InterfaceEngineeringofCarbon- FiberReinforcedMg–AlAlloys**Feldhoff, Pippel, Woltersdorf/Interface Engineering of Carbon-Fiber Reinforced Mg–Al Alloys A further parameter influencing

Interface Engineering of Carbon-Fiber Reinforced Mg±Al Alloys**By Armin Feldhoff, Eckhard Pippel, and Jörg Woltersdorf*

1. IntroductionMetal matrix composites (MMCs) of the system carbon-fi-

ber/magnesium-matrix allow the outstanding combination ofi) the high strength (3±4 GPa) and Young's modulus (some100 GPa) of the carbon fibers, ii) the low density of both part-ners (

Feldhoff, Pippel, Woltersdorf/Interface Engineering of Carbon-Fiber Reinforced Mg±Al Alloys

croscope techniques offer powerful tools to gain structuraland chemical information of the fiber/matrix interfaces downto the atomic scale: High voltage electron microscopy(HVEM), high-resolution electron microscopy (HREM),energy-filtered transmission electron microscopy (EFTEM),and scanning transmission electron microscopy (STEM) incombination with energy dispersive X-ray spectroscopy(EDXS) and electron energy loss spectroscopy (EELS) withparticular emphasis on energy loss near edge structures(ELNES). The latter method allows one to characterize thechemical bonding state in the interlayers on the nanometerscale by analyzing the fine structures of the relevant ioniza-tion edges, which can be attributed to transitions of core±shellelectrons into unoccupied states above the Fermi level.[20±24]

As characteristic ELNES details the edge onset as well as theshape, the position and the intensity of individual peaks inthe fine-structure can be employed. Furthermore, using an

energy-selective device, transmitting only electrons with apredetermined energy, EFTEM becomes possible, allowingthe imaging of elemental distribution in the fiber/matrix in-terregion with a spatial resolution of about one nano-meter.[25±27]

2. Alloying the Matrix with a Carbide FormerAluminum, which is the technologically most important

alloying element for magnesium,[28±31] can be used to increasethe fiber/matrix interactions via the formation of carbides.Some of the Mg±Al alloys most frequently used are AM20,AM50, and AZ91 with Al contents of approximately 2, 5, and9 wt.-%, respectively.[28±31] In contact with carbon fibers, alu-minum may form the stable binary carbide Al4C3,

[32±34] or itmay cause the formation of stable ternary carbides of approx-imate composition Al4Mg2C3

[35] or Al2MgC2.[36±39]

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A. Feldhoff studied physics at the Westphalian-Wilhelms-University, Münster until 1994. Since then heworked as a research scientist at the Max-Planck-Institute for Microstructure Physics in Halle. In 1997he graduated at the Martin-Luther-University Halle-Wittenberg with a Ph D thesis on interface optimi-zation of carbon-reinforced magnesium alloys.

E. Pippel is working as a research scientist in the group ªInterfaces/ New Materialsº at the Max-Planck-Institute for Microstructure Physics, Halle. He graduated 1982 with a thesis on pseudomorphousgrowth in epitaxial bicrystal systems and won the Bühler award in 1993. He authored 130 scientific pub-lications and is also co-author of several technical books.

J. Woltersdorf is head of the research group ªInterfaces/ New Materialsº at the Max-Planck-Institute forMicrostructure Physics, Halle, and lecturer at the Martin-Luther-University Halle-Wittenberg. He re-ceived his habilitation in 1988 and was in 1992 appointed as professor of surface technology at the Tech-nical University Dresden. He also won the Bühler award in 1993 and authored 200 scientific publica-tions as well as several technical books and monographs.


A further parameter influencing the fiber/matrix reactivityis the surface microstructure of the carbon fiber itself. As aconsequence of their production process, which is a con-trolled pyrolysis of a polyacrylonitrile (PAN) polymer precur-sor fiber,[40±43] commercially available carbon fibers are classi-fied in two types that differ mainly in the degree ofgraphitization of the turbostratic carbon[44]: The more graphi-tized high-modulus fibers possess a relatively smooth sur-

face, with the graphitical basal planes in preferen-tially parallel orientation to the fiber surface. Incomparison the surface of the less graphitized high-tensile-strength fibers is rougher, with many basalplanes ending freely on it.[5,45±48] As the free-endinggraphite layers show a higher chemical reactivitythan the graphitic basal planes,[48,49] the high-ten-sile-strength fibers are generally more reactive thanthe high-modulus ones.

In general, by combining carbon fibers of differ-ent surface microstructures with Mg±Al matrices ofdifferent Al contents, the fiber/matrix reactivity ofC/Mg±Al composites can be varied over a widerange. With the interface reactivity increasing threecharacteristic failure mechanisms occur, which cor-relate with the strength of the composites.[5,6,50±57]

They are illustrated in Figure 1, showing results ofin-situ bending tests[5,53±58] on three composites ofincreasing fiber/matrix reactivity:

At the beginning of the bending test with aC/Mg±Al composite of low interface reactivity(Fig. 1a), the load increases linearly with increasingdeflection. With accumulation of damage events,the slope decreases. After the load maximum atFmax = 240 N (rB » 544 MPa) the curve declines con-tinuously. The fracture surface in Figure 1a is char-acterized by pull-out and fracture of single fibersdue to a very weak fiber/matrix bonding. This is af-fected by the absence of any carbides at the fiber/matrix interface as is proven by transmission elec-tron microscopy (TEM) investigations.[4,5,54]

In the bending test with a C/Mg±Al compositeof medium interface reactivity (Fig. 1b), the load in-creases linearly with increasing deflection until theload maximum is reached at Fmax = 410 N (rB »929 MPa). The subsequent decrease in load showscharacteristic steps, which can be attributed to thefailure of individual fiber bundles. Bundles of 30 to80 fibers, which are joined by matrix metal emergefrom the fracture surface. Such a bundle fracturebehavior implies a bending strength of the compos-ite that is approximately 1.7 times higher than thatof the first material.

A composite of high interface reactivity (Fig. 1c)shows a much lower strength. In the correspondingbending test the load increases linearly up to itsmaximum, which is reached already at Fmax = 70 N

(rB » 158 MPa). The flat fracture surface hints at a brittle frac-ture as known from monolithic ceramics.

Figure 1 illustrates that the interface reactivity shouldneither be too low nor too high to enable the maximum use ofthe fiber strength. With increasing fiber/matrix reactivity thecomposite strength correlates with three characteristic failurepatterns: single-fiber fracture, bundle fracture, and brittlefracture.

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Fig. 1. Load-deflection diagrams and SEM images of the fracture surfaces: a) High-modulus fiber/AM20-matrix composite with the load maximum at 240 N (rB » 544 MPa) showing single-fiberfracture, b) high-tensile-strength fiber/AM20-matrix composite with the load maximum at 410 N(rB » 929 MPa) showing bundle fracture, c) high-tensile strength fiber/AZ91-matrix composite withthe load maximum at 70 N (rB » 158 MPa) showing brittle fracture.


Depending on their fiber/matrix reactivity, C/Mg±Alcomposites show extents of plate-shaped precipitates at thefiber/matrix interfaces.[4±6,54,56] With the fiber/matrix reactiv-ity being rather high, many mm-sized precipitates are formed(Fig. 2). Amount and size of the precipitates characterize therespective interface reactivity and control the mechanicalproperties.[5,6,50±57] Therefore, the optimization of the C/Mg±Al composite materials requires the almost comprehensivecharacterization of the nature of the precipitates.

EDXS and EELS investigations revealed that the precipi-tates are composed of the three elements Mg, Al, andC.[35,54,57] The fine-structures of the ionization edges in theEEL spectra contain further information that goes beyond thecomposition as the ELNES peculiarities are determinedmainly by the bonding states, thus reflecting the chemicalneighborhood (cf. Sec. 1). In the following, the ELNES fea-tures of the interface precipitates are discussed referring toFigure 3, where the fine-structures of the Mg±L23, Al±L23, andC±K ionization edges are compared to those of standard sub-stances.

The Mg±L23 ELNES curves of Figure 3a show a significantdifference in the chemical neighborhood of the magnesiumatoms of a precipitate compared to that of metallic (top) oroxidic magnesium (bottom). In addition, for the precipitatephase the ELNES shows a peak structure at about 73 eV (on-set), which can be attributed to the L23 ionization edge of alu-minum. In Figure 3b, this peak structure is compared to theL23 edge of different Al compounds.

The Al±L23 ELNES (Fig. 3b) of the Al4C3 standard and ofthe precipitate phase clearly differ from those of metallic andoxidic aluminum. Between each other they show similar pro-files. In both cases, the edge-onset energy is at about 73 eV,followed by a small peak at about 77.5 eV after a steep rise ofthe signal.

At the C±K ionization edge (Fig. 3c), the spectra of theAl4C3 standard and of the precipitate phase show a steep risein the signal, with its maximum peak at about 291 eV. For theprecipitate phase, however, the slope decreases slightly onthe left of this peak, viz. at about 287 eV. Furthermore, atabout 30 eV above the onset energy, another broader peak oc-curs, which differs by about 3 eV in the energy position of itsmaximum (at about 305 eV for Al4C3, and at about 302 eV forthe precipitate).

Because of the strong similarities in the Al±L23 and C±KELNES (Figs. 3b and 3c), these ELNES peculiarities suggestthe precipitate phase to be considered as an aluminum±mag-nesium-carbide, the crystal chemistry of which is closely re-lated to that of the binary carbide Al4C3.

Indeed, detailed crystallographic and morphological in-vestigations[38,54,56] including the comparison with X-ray pow-der diffractometry studies[36] identify the observed precipitatephase as the ternary carbide Al2MgC2, which has an (0002)lattice fringe distance of 0.62 nm. The proposed crystal struc-ture describes the ternary carbide Al2MgC2 as an interstitialcarbide, analogous to the well-known binary carbide

Al4C3.[38,54±56] According to Figure 4, showing a projection on

the (1120) plane, the carbide Al2MgC2 consists of a metalatom arrangement, with the Al atoms in a cubic close-packed(c) and the Mg atoms in an hexagonal close-packed (h) stack-ing. The carbon atoms fill two kinds of interstices of the metalatom host lattice: octahedral interstices between two adjacentaluminum layers, and trigonal-bipyramidal ones within themagnesium layers.

The ternary carbides having formed in the C/Mg±Al com-posites have the morphology of hexagonal-shaped plates, i.e.,their lateral extension is in perpendicular orientation to theirc-axis. Figure 5 demonstrates that the interfaces between theside faces of the Al2MgC2 plates and the metal matrix exhibitmany growth ledges, whereas the interfaces between the


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Fig. 2. HVEM bright-field image of a high-tensile strength fiber in a composite with anAZ91 matrix (middle), revealing plate-shaped precipitates in the fiber/matrix interre-gion.

Fig. 3. ELNES profiles of carbidic precipitates in a C/Mg±Al composite and of standardsubstances at different ionization edges: a) Mg±L23, b) Al±L23, c) C±K.


(0001) habit planes and the matrix are flat on the atomic scale.These two distinct interface structures between the Al2MgC2plates and the Mg matrix indicate the action of two differentgrowth mechanisms,[54±57,59] which were observed also forAl4C3 in Al matrixes:

[60,61]

l The atomically-rough interface moves in a continuous dif-fusion-controlled growth mode, which is determined bythe diffusion rates of the reaction partners Al and C.

l In directions normal to the atomically-flat interface (i.e.,along the c-axis of the carbide), the crystal grows by an in-terface-controlled ledge mechanism, which requires amultiple two-dimensional nucleation or a spiral growthmechanism.

With the driving force being low, the diffusion-controlledgrowth process is several times faster than the interface-con-trolled one,[62,63] yielding the observed plate-shaped morphol-ogy of the ternary carbides. Of course the carbides nucleateand grow preferentially in the liquid matrix state, when allthe species have the highest mobility.[39,54,64]

The velocity of the growth of the Al2MgC2 plates inC/Mg±Al composites can be estimated by comparing it withthe analogous growth of Al4C3 in C/Al composites: Formelt/fiber contact times of 0.3 s, the C/Al composites showed

an amorphous band of approximately 5 nm in thickness tohave formed around the fibers, which obviously was alumi-num with dissolved carbon.[65,66] This amorphous band con-tained nuclei of Al4C3. After raising the melt/fiber contacttime to 1.3 s, mm-sized Al4C3 plates have formed. In C/Mg±Al composites, the Al content is minor, and Al has to diffuseto the fiber/matrix interface prior to its incorporation intocarbides. Consequently, the growth of Al2MgC2 in C/Mg±Alsystems is expected to be slower than that of Al4C3 in C/Alsystems. Nevertheless, if the Al content of the Mg matrix ishigh enough, mm-sized Al2MgC2 plates should form within afew seconds.

In C/Mg±Al composites with a high fiber/matrix reactiv-ity, the carbidic plates are grown directly on the fiber surfaces(Fig. 6). Hence, if such a composite is loaded in the directionof the fiber axes, the fibers may be notched by the carbides,revealing the observed brittle fracture behavior (cf.Fig. 1c).[5,53±58]

The studies reviewed have shown that it is possible to con-trol the interface reactions in C/Mg±Al composite materialsby varying the aluminum content of the matrix and by usingcarbon fibers of different surface microstructures. For a medi-um interface reactivity, an MMC of optimum properties canbe produced, which, in the three-point bending test, is charac-terized by a bundle fracture behavior. Contrary to the wide-spread assumption of the major role of the binary carbideAl4C3 in the fiber/matrix chemistry of C/Mg±Al composites,the properties of the composites proved to be affected by theformation of plate-shaped ternary carbides (Al2MgC2).

Furthermore, note that in some manufacturing processeslarge amounts of oxygen are incorporated into the MMC, in-ducing the formation of thick, passivating MgO interlayers(100 nm, or thicker) between fiber and matrix.[67±71] In thoseC/Mg±Al composites, the remaining aluminum occurs asMg±Al intermetallic in the matrix, but not as carbide at the fi-


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Fig. 4. Projection of the atomic arrangement of Al2MgC2 on the (1120) plane.

Fig. 5. Ternary carbide in the matrix of a C/Mg±Al composite.

Fig. 6. HREM image showing an Al2MgC2 plate grown directly on the surface of ahigh-tensile-strength fiber in a composite with an AZ91 matrix (fiber below).


ber/matrix interface, leaving a weak fiber/matrix bond-ing.[67±74] This kind of passivation of the reaction front byoxide interlayers is well-known for C/Al systems, where wet-ting usually occurs if the oxide interlayer loses its compact-ness (at temperatures above 900 �C, or due to some alloyingwith elements of higher affinity to oxygen than to alumi-num).[19,75]

The next section discusses another possibility of achievingan appropriate interface bonding, which, at the same time,prevents the formation of carbides in C/Mg±Al composites.

3. Modification of the Fiber Surface by CoatingQuite a different way to generate an appropriate interlayer

is the direct fiber coating prior to the matrix infiltration. Thecoating should meet the following conditions: i) to provide anadequate fiber/matrix adhesion and ii) to act as a diffusionbarrier. The latter results because a potential application ofC/Mg composites, with a strength of 1 to 2 GPa, is the partialreinforcement of highly loaded parts in mechanical compo-nents (hybrids).[76±79] In those parts of the components wherelower stresses are applied, the strength of Mg alloys with 5 to9 wt.-% Al (approx. 250 MPa at room temperature) is suffi-cient, which is also considerably higher than that of pure Mg(150 to 180 MPa). However, as described in Section 2, in con-tact with carbon fibers (high-strength part of the hybrid),Mg±Al alloys with a high Al content will cause an extensiveformation of Al2MgC2 leading to an embrittlement of thecomposite. Coating the fiber surface to slow down the detri-mental mass transfer across the interface is a promising wayto suppress the carbide formation and to provide a mechani-cal protection of the fibers.

As coating materials interstitial compounds seem to besuitable for this purpose as they consist of close-packed metalatom sublattices, with their interstices (octahedral and/or tet-rahedral) being filled with smaller non-metallic atoms (hy-drogen, carbon, nitrogen, boron, etc.) thus having a low per-meability.[80] In general, interstitial compounds are hard andbrittle, and therefore the coating should not exceed a criticalthickness of about 1 % of the fiber diameter (df = 7 mm) toavoid notch effects, which might arise from microcracks inthe coating and which would lower the fiber strength.[81,82] Asit is necessary to achieve appropriate films each filament ofthe multifilamentous carbon fiber yarn, chemical vapor de-position (CVD) seems to be the appropriate technique.

As the coating has to be compatible to both fiber and ma-trix, the number of possible interstitial compounds to be usedfor the system C-fiber/coating/Mg-matrix is restricted. In ad-dition, a melt infiltration of the fiber prepreg is favorable (cf.Sec. 1), and hence a good wetting of the metal on the coatingis necessary.[19] In general, the wettability of metal-like com-pounds by metals is better than that of covalent ones as a highpercentage of delocalized electrons in the solid phase favorsthe electron exchange, which is necessary for the formation ofstable chemical bonds with the liquid-phase metal.[83]

Regarding the above, for the use in C/Mg±Al composites,TiN is an attractive coating material among the others[84±90]

that have been deposited on carbon fibers by CVD routes. Ti-tanium nitride has the cubic rock-salt structure, which is thesimplest structure of interstitial compounds, and it can bebrought onto the carbon fibers via CVD from TiCl4±N2±H2precursor gas mixtures,[91±93] even in thin films (i.e., some10 nm in thickness):[59,91±94]

Using a TiN coating of some 10 nm in thickness between ahigh-tensile-strength fiber and a cp-Mg matrix (commerciallypure Mg), a relatively high tensile strength perpendicular tothe fiber axis (approx. 25 MPa) was measured, indicating agood adhesion of the matrix to the coated carbon fibers.[92,93]

In addition, the tensile strength parallel to the fiber axes (ap-prox. 1200 MPa) as well as the bending strength (approx.1900 MPa) of this composite are extremely high.[92,93] In thismaterial, the interfacial structure is mainly characterized byan almost stringent separation of matrix and fiber by the coat-ing as it is shown in Figure 7, in the bright-field image (a)and corresponding EFTEM images taken at the ionizationedges Ti±L23 (456 eV) (b), C±K (284 eV) (c), and Mg±K(1305 eV) (d), with the fiber on the left and the matrix on theright. The elemental distribution of titanium (Fig. 7b) revealsthe continuous covering of the fiber with the coating, here 10to 20 nm in thickness. The distribution of carbon (Fig. 7c) rep-resents mainly the fiber, and that of magnesium (Fig. 7d) thematrix. In Figure 7d, some bright contrast features on the leftof the coating indicate that Mg penetrates it (the coating)slightly towards the fiber. The HREM image of Figure 8shows the turbostratic carbon of the fiber (left) and the (1010)lattice fringes of the Mg matrix (right) ending at the coating(middle part). Clearly revealed is the polycrystalline nature


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Fig. 7. TiN-coated high-tensile-strength fiber in a cp-Mg matrix: a) bright-field image,b±d) EFTEM at ionization edges.


of the coating, with grain sizes comparable to the film thick-ness.

Figure 9a show a three-dimensional plot of EEL spectra re-corded across the fiber/matrix interregion at 30 equidistantpoints at spacings of about 1.3 nm along the white marker inthe STEM bright-field image (Fig. 9b). In the energy intervalchosen between 250 and 750 eV, there appear the C±K, N±K,Ti±L23, and O±K ionization edges with their associated fine-structures. In the region of the fiber (spectra in the back-ground, Fig. 9a), only the C±K ionization edge at 284 eVarises from the exponentially decreasing background. At thefiber/coating interface, a Ti±L23 ionization edge appears, andthe C±K ELNES changes. Within a gradient of approximately5 nm, an N±K ionization edge arises and the C±K edge disap-pears. The Ti±L23 ELNES remains very clear above the back-ground due to the presence of sharp white lines. The latter oc-cur at the L23 ionization edges of the transition metals and atthe M45 ionization edges of the rare-earth elements, becausethe Fermi energy is located in the narrow 3d or 4f valence

band, respectively.[21] At the coating/matrix interface, oxygenis present, which is indicated by an O±K ionization edge (cf.Fig. 9a). Reaching the first matrix areas, again carbon occurs.

A detailed analysis of the EEL spectra of Figure 9 showedthat within the coating the C±K and N±K edges have similarELNES profiles,[94] resembling those of stoichiometric TiCand TiN,[95±97] which have a similar crystal chemistry.[95] Bothcompounds exhibit the rock-salt structure, where the Tiatoms form an face-centered cubic (fcc) lattice, with the octa-hedral interstices being filled completely with C and N atoms,respectively. The occupation of some octahedral intersticeswith carbon and others with nitrogen, leads to the formationof titanium carbonitrides TiCxNy (x + y £ 1) substitutional sol-id solutions with a wide range of the C and N solubility.[98]

Changes in the stoichiometry affect the C±K and N±K ELNESof these carbonitrides only slightly.[99]

Including quantification of the EEL spectra of Figure 9, thecoating was proven to consist of a substoichiometric titaniumcarbonitride (TiCxNy), with a high carbon content at the fi-ber/coating interface, and gradually becoming richer in nitro-gen to the middle of the film.[94] The compositional changeswithin the coating indicate that the reactor gas mixture (TiCl4,N2, and H2)

[91±93] reacts with carbon from the fiber at the be-ginning of CVD. Thus, the term of ªreactive chemical vapordepositionº (RCVD)[85] seems to be suitable here.

Locally, carbon is transported from the fiber to the matrix(cf. Figs. 7c and 9). It is assumed that, first of all, liquid Mg pe-netrates through fissures in the film towards the fiber (cf.Fig. 7d):[94] In successive steps, carbon from the fiber dissolvesand diffuses in the liquid Mg through the coating and precipi-tates as graphitic ribbons at the coating/matrix interface.

Of course, in presence of Al in the matrix carbides wouldform at these locations. Inherently, in the composites, bothgraphitic ribbons[4,5,54,94] and Al2MgC2

[39,94] form via isother-mal dissolution±diffusion±precipitation mechanisms, whichare driven by a gradient of the chemical potential of car-bon.[100,101]

Conclusively, when the Mg matrix is alloyed with 5 wt.-%Al (AM50), locally isolated carbidic plates are observed, withdimensions above 1 mm, which may pierce the TiCxNy filmand reach underneath the fiber surface.[94] This effect is notvery pronounced, and the microstructure of the C/TiN/AM50 composite, which is shown in Figure 10, resembles thatof the corresponding C/TiN/cp-Mg composite (cf. Fig. 7a),with its fiber/matrix interfaces being mainly free of carbides.This is due to the fact that the TiCxNy coating serves well asan inhibitor of deleterious carbide formation, even if Mg al-loys with large amounts of Al as matrices are used.

However, combining the same Mg±Al matrix with the un-coated fiber results in the substantial formation of many car-bidic Al2MgC2 plates of different size at the fiber/matrix in-terface (Fig. 11). Here, many carbides extend up to 1.5 mminto the matrix[94] and embrittle the composite (cf. Sec. 2).

Of course, in the C/TiN/AM50 composite at those loca-tions where the carbidic plates have formed notching may oc-


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Fig. 8. TiN-coated high-tensile-strength fiber in a cp-Mg matrix: HREM image show-ing the (1010) lattice fringes of the Mg matrix ending at the coating (fiber left, matrixright).

Fig. 9. TiN-coated high-tensile-strength fiber in a cp-Mg matrix: a) three-dimensionalplot of EEL spectra taken at 30 equidistant points across the fiber/matrix interregion,arranged from fiber (back) to matrix (front), b) STEM bright-field image showing theanalysis distance as a white line.


cur, which can initiate fiber cracking. But, as there are only afew of these large carbidic plates formed in the C/TiN/AM50composite, the effect is very limited here. Accordingly, thestrength of this composite turned out to be only slightly lowerthan that of the corresponding C/TiN/cp-Mg composite.[92,93]

Thus, the manufacturing of hybrid components based onMg±Al matrices is enabled by the diffusion retarding effect ofthe chemical vapor deposited TiCxNy coating.

Besides suppressing deleterious carbide formation theTiCxNy interlayer offers a further benefit: For uncoated high-tensile-strength carbon fibers, the penetration of significantamounts of magnesium into the body of the fibers has beenreported, which seems to cause a degradation of the fiberstrength.[11,54,94,102±105] A TiCxNy interlayer has been proven toreduce the magnesium penetration remarkably.[94] Thus, fiberdegradation as well as carbide formation can be avoided si-multaneously by a proper interlayer.

4. ConclusionsDepending on the intended purpose of the special C/Mg±

Al composites two strategies of interface optimization are ap-plicable:

Firstly, the fiber/matrix reactivity of C/Mg±Al compositescan be controlled by varying i) the Al content of the Mg ma-trix and ii) the carbon fiber type. Appropriate interface prop-erties can be achieved with small amounts of Al (approx.

2 wt.-%) and carbon fibers of the high-tensile-strength typevia moderate formation of the ternary carbide Al2MgC2 in thefiber/matrix interlayer, thus adjusting the interfacial bond-ing.

Secondly, for hybrid components higher Al contents of thematrix are needed. Here, the formation of Al2MgC2 becomesstronger and embrittles the composite material. An appropri-ate TiCxNy coating on the carbon fiber prior to the compositemanufacturing prevents this deleterious carbide formationand enables the use of Mg±Al matrices with high Al contents.

A further improvement of C/Mg±Al composites with largeamounts of Al in the matrix alloy requires a tailoring of theTiCxNy interlayer with respect to compactness and grain sizedistribution, which is one challenging topic for additional at-tempts of interface engineering in metal matrix composites.

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Fig. 10. Bright-field image of a TiN-coated high-tensile-strength fiber in an AM50 ma-trix showing the fiber/matrix interfacial region to be mainly free of carbide precipitates.

Fig. 11. Bright-field image of an uncoated high-tensile-strength fiber in an AM50 ma-trix showing many platelets of Al2MgC2 in the fiber/matrix interregion.


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