltration time of 45 s; die temperature - bottom/ top: 220/190°C. Initial pressure ratio was 5-10 MPa during first 10 s of infiltration. Properties: improved mechanical properties at room and at elevated temperature, high resistance to thermal shock, decreased thermal expansion, electrical conductivity, and thermal conductivity, and increased resistance to wear. Structural characteristics (Fig. II.1.2): porous ceramic preform is characterized by a uniform arrangement of fibers in its entire volume (Fig. II.1.1). The application of external pressure for the infiltration (Fig. II.1.2) ensures complete filling of the preform cavities with the liquid metal and aids in establishing permanent and intimate contact between the reinforcing phase and the metal-matrix. At the same time, the external pressure does not cause distortion of the preform, which is evidenced by a homogenous spacing of the fibers in the matrix and a lack of directional orientation. The characteristic features of the microstructure of such composites are: 1) uniform infiltration of the preform by the liquid metal, 2) lack of structural discontinuities, and 3) growth of silicon phase precipitates (both eutectic silicon and primary silicon) on the surface of the aluminum oxide fibers; this is a very beneficial phenomenon because it enhances the interfacial bonding between matrix and reinforcing phase, thus favouring improved mechanical properties of the composite material.

Fig. II.1.1. SEM image of the preform made of short fibers based on Al2O3

84 Particulate dispersion-reinforced composites

a

b

c

Fig. II.1.2. Structure of AK12/78 vol. % Al2O3 composite produced by squeeze infiltration of porous preform (made from

short alumina fibers with SiO2 binder) with AlSi12CuNiMg alloy. Non-etched sample. Conventional light (a, b) and phase contrast (b)

Short fiber-reinforced composites 85

II.2. Carbon felt-reinforced aluminum alloys Al/carbon felt composite Matrix: pure aluminum (99.5 wt. % Al). Reinforcing phase: carbon felt preform of estimated initial porosity 70 vol. %, manufactured from fibrous organic raw material by POLGRAPH Company, Nowy Sącz (Poland). The preform consists of individual carbon fibers of 10-13 µm thickness with a random spatial distribution and in-plane orientation. The felt preforms alone were of 80 mm diameter and 40-50 mm height. Manufacturing method: squeeze infiltration of the thermally-activated carbon preform by liquid aluminum at 780°C, under a pressure of 230 MPa applied for 30 s; die temperature - bottom/top: 130-180/120-160°C. Properties: lower density, increased hardness, lower strength and plasticity, decreased thermal expansion, decreased thermal and electrical conductivity, good machinability, and advantageous economic and ecological manufacturing parameters. Structural characteristics: carbon felt in its initial state (Fig. II.2.1) forms entangled continuous fibers arranged in one direction and characterized by an inhomogeneous diameter and varying cross-section. Within the carbon felt-reinforced aluminum-matrix composite structure (Fig. II.2.2), the majority of the fibers retain their primary orientation, while the cavities between the fibers become completely filled with liquid matrix. Structural discontinuities, in the form of pores between agglomerates of many fibers occur rather sporadically (Fig. II.2.2 c). However detailed microscopic examination under polarized light leads to the conclusion that such pores are actually primary defects, which exist in the original felt prior to infiltration. These defects originate during the process of preform manufacture when some fibers may form permanent bond with some or even umpteen fibers in such a way that closed porosity forms inside this agglomerate. This closed porosity remains inaccessible to the liquid metal during infiltration even under large pressures. Aluminum carbide (Al4C3) precipitates were not found in the composite under optical microscopy magnifications. A characteristic feature of the carbon felt-reinforced composites is the presence of porosity inside individual fibers. This is primary porosity of the fibers, spread over their entire length, because fibers themselves are carbon tubes of different shapes and wall thicknesses (not necessarily coaxial).

Fig. II.2.1. SEM image showing carbon felt structure

86 Particulate dispersion-reinforced composites

a

b

c

d

e

Short fiber-reinforced composites 87

f

g

h i Fig. II.2.2. Structure of squeeze infiltrated Al/carbon felt composite. Non-etched sample. Conventional light (a, b, d, f, h) and

phase contrast (c, e, g, i)

88 Particulate dispersion-reinforced composites

AlTi4/carbon felt composite Matrix: AlTi4 aluminum alloy (3.8 wt. % Ti). Reinforcing phase: carbon felt preform of estimated initial porosity 70 vol. %, manufactured from fibrous organic material by POLGRAPH company, Nowy Sącz (Poland). The preform consists of individual carbon fibers of the 10-13 µm thickness with a random spatial distribution and in-plane orientation. The felt preforms alone were of 80 mm diameter and 40-50 mm height. Manufacturing method: squeeze infiltration of the thermally-activated carbon preform by liquid aluminum alloy at 1000°C, under a pressure of 230 MPa, for 30 s; die temperature - bottom/top: 130-180/120-160°C. Properties: lower density, increased hardness, lower strength and plasticity, decreased thermal expansion, decreased electrical and thermal conductivity, good machinability, and advantageous economic and ecological manufacturing parameters. Structural characteristics (Fig. II.2.3): the character of carbon fiber arrangement in the AlTi4 metal matrix is similar to that in pure aluminum. The distinguishing feature is the presence of the Al3Ti phase precipitates in the metal matrix alone, especially on the preform edges (Fig. II.2.3 a). Their rounded shape indicates that they are primary precipitates, which have not quite been dissolved in the matrix during its melting. Within the preform, the Al3Ti phase appears in the form of fine needles, often growing on the surface of the carbon phase, including the surfaces of the carbon flakes (Fig. II.2.3 b) as well as the carbon fibers (Fig. II.2.3 c). Figures II.2.3 b-d show structural defects, in the form of cracks, that occur occasionally in the carbon phase of the composite. An examination of the regions under polarized light revealed that the origin of these defects is linked with the carbon felt manufacturing process. In particular, these discontinuous form as a result of wetting of the whole group of fibers by a large amount of binder in such a way as to fuse the fibers inside large flakes of binder while maintaining the same uniform and directionally oriented arrangement as in remaining part of the felt. Also, because during the felt manufacturing process the binder gets carbonized, the homogeneity of the carbon phase areas is maintained. The fractured areas (cracks) appearing in such flakes are not filled with metal, which proves that they are formed after the infiltration process is complete, probably during cooling of the composite as a result of considerable solidification shrinkage, which is characteristic for the matrix alloy. The phenomenon of cracking of large flakes of the carbon phase does not occur in the case of other matrix alloys (e.g., AK12 alloy or pure Al).

a

Short fiber-reinforced composites 89

b

c

d

Fig. II.2.3. Structure of squeeze infiltrated AlTi4/carbon felt composite. Non-etched sample. Polarized light

90 Particulate dispersion-reinforced composites

AK12 / carbon felt composite Matrix: AK12 aluminum alloy (AlSi12CuNiMg) - 12.2 Si, 1.15 Cu, 0.95 Ni, 0.90 Mg, 0.60 Fe, 0.20 Zn 0.19 Mn, 0.09 Ti (wt. % ). Reinforcing phase: carbon felt preform of estimated initial porosity 70 vol. %, manufactured from fibrous organic material by POLGRAPH company, Nowy Sącz (Poland). The preform consists of individual carbon fibers of the 10-13 µm thickness with a random spatial distribution and in-plane orientation. The felt preforms alone were of 80 mm diameter and 40-50 mm height. Manufacturing method: squeeze infiltration of the thermally-activated carbon preform by the liquid aluminum alloy at 760°C, under a pressure of 230 MPa for 30 s; die temperature - bottom/top: 130-180/120-160°C. Properties: lower density, increased hardness, lower strength and plasticity, decreased thermal expansion, decreased electrical and thermal conductivity, good machinability, and advantageous economic and ecological manufacturing parameters. Structural characteristics (Fig. II.2.4): The characteristic feature of carbon felt-reinforced AK12 alloy-matrix composite is a complete lack of structural discontinuous inside the carbon preform and even very large flakes of the carbon phase do not show any minute cracks (Fig. II.2.4 a-f).

a

b

c d

Short fiber-reinforced composites 91

e f

g

h

Fig. II.2.4. Structure of squeeze infiltrated AK12/carbon felt composite. Non-etched sample. Polarized light (a-f) and

conventional light (g-h)

92 Particulate dispersion-reinforced composites

II.3. Carbon felt-reinforced zinc composites Zn/carbon felt composite Matrix: pure zinc (99.6 wt. % Zn). Reinforcing phase: carbon felt preform of estimated initial porosity 70 vol. %, manufactured from fibrous organic material by POLGRAPH company, Nowy Sącz (Poland). The preform consists of individual carbon fibers of the 10-13 µm thickness with a random spatial distribution and in-plane orientation. The felt preforms alone were of 80 mm diameter and 40-50 mm height. Manufacturing method: squeeze infiltration of the thermally-activated carbon preform by liquid zinc at 550°C, under a pressure of 230 MPa for 30 s; die temperature - bottom/top: 130-180/120-160°C. Properties: lower density, increased hardness, lower strength and plasticity, decreased thermal expansion, decreased electrical and thermal conductivity, good machinability, and advantageous economic and ecological manufacturing parameters. Structural characteristics (Fig. II.3.1): under conventional light, the composite structure (Fig. II.3.1 a-e) shows defects in the form of structural discontinuities (due to the lack of metal infiltration into the enclosed space between aggregates of permanently bonded fibers) (Fig. II.3.1 b-d). In contrast to the Al matrix, in the case of Zn matrix, the minute cracks were found in large carbon phase flakes, and some of these cracks were filled with metal (Fig. II.3.1 e). The use of polarized light allows observation of the zinc matrix structure, i.e. its directional solidification and twinning (Fig. II.3.1 f-o). Within the area around a single grain of the matrix phase, the directional arrangement of twins is the same as the direction of fracture in the carbon phase flake in this grain (Fig. II.3.1 o).

a

b

c d

Short fiber-reinforced composites 93

e

f

g

h

i

j k

94 Particulate dispersion-reinforced composites

l

m

n

Fig. II.3.1. Structure of squeeze infiltrated Zn/carbon felt composite. Non-etched sample. Conventional light (a-e) and

polarized light (f-n)

Short fiber-reinforced composites 95

II.4. Carbon felt-reinforced lead composites Pb / carbon felt composite Matrix: pure lead (99.6 wt. % Pb). Reinforcing phase: carbon felt preform of estimated initial porosity 70 vol. %, manufactured from fibrous organic material by POLGRAPH company, Nowy Sącz (Poland). The preform consists of individual carbon fibers of the 10-13 µm thickness with a random spatial distribution and in-plane orientation. The felt preforms alone were of 80 mm diameter and 40-50 mm height. Manufacturing method: squeeze infiltration of the thermally-activated carbon preform by liquid lead at 480°C, under a pressure of 230 MPa for 30 s; die temperature: bottom - 130-180°C, plunger 120-160°C Properties: substantially lower density, increased hardness and strength, lower plasticity, decreased thermal expansion, decreased electrical (insignificantly) and thermal conductivity, good machinability, and advantageous economic and ecological manufacturing parameters. Structural characteristics: selected technological process ensures good infiltration of the fibrous preform with liquid lead together with the establishment of intimate contact between the reinforcement and the matrix as well as their good interfacial bonding.

Fig. II.4.1. SEM images of fracture of Pb/carbon felt composite

96 Long fiber-reinforced composites

III. Long fiber-reinforced composites III.1. Carbon fiber-reinforced aluminum composites AlTi4/carbon fiber composite Matrix: AlTi4 aluminum alloy (3.8 wt. % Ti). Reinforcing phase: long (continuous) carbon fiber of PAN type, shaped into porous preform. Individual carbon fibers of about 8 µm diameter, spatially distributed and oriented in one selected direction in the preform. Manufacturing method: squeeze infiltration of the thermally-activated carbon fiber preform by the liquid alloy at 1000°C, under a pressure of 230 MPa for 30 s; die temperature - bottom - 180°C, plunger: 170°C. Properties: substantially lower density (down to 2.15 g/cm3), increased stiffness and strength, and significantly lower thermal expansion (especially in the direction of fiber orientation). Structural characteristics: (Fig. III.1.1): in spite of the presence of a large volume fraction of the reinforcing phase in the preform, the use of squeeze infiltration causes complete filling of the carbon fiber preform by liquid metal matrix. This is the reason for the complete absence of structural discontinuities, such as porosity, which would form due to impossibility to locally wet the reinforcing phase (Fig. III.1.1). An additional factor, positively influencing the homogenous arrangement of the fibers in the metal matrix, is the use of titanium as an alloy additive to aluminum matrix. Titanium increases chemical interaction in the fiber/matrix system. It also contributes to improving physiochemical compatibility of the composite components by slowing down the formation of undesirable aluminum carbide (Al4C3) while at the same time ensuring an improvement in wettability of carbon fibers by the liquid matrix.

a

Long fiber-reinforced composites 97

b

c

d

e

f g

Fig. III.1.1. Structure of squeeze infiltrated AlTi4/carbon fiber composite. Non-etched sample. Conventional light (a, b, d, f) and phase contrast (c, e, g)

98 Long fiber-reinforced composites

Al/carbon fiber composite (nickel-coated) Matrix: pure aluminum (99,99 wt. % Al). Reinforcing phase: long (continuous) carbon fiber from CYANAMID (USA) nickel coated (50 wt. % Ni), formed into a preform. The preform consists of individual nickel-coated carbon fibers of about 10-18 µm thickness spatially distributed and oriented in one selected direction. Manufacturing method: squeeze infiltration of the thermally-activated carbon fiber preform by liquid aluminum at 1000°C, under a pressure of 230 MPa for 30 s; die temperature – bottom/top: 180/170°C. Properties: increased rigidity and strength, lower thermal expansion (especially in the direction of fiber orientation) Structural characteristics (Fig. III.1.2): the Al-matrix composite, reinforced with nickel-coated carbon fibers (Cf(Ni)), is characterized by relatively homogenous arrangement of fibers in the metal matrix on a macro-scale (Fig. III.1.2 a). However, on a micro-scale, in spite of a very good wettability of nickel coated carbon fibers by the liquid Al, there are agglomerates of fibers in the composite (Fig. III.1.2 b) as well as structural discontinuities in the form of porosities inside these agglomerates. Structural analysis under higher magnifications (about 500x) shows that these are primary flaws occurring in the preform in the initial state before infiltration and preform solutioning with liquid metal. These primary flaws originate in the process of nickel coating, which is usually carried out on a bundle of a few hundred fibers. Additionally, occasional “bonding” of a few to a few dozens of fibers into one conglomerate may take place, which then gets coated with one common nickel layer. This creates enclosed porosity between fibers, which is impossible to be filled by liquid metal even under high pressure. Interestingly, in areas where heterogeneous arrangement of fibers can be observed, there is also heterogeneous thickness of the nickel coating, with the coating thickness even exceeding the fiber diameter itself. Under higher magnifications, eutectic Al-Al3Ni is revealed in grain boundaries, which results from the dissolution of the nickel coating in the liquid Al.

a

b

Fig. III.1.2. Structure of squeeze infiltrated Al/nickel coated carbon fiber composite. Non-etched sample. Conventional light

Long fiber-reinforced composites 99

III.2. Carbon fiber-reinforced copper composites Cu/carbon fiber composite Matrix: pure copper (99.7 wt. % Cu). Reinforcing phase: long (continuous) carbon fiber of PAN type, shaped into a preform. Individual carbon fibers of about 8 µm diameter, spatially distributed and oriented in one selected direction in the preform. The initial preform thickness was 5 mm, with relatively loosely arranged, but directionally oriented fibers. Manufacturing method: infiltration of the thermally-activated carbon fiber preform by liquid copper at 1200°C under a pressure of 230 MPa for 60 s; die temperature - bottom/top: 300/270°C. Properties: substantially lower density, increased rigidity and strength, significantly lower thermal expansion (especially in the direction of fiber orientation) Structural characteristics (Fig III.2.1): due to a relatively loose character of fiber arrangement in the preform and initially, during pressure application, the metal infiltration is accompanied by thickening of the preform in such a way as to obtain local reinforcement of the cast surface, up to 0.5 mm. The areas between the fibers are completely filled with liquid matrix; however, in areas where the fibers were not sufficiently stiffened, the loss in the homogeneity of fiber arrangement may occasionally occur in the metal matrix because of their displacement during infiltration with liquid metal and bunching together into agglomerates of several fibers (Fig. II.2.1 a, b). For comparison, in areas where the fibers are immobilised, their arrangement in the metal matrix is homogenous (Fig. III.2.1 c). Under large magnifications, micro-cracks are revealed inside individual fibers (visible in figures III.2.1 b, c). Their location and shape suggest that they are formed during preparation of the sample itself for the structural examination, most probably, during cutting of the composite casting.

a

100 Long fiber-reinforced composites

b

c

d e

Fig. III.2.1. Structure of squeeze infiltrated Cu/carbon fiber composite. Non-etched sample. Conventional light (a, b, d) and

phase contrast (c, e)

Long fiber-reinforced composites 101

III.3. Sapphire fiber-reinforced nickel composites NiAl/saphire fiber composite Matrix (at. %): nickel-base intermetallic NiAl (49.9-50.3Ni, 49.7-50.1Al), NiAl(Yb) (62.1-62.6Ni, 36.8-37.3Al, 0.36-0.11 Yb), NiAl(Cr) (60.2Ni-28.2Al-11.4Cr), and NiAl(W) (67.4Ni-31.4Al-1.5W). Reinforcing phase: long (continuous) single crystal sapphire fibers (SaphikonTM) with average diameter 178 µm. Manufacturing methods: two methods of composite manufacture were utilized: 1. NASA’s patented ‘Powder-cloth’ (PC) technique in which a cloth or sheet of atomized matrix

powders is created by mixing them with Teflon binders and a solvent to a dough-like consistency; fiber mats of specified thickness are produced by filament winding and by application of a PMMA coating; matrix cloth and fiber mats are then stacked in layers and vacuum hot pressed during which binders are removed; the orientation of individual plies and stacking sequence provide flexibility in designing the composite properties.

2. Floating-zone directional solidification (DS), which is a container-less solidification technique to

produce the composites. The DS was done on either powder-cloth feedstock with ~25 vol. % fibers, or on vacuum induction melted and chill cast (in Cu-mould) composites (the chill cast composites had low, <0.5%, fiber volume fraction and were fabricated by laying sapphire fibers in a Cu-mould with fiber ends rigidly supported within holes drilled in steel end caps). The DS was done under 3 psi (gauge) pressure of ultrahigh purity argon gas in a 25 kW radio frequency (RF, 400 kHz) vacuum induction furnace using a water-cooled copper current concentrator; a current adjusting type three-mode proportional controller was used to set the temperature and control the power input to the coil; DS was carried out by traversing a 8 mm long molten zone along the length of the feedstock at a constant speed (~6.0 cm/h).

Properties: substantially lower density, increased rigidity and strength at room-temperature and elevated temperatures, significantly lower thermal expansion (especially in the direction of fiber orientation), increased hardness, decreased ductility, and improved wear resistance are expected. With W and Cr alloying of NiAl and incorporation of sapphire in the composite, dual-phase toughening potential together with high-temperature strength is expected in the composite.

a b Fig. III.3.1. Structure of NiAl/saphire fibers composite made by “powder-cloth (PC)” method. The arrows in (b) denote the

residual micro-porosity from incomplete densification of the NiAl matrix powders during composite fabrication Structural characteristics: in the powder-cloth (PC) composite, polycrystalline NiAl matrix contains rows of sapphire fibers distributed in the matrix (Fig. III.3.1 a), and there is some residual porosity (marked by arrows in Fig. III.3.1 b). After hot pressing there is no evidence of interfacial reaction. The directional solidification (DS) of polycrystalline PC feedstock reduced the number of NiAl grain boundaries coincident on the fiber surface (Fig. III.3.2), which led to an increase in the room-temperature fiber/matrix interfacial shear strength. This is because the NiAl grain boundaries coincident on the fiber surface act as crack initiation sites during loading of the fiber. In a manner

200 µm

25 µm

102 Long fiber-reinforced composites

similar to the PC composites, there was no evidence of interfacial reaction layers in the DS composites; however, porosity in the matrix persisted after the DS. The surface appearance of sapphire fibers extracted the sapphire-NiAl matrix composites synthesized by different fabrication techniques is shown Fig. III.3.3; there is no evidence of extensive surface degradation or damage. The original fiber distribution of the PC feedstock was slightly altered after DS (Fig. III.3.2), but the overall uniformity of fiber distribution was not appreciably affected. The matrix structure of the DS composites (Fig. III.3.2) exhibited coarser grains, as compared to the NiAl matrix in the PC composites (Fig. III.3.1). Unlike Yb in NiAl, chromium alloying of NiAl only moderately attacks the sapphire; Cr preferentially precipitates on to the sapphire fiber, as revealed by the SEM examination (Fig. III.3.4) of fibers separated from the NiAl(Cr) matrix by etching. In addition, Cr bonds to the sapphire fiber (Fig. III.3.5). There is, however, no evidence of reaction product layer formation. The DS structure of the NiAl(Cr)/sapphire composite displays NiAl cells with fine Cr precipitates distributed within the cells and the eutectic in the intercellular colonies (Fig. III.3.6). Depending upon whether the fiber resides predominantly in the intercellular colonies or within the NiAl grains, preferential segregation of either eutectic Cr or primary Cr precipitates, respectively occurs at the fiber/matrix interface (Fig. III.3.7). In contrast to both Yb and Cr alloying of NiAl, with 1.5 at. % W in NiAl, the fiber/matrix interface is completely devoid of any reaction products and secondary precipitates (Fig. III.3.8.). The addition of W and Cr to NiAl and incorporation of sapphire in the composite creates fiber-reinforced composites with dual-phase toughening potential together with attainment of high-temperature strength.

a b Fig. III.3.2. Structure of NiAl/saphire fibers composite after directional solidification (DS) from PC feedstock. The arrows in

(a) indicate the porosity in the matrix. The DS was done in the floating-zone mode. DS reduced the number of NiAl grain boundaries coincident on the fiber surface, which led to an increase in the fiber/matrix interfacial shear strength. No interfacial reaction layers at optical magnifications. Non-etched sample

a b

Fig. III.3.3. SEM views of the surface of fibers extracted from NiAl/sapphire fibers composites processed under the following

conditions: (a) “powder-cloth” (PC) composites that were subsequently directionally solidified using the “floating-zone directional solidification” (DS) technique, (b) vacuum induction melted and chill cast (in copper mould) composites after DS in the floating-zone mode. The fibers were extracted from the composite by dissolving the matrix in a boiling solution of 75% acetic acid, 23% nitric acid, and 2% HCl acid. Average fiber diameter 178 µm. The matrix contained equi-atomic fractions of Ni and Al

200 µm

25 µm

100 µm

Long fiber-reinforced composites 103

a b Fig. III.3.4. SEM view of a sapphire fiber extracted from a sapphire-reinforced NiAl(Cr) composite at (a) low- and (b) high

magnifications. The composite was fabricated using the “floating-zone” directional solidification (DS) of vacuum induction melted and chill cast composite (fiber volume fraction ~0.5 wt. %, average sapphire fiber diameter 178 µm). The fiber surface shows good bonding between the chromium precipitates and the oxide fiber. The fibers were extracted from the composite by dissolving the matrix in a boiling solution of 75% acetic acid, 23% nitric acid, and 2% HCl acid

a b

c

Fig. III.3.5. SEM views of the fiber/matrix interface in a directionally solidified NiAl(Cr)/sapphire fibers composite: (a)

chromium-rich layer around a sapphire fiber embedded in the composite matrix, (b) a higher magnification view of a region of Fig. (a), and (c) the sapphire/NiAl(Cr) interface region showing intimate bonding between eutectic Cr particles and the sapphire fiber. Also noticeable is the cellular NiAl matrix structure and the fine secondary Cr precipitates that form from supersaturated NiAl matrix during cooling at the conclusion of solidification. The composite was produced by vacuum induction melting and chill casting in a copper mould followed by directional solidification (DS) in the “floating-zone” mode. The fiber volume fraction in the composite: ~0.5 wt. % with average fiber diameter of 178 µm

Sapphire fiber

Eutectic chromium

Eutectic chromium

Sapphire/NiAl(Cr) interface region

Cr precipitates

7 µm

1 µm

30 µm

2 µm

104 Long fiber-reinforced composites

a b Fig. III.3.6. SEM photomicrographs of a directionally solidified NiAl(Cr)/sapphire fibers composite in a fiber-free region

showing the cellular NiAl structure with intercellular NiAl-Cr eutectic (a), and a higher magnification view of a eutectic region (b). The DS was carried out in a floating-zone mode. The SEM views in (a) and (b) are transverse to the growth direction. The composition of the matrix alloy after DS (in at. %) is: 60.2Ni-28.2Al-11.4Cr and ~300 ppm oxygen

a b Fig. III.3.7. SEM views of de-bonded interfaces in DS NiAl(Cr)/sapphire fiber composite after the fiber push-out test. The DS

was done on vacuum induction melted and chill cast feedstock (fiber volume fraction < 0.5%). Front-face of a composite wafer showing secondary Cr precipitates within the NiAl cells and microplastic grooving of the NiAl matrix during sliding of the fiber (a), and preferential segregation of eutectic Cr at the fiber/matrix interface in the region of the displaced fiber (b). The composition of the DS matrix (in at. %) is 60.2Ni-28.2Al-11.4Cr and ~300 ppm oxygen

a b Fig. III.3.8. Fiber/matrix interface in NiAl(W)/sapphire composite produced by vacuum induction melting and chill casting in

a copper mould followed by directional solidification (DS) in the “floating-zone” mode. Reaction-free fiber/matrix interface on a polished surface (a) and within the region of a pushed fiber in a fiber push-out test (b). The fiber volume fraction in the composite: ~0.5 wt. %, average fiber diameter is 178 µm, and the composition of the matrix (in at. %) is 67.4Ni-31.4Al-1.5W and ~300 ppm oxygen

Fiber

Secondary Cr precipitates and

wear tracks

Eutectic Cr

Fiber

4 µm

2 µm

Fiber

Long fiber-reinforced composites 105

Nickel-base superalloy (Hastealloy)/sapphire fiber composite Matrix: Nickel-base superalloy of composition (in wt. %): 47.5Ni-21.5Cr-17.8Fe-8.3Mo-1.7Co-0.3Mn-0.4Si-0.2Al-0.1Cu-0.4Nb-0.06Ti-0.08C-0.020. Reinforcing phase: single crystal sapphire fibers (SaphikonTM), average fiber diameter: 130 µm. Manufacturing method: pressure infiltration casting in a pressure vessel that contained a pair of induction coils, powered with a RF generator, to independently heat preform and metal; preforms of uni-directionally aligned sapphire fibers in recrystallized alumina moulds; top-pour infiltration configuration; melt temperature: 1400-1500°C; infiltration time: 3-5 min; pressurizing gas: ultrahigh-purity argon; samples cooled at the conclusion of infiltration with argon gas inside the pressure chamber. Properties: increased hardness and strength at room-temperature and elevated-temperatures, strong interfacial bond, decreased ductility, improved wear resistance, and decreased thermal expansion are expected. Structural characteristics: good penetration of inter-fiber channels occurred (Fig. III.3.8); some fiber-to-fiber contact was unavoidable because of high fiber volume fractions. Short infiltration times and large fiber contents limited the extent of fiber attack by the reactive solutes (e.g., Cr) of the matrix. Structural observation of extracted fibers, however, showed evidence of chemical interaction.

a

b

c d

Fig. III.3.8. Structure of Hastealloy/sapphire fibers composite produced by pressure infiltration showing fibers distribution in

the cross-section of the infiltrated composite bar (a) and (b), and higher magnification views of the fiber arrangement, and capillary channels between fibers (c) and (d). Fiber diameter 130 µm

50 µm

10 µm

200 µm

500 µm

106 Gas „reinforced” composites (gasars)

IV. Gas "reinforced" composites (gasars) Copper and nickel gasars Matrix: pure copper, pure nickel Reinforcing phase: gas in the form of bubbles of diversified morphology, dimensions and contents in the metal matrix. Manufacturing method: directional crystallization under pressure from liquid metal saturated with gas mixture (hydrogen and argon or nitrogen) according to the know-how of a protected manufacturing technique, known as gasar process or lotus-like structures process. Properties: substantially lower density, increased mechanical properties, especially tensile strength and plasticity (generally in case of the bubbles of diameter below 1-10 µm range and their contents in the structure up to 30 vol. %) in the direction of bubble growth, increased damping of mechanical energy and sound, possibility to obtain continuous bubbles, i.e. by creating open porosity (permeable structures).

a

b

c d

Fig. IV.1. SEM images of the fractured copper gasars under various magnifications

300 µm

300 µm

70 µm

20 µm

Gas „reinforced” composites (gasars) 107

a

b

c

d

e

f

300 µm

300 µm

300 µm

300 µm

70 µm

70 µm

108 Gas „reinforced” composites (gasars)

g h Fig. IV.2. SEM images illustrating the structural changes in the nickel gasar, depending on the technological parameters of

the manufacturing process. Transverse cross-section (a, c, e, g) and lengthwise section (b, d, f, h) Structural specifics: gasars can be classified as a separate, specific group of metal composite materials manufactured using liquid-phase techniques. Gas being released from liquid metal (during directional solidification, under variable ambient pressure) is regarded as a reinforcing phase. For the purpose of practical applications, the most desirable bubble morphology is the one that closely resembles cylindrical, parallel and evenly arranged capillaries, allowing full permeability of gasars in a direction parallel to crystallization front to be achieved. Fig. VI.3. Proposed scheme of gasar growth kinetics based upon structural examination, during different time of directional

solidification (τ) at a constant velocity. The fracture surface of a copper gasars is shown as a final structure under original magnification of 35x (SEM)

Fig. IV.4. SEM micrographs and corresponding evaluation schemes of the nickel gasar structure formation beginning from

“quasiboiling” process (blister-like structure) (a) through the increase of crystallization velocity Vs (b, c, d) and development of required dendritic structure (e)

τ1<τ2<τ3

Solid phase

Gas phase Liquid Gas

Cooling direction

τ1 τ2 τ3 Final structure

Solid phase

Gas phase

Metallic liquid

Gas

Non-metallic inclusions

Cooling direction

20 µm

20 µm

Nano-composites 109

V. Nano-composites AlSi7Mg/ carbon nano-fibers composite Matrix: AK7 aluminum alloy (AlSi7Mg) - 7.24 Si, 0.35 Mg, 0.46 Fe, 0.20 Mn. Reinforcing phase: pipe-like carbon nano-fibers (Fig. V.1), thermally treated and known by the trade name - Pyrograf® - III (PR-19-LHT); these nano-fibers have high graphite content and unique combination of properties, especially high thermal conductivity, the highest out of all known reinforcing phases (Table V.1).

Fig. V.1. Structure of the Pyrograf® - III nano-fiber after catalytic elongation and thickening (CVD - chemical vapour deposition). Thickness of the nano-pipe wall in the initial stage is about 20 nm (± 20%)

Table V.1. Selected physical, mechanical and thermal properties of the Pyrograf® - III nano-fibers in relation to heat treatment

Properties Nano-fibers as-received Nano-fibers after heat treatment Thickness, g/cm3 1.8 2.1 Tensile strength, GPa 2.7 7.0 Young module, GPa 400 600 Thermal expansion coefficient, µm/m·K - -1.0 Electrical resistance, µΩ·cm 1000 55 Thermal conductivity, W/m·K 20 1950 Properties: lower density, significantly increased stiffness and strength, significantly lower thermal expansion, drastically increased thermal conductivity. Manufacturing method I: squeeze infiltration of the loose carbon fibers deposit by liquid alloy at 740°C, under a pressure of 120 MPa for 30 s; die temperature - bottom/top: 180/220°C. Structural characteristics (Fig. V.2): optical microscopy examination of nano-composites at the micro-scale has only illustrational value due to the very fine dimensions of individual nano-fibers, which are impossible to be. In the next stage of structural examinations, it is necessary to use high resolution electron microscopy (HREM). Nevertheless, even observations under an optical microscope show that arrangement of the reinforcing phase is non-uniform. There are areas of varying degree of nano-fibers density, which may not be properly described even as nano-fiber colonies.

110 Nano-composites

a

b c

d e

f g Fig. V.2. Structures of AK7/carbon nano-fibers composite obtained by squeeze infiltration of loose-bed nano-fibers

deposits. Transition layer: alloy-composite as well as structural areas illustrating arrangement of reinforcing phase under various magnifications. Non-etched sample. Conventional light (a, b, d, f), phase contrast (c, e, g)

Nano-composites 111

Manufacturing method II: squeeze infiltration of the thermally-activated carbon nano-fiber preform bonded with hydrated sodium silicate (about 75 vol. % porosity) with liquid alloy at 740°C under a pressure of 120 MPa for 30 s; die temperature - bottom/top: 180/220°C. Structural characteristic (Fig. V.3): observations under the optical microscope, especially with 500-1000x magnifications, show colonies of nano-fibers located in the preform walls, and being divided by a significant amount of unburned (unremoved) bonding compound. In this case, a heterogeneous arrangement of fibers in the infiltrated preform results from the specifics of the preform preparation technology itself.

a

b

c

112 Nano-composites

d e

f g

Fig. V.3. Structure of AK7 alloy reinforced with carbon nano-fiber preform by squeeze infiltration. Non-etched sample. Conventional light (a, b, d, f) and phase contrast (c, e, g)

Nano-composites 113

Al/50 wt. % carbon nano-fibers composites Matrix: pure aluminum. Reinforcing phase: pipe-like carbon nano-fibers (Fig. V.1), thermally treated and known by the trade name - Pyrograf® - III (PR-19-LHT); these nano-fibers have high graphite content and unique combination of properties, especially high thermal conductivity, the highest out of all known reinforcing phases. Manufacturing method: powder metallurgy (know-how) Structural characteristics (Fig. V.4): microscopic observations show relatively uniform arrangement of the nano-fiber colonies in comparison to liquid-phase methods. A characteristic feature of the microstructure is the arrangement of the fiber colonies along the direction of deformation.

a b

c

d

Fig. V.4. Structure of Al/50 wt. % carbon nano-fibers composite obtained by powder metallurgy. Cross section (a, b) and longitudinal section (c, d) related to the deformation direction. Non-etched sample. Conventional light

114 Typical structural defects in metal-matrix composites

VI. Typical structural defects in metal-matrix composites VI.1. Particulate dispersion-reinforced composites VI.1.1. Macro-scale nonuniform arrangement of the reinforcing phase (macro segregation) in

gravity cast composites, linked with the stability loss of initially homogenous metal-ceramic slurry obtained, due to a significant density difference between reinforcing phase (dzb) and metal matrix (dMe), resulting in either sedimentation phenomenon of the reinforcing phase α dzb > dMe or floating phenomenon of the reinforcing phase at dzb < dMe (Fig. VI.1.1).

a

b

Fig. VI.1.1. Macrostructures of the lengthwise cross-sections of gravity cast composite ingots illustrating segregation of reinforcing phase as a result of: a) sedimentation of the silicon carbide particulates in the A356/20 vol. % SiC composite, for which dSiC>dA356; b) graphite floatation in the AK12/5.7 wt. % nickel coated graphite (50 wt. % Ni) composite, for which dGr<dAK12

Typical structural defects in metal-matrix composites 115

VI.1.2. Micro-scale nonuniform arrangement of the reinforcing phase (micro segregation) in gravity cast composites, as a result of particulates agglomeration phenomenon (gathering together of the particulates, contacting each other, into agglomerates, containing from few to several dozens of particulates) (Fig. VI.1.2.1).

Three types of agglomeration are distinguished in metal matrix composites: 1) Primary agglomeration occurrs in the reinforcing phase before it gets introduced into the liquid

metal matrix and characterized by varying stability.

Particulates in unstable agglomerates, called soft agglomerates, are held by relatively week Van der Waals forces or electrostatic ones, and whose de-agglomerating can be achieved by additional sifting, mixing themselves or mixing with other powder, directly prior to the introduction of reinforcement into the metal matrix. In stable agglomerates, called hard agglomerates, the particulates are linked with each other by permanent bridges, which do not get completely destroyed even during intensive mixing in molten metal bath and in the same undestroyed state they are then permanently imbedded in solidifying composite casting. In this case, the phenomenon of inheriting primary heterogeneity of the reinforcing phase takes place, which can be effectively prevented by additional powder milling operation coupled with its de-agglomeration using other powder material. 2) Secondary agglomeration occurs during liquid-state processing, i.e. at some stages in the

formation of the metal-ceramic slurry or solidification of the composite casting. It is directly linked with the particulate-liquid matrix or particulate-solid matrix interactions, which determine the character of a particulate arrangement in the liquid metal bath or in the solidified composite casting. Even initial de-agglomeration of the reinforcing phase prior to its introduction to the bath, does not ensure homogenous arrangement of the reinforcing phase in the final product.

The following phenomena contribute to secondary agglomeration formation:

‐ Interaction between reinforcing phase particulates and gas bubbles in the metal bath or as a result of gas adsorption from the bath surface, which contributes to creation of clusters containing few or a dozen of particulates linked with common blister (Figs. VI.1.2.2 a, b) or gas layer (Fig. VI.1.2.2 c).

‐ Local non-wetting of the reinforcing phase by liquid matrix, which contributes to a particle-particle interaction in metal-ceramic slurry and facilitates gathering of the particulates in agglomerates in such a way, that they touch each other but some voids remain between these particles because they are non-wettable by the liquid matrix (Fig. VI.1.2.3).

‐ Pushing out of the reinforcing phase particulates in front of the solidifying metal matrix (pushing phenomenon), causing their forced undesirable redistribution. Significant fact is, that after creating very homogenous metal-ceramic slurry using intensive mechanical stirring method, the gravity castings of such composites have relatively homogenous distribution of the reinforcing phase in the metal matrix on the macro-scale. However, the observation of the structure in the micro-scale shows considerable segregation of the reinforcing phase particulates. This is linked with the phenomenon taking place in the composite casting already during its solidification, thus not really being a proof of the lack of homogeneity of the suspension formed. In case of aluminum alloys matrix composites, the solidification begins from the formation of primary precipitates of the solid α−phase solution. Further growth of its dendrites causes pushing reinforcement phase particulates out, which are gathered, following solidification process, as agglomerates in the eutectic between α-phase dendrites (Fig. VI.1.2.4). The use of technological operations decreasing grain size of primary phase precipitates, effectively enhances homogeneity of the reinforcement phase distribution in the aluminum matrix of the gravity cast composites (e.g. refining, use of the external pressure, fast cooling).

3) Ternary agglomeration occurs at some stages in creation of metal-ceramic slurry or during

formation of composite casting, due to the action of some external factors (e.g. centrifugal or electromagnetic force), forcing a violent particulates movement in a particular direction, their “congestion” and formation of agglomerates containing large amount of particulates. The difference between secondary and ternary agglomerates based upon structural examinations is difficult, because visually they look very much alike. In this case decisive significance has the comparison of the same composite material structure, cast in two methods - gravitationally and with the presence of external factor (i.e. two composite samples are manufactured in one experiment using two different methods, but out of the same composite slurry). Such example of the copper alloy matrix composite containing graphite particulates, is shown on Figs. I.6.2 and I.6.3.

116 Typical structural defects in metal-matrix composites

Following gravity casting, the composite has very homogenous structure, which results from uniform arrangement of graphite phase in the matrix (Fig. I.6.2). This is a testimony to not only a very high degree of homogeneity of the composite slurry obtained, but also of its sufficient stability, allowing to permanently form such a structure during further solidification of matrix in the gravity casting. The same homogenous and sufficiently stable in gravity conditions slurry becomes unstable under the centrifugal force, which manifests itself by the formation of ternary agglomerates occurring only in the external layer of the centrifugal casting (Fig. I.6.3).

a b

c

d

Fig. VI.1.2.1. Structure of squeeze cast AG10/5.0 wt. % fly ash composite, illustrating occurrence of primary agglomeration;

(a, b) depict stable (hard) agglomerates, while (c, d) depict the agglomerate of the particles, which partially becomes destructed, because fly ash particles are separated with a metal-matrix layer, maintaining their initial position though, in relation to each other (structure heredity phenomenon). Non-etched sample. Phase contrast (a, c, d) and conventional light (b)

Typical structural defects in metal-matrix composites 117

a b

c

Fig. VI.1.2.2. Structure of gravity cast AK12/3.7 wt. % fly ash composite (Dayton Power and Light Co., USA), illustrating secondary agglomeration due to particle-gas bubble interaction. Non-etched sample. Conventional light

Fig. VI.1.2.3. SEM image of gravity cast F3S.15S composite fracture (A359/15 vol.% SiC), depicting secondary agglomeration of a few SiC particles due to local non-wetting of reinforcing phase by the liquid matrix evidenced by the absence of continuous contact between SiC particles and metal matrix)

118 Typical structural defects in metal-matrix composites

Fig. VI.1.2.4. Structure of gravity cast F3S.15S composite (A359/15 vol. % SiC). Arrows show secondary agglomeration effect due to particle pushing phenomenon in such a way that the reinforcing particles are being gathered in low melting temperature matrix component (in this case, the Al-Si eutectic). Non-etched sample. Polarized light

Typical structural defects in metal-matrix composites 119

VI.1.3. Porosity - gas bubbles and shrinkage This type of structural macro discontinuity is characteristic to gravity casting. The main reason for the gas bubbles to occur is excessive gassing of the composite slurry during melting of metal-matrix or during intensive mechanical stirring of metal bath itself while introducing reinforcing phase. The presence of ceramic particles in the bath contributes to increasing stability of gas bubbles formed, because particulates gather on the bubble surface, creating rigid skeleton like three-dimensional construction (Fig. VI.1.3.1). The occurrence of shrinkage porosity is liked with the tendency of the metal matrix to shrink during solidification and further cooling of the casing. It is possible to differentiate between these two types of porosity by the shape of the defect itself, because the gas bubbles naturally have the spherical shape, while shrinkage porosity has irregular shape, often of very developed area (Fig. VI.1.3.2).

a b

Fig. VI.1.3.1. Structure of gravity cast F3S.15S composite (A359/15 vol. % SiC), depicting the gas bubbles and shrinkage

porosity. None-etched sample. Polarized light

120 Typical structural defects in metal-matrix composites

Fig. VI.1.3.2. Structure of A356/20 vol. % alumina composite (COMRAL™) cast into the metal mould, depicting

micro-shrinkage (marked with arrows). Non-etched sample. Polarized light

Fig. VI.1.3.3. Structure of AK12/fly ash (microspheres) composite obtained by squeeze infiltration, depicting gas-shrinkage

porosity (marked with arrow). Non-etched sample. Phase contrast

Typical structural defects in metal-matrix composites 121

VI.1.4. Excessive reactivity between the metal matrix and reinforcement Reactivity in the metal-matrix reinforcement phase system ensures the formation of a good bonding between matrix and reinforcement, but excessive reactivity can lead to undesirable effects including structural defects.

1) Occurrence of the structural discontinuities as a result of cracking of the reinforcement material

Example 1 High reactivity in Al – fly ash system (Figs. VI.1.4.1, VI.1.4.2) may lead to intensive oxy-redox reaction of the FexOy, SiO2 and mullite (main ingredients of the fly ash) with Al, which lead to the formation of aluminum oxide, Si and Fe, and which are accompanied by cracking of fly ash particles. Example 2 Cracking of the graphite phase in the aluminum matrix after heating of the Al-graphite slurry to the temperature of about 950ºC, occurring due to excessive reactivity, stemming from the creation of the aluminum carbide (Fig. VI.1.4.3). Similar phenomenon, in the case of the composites reinforced with carbon fibers or graphite, leads to weakening fibers and falling level of composite mechanical properties, and even their complete loss.

2) Increase of the composite structural heterogeneity High reactivity in the Al - fly ash system (Figs. VI.1.4.1, VI.1.4.2) leads to rapid increase of the amount and size of the primary Si precipitates, which have a tendency to grow on the fly ash particles, thus contributing to the increase of the structural heterogeneity of such composite. In the case of reinforcement in the shape of microsphere (Fig. VI.1.4.2) their cracking contributes to filling cavity inside the microsphere with liquid metal (particularly in the case of casting under external pressure). Inside the sphere, the content of Si and Fe is significantly higher than in the outside areas, which also contributes to increasing structural homogeneity of the composite.

3) Formation of undesirable compounds Example 1 Excessive reactivity in Al/C system leads to the formation of undesirable aluminum carbide (Al4C3), which, actively reacting with water vapor, contributes to the creation of a gas product (methane), thus weakening bonding between the matrix and the reinforcement phase and contributing to the formation of micro-cracks in the ready product (Fig. VI.1.4.4). Example 2 Similar to Al/C system, excessive reactivity in Al/SiC system may cause the formation of undesirable Al4C3 compound (Fig. VI.1.4.5). Example 3 Unproper conditions of manufacturing cast composites reinforced with metal-coated ceramic phase lead to the excessive dissolving of the coating in the metal matrix, cause weakening of the bonding between matrix and reinforcement, contributing to undesirable changes of the end product properties. Widely used technological coating of the ceramic phase (e.g. carbon fibers, graphite powder and aluminum oxide) is nickel coating, whose dissolving in the aluminum matrix causes the formation of the excessive amount of brittle Al3Ni compound precipitations (Fig. VI.1.4.6).

122 Typical structural defects in metal-matrix composites

Si

Si

Si

Fig. VI.1.4.1. Structure of AK12/3.7 wt. % fly ash composite (Dayton Power & Light Co., USA), depicting excessive reactivity of

the fly ash particulates in contact with liquid matrix, which leads to reduction of such type of oxides as SiO2 and mullite by Al (main fly ash ingredients) and the formation of Si. In the saturated with silicon metal-matrix, large primary precipitations of the Si crystals form, growing on the surface of the ash particulates (arrows indicate cracks in the particulates). Non-etched sample. Conventional light

Fig. VI.1.4.2. Structure of 52K/15 vol. % fly ash composite, depicting cracking of the microsphere and its filling with the liquid

matrix as well as the formation of large amounts of Si and iron-reach precipitates inside the microsphere as a result of chemical reactions taking place between liquid matrix and oxides. Sample etched with Kevlar agent. Polarized light

Typical structural defects in metal-matrix composites 123

Fig. VI.1.4.3. Structure of Al-graphite composite after heating to about 950°C, depicting cracking of the large graphite particle,

due to excessive reactivity caused from aluminum carbide formation. Non-etched sample. Polarized light

← Al

← Al4C3 ← graphite

Fig. VI.1.4.4. Structure of the Al/graphite interface after heating up to 950°C for 30 min and weathering out in the open for

1 month. Al4C3 precipitates are visible at the metal graphite interface as well as deformation lines and cracks in Al as a result of gas product formation (methane) due to reacting Al4C3 with water vapor from the atmosphere. Non-etched sample. Polarized light

Al4C3

124 Typical structural defects in metal-matrix composites

Fig. VI.1.4.5. Structure of gravity cast F3S.20S composite structure (A359/20 vol. % SiC) after heating in the liquid state with

visible needle-like Al4C3 precipitates formed on the surface of SiC particles. Effective methods to prevent occurrence of aluminum carbide in DURALCAN composites are: thorough temperature control of the liquid suspension as well as increased silicon contents in the matrix (usually Si ≥ 7 wt. %). Non-etched sample. Polarized light

Typical structural defects in metal-matrix composites 125

VI.1.5. Contaminations and non-metallic impurities Typical example of the contamination of the cast composites that may take place during improper melting process is presence of continuous oxide films (Fig. VI.1.5).

a b

c

Fig. VI.1.5. Structures of gravity cast F3S.20S composite (A359/20 vol. % SiC), illustrating inhomogeneous distribution of reinforcing phase due to the formation of the oxide films and their getting through from the areas inside the composite casting. Non-etched sample. Polarized light

126 Typical structural defects in metal-matrix composites

VI.2. Porous ceramic preform reinforced composites VI.2.1. Structural discontinuities due to insufficient preform infiltration, caused, amongst the

others by low metal temperature, local non-wetting, insufficient infiltration pressure, capture of gases inside the preform and lack of possibility to extract them (Figs. VI.2.1.1 - VI.2.1.3).

Fig. VI.2.1.1. SEM images of the AK12/22 vol. % Al2O3 composite fracture, depicting the solidified infiltration front inside short

fiber preform (MORGAN Company) due to its unfinished infiltration with AK12 alloy (time of pressure being exerted about 3 s)

Typical structural defects in metal-matrix composites 127

Fig. VI.2.1.2. SEM images of the AK12/20 vol. % Al2O3 composite fracture, depicting the solidified infiltration front inside short

fiber preform (DIDIER Company) due to its unfinished infiltration with AK12 alloy

128 Typical structural defects in metal-matrix composites

Fig. VI.2.1.3. Structure of AK12/fly ash composite (microspheres) obtained by squeeze infiltration of porous preform with AK12 alloy, illustrating its incomplete filling with liquid matrix due to insufficient wettability in the metal-ceramic system (marked with arrow). Non-etched sample. Phase contrast

Typical structural defects in metal-matrix composites 129

VI.2.2. Structural heterogeneity and discontinuities resulting from the improper preform preparation

Both heterogeneous thickness of the nickel coating layer on the fibers as well as their heterogeneous distribution contribute to the formation of significant Ni segregation in the matrix, where the areas between the fibers are used up by eutectic (Al+Al3Ni), while other area – by Ni solid solution in Al (Fig. VI.2.2.1).

a

b

c

130 Typical structural defects in metal-matrix composites

d e

f g

Fig. VI.2.2.1. Structure of Al/nickel coated carbon fibers composite, depicting structural heterogeneities caused from nonuniform distribution of reinforcing phase in Al matrix, nonuniform thickness of nickel coating and matrix segregation due to the dissolution of Ni coating in liquid Al. Non-etched sample. Conventional light (a, b, d, f) and phase contrast (c, e, g)

Typical structural defects in metal-matrix composites 131

VI.2.3. Excessive reactivity between composite components

1) Cracking of the reinforcing phase (Fig. VI.2.3.1). 2) Weakening of the ceramic preform.

In the case of an intensive interaction between liquid matrix and binder component, used to form a rigid construction of the ceramic preform, the mechanical weakening of the inter-fiber bonding occurs in the preform. This is accompanied by changes of its dimensions and shape, even crack appearance, thus contributing to undesirable changes of the end product properties (Fig. VI.2.3.2).

3) Formation of undesirable compounds (Fig. VI.2.3.3). 4) Disadvantageous changes of the matrix chemical composition (Fig. VI.2.3.4).

a

b

c d

Fig. VI.2.3.1. Structure of Al matrix composite reinforced with porous preform made from short Al2O3 fibers and SiO2, depicting

cracking of the reinforcing phase as a result of its interaction with metal matrix. Non-etched sample. Phase contrast (a-c) and conventional light (d)

132 Typical structural defects in metal-matrix composites

a

SiO2 + Mg MgSiO2 SiO2 + 2Mg 2MgO + Si Al(Mg) + Si [Al(Mg) + Si]eut

b Fig. VI.2.3.2. Structure of MgAl8Zn2 alloy locally reinforced with ceramic preform (MORGAN, UK), consisting of 95 wt. % short

fibers (96% Al2O3 and 4% SiO2) as well as 5 wt. % of SiO2-based binder, showing the dimensional and shape changes of the preform, its deformation and cracking (a). Scheme (b) shows changes in the preform structure caused from two effects, i.e. applied external pressure and weakening of interfiber connections in the preform that takes place due to consumption of binder by chemical reactions liquid.

Mg

200 µm

Typical structural defects in metal-matrix composites 133

a b

c

Fig. VI.2.3.3. Structure of Al/carbon fiber composite obtained by gas pressure infiltration in the autoclave (p = 7 MPa, T = 800°C, τ = 30 min), depicting the formation of the aluminum carbide at the fiber/matrix interface. Non- etched sample. Polarized light

a b

Fig. VI.2.3.4. Structure of the composite obtained with initially pure Al matrix and reinforced with nickel coated carbon fibers using gas pressure infiltration in the autoclave (p = 7 MPa, T = 800°C, τ = 30 min). As a result of intensive interaction between nickel coating and Al, a complete dissolving of Ni in the metal matrix takes place, and thus exposing the fibers. This causes further chemical reaction between Al and nickel and formation of large amount of aluminum carbide (Al4C3). At the same time, a disadvantageous change of chemical composition of the matrix itself, takes place, caused by rapid growth of Ni content as well as contributing to the formation of large amount of brittle Al3Ni precipitates. Non-etched. Polarized light

134 Typical structural defects in metal-matrix composites

VI.2.4. Non-metallic inclusions Just like with particulate dispersion-reinforced composites, in case of continuous fibers reinforced composites, there may be oxide films, especially appearing in the areas close to the castings’ surface (Fig. VI.2.4.1).

a

b

c

Fig. VI.2.4.1. Structure of AK11/carbon felt composite in the vicinity of the casting surface, depicting structural discontinuities due to occurrence of oxide films. Non-etched sample. Conventional light

Typical structural defects in metal-matrix composites 135

VI.3. Hybrid composites In the case of hybrid composites, introduction of the second reinforcing phase may be aimed at obtaining additional positive effect of increased structural homogeneity through preventing segregation of the reinforcing phase, using the effect of the neutral buoyancy. To this effect, two phases are selected in such a way, so that one of them had a density higher than the other, in comparison to liquid matrix, thus ensuring significant limiting of the undesirable movement of the particulates in the metal bath. To make use of the neutral buoyancy effect it is required to strictly control the conditions of the process, both of reinforcing phase preparation as well as manufacturing of the composite, while not following these requirements contributes to increasing structural heterogeneity, in comparison to other composites. Example 1 In the Al/SiC/graphite type of composites, too thick nickel coating on the graphite particles causes disadvantageous increase of their density, which contributes to speeding up of the sedimentation process of both SiC and graphite reinforcing phases (Fig. VI.3.1). Example 2 For excessive dissolution of the Ni in the liquid metal in case of the Al/SiC/graphite type composites containing nickel coated graphite, the excessive increase of the matrix density itself takes place, above that of the SiC phase. This causes increase of structural heterogeneity due to an undesirable floating of particles, both SiC and graphite ones.

Fig. VI.3.1. Structure of lower part of the gravity cast hybrid composite made from A356 matrix reinforced with 15 vol. % SiC

and 5 vol. % nickel coated graphite particles. Non-etched sample. Conventional light

136 Typical structural defects in metal-matrix composites

Fig. VI.3.2. Structure of hybrid composite produced from A356 alloy reinforced with 15 % silicon carbide and 5 % graphite,

gravity cast into the metal die. Sample etched with Kevlar agent. Conventional light

SiC

SiC

Al3Ni

Al3Ni

Al3Ni

Al3Ni Al3Ni

Graphite particle Ni coating

Typical structural defects in metal-matrix composites 137

VI.4. High volume faction particulate-reinforced composites 2014/52 vol. % SiC composite Matrix: 2014 aluminum alloy (AlCu4SiMg) – 4.40Cu, 0.80Si, 0.8Mn, 0.60Mg, 0.30 Fe (wt. %). Reinforcing phase: single crystal silicon carbide platelets (hexagonal, α-SiC) (from Third Millennium Technologies, Inc., USA) of three average platelet sizes - 250-500 µm diameter × 25-50 µm thickness; 50-250 µm diameter × 5-25 µm thickness; and 20-70 µm diameter × 0.50-5 µm thickness; both as-received and Cu-coated SiC platelets were utilized. Manufacturing method: counter-gravitational infiltration of packed binder-less silicon carbide platelets (packing density ~52 %) by the molten 2014 aluminum alloy under nitrogen gas pressure; SiC platelets were packed in quartz tubes via multistage compaction of loose powders; infiltration was done in a stainless steel pressure vessel with a melting furnace, consisting of two semi-cylindrical heating elements, housed inside the pressure vessel; quartz tubes containing packed SiC were attached using leak-tight fittings to the lid of the pressure vessel which, upon closure, led to partial immersion of the tube in the molten alloy held in a graphite crucible inside the pressure vessel; an argon blanket was maintained over the SiC inside the quartz tube to reduce the oxidation of the liquid front; preheating time of SiC beds: 3-5 min; melt temperature: 800°C; nitrogen gas pressure: variable, 525-850 kPa; infiltration time: 3-5 min; slow cooling at the conclusion of infiltration in ambient air. Both as received and copper-coated SiC platelets were infiltrated. Cu-coating was done to promote the wettability and facilitate the infiltration. Copper coating was deposited on SiC using an electroless plating technique. For this, SiC platelets were stirred for 3 h in a solution of 10 g/l of SnCl2 and 40 ml/l of concentrated HCl. The platelets were then stirred for 15 min in a solution 0.5 g/l PdCl2 and 1 ml/l concentrated HCl. After stirring, the solution was decanted, and platelets washed twice with distilled water. Finally, the platelets were stirred for 15 min in the plating bath of following composition: 10 g/l CuSO4

.5H2O, 10 g/l NaOH, 50 g/l sodium potassium tartrate, and 10 ml/l formaldehyde. The coating deposition was confirmed with the SEM examination. Properties: considerable improvement of tribological characteristics expected, significant reduction in thermal expansion coefficient and increase in the thermal conductivity, increased hardness and strength, and decreased ductility. Structural characteristics: A structural imperfection sometime observed in infiltrated composite bars made from loosely-packed SiC beds is the disruption and channeling of the packed bed under melt shear. This leads to the formation of clearly demarcated SiC-impoverished and SiC-enriched regions (Fig. VI.4.1). Normal solidification microstructure (cored dendrites and interdendritic eutectic) of the matrix alloy develops in the SiC-impoverished region of the disrupted compact (Fig. VI.4.2) Another imperfection is the spatial variation of the composite microstructure with the infiltration height. A greater amount of porosity and secondary phases precipitate near the top of the infiltrated bar and less porosity and secondary phases form near the bottom (Fig. VI.4.3). In addition, a greater preferred orientation (stratification) of the SiC platelets due to higher melt shear occurs at the bottom. Thus, both the matrix solidification features (secondary phases, porosity) and SiC distribution vary with the infiltration height, resulting in an undesirable spatial variation of the composite properties. Finally, the chemical interactions between the reinforcement particulates and matrix after prolonged contact at elevated temperatures gives rise to extensive degradation of the SiC platelets (Fig. VI.4.4) due to the formation of a large amount of the brittle reaction product, aluminum carbide (Al4C3).

138 Typical structural defects in metal composites

Fig. VI.4.1. Structure of 2014/xSiC composite showing channeling near the periphery of specimen transversed to the infiltration direction. Sharply demarcated SiC-starved and SiC-enriched regions form because of channeling in the packed SiC network under melt shear; pressure: 825 kPa, melt temperature: 800°C, infiltration time: 3 min, SiC size: 50-250 µm diameter and 5-25 µm thickness, as-received SiC. Non-etched sample. Conventional light

Fig. VI.4.2. Structure of 2014/xSiC composite at two magnifications of cored dendritic structure in the channeled region (SiC

impoverished region) of sample of I.3.12; pressure of 825 kPa, melt temperature: 800°C, infiltration time: 5 min, SiC size: 50-250 µm diameter and 5-25 µm thickness, as-received SiC; the sample was color etched with a solution of 4 g KMnO4 and 2 g NaOH in 1 l distilled water. Conventional light

20 µm

10 µm

200 µm

Typical structural defects in metal-matrix composites 139

a b Fig. VI.4.3. Structure of 2014/xSiC composite showing the variation of microstructure and SiC distribution with height in

composite produced by infiltration counter to gravity. The microstructures are shown for samples cut lengthwise across their diameters (i.e., views are parallel to the infiltration direction): (a) shows greater porosity and second phase precipitation near the top of the infiltrated bar, and (b) shows less porosity and secondary phases at the bottom (melt entrance) but also greater directional stratification of the platelets due to greater melt shear at the bottom; pressure of 825 kPa, melt temperature: 800°C, infiltration time: 5 min, SiC size: 250-500 µm diameter and 25-50 µm thickness, Cu-coated SiC; the sample was color etched with a solution of 4 g KMnO4 and 2 g NaOH in 1 l distilled water. Conventional light

Fig. VI.4.4. Transverse section of 2014/xSiC infiltrated counter-to-gravity with a 30 min hold at 800°C. The image, taken from

a section near the bottom of the infiltrated bar, shows extensive degradation of the Cu-coated SiC platelets due to chemical reaction of SiC with molten Al upon prolonged contact; infiltration pressure: 525 kPa, melt temperature: 800°C, infiltration time: 30 min, SiC size: 50-250 µm diameter and 5-25 µm thickness, Cu-coated SiC. Conventional light

10 µm

10 µm

10 µm

140 Typical structural defects in metal-matrix composites

VI.5. Long fiber-reinforced composites NiAl/saphire fiber composite Matrix: nickel-base intermetallic NiAl (49.9-50.3Ni, 49.7-50.1Al, at. %) and NiAl(Yb) (62.1–62.6Ni, 36.8–37.3Al, 0.36–0.11 Yb at. %). Reinforcing phase: long (continuous) single crystal sapphire fibers (SaphikonTM) with average diameter 178 µm. Manufacturing methods: two methods of composite manufacture were utilized: 1. NASA’s patented ‘Powder-cloth’ (PC) technique in which a cloth or sheet of atomized matrix

powders is created by mixing them with Teflon binders and a solvent to a dough-like consistency; fiber mats of specified thickness are produced by filament winding and by application of a PMMA coating; matrix cloth and fiber mats are then stacked in layers and vacuum hot pressed during which binders are removed; the orientation of individual plies and stacking sequence provide flexibility in designing the composite properties.

2. Floating-zone directional solidification (DS) which is a container-less solidification technique to

produce the composites. The DS was done on either powder-cloth feedstock with ~25 vol. % fiber, or on vacuum induction melted and chill cast (in Cu-mould) composites (the chill cast composites had low, <0.5%, fiber volume fraction and were fabricated by laying sapphire fibers in a Cu-mould with fiber ends rigidly supported within holes drilled in steel end caps). The DS was done under 3 psi (gauge) pressure of ultrahigh purity argon gas in a 25 kW radio frequency (RF, 400 kHz) vacuum induction furnace using a water-cooled copper current concentrator; a current adjusting type three-mode proportional controller was used to set the temperature and control the power input to the coil; DS was carried out by traversing a 8 mm long molten zone along the length of the feedstock at a constant speed (~6.0 cm per hour).

Properties: substantially lower density, increased rigidity and strength at room-temperature and elevated temperatures, significantly lower thermal expansion (especially in the direction of fiber orientation), increased hardness, decreased ductility, and improved wear resistance are expected. With W and Cr alloying of NiAl and incorporation of sapphire in the composite, dual-phase toughening potential together with high-temperature strength is expected in the composite. Structural characteristics: in the powder-cloth (PC) composite (Fig. VI.5.1), polycrystalline NiAl matrix exhibits some residual porosity due to incomplete densification during hot pressing to fabricate the composites. There is, however, no evidence of interfacial reaction. In the sapphire-reinforced NiAl(Yb)-matrix composites made using the powder-cloth process (Fig. VI.5.2), besides matrix porosity (marked by arrows in the figure), a considerable amount of interfacial reaction occurred between the sapphire fiber and the ytterbium, which is a strong oxide former, and Yb2O3 is more stable than Al2O3 (the free energies of formation of Yb2O3 and Al2O3 at the DS are ~1727.5 kJ.mol-1 and ~1583.1 kJ.mol-1, respectively). The principal reaction products, independently identified using EDS and x-ray diffraction, consist of Yb2O3 and Yb3Al5O12. The grain boundaries in the NiAl(Yb) matrix are enriched in Ni3Al phase (arrows in Fig. VI.5.2 b). The directional solidification (DS) of the powder-cloth sapphire-reinforced NiAl(Yb) composite feedstock led to an even more extensive fiber-matrix interfacial reaction and formation of a thicker reaction zone as compared to the powder-cloth composite (Fig. VI.5.3). This together with the presence of porosity and shrinkage cracks in DS composites decreased the interfacial shear strength. Fig. VI.5.4 shows the SEM images of the surface of fibers extracted from the PC and DS composite. The fiber surface shows extensive chemical attack of the sapphire by the ytterbium in the NiAl(Yb) matrix. The fibers were extracted from the composite by dissolving the matrix in a boiling solution of 75% acetic acid, 23% nitric acid, and 2% HCl acid. There was some loss of Yb from the matrix in the fiber/matrix interface reactions; as a result, the Yb content (in at. %) decreased from 0.36 in the PC composite to 0.19 in the DS composite. The appearance of the sapphire fibers distributed in the matrix is shown for the various composites in Fig. VI.5.5. The hot pressing step in the powder-cloth process leads to appreciable fiber breakage during composite fabrication. This fiber damage is further accentuated when the PC composite is used as a feedstock for floating-zone directional solidification (DS); this is because of the additional damage caused due to the vigorous melt recirculation within the melt zone under a strong electromagnetic

Typical structural defects in metal-matrix composites 141

field. In contrast, the composite made by vacuum induction melting and chill casting in copper mould do not show any appreciable fiber breakage.

a b

Fig. VI.5.1. (a) Matrix region around a debonded sapphire fiber in a “powder-cloth” (PC) sapphire/NiAl composite showing

porosity due to incomplete densification during composite fabrication, and (b) shows wear tracks on NiAl grains that formed due to sliding of debonded fibers during fiber push-out test. The matrix contains approximately equi-atomic fractions of Ni and Al. Reinforcement: single crystal sapphire fibers (SaphikonTM), average fiber diameter: 178 µm

a b

Fig. VI.5.2. Structure of NiAl(Y)/saphire fibers composite produced by the “powder-cloth (PC)” technique. Arrows in (a) denote the porosity in the matrix due to incomplete densification during hot pressing, and arrows in (b) denote the Ni3Al phase that formed at matrix grain boundaries in the NiAl(Yb) matrix. A reaction layer at the fiber/matrix interface in (b) is also visible. Non-etched sample

a b

Fig. VI.5.3. Low (a) and high (b) magnification views of a sapphire-reinforced NiAl(Yb) matrix composite after floating-zone directional solidification. The DS was done on ‘Powder-Cloth (PC)’ feedstock specimen. Extensive fiber/matrix interfacial reaction and formation of shrinkage cracks is observed. Reinforcement: single crystal sapphire fiber (SaphikonTM), fiber diameter: 178 µm

200 µm

10 µm

142 Typical structural defects in metal-matrix composites

a b

Fig. VI.5.4. SEM images of the surface of fibers extracted from the PC composite (a) and the DS composite (b). The fiber surface shows extensive chemical attack of the sapphire by the ytterbium in the NiAl(Yb) matrix. The fibers were extracted from the composite by dissolving the matrix in a boiling solution of 75% acetic acid, 23% nitric acid, and 2% HCl acid. Reinforcement: single crystal sapphire fibers (SaphikonTM), average fiber diameter: 178 µm. The composition of the powder-cloth (PC) feedstock (in at. %) is 62.1Ni-36.8Al-0.36Yb. There was some loss of Yb in the fiber/matrix interface reactions; as a result, there was less Yb in the DS material. The composition of the DS composites (in at. %) is 62.5Ni-36.4Al-0.19Yb

a b

c d

e

Fig. VI.5.5. Structure of various NiAl/sapphire fibers composites: (a) “powder-cloth (PC)” composite showing fiber breakage during composite fabrication by hot pressing, (b) directionally solidified (DS) composite made from a PC feedstock, (c) composite made by vacuum induction melting and chill casting in copper mould, (d) DS composite showing fiber fragmentation due to matrix recirculation in the molten zone, and (d) fibers in a sapphire-NiAl(Yb) composite

143

Acknowledgments In the course of the preparation of this atlas, the results of experimental research were used, which in turn were obtained during the following research projects:

1. Infiltration processing of silicon carbide platelet-reinforced aluminum alloy composites. Doctoral Dissertation by R. Asthana, 1991, University of Wisconsin-Milwaukee, USA (Adviser: P.K. Rohatgi)

2. Complex investigation of the interfacial phenomena of the liquid metals in contact with solid bodies, Project No. MP/NIST 92-90, conducted in 1992-1996 by the Foundry Research Institute and University of Wisconsin - Milwaukee, USA, and financed by Polish-American M. Sklodowska-Curie Fond (Principal Investigators: N. Sobczak, P.K. Rohatgi)

3. Physical and chemical mutual interaction on the solid body – liquid metal contact border, as a factor shaping the structure and properties of heterogenic materials of aluminum alloys matrix. Research project No. 773729203, conducted in 1992-1995 by the Foundry Research Institute and financed by the Committee for Scientific Research, Poland (Project Manager: N. Sobczak)

4. Working out comprehensive technology to manufacture new generation castings reinforced with ceramic or composite insertions using squeeze casting technology. Research project No. 7S20201804, conducted in 1993-1996 by the Foundry Research Institute, financed by the Committee for Scientific Research, Poland (Principal Investigator: J. Sobczak)

5. Zone directional solidification of sapphire-reinforced NiAl-matrix composites, U.S. National Academy of Sciences (through National Research Council), 1994-95, USA (Principal Investigator: R. Asthana)

6. Effect of external pressure during solidification on the production of metal matrix composites and use of such products in automotive industry, Project No. MP/NIST-95-231, conducted in 1995-1998 by the Foundry Research Institute and University of Wisconsin - Milwaukee, USA, financed by Polish-American M. Sklodowska-Curie Fond (Principal Investigators: J. Sobczak, P.K. Rohatgi)

7. Settling behaviour of ceramic reinforcements in composite slurries, Faculty Research Initiative Grant, University of Wisconsin-Stout, 1996, USA (Principal Investigator: R. Asthana)

8. Graded products with local composite reinforcement manufactured using centrifugal casting. Research project No. 7T08B02211, conducted in 1997-1999 by the Foundry Research Institute and financed by the Committee for Scientific Research, Poland (Principal Investigator: J. Sobczak)

9. New generation of pistons for combustion engines as a result of optimising their properties by formation of composite material structure. Ph.D. project No. 9T12C0131, conducted in 2000-2001 by the Motor Transport Institute and financed by the Committee for Scientific Research, Poland (Principal Investigator: J. Sobczak, Ph.D. student: D. Rudnik)

10. Metallographic quantitative criteria of the selected service properties of composite materials made brake discs. Ph.D. project No. 7T08B02718, conducted in 2000-2001 by the Motor Transport Institute and financed by the Committee for Scientific Research, Poland (Principal Investigator: J. Sobczak, Ph.D. student: A. Wojciechowski)

11. Working out the foundation of gasars technology – new highly porous materials with predetermined spatial lay out of porosity. Research project No. 7T08B01419, conducted in 2000-2003 by the Foundry Research Institute in cooperation with the Motor Transport Institute and financed by the Committee for Scientific Research, Poland (Principal Investigator: J. Sobczak)

12. Physicochemical aspects of in-situ synthesis of Al-Al2O3 composites by liquid-phase processing. Research project No. 7T08B00320, conducted in 2001-2004 by the Foundry Research Institute and financed by the Committee for Scientific Research, Poland (Principal Investigator: N. Sobczak)

13. Starting up production of brake discs made of composite materials for the automotive industry needs: Working out comprehensive technology of manufacturing brake discs out of Al/SiC composites. Target project No. 10T120132000C/5220 conducted in 2001-2004 by the Motor Transport Institute with the participation of the Foundry Research Institute, commissioned by the Institute of Implementations and Polish Foundries Technologies Ltd. Project co-financed by the Committee for Scientific Research, Poland

144

14. Optimisation of car component properties by formation of the new composite material structure. Research project No. 9T12C01218 conducted in 1999-2002 by the Motor Transport Institute and financed by the Committee for Scientific Research, Poland (Principal Investigator: K. Pietrzak)

15. Development of the synthesis of the light and economically justified construction materials utilising waste products (fly ashes). Research project No. 3T08B06426, conducted in 2004-2007 by the Foundry Research Institute in co-operation with the Motor Transport Institute and financed by the Ministry of Science and Information Technology, Poland (Principal Investigator: J. Sobczak)

16. Effect of technological parameters on structure and properties of fly ash reinforced aluminum matrix composites. Ph.D. project No. 3T08B05728, conducted in 2005-2007 by the Foundry Research Institute and the Motor Transport Institute, and financed by the Committee for Scientific Research, Poland (Principal Investigator: J. Sobczak, Ph.D. student: P. Darlak)

17. Innovative materials and processes for use in foundry industry. Commissioned research project No. PBZ-KBN-114/TO8/2004, conducted in 2004-2008 by the Foundry Research Institute in co-operation with the Motor Transport Institute, and financed by the Ministry of Science and Information Technology, Poland (Principal Investigator: J. Sobczak)

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