Open Ceramics 6 (2021) 100090

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Open Ceramics

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Processing and oxidation response of Cr2AlC MAX-phase compositescontaining ceramic fibers

Teresa Go a,b,*, Robert Vaßen a, Olivier Guillon a,b,c, Jesus Gonzalez-Julian a,b

a Forschungszentrum Jülich GmbH, Institute of Energy and Climate Research, Materials Synthesis and Processing (IEK-1), 52425, Jülich, Germanyb Department of Ceramics and Refractory Materials, Institute of Mineral Engineering, RWTH Aachen University, 52064, Aachen, Germanyc Jülich Aachen Research Alliance, JARA-Energy, 52425, Jülich, Germany

A R T I C L E I N F O

Keywords:MAX phasesCr2AlCOxidation responseCeramic fibersCMC

* Corresponding author. Forschungszentrum JüliGermany.

E-mail address: t.go@fz-juelich.de (T. Go).

https://doi.org/10.1016/j.oceram.2021.100090Received 18 December 2020; Received in revised fAvailable online 24 March 20212666-5395/© 2021 The Author(s). Published by El(http://creativecommons.org/licenses/by-nc-nd/4.0/).

A B S T R A C T

Three different ceramic matrix composites (CMCs) were produced using Cr2AlC as a matrix, and carbon, SiC andAl2O3 short fibers as a secondary phase. Cr2AlC powders were synthesized by solid-state reaction, followed bymixing with the fibers, and full densification using a field-assisted sintering technique. Of the three different fibertypes, Carbon fibers reacted strongly with Cr2AlC, while the reaction with SiC fibers was more limited andalumina fibers didn’t show any reaction. Oxidation tests of the monolithic Cr2AlC and the composites wereperformed by thermogravimetric analysis. An alumina layer formed at 1000 �C on every sample, well attachedand worked as a good oxidation barrier. Under realistic conditions using a burner rig for cyclic oxidation at 1200�C for 500 cycles, the oxidation resistance of the alumina fiber CMC is good, as no defects or degradation arevisible and the alumina layer is well attached.

1. Introduction

MAX phases are a relatively new family of materials with more than150 different compositions, which have gained considerable attentiondue to their combination of metallic and ceramic properties [1–3]. Asceramics, MAX phases are stiff and lightweight, while as metals they arereadily machinable, damage-tolerant, good thermal and electrical con-ductors, and thermal shock resistant [4,5]. Additionally, some composi-tions that contain aluminum (Al-based MAX phases that are mostlyTi2AlC, Cr2AlC, and Ti3AlC2) exhibit good oxidation and corrosionresistance and crack self-healing behavior up to 1400 �C under aggressiveconditions thanks to the in-situ formation of an external, dense, andadherent α-Al2O3 layer [6–9]. Among these Al-basedMAX phases, Cr2AlCdemonstrated great potential for high-temperature applications,although it has been less investigated than the other two phases. It alsoshowed an excellent oxidation response in air up to 1000 �C (oxidationrate of 1.1⋅10�11 kg2m�4s�1 at 1000 �C) and a good response up to 1200�C (oxidation rate of 5.6 10�10 kg2m�4s�1 at 1200 �C) [10,11]. However,the formation of a porous Cr7C3 layer underneath the Al2O3 layer limitsits performance at higher temperatures. In addition, secondary phasessuch as carbides, which are typically formed during the synthesis process,play a crucial role in the oxidation response. In 2007, Lee and Park

ch GmbH, Institute of Energy an

orm 15 March 2021; Accepted 15

sevier Ltd on behalf of European

proposed that the primary formation of the α-Al2O3 layer on the surfaceof Cr2AlC is based on the inward diffusion of oxygen [12]. One year later,based on the microstructure and phase composition of the oxidizedsurface, Tian et al. additionally demonstrated an outward diffusion ofaluminum to form the α-Al2O3 layer. Below this alumina layer, a porouslayer containing Cr7C3 and Cr3C2 formed [13]. However, Cr2AlC is theonly known MAX phase composition that forms an intermediate porouslayer between the protective alumina scale and the unreacted material[6]. Lee et al. investigated the long-term (360 days) oxidation behavior ofCr2AlC at temperatures from 700 �C to 1000 �C [14]. They foundmetastable θ-Al2O3 platelets formed on the α-Al2O3 layer, which almostcompletely disappeared at 1100 �C due to θ → α transformation. This isalso known from alumina scale formation on alloys [15,16]. In addition,at every temperature, an Al2O3 layer formed with a Cr7C3 sublayer, with0.7–8.3 at.% Cr detected in the Al2O3 layer. In that study, the oxidationbehavior was described as parabolic in the temperature range between700 �C and 1100 �C, with a parabolic rate constant of 7.5⋅10�11

kg2m�4s�1 at 1100 �C. Lin et al. also reported parabolic rate constants of1.08⋅10�12 kg2m�4s�1 at 800 �C and 2.96⋅10�9 kg2m�4s�1 at 1300 �C,respectively, measured by thermogravimetric analysis [10]. These valueswere faster as compared to the rate constants of alumina scale formingalloys [17].

d Climate Research, Materials Synthesis and Processing (IEK-1), 52425, Jülich,

March 2021

Ceramic Society. This is an open access article under the CC BY-NC-ND license

T. Go et al. Open Ceramics 6 (2021) 100090

In 2013, Tallman et al. reviewed the oxidation response of Ti2AlC,Ti3AlC2, and Cr2AlC in air [6]. The activation energy for the oxidation ofCr2AlC is 507� 90 kJ/mol, which is roughly double that of the activationenergy of Ti2AlC and Ti3AlC2. The reason for this higher activation en-ergy remained unclear. Furthermore, they concluded that the oxidationbehavior of Cr2AlC as parabolic was not precise and that also a cubic rateconstant was only an approximation. The oxidation kinetics of Cr2AlC issensitive to several factors, such as purity, orientation, and the size of thegrains. As a result, the oxidation kinetics of Cr2AlC cannot be fittedproperly using a simple model. Smialek concluded from comparingdifferent datasets of TGA investigations for Cr2AlC that the higher weightloss above 1200 �C is accompanied with higher amounts of Cr7C3 [18]. Insome of the materials, Cr7C3 oxidized to Cr2O3. It is possible that athigher temperatures, volatile species such as CrO3 accelerated thedegradation mechanism. In these cases, the oxidation rates no longerhave any comparative value. In addition to these experiments underwell-defined and mild conditions, oxidation tests were performed underrealistic operating conditions similar to the corrosive gas turbine envi-ronments using a burner rig [19]. Cr2AlC was able to withstand at least500 short thermal cycles at 1200 �C due to the formation of a denseα-Al2O3 layer. This protective layer exhibited strong adhesion to theCr2AlC sample, which showed no visible damage. The thermal stress isreduced due to a good matching CTE between the Cr2AlC sample(12.0⋅10�6 K�1) and the alumina scale (8.5⋅10�6 K�1), which is onereason for the strong adhesion and the convoluted interface [20].

Ceramic matrix composites (CMCs) are the most common approachfor facilitating a pseudo-plastic deformation and increasing resistance tothe crack propagation of brittle materials. In addition to crack deflection,other reinforcing mechanisms, such as fiber pull-out, enhance the me-chanical response of the final composites. The fracture energy isincreased by the dissipative fiber pull out mechanism, giving the com-posite an enhanced strain tolerance. In the case of low-strength matrixmaterials with low elastic modulus (e.g. carbon, mullite), mechanicalresilience is increased by transferring the load to the reinforcing fibers.For maximum load transfer, the fibers must be firmly integrated into thematrix [21,22]. Ideally, damage-tolerant behavior is achieved, meaningthat internal defects or external damage under load do not lead to asudden total failure of the component [21]. In the particular case of MAXphases, the reinforcement with ceramic fibers is only investigated to aminor extent despite its high potential [23–31], in particular forCr2AlC-based materials. Lenz and Krenkel fabricated Ti3SiC2 compositescontaining carbon fibers by liquid silicon infiltration, using carbon andTiC preforms as starting materials [23]. The Ti3SiC2 matrix was formedafter infiltration and a thermal process at 1350 �C for 8 h. Dash et al.investigated the creep behavior of Ti3SiC2 reinforced with SiC whiskersbetween 1100 �C and 1300 �C [24,25]. The creep rates decreased byaround two orders of magnitude with 10 vol.-% SiC whiskers. The acti-vation energy for creep was constant at 650–700 kJ/mol forsmall-grained Ti3SiC2, and increased – compared to the monolithic ma-terial – from 454� 29 kJ/mol to 576� 33 kJ/mol for the coarse-grainedMAX phase material. Spencer et al. examined the reaction between SiCfibers (from Nippon Carbon Co. Ltd, Tokyo, Japan) and commercialTi2AlC or Ti3SiC2 powder (both from 3-ONE-2,Voorhees, NJ) [26]. Inboth cases, the production of fully dense composites was successful. Theydetected a reaction between the SiC fibers and Ti2AlC as a matrix ma-terial, while no reaction was observed for Ti3SiC2. Interfacial character-ization and the mechanical behavior of composites consisting of Ti3AlC2and continuous SiC fibers were studied by Guo et al., in 2014 [27]. Athick interface reaction layer formed between the SiC fiber and the ma-trix during hot pressing at temperatures of 1300 �C. The dominantdiffusing species in the interface reaction between the SiC fiber and theTi3AlC2 matrix was aluminum. Due to the presence of a thin Ti foilcoating on the fibers, which acts as a diffusion barrier, the compositesexhibited non-catastrophic fracture behavior and thus the mechanicalproperties were improved [28]. Parrikar et al. used Al2O3 fibers toreinforce Ti2AlC and thus improved the compressive fracture strength by

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up to 39.7 % under static load [29], although no information about theincorporation of the fibers or the reaction between fibers and matrix wasreported. In another study, Spencer et al. investigated composites con-sisting of Ti2AlC and alumina fibers [30]. Fully dense composites wereobtained at temperatures below 1300 �C, while at higher temperatures,the fibers sintered together. As mentioned above, there have been limitedinvestigations of Cr2AlC reinforced with ceramic fibers and only onestudy was published in 2017. In that study, Gonzalez et al. developedCr2AlC composites containing short SiC fibers, increasing wear resistanceup to 70–80 % due to the incorporation of SiC fibers [31].

Especially for very harsh environments, the existing strength of theMAX phases is insufficient, so that reinforcement is necessary. Ceramicfibers can provide this reinforcement and, due to the various reinforce-ment mechanisms, work better than, for example, particle reinforcement.In this study, ceramic matrix composites based on Cr2AlC containing C,SiC, or Al2O3 short fibers were developed. This three different kinds offibers with a wide range of costs are chosen to find the best reinforcementfor Cr2AlC. The composites were fully densified using the field-assistedsintering technique (FAST/SPS), and the potential reactions betweenthe fibers and the matrix analyzed. The oxidation behavior was studiedby thermographic analysis up to 1300 �C at different times and comparedto the monolithic response as well as the literature. Additionally, thecomposites were tested under realistic thermal conditions using a burnerrig at 1200 �C to explore the potential of these materials as componentsin gas turbines.

2. Experimental

Cr2AlC powder was synthesized using chromium carbide (Cr3C2),aluminum, and chromium powders (all from Alfa Aesar, Germany) bysolid–liquid-state reaction at a high temperature. These powders – Cr3C2(�325 mesh, d50 ¼ 6.0 μm, 99.5% pure), Al (�325 mesh, d50 ¼ 5.6 μm,99.5% pure), and Cr (�325 mesh, d50 ¼ 28.0 μm, 99% pure) – weremixed according to the weight ratio 1:2.1:1. Compared to the stoichio-metric composition, an additional 10 wt% of aluminum was added tocompensate its loss during the thermal synthesis process. The powdermixture was treated at 1300 �C for 2 h under an argon atmosphere with aheating rate of 10 �C/min. The resulting Cr2AlC porous pellet wascrushed and milled using a planetary mill (PM400, Retsch GmbH, Ger-many). ZrO2 milling balls were used with a 5 mm diameter, a powder toball ratio of 1:4 in weight, ethanol as liquid media, and 300 rpm for 1 h.The milling balls were removed and the Cr2AlC powder was dried in anoven at 70 �C overnight. A Cr2AlC powder was obtained with a unimodalparticle-size distribution and a mean particle size of 7.0 μm, measured bylaser diffraction (Horiba LA950V2, Retsch, Germany).

The Cr2AlC powder was mixed with three different kinds of short fi-bers: carbon fibers, silicon carbide fibers, and alumina fibers. The carbonfibers (Tenax®, Teijin Carbon Europe GmbH) were approx. 100 μm inlength, had a diameter of 5–7 μm, and a density of 1.8 g/cm3 [32]. TheSiC fibers (Hi-Nicalon™, COI Ceramics, Inc.) were approx. 1 mm long,had a diameter of approx. 14 μm, and a density of 2.65 g/cm3 [33]. Thealumina fibers (Nextel™, 3M™ Deutschland GmbH) were continuousfibers with a diameter of approx. 12 μm and a density of 3.9 g/cm3 [34].These alumina fibers were chopped by hand into a length of approx. 1mm. The alumina fibers were desized at 700 �C for 30 min, while thecarbon and SiC fibers were washed three times with boiling water fol-lowed by drying at 80 �C in an oven overnight.

The Cr2AlC powder was then mixed with the different short fibers,using ad-hoc developed different methods for each fiber type. Eachpowder batch was targeted at producing 15 g (20 g) of powder, so toproduce a 20 mm (30 mm) diameter and 6 mm (4 mm) height sinteredpellet. The Cr2AlC powder and carbon fibers were dispersed in ethanol bymechanical stirring for 30 min, followed by drying in a rotary evaporator.The Cr2AlC powder were weighed out with the SiC fibers, and themixture was dispersed by mechanical agitation for 5 min. For the lastcomposition, the alumina fibers were dispersed in ethanol by mechanical

Fig. 1. a) powder particles and b) XRD pattern of Cr2AlC powder after milling; c) low magnification and d) high magnification of sintered Cr2AlC.

T. Go et al. Open Ceramics 6 (2021) 100090

stirring for 10 min and the Cr2AlC powder was subsequently added andmixed for further 10 min. Each suspension was heated during the me-chanical stirring to evaporate the ethanol until 50 ml were left. Thepowder mixture was then completely dried in an oven at 70 �C. Thecompositions are named according to the type and weight content of thefibers, resulting in the following names: C-10, SiC-5, SiC-10, SiC-15, SiC-20, Al2O3-5, Al2O3-10, Al2O3-15, and Al2O3-20. In addition, one mono-lithic sample was prepared without fibers with the name Cr2AlC. 10 wt%C fibers correspond to 25,8 vol%. For SiC fibers 5, 10, 15 and 20 wt%correspond to 11.0, 22.0, 33.0 and 43,9 vol%, respectively. 5, 10, 15 and20 wt% of Al2O3 fibers correspond to 7.4, 14.7, 22.1 and 29.5 vol%,respectively.

All the powder compositions were sintered by field-assisted sinteringtechnology/spark plasma sintering (FAST/SPS, FCT-HPD5, FCT SystemeGmbH, Germany) using standard graphite dies with a diameter of 20 mmor 30 mm. A graphite foil was used to separate the powders from thegraphite die and the punches. The sintering conditions were as follows: apressure of 50 MPa, a temperature of 1200 �C, a heating rate of 100 �C/min, a dwell time of 10 min, and a vacuum atmosphere. After sintering,the graphite foil was removed from the samples and the pellets were cutinto smaller equilateral triangles with an edge length of approximately 9mm and a height of approximately 2 mm for characterization. A mono-lithic Cr2AlC sample was sintered under the same conditions as referencespecimen.

Crystal phases of the powders and sintered samples were analyzed byX-ray diffraction (XRD, D4-Endeavor, Bruker AXS, Germany) with Cu-Kαradiation over the range of 10–80 2θ degree and a step size of 0.02�. Withthe TOPAS V 4.2 software (Bruker AXS, Karlsruhe, Germany), a Rietveldrefinement is carried out and thus the mass fractions of the detectedphases are determined [35]. For crystallographic files the InorganicCrystal Structure Database (ICSD, FIZ Karlsruhe, Germany) is used withthe following ICDS numbers: ICSD-42918 (Cr2AlC) [36], ICSD-24170

3

(SiC) [37], ICSD-88028 (Al2O3) [38], ICSD-52230 (C) [39],ICSD-44731 (Cr) [40], ICSD-57009 (Cr3C2) [41], ICSD-76799 (Cr7C3)[42], ICSD-617535 (Cr5Si3C) [43], ICSD-66751 (Al4C3) [44],ICSD-57651 (AlCr2) [45]. Cross sections of the samples were preparedand observed using a scanning electron microscope (SEM, Zeiss Ultra 55,Germany). In addition, some of the samples were analyzed by energydispersive X-ray spectroscopy (EDX). The density was determined byArchimedes’ method using the rule of mixture for the composites andtheoretical values of 5.24 g/cm3, 2.26 g/cm3, 2.65 g/cm3, and 3.95g/cm3 for Cr2AlC, C, SiC, and Al2O3, respectively. The samples used forthe thermogravimetric analysis (TG) and oxidation tests weremirror-polished using sandpapers and diamond paste. For TG tests, fivedifferent samples – Cr2AlC, SiC-10, SiC-20, Al2O3-10, and Al2O3-20 –

were measured at five different temperatures – 800 �C, 1000 �C, 1100 �C,1200 �C, and 1300 �C – with a heating rate of 10 �C/min. The isothermalholding time was 20 h at 800 �C, 1000 �C, and 1300 �C and 68 h at 1100�C and 1200 �C. During the TG measurements the mass gain Δm ismeasured. Together with the corresponding oxidation time t and thesurface area of the samples A, the parabolic (kp) or cubic (kc) oxidationconstant can be determined using the following formulas (1) and (2):

kp ¼ðΔm=AÞ2t

(1)

or

kc ¼ðΔm=AÞ3t

(2)

For oxidation, Cr2AlC, SiC-10, and Al2O3-10 were heated for 50 h at1000 �C. Cross sections were analyzed by SEM. In addition, cycling testswere carried out with a burner rig and round specimens with a diameterof 30 mm. Detailed information about the facility can be found elsewhere

Table 1Phase compositions based on Rietveld refinement after sintering and relativedensities recalculated accordingly for all produced composites.

Phase composition [wt.%]

Cr2AlC SiC-10 Al2O3-10 C-10

Cr2AlC 94 Cr2AlC 80 Cr2AlC 88 Cr2AlC 73

AlC2 3 Cr5Si3C 8 Al2O3 11 C 12Al2O3 2 Al4C3 4 Cr 1 Cr7C3 6Cr7C3 1 SiC 3 AlCr2 4

Cr3C2 3 Cr3C2 3Al2O3 2 Al2O3 2

Relative density [%]Cr2AlC SiC-composites Al2O3- composites C-1097,5 � 0,4 SiC-5 97,6 � 0,3 Al2O3-5 98,2 � 0,4 84,9 � 0,1

SiC-10 96,6 � 0,2 Al2O3-10 97,0 � 0,2SiC-15 94,3 � 0,5 Al2O3-15 97,2 � 0,8SiC-20 90,3 � 1,3 Al2O3-20 95,6 � 0,3

T. Go et al. Open Ceramics 6 (2021) 100090

[46]. In brief, natural gas (CH4) and oxygen are used in a volume ratio of1:2.2 (over-stoichiometric) to operate the gas burner rig. The wholesystem is PLC-controlled and all gas fluxes were adjusted by mass flowcontrollers. The surface temperature was measured with a long-wavepyrometer. As the emissivity of the examined materials is unknown,and thus the measurement of the temperature is not precise, a magnetite(Fe3O4) coating was applied to the sample. This coating was placedprecisely in the center of the sample, where the temperature wasmeasured. CMCs with 10 wt% SiC or Al2O3 fibers were heated with aburner for 5 min and then cooled for 2 min. A total of 500 such cycleswere performed achieving a sample surface temperature of 1200 �C. Theoverall heating time was thus approximately 41 h for each sample. Afterthe thermal cycling tests, the surface of the samples was observed bySEM. The samples were subsequently cut and polished to analyze thecross section. The chemical composition of the different phases wasidentified by EDX.

Fig. 2. Overview and detailed view of the composites with 10 wt% a

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3. Results and discussion

3.1. Microstructural characterization of the composites

The synthesized and milled Cr2AlC powder presents a unimodalparticle-size distribution with a mean particle size of 7.0 μm (d50), and ad10 and d90 of 4.1 μm and 13.2 μm, respectively. The typical layeredstructure of the MAX-phase particles is shown in Fig. 1a. The XRD pattern(Fig. 1b) reveals the high purity of the powder, where most of the reflexescorrespond to the Cr2AlC phase. Nevertheless, traces of Al2O3 and Al8Cr5are also detected. Rietveld refinement determined a composition of 96 wt% Cr2AlC, 2 wt% Al2O3, and 2 wt% Al8Cr5. This powder was sintered as ablank specimen, and an overview of the surface at two different magni-fications is shown in Fig. 1c and d. In the sintered reference sample, fourdifferent phases were detected by XRD and the content of each phase wascalculated by Rietveld refinement (Table 1). The amount of Cr2AlCdecreased during sintering, but a 94 wt% was preserved, the amount ofthe other phases resulted to be Al2O3 (2 wt%), AlCr2 (3 wt%), and Cr7C3(1 wt%). The change in the phase composition and the content of phasesafter the sintering process indicates some reactions. Cr7C3 is typicallyformed due to the degradation of Cr2AlC [47], while the Al8Cr5 present inthe powder might have reacted with traces (below the detection limit ofXRD) of C or Cr to form Cr2AlC and/or AlCr2. The relative density aftersintering was 97.0 % using the rule of mixture re-adapted to the actualpresent phases.

In a first phase, the three different composites containing 10 wt% offibers were sintered by FAST/SPS in order to evaluate the dispersion andpotential reactions. Fig. 2 shows the polished cross sections at differentmagnifications of the sintered composites. Carbon fibers were homoge-nously dispersed into the matrix (Fig. 2a), although reactions betweenthe fibers and matrix can be observed at a higher magnification (Fig. 2b).The contour of the fibers is uneven, and Rietveld analysis indicatedCr2AlC and C as well as Al2O3, Cr7C3, AlCr2, and Cr3C2 (Table 1). Incomparison to the monolithic sample, the amount of Cr7C3 increasedfrom 1 wt.-% to 6 wt%. AlCr2 increased from 3 wt% to 4 wt%, and 3 wt%of Cr3C2 formed in C-10 compared to the monolithic sample. The amountof Al2O3 remained constant with a content of 2 wt%. In some carbonfibers, a dark core is visible, as shown in Fig. 3. Here, some almost purecarbon remained, as EDX measurements show. In the nearby mixed greyand black area, chromium peaks increased and the carbon peak

) and b) C fibers; c) and d) SiC fibers; and e) and f) Al2O3 fibers.

Fig. 3. SEM micrograph and EDX measurements of C-10.

Fig. 4. SEM micrograph and EDX measurements of SiC-10.

T. Go et al. Open Ceramics 6 (2021) 100090

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Fig. 5. Diffraction patterns of a) Cr2AlC, SiC-10, and SiC fibers, and b) Cr2AlC, Al2O3-10, and Al2O3 fibers.

T. Go et al. Open Ceramics 6 (2021) 100090

decreased suggestion Chromium diffusion from the matrix. There is anincreased amount of aluminum in the darker and homogenous border ofthe carbon fiber (see EDX measuring point 3 in Fig. 3). The light greyphase (point 5) corresponds to chromium carbide and was manly foundaround the carbon fibers. The former carbon fibers consisted of a mixtureof different phases after sintering; with only some fibers displaying asmall carbon core. In addition, the density of C-10 was only 84.9� 0.1%,which is much lower than the blank specimen (97.5� 0.4%), Table 1. Asa result, composites containing carbon fibers were discarded due to thereaction of the fibers and the low density.

SiC fibers were well distributed in the matrix without any visibleagglomeration (Fig. 2c). As expected, the fibers are orientated due to theuniaxial pressure during sintering. A reaction between the SiC fibers andthe matrix material is visible at higher magnification (Fig. 2d). Thecontour of the fibers is irregular and the secondary phases formed at theinterface. Six different phases were revealed by Rietveld analysis: Cr2AlC(80 wt%), Cr5Si3C (8 wt%), Al4C3 (4 wt%), SiC (3 wt%), Cr3C2 (3 wt%),and Al2O3 (2 wt%) (Table 1). SiC fibers might react with the Cr2AlCmatrix according to equation [3].

3 SiC þ 4 Cr2AlC → Cr5Si3C þ Al4C3 þ Cr3C2 þ C (3)

EDX measurements were performed to identify the different phasesand understand the different reaction paths (Fig. 4). As expected, point 1

Fig. 6. Fiber dispersion in a) SiC-5, b) SiC-10, c) SiC-15, d) SiC

6

at the core of the fiber mainly exhibits Si and C. Point 2 is situated in thelighter grey area of the fiber, close to the interface. There, aluminum isdetected in addition to silicon and carbon, indicating diffusion from thematrix into the fiber. In point 3, the dark grey phase with an irregularshape, the aluminum peak increases strongly while the silicon peak de-creases considerately. Identification of this phase is complicated due tothe size and the number of phases around, but it might be correlated toAl4C3. In addition, some chromium and oxygen peaks were present, too.This dark phase is surrounded by a light grey phase (point 4), in whichmainly silicon and chromium are detected. The amount of siliconincreased again after an area with less silicon. The peaks in point 5correspond to the matrix material Cr2AlC with a small peak of silicon.This silicon is correlated with the diffusion from the fiber to the matrixduring sintering or due to the proximity of the silicon carbide fibers.Using EDX, a small area was investigated to determine the reaction be-tween the matrix and the fiber. Thus, the matrix and possible secondaryphases in this matrix are not taken into account. Besides the main phaseCr2AlC, 8 wt% of Cr5Si3C is detected by XRD (Fig. 5). This might be thelight grey phase surrounding the fibers, measuring point 4 in Fig. 4.These results do not coincide with the literature, where no reactionappeared between the SiC fibers and the Cr2AlC [31]. The main differ-ence between the previous results in the literature and this study is thetype of SiC fibers and the purity of the Cr2AlC. Both types of fibers – SF-7(SI-TUFF SF-7, Advanced Composite Materials, LLC, US) and

-20, e) Al2O3-5, f) Al2O3-10, g) Al2O3-15 and h) Al2O3-20.

Fig. 7. Fracture surfaces of a) SiC-10 and b) Al2O3-10.

T. Go et al. Open Ceramics 6 (2021) 100090

Hi-Nicalon™ (Hi-Nicalon™ from COI Ceramics, Inc.) – are composed offine-grained β-SiC. Other studies on non-oxide ceramics with Hi-Nica-lon™ fibers showed a reinforcing effect up to 1700 �C sintering tem-perature as long as the majority of the fiber core is present, even if there isa strong reaction between fiber and matrix [48–50]. In this study, XRDanalysis showed only SiC reflexes for the Hi-Nicalon™ fibers (Fig. 5). Asthe detection limit of XRD is about 1–2 wt.% of each phase, it is notpossible to detect small amounts of additives/secondary phases with thismethod. Nevertheless, differences in the composition of the fibers couldlead to the differences in the reaction.

As last composite, the dispersion of Al2O3 fibers was not completelyhomogenous, as some agglomerations were detected at some locations inAl2O3-10 (Fig. 2e). Interestingly, the fibers did not react with the matrix,and consequently clean interfaces were observed within the resolutionlimits (Fig. 2f)). Three phases were detected by XRD: Cr2AlC, Al2O3, andpure Cr (Table 1, Fig. 5). The amount of chromium (1 wt%) is near thedetection limit of XRD and thus the analysis must be considered withcaution.

In general, the distribution of SiC fibers was better than the distri-bution of Al2O3 fibers, Fig. 6. In both cases, some cracks appeared in thecomposite material when the fiber contents exceeded 10 wt%, Fig. 6c),d), g) and h), which might be correlated to differences in the CTE’s of thematerials. In addition, the agglomeration of the fibers increased stronglywith the content and, consequently, the trapped porosity. Fiber contentsof 5 wt% and 10 wt% showed a good distribution and the changes in thecomposite density are modest, as shown in Fig. 6 and Table 1. With SiCfibers, the density varied from 97.0% (without fibers) to 90.3% of thetheoretical density (with 20 wt% fibers). The evaluation of the densitywas not linear: with 10 wt% SiC fibers; the relative density was about96.6%. The high variation of the relative density of the composite with

Fig. 8. Mass gain over time between 800 �C and 13

7

SiC fibers might be due to the reactions between the fibers and the matrixmaterial. The relative density of composites with Al2O3 fibers variedfrom 98.2� 0.4% for Al2O3-5 to 95.6� 0.3% for Al2O3-20. This variationof almost 3% might be caused by increasing trapped porosity within thehigh aspect ratio fibers, which tended when added in amount above 10wt%. Based on this results, the composites SiC-10 and Al2O3-10 werechosen for the burner rig experiments. The amount of 10 wt% fibers waschosen as compromise of between the highest possible fiber content andthe highest possible density.

The fracture surfaces of SiC-10 and Al2O3-10 showed different frac-ture behavior (Fig. 7). SiC fibers were broken evenly with the matrixmaterial and some cracks run from the matrix into the fibers. These fiberswere strongly adherent to the matrix, likely due to the reaction betweenthe fibers and the matrix. In contrast, the Al2O3 fibers and the Cr2AlCmatrix exhibited weak interfaces. Some fibers detached from the MAX-phase matrix displayed limited pull-out, Fig. 7b). This fiber pull-outmight help to enhance the mechanical properties. Nevertheless, the fi-bers broke and their fracture surfaces were as fine-grained as beforesintering within Cr2AlC. Overall, the MAX-phase grains in the fracturesurface of Al2O3-10 were coarser those that of SiC-10 with the typicalMAX-phase layered structure is visible. This is another indication that noreaction between the Al2O3 fibers and the Cr2AlC matrix occurred.

3.2. Oxidation response of Cr2AlC composites

The two composites containing 10 wt% SiC and alumina fibers as wellas the reference Cr2AlC sample were oxidized at temperatures of up to1300 �C for 68 h. Fig. 8 shows the TG curves for the fiber free specimenand the composites. As expected, the higher the temperature, the higherthe mass gain for each material. At temperatures up to 1000 �C, the mass

00 �C for a) Cr2AlC, b) SiC-10, and c) Al2O3-10.

Fig. 9. Surface and cross section after 50 h at 1000 �C of a) Cr2AlC, b) SiC-10, and c) Al2O3-10.

T. Go et al. Open Ceramics 6 (2021) 100090

gain for all the materials is below 0.05 %/mm2. At 1100 �C and 1200 �C,the mass gain is moderate, with higher values for the monolithic samplethan for the composites, particularly for Al2O3-10. Nevertheless, the massgain is below 0.5 %/mm2 and 0.25 %/mm2 for Cr2AlC and Al2O3-10,respectively. However, at 1300 �C, the mass gain increased rapidly for allsamples, therefore the oxidation was stopped after 24 h. As a result, 1200�C seems to be the maximum operating temperature for Cr2AlC materialsfor applications operating for at least several days. These results are inagreement with other publications [11,51].

Representative polished cross sections of the post-mortem samples at1000 �C (50 h) and 1200 �C (68 h) are shown in Fig. 9 and Fig. 10,respectively. A dense Al2O3 layer with an even and uniform surface and athickness of approximately 5 μm forms without fibers at 1000 �C. Un-derneath this alumina layer, a Cr7C3 layer with a comparable thicknessforms as reported in the literature [14,19]. Incorporation of fibers alteredthe formation of the alumina layer. With SiC fibers, a very thin (~0,5-1μm) continuous alumina layer formed during oxidation. Fibers located atthe surface of the sample could still be observed though this thin layer(Fig. 9c)). In addition, the reaction zone between the fibers and thematrix increased. This reaction is thus an ongoing process when the

8

material is exposed to high temperatures. In addition, this reaction can bea reason for the thin oxide layer, as the aluminum that is weakly bound inthe matrix is not available for oxidation. In contrast, Al2O3 fibers did notreact with the matrix material (Fig. 9f)). With incorporated alumina fi-bers, a very thin uniform layer formed with some bigger clusters of Al2O3(Fig. 9e)). These clusters are 6 μm–7 μm in thickness. Some Cr7C3 alsoformed, but not as a continuous layer. The thicknesses of these Cr7C3clusters corresponds to the thickness of the Al2O3 on top, regardlesswhether it is the continuous layer or the cluster. The incorporated fibersmight therefore hinder the formation of a continuous Cr7C3 layer. Inmonolithic material, the element with the weakest bond, aluminum,diffused to the surface and formed an alumina layer with oxygen [6,12,13]. Due to the decrease of aluminum, a Cr7C3 layer formed conse-quently. Here, with alumina fibers, the amount of aluminumwith a weakbond is lower and not as well distributed than in the monolithic material.Regions with greater amounts (Cr2AlC) and lesser amounts (Al2O3 fiber)of aluminum are distributed in this kind of CMC. As a result, the for-mation of the Cr7C3 layer is inhomogeneous.

After oxidation at 1200 �C for 68 h, the monolithic sample formed anexternal dense but uneven alumina layer. This layer was thicker (5–10

Fig. 10. Oxidation response of a) and b) monolithic; c) and d) SiC-10; and e) and f) Al2O3-10 after TG at 1200 �C with an isothermal holding time of 68 h.

T. Go et al. Open Ceramics 6 (2021) 100090

μm) than the layer formed at 1000 �C and presented waves (see Fig. 10a)surface view and b) cross section). This result is in good agreement withother oxidation tests of Cr2AlC reported in the literature [6,10,51]. Here,spalling of the alumina layer started at 1200 �C. As expected, a Cr7C3layer with a comparable thickness (8–14 μm) formed underneath this

Fig. 11. mass gain over isothermal holding time

9

alumina layer, as reported in the literature [14,19,51]. Badie et al.observed a phenomena of rumpling at their oxidized Ti2AlC MAX phasesamples [52] and proposed a model to estimate deformation andrumpling of the oxide layer as a consequence of stresses in that layer.Thus, the roughness of the sample surface has a strong influence on the

for a) 1000 �C, b) 1100 �C, and c) 1200 �C

Fig. 12. X-ray diffraction pattern of the surface of monolithic Cr2AlC material and CMCs before and after TG at 1200 �C for 68 h.

Table 2Calculation of parabolic and cubic oxidation rates with corresponding R2 values.

T and Sample kp [kg2 m�4 s�1] R2 of kp kc [kg3 m�6 s�1] R2 of kc

1300 �C, Cr2AlC 5.58E-04 0.99732 3.42E-03 0.960461300 �C, Al2O3 -10 8.36E-04 0.92844 5.93E-03 0.825311300 �C, Al2O3 -20 2.98E-04 0.98912 1.33E-03 0.933841300 �C, SiC-10 1.29E-04 0.99671 3.77E-04 0.959191200 �C, Cr2AlC 4.76E-05 0.98871 1.76E-04 0.998511200 �C, Al2O3 -10 1.75E-05 0.99859 3.62E-05 0.967521200 �C, Al2O3 -20 1.20E-05 0.99773 2.12E-05 0.988721200 �C, SiC-10 7.33E-06 0.99892 9.89E-06 0.982671100 �C, Cr2AlC 4.88E-06 0.99020 5.15E-06 0.946531100 �C, Al2O3 -10 8.33E-06 0.97888 4.08E-06 0.885681100 �C, Al2O3 -20 2.09E-06 0.98427 2.29E-06 0.941651100 �C, SiC �10 1.74E-07 0.90650 3.66E-08 0.832101000 �C, Cr2AlC 7.86E-07 0.99463 1.97E-07 0.969171000 �C, Al2O3 -10 2.02E-06 0.96862 7.92E-07 0.991461000 �C, Al2O3 -20 8.25E-07 0.98798 2.11E-07 0.983081000 �C, SiC �10 4.46E-08 0.82341 2.69E-09 0.78023

T. Go et al. Open Ceramics 6 (2021) 100090

rumpling, as it caused higher grow stresses in the oxide layer, promotingrumpling and finally interfacial decohesion. In the case of rumpling, theprobability increases that the oxide layer gets cracks, whereby the pro-tective effect of this oxide layer is reduced. SiC-10 formed a continuousalumina layer with a thickness of 6–9 μm, well attached to the substrate.In addition, the reaction zone between the fibers and the matrixincreased (Fig. 10d)) as compared to the sintered sample and the sampleunder 1000 �C oxidation. Thus, this reaction, which was first observedduring sintering, is facilitated at high temperatures. A possible furtherdevelopment under these oxidation conditions is the complete reaction ofthe Cr2AlC MAX phase with SiC fibers into the phases found in Table 1.

Al2O3-10 behaved in a similar way to the monolithic Cr2AlC, at leastin terms of alumina scale adhesion. Furthermore, Al2O3-10 presented asurface similar to that of the wrinkled surface of the Cr2AlC sample(Fig. 10e)). Cr7C3 was also formed under the alumina scale, although inthis case it formed a discontinuous layer and the content is lower(Fig. 10f)). The alumina fibers themselves did not react with the matrixmaterial, but the alumina grains grew during oxidation at 1200 �C,resulting in an uneven appearance. As the continuous use temperature ofthese fibers (Nextel 610) is 1000 �C, this kind of grain growth is expected.So far, grain growth has not influenced the oxidation response of Al2O3-10.

In the following section, the measurement data of the TG experimentsare compared at a constant temperature for the different compositions(Fig. 11), excluding the values at 800 �C due to the low weight gain(<0.0125 %/mm2). Interestingly, SiC-10 exhibited the lowest mass gainat every temperature. Compared to pure Cr2AlC, Al2O3-10 showed ahigher mass gain at the lower temperatures (Fig. 11a) and b)). The higherthe oxidation temperature, the greater the advantage of alumina fibers inthe composite; the mass gain was clearly reduced compared to mono-lithic Cr2AlC at 1200 �C, not only for Al2O3-10 even for Al2O3-20. Here,the behavior of the oxidation curves changes; there is a higher oxidationof monolithic Cr2AlC than of Al2O3-10. Through the incorporation ofalumina fibers, the amount of Cr2AlC at the surface decreased, lessaluminum could diffuse to the surface, and thus less Al2O3 is formed.

10

However, this alumina still formed a wavy but continuous layer, asshown in Fig. 10f). It is not only the fibers that might change theoxidation response of a material, but also the secondary phases, whichare present before oxidation (see Table 1).

The different phase compositions before and after oxidation at 1200�C for 68 h of monolithic Cr2AlC, SiC-10, and Al2O3-10 are shown inFig. 12. As expected, alumina and Cr7C3 were detected on the surface ofthe monolithic material after exposure to air at a high temperature.Additionally, some reflexes of the MAX phase Cr2AlC appeared. Cr2AlCreflexes could appear when the formed alumina layer is thinner than themeasuring depth of the XRD or the alumina layer shows some partialspallation. There was a clear difference between the oxidation behaviorof the monolithic material and SiC-10. Here, different phases were visibleafter the TG experiments (red XRD pattern, Fig. 12), but no Cr2AlC can beseen. This is an indication of an ongoing reaction or internal oxidation.The patterns after TG experiments of the monolithic Cr2AlC material and

Fig. 13. Formation of oxide layer and reaction of fibers after 500 cycles in an oxidation at burner rig achieving a surface temperature of 1200 �C for a) SiC-10 and b)Al2O3-10.

T. Go et al. Open Ceramics 6 (2021) 100090

Al2O3-10 are very similar. The two main phases are Al2O3 and Cr7C3,which was caused by the formation of the outer layers described above.Accordingly, the incorporation of alumina fibers had no influence on thesurface of Cr2AlC after oxidation at 1200 �C. Overall, the presence ofsecondary phases in the Cr2AlC MAX phase material changed theoxidation response, but the exact nature of this has yet to be examined. Inparticular, the presence of Cr7C3 might alter the oxidation response, sincethese particles prevent the formation of a continuous alumina layer.

The oxidation kinetics of each sample and temperature were calcu-lated, including the parabolic (kp) and cubic (kc) rates (Table 2) with thecorresponding R2 values; k value with the highest R2 value is highlightedin bold. The oxidation rates fit better for parabolic behavior than for acubic response for most of the calculations. Nevertheless, both R2 values(for kp and kc) are rather similar, and as such it cannot be confirmed thatthese Cr2AlC materials follow a specific single trend. This might berelated to the content of secondary phases, which alters the oxidationbehavior of the Cr2AlC compounds [6,18,53]. Nevertheless, in this study,the oxidation rates are higher than those reported in the literature. Forexample, Lee et al. found a parabolic oxidation rate for Cr2AlC at 1100 �Cof 7.5 � 10–11 kg2m-4s-1 [14] compared to 4.88 � 10-6 kg2m-4s-1 foundhere. Lin et al. reported a parabolic rate of 1.08 � 10–12 kg2m-4s-1 at 800�C and 2.96 � 10-9 kg2m-4s-1 at 1300 �C [10] while here, 5.58 � 10-4

kg2m-4s-1 is calculated at 1300 �C. The actual phase composition of theMAX-phase material plays a paramount role, as it is not 100% pureCr2AlC. At the used temperatures, the formation of SiO2 should befavored, but with the used analyzing techniques (XRD, SEM, EDX) itcould not be detected. Thus, there is no SiO2 in the material or theamount of it is below the detection limit. It is unlikely that the amount ofAl2O3 (2%) has any influence, since Al2O3 was formed during oxidation.However, AlCr2 (3%) and/or Cr7C3 (1%) facilitated oxidation as well asthe SiC fibers. Cr7C3 and oxygen formed CO2 and chromium oxidefollowing this reaction:

4 Cr7C3(s) þ 33 O2(g) → 14 Cr2O3(s) þ 12 CO2(g) (4)

The formation of a gas phase could lead to pores and channels in thematerial. With a bigger surface area due to the walls of these defects, theoxidation of Cr2AlC, and thus the whole material/composite, might befaster.

The oxidation response of the two composites was tested under morerealistic thermal conditions using a burner rig. These experiments wereperformed in a previous study by most of the authors of the current paperon a similar Cr2AlC monolithic material and at a surface temperature of1200 �C [19]. In that study, no visible damage to the sample wasobserved after 500 thermal cycles, and the external and protectivealumina layer with a thickness of around 7 μm was well adhered to the

11

sample. The porous Cr7C3 layer underneath showed a similar thicknessthan the oxide layer The good response during thermal cycling, includingno cracks and no delamination at the interphase of the phases, is causedby the good match in the coefficient of thermal expansion (CTE) ofCr2AlC (11.0⋅10�6 K�1), Al2O3 (8.5⋅10�6 K�1), and Cr7C3 (10⋅10�6 K�1),and the convoluted interface. Encouraging results by Gibson et al.showed that a Cr2AlC sample treated under same thermal conditions didnot undergo mechanical properties variation afterwards [20]. Aftermicro-cantilever fracture tests at the interfaces between the differentlayers (Al2O3, Cr7C3 and Cr2AlC), the fracture toughness remained below4 MPa m1/2, which are typically associated with ceramic materials andthus ceramic properties. Fig. 13 shows the polished cross sections ofSiC-10 and Al2O3-10 after 500 thermal cycles at a surface temperature of1200 �C using the burner rig. For both composites, a dense, continuousand well-adherent alumina layer was formed in situ with a thickness ofabout 5 μm. These results are similar to those reported for Cr2AlC underthe same conditions. For SiC-10, a strong reaction occurred between thefibers and the Cr2AlC matrix. EDX measurements (not shown here)exhibited a large content of Cr7C3 underneath the alumina layer. As seenpreviously, this combination of materials is not suitable for long-termhigh-temperature applications. However, the results of the oxidationtest of Al2O3-10 are promising. In Fig. 13b), a well-adhered alumina layerwith a Cr7C3 layer underneath can be observed in the cross section. Evenafter 500 cycles at 1200 �C, the Al2O3 fibers did not react with the matrixmaterial. Fibers close to the surface of the CMC did not hinder the for-mation of a continuous alumina layer, which is required for protectionfrom oxidation. In comparison with the results of the TG experiments,there is no visible sign of grain growth after the burner rig experimentand the alumina layer is not wrinkled, which leads to the conclusion thatthe temperature inside the sample did not reach 1200 �C, as it wasmeasured by a pyrometer at the surface of the sample. In summary, theoxidation response of Al2O3-10 in a realistic oxidizing environment wasgood and promising. A protective, dense, continuous, and well-adheredalumina layer formed and the fibers survived this test without anyvisible damage, grain growth, reactions, and/or cracks.

4. Conclusion

Cr2AlC MAX-phase powder was successfully synthesized and subse-quently mixed with three different kinds of ceramic fibers, added inamount from 5 to 20 wt%. Carbon fibers resulted not suitable for thereinforcement of Cr2AlC because of their reaction with the matrix ma-terial during sintering at 1200 �C. SiC fibers also reacted with the matrix,but not completely. Al2O3 fibers did not exhibit any reaction with theMAX-phase matrix.

T. Go et al. Open Ceramics 6 (2021) 100090

The formation of the oxide layer during TG experiments was influ-enced by the type of fiber. The incorporation of Al2O3 fibers led to adiscontinuous, wrinkled alumina layer. The response of the materialduring the TG experiments was not uniform. Each CMC materialexhibited the highest mass gain at 1300 �C and the lowest mass gain at800 �C. Regarding the type of fiber, SiC-10 displayed the lowest massgain and when the oxidation temperature exceeded 1200 �C, also Al2O3-20 displayed reduced the mass gain as compared to pure Cr2AlC. Overall,the oxidation rates were higher than those reported in the literature,possibly due to small traces of secondary phases in thematerial that couldpromote internal oxidation, such as Cr7C3. The most important result ofthis study is the very good oxidation response of Al2O3-10 in a realisticoxidizing environment. A protective, well-adhered alumina layer formedupon 500 cycles at 1200 �C and the fibers survived this test without anydamage or visible grain growth. Underneath the alumina layer a Cr7C3layer formed. Overall, the results of these experiment are very promisingfor the further development of MAX-phase material with ceramic fibers.

Declaration of interest

The authors declare that they have no known competing financialinterests or personal relationships that could have appeared to influencethe work reported in this paper.

Acknowledgment

This study has been funded by Germany’s Federal Ministry of Edu-cation and Research (“Bundesministerium für Bildung und Forschung”)as part of the MAXCOM project (03SF0534). The authors would like tothank Dr. Doris Sebold for her valuable assistance during SEM analysis,Marie-Theres Gerhards for her assistance during TG measurements, Dr.Yoo Jung Sohn for her support with the Rietveld refinement analyses, Dr.Daniel Mack and Martin Tandler for performing the burner rig experi-ments, and Sylvain Badie for his support in manufacturing the magnetitecoating.

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