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ARTICLE IN PRESS+ModelECS-9587; No. of Pages 12

Available online at www.sciencedirect.com

ScienceDirect

Journal of the European Ceramic Society xxx (2014) xxx–xxx

Laser-supported joining of SiC-fiber/SiCN ceramic matrix compositesfabricated by precursor infiltration

Marion Herrmann a,∗, Katrin Schönfeld b, Hagen Klemm b, Wolfgang Lippmann a,Antonio Hurtado a, Alexander Michaelis b

a Technische Universität Dresden, Institute of Power Engineering, Chair of Hydrogen and Nuclear Engineering, D-01062 Dresden, Germanyb Fraunhofer Institute for Ceramic Technologies and Systems (IKTS), D-01277 Dresden, Germany

Received 5 December 2013; received in revised form 12 March 2014; accepted 16 March 2014

bstract

iC-fiber/SiCN ceramic matrix composites were manufactured by means of polymer infiltration and pyrolysis. The fiber preform was made bylurry infiltration and winding using a computer-controlled winding module. Multiple infiltration steps using a Si–C–N precursor were includedo increase the density. The influence of the sintering conditions on the microstructure of the CMC was demonstrated.

Pipe sections made of the CMC materials were joined using a laser-supported heating technology with an Y–Al–Si–O glass–ceramic filler. Thehermal response of the CMC components was controlled by the anisotropic thermal conductivity. Fast heating by laser beam was achieved forlements rotating in the direction of the fiber winding. SEM micrographs of the joints showed the good wettability of the CMC by the glass–ceramicller. Nearly defect-free joints were obtained using a nitrogen process atmosphere. The laser-supported technology was shown to be promising for

he joining of CMC components. 2014 Elsevier Ltd. All rights reserved.

eywords: CMC; SiC fiber; PIP; Laser joining; Glass–ceramic filler

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. Introduction

Rising needs for high-temperature components in the energyndustry cannot be met with conventional metallic material-ased solutions alone. Ceramic matrix composites (CMCs) have

high potential for applications in ground-based and automo-ive gas turbine components such as combustors, liners, turbineanes, and blades as well as aerospace engines and other indus-rial applications such as heat exchangers, hot gas filters, andadiant burners.1

CMCs can be divided into oxide and nonoxide compos-

Please cite this article in press as: Herrmann M, et al. Laser-supported jprecursor infiltration. J Eur Ceram Soc (2014), http://dx.doi.org/10.1016/

tes. In oxygen-containing atmospheres at high temperatures,xide composites exhibit high stability but inadequate mechan-cal properties (e.g., strength and creep resistance) for structural

∗ Corresponding author. Tel.: +49 351 463 32371; fax: +49 351 463 37161.E-mail address: [email protected] (M. Herrmann).

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ttp://dx.doi.org/10.1016/j.jeurceramsoc.2014.03.016955-2219/© 2014 Elsevier Ltd. All rights reserved.

pplications. In comparison, nonoxide composites are moreusceptible to oxygen environments at elevated temperaturesut exhibit considerably better mechanical properties at highemperatures.2

Nonoxide fiber composites based on SiC fibers were used inhis study. CMCs can be manufactured using a variety of pro-esses, including melt infiltration, chemical vapor infiltration,nd polymer infiltration and pyrolysis (PIP).3 Because of theow complexity of the required facilities and the low costs, theIP process was used in this work.

However, the processes available only allow fiber-reinforcedomponents of limited size and with simple geometries toe produced. Suitable joining technologies needed for fabri-ation of complexly shaped CMC components are currentlyeing developed, with the main emphasis being placed on tech-

oining of SiC-fiber/SiCN ceramic matrix composites fabricated byj.jeurceramsoc.2014.03.016

ologies using special materials for the joining of SiCf/SiC astructural materials for fusion power reactors and for the Veryigh Temperature Reactors (VHTR) within the framework of

dx.doi.org/10.1016/j.jeurceramsoc.2014.03.016

http://www.sciencedirect.com/science/journal/09552219


mailto:[email protected]


ARTICLE IN PRESS+ModelJECS-9587; No. of Pages 12

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M. Herrmann et al. / Journal of the Eu

he international Gen-IV program.4,5 The joining materials forpplication at high temperatures can be divided into

reactive materials, preceramic polymers, and

glass–ceramic fillers.

Examples of reactive materials are technologies using solid-tate displacement reactions for the formation of joints. Fornstance, mixtures of TiC and Si react to form an epitaxial layerf secondary SiC at the interface to the SiC matrix, therebyncreasing the joint shear strength.6,7

Different authors described the application of preceramicolymers for the joining of SiCf/SiC.6,8,9 The intrinsic characterf the joint was found to be advantageous; however, the pyrol-sis process itself led to material shrinkage and gas formation.hese effects could be partially reduced through the addition of

nert filler materials.Katoh et al.10 and Ferraris et al.11 proposed the use of

lass–ceramic fillers for joining. They identified the advan-ages as being the possibility of tailoring the properties andhe chemical resistance of the glass–ceramics. The mostrequently investigated glass–ceramic filler systems for SiCaterials were Y2O3–Al2O3–SiO2,12 MgO–Al2O3–SiO2,11,13

nd CaO–Al2O3.10,14,15 The joints exhibited superior thermo-echanical stability when thermally stable crystalline phasesere formed. The coefficients of thermal expansion (CTEs) of

he crystalline phases should be adjusted to the CTE of the matrixaterial for a high joint quality to be achieved. The residual glass

hase shows self-sealing behavior if the glass–ceramic is appliedt temperatures higher than the glass transition temperature.10

s a consequence of their high glass transition temperatures,lasses in the Al2O3–Y2O3–SiO2 system play an important rolemong glass–ceramic fillers for the joining of SiC.16 These com-ositions were also extensively investigated in connection withiquid phase sintering of the nonoxide ceramic materials SiC andi3N4.17,18

NITE (nano-infiltration and transient eutectic) sintering tech-ology was shown to be promising for joining CMCs.19

his technology requires temperatures of at least 1700 ◦C andressures between 10 MPa and 20 MPa. The composition ofhe joining zone (mix of Y2O3, Al2O3, SiO2, and �-SiC nano-owder) was adjusted to be close to that of the matrix material.20

he joints exhibited very high thermal and mechanical stabil-ties, especially under conditions of irradiation (investigated at

6 dpa at 800 ◦C).5 For this reason, the technology was rec-mmended for joining of SiCf/SiC components under neutronrradiation.

With all of these joining methods, the components are heatednd joined, in some cases under high pressure, in a furnace.n contrast, the laser-supported technology enables heat inputo be localized. Thus, the dimensions of the components to be


oined are not restricted through the dimensions of the heat-ng furnace.21 Superior results in joining of monolithic SiComponents with laser-supported heating were obtained in pre-ious studies.21–23 A glass–ceramic filler composition in the

swii

n Ceramic Society xxx (2014) xxx–xxx

2O3–Al2O3–SiO2 system was used for these experiments.hort processing times and localized heating of componentsere the main advantages of the laser brazing method. Fur-

hermore, the high thermal conductivity and the low thermalxpansion coefficient of SiC were found to be beneficial foriminishing temperature gradients and transients introduced byhe laser energy input. In comparison with other technologies,he laser-supported joining technology proved to be very energy-fficient.

The goal of the current study was to develop a laser-supportedechnology for joining CMC components. This involved twoifferent tasks:

investigation of the thermal response of the CMC componentsto laser heating and

testing of glass–ceramic fillers for joining of the CMC mate-rial.

Focus was on the development of materials with propertiesuitable for application under corrosive conditions at tem-eratures of above 1000 ◦C and the implementation of thesedvanced materials in components and functional prototypes fornergy technology (e.g., heat pipes).

. Materials and methods

.1. CMC fabrication process

Polycrystalline SiC fibers of type “Tyranno SA3” (UBEndustries Ltd.) were used to produce the fiber preforms. Aftereing desized at 800 ◦C the fiber roving was wound onto differ-ntly sized mandrels at a winding angle of 85◦ according to austomized winding program. During this winding process theoving was infiltrated with a ceramic slurry comprising a SiC ori3N4 powder (SNE-3, SNE-10, UBE Industries), a binder, and

sintering aid. The powder loading in the slurry was 20 wt%.After the preform was dried PIP was performed using

he commercially available polysilazane Si–C–N precursorHTT1800” (Clariant Advanced Material GmbH). Pyrolysisas subsequently performed at 900 ◦C in argon. Conversion

rom the liquid precursor to the amorphous SiCN was accompa-ied by shrinkage of the matrix, which led to formation of poresnd cracks in the matrix. A defined fiber spacing was obtainedhrough the presence of powder particles between the individualbers introduced during slurry infiltration in the winding pro-ess. Several reinfiltration and pyrolysis steps were performedn order to minimize the porosity of the matrix. At least a sinter-ng process was performed in a gas pressure sintering furnaceetween 1500 ◦C and 1700 ◦C in a nitrogen atmosphere. Thereparation route used is shown in Fig. 1.

The components for the laser joining experiments underwenthree infiltration/pyrolysis cycles and were sintered at 1700 ◦C.wo sample configurations were prepared. Fig. 2a shows pipe


ection samples of outer diameter 30 mm, length 20 mm, andall thickness 1 mm for the heating and simple joining exper-

ments. The softening and wetting behavior of the filler wasnvestigated on these samples; the joints were realized as butt



M. Herrmann et al. / Journal of the European Ceramic Society xxx (2014) xxx–xxx 3

Fig. 1. Preparation route for SiC fiber-reinforced CMC.

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The magnitude of the laser power was adjusted according tothe filler softening temperature. The joining process was stopped

Fig. 2. CMC test samples for joining experiment

oints. Microstructural analysis of the joints was performed onolished cross sections prepared from these samples. Planaramples were fabricated for preparation of bending bars. Theest sample geometry (Fig. 2b) was 5 mm × 5 mm × 25 mm,hereby the joint had an overlap (5 mm × 5 mm) to ensure a

ufficient joining surface.

.2. Filler materials

A glass with the composition 16.55% Y2O3, 26.05% Al2O3,nd 57.40% SiO2 (mol%) was used as the filler. Glass synthe-is and characterization were described by Rohbeck et al.24 Thelass transition temperature was approx. 890 ◦C24; the workingemperature was in the range of 1420 ◦C. The suitability of thisller for laser-supported joining of silicon carbide was demon-trated for pressureless-sintered SiC (SSiC) by the authors.23

he filler was ground to a grain size of less 10 �m, mixed withthanol, and placed on the surfaces to be joined. For somexperiments, SiC powder with a grain size of 0.1–1 �m (�-iC, Goodfellow) was added in the amount of 10 vol% to thelass–ceramic filler.

.3. Laser-supported joining technology

A diode laser (DL 031 Q, ROFIN) with a power of 3100 Was used for the joining tests; both wavelengths (808 ± 10 nm

nd 940 ± 10 nm) were used simultaneously. The laser beam was


uided via a scanner (powerSCAN 33, Scanlab) with a lens ofocal width 400 mm so that the laser energy could be distributedptimally over the sample surface. The test setup is shown inig. 3.

e sections (a) and test bars for bending tests (b).

The laser process was performed in a special box thatemained open for the experiments in ambient air. For the exper-ments in nitrogen, the box was closed, and the laser beam wasassed through a quartz glass window.

The temperature was measured during the joining process byeans of an infrared camera (VarioCAM hr head, InfraTec). A

rotective filter was positioned in front of the camera to absorbhe scattered laser radiation and prevent artifacts in the measure-

ents. The overall transparency of the filter (0.88) and that of thenSe window in front of the vacuum chamber (0.9) were taken

nto account during temperature measurement. The emissivityf the CMC material was assumed to be 0.86. The temperatureas taken to be the average value yielded in a 2-mm-diameter


Fig. 3. Experimental setup for the laser-supported joining process.



4 M. Herrmann et al. / Journal of the European Ceramic Society xxx (2014) xxx–xxx

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The thermophysical properties of the CMC material are

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Fig. 4. Test setup for the four-point bending test.

fter 70 s through abrupt cutoff of the laser power. Subsequentooling was realized by free convection in the atmosphere usedor heating. During the tests the samples were rotated at a ratef two revolutions per second. The volume reduction duringoftening of the filler was compensated for through applicationf an axial contact pressure of approx. 0.5 MPa.

.4. Characterization

During the fabrication of SiC fiber/SiCN composites theorosity of the material was detected after each preparationtep according to Archimedes’ principle. Qualitative analysisf crystalline phases was performed by XRD (D8, Bruker).he mechanical properties were evaluated by means of circu-

ar ring tests with an INSTRON 8562 according to DIN EN95-3. Four-point bending tests were also performed on sam-


les of dimensions 3 mm × 4 mm × 50 mm according to DINN 843-1.

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ig. 5. Microstructures of SiC/Si3N4 composites after 3 pyrolysis steps (a and c) and

ig. 6. Open porosities after different numbers of pyrolysis steps and sinteringrocesses.

Four-point bending tests were performed5 mm × 5 mm × 50 mm) according to DIN EN 843-1 toield information about joint strength. The positioning of theoined samples during the test is shown in Fig. 4. For compari-on purposes, tests on CMC bars with the same configurationut without joints were conducted. The support and load spansere 40 mm and 20 mm, respectively. A 10-kN load cell and aelocity of 0.5 mm s−1 were used for the measurements.


mportant for the laser-supported joining process. The coeffi-ient of thermal expansion and the thermal conductivity were

after sintering at 1600 ◦C (b and d) perpendicular and parallel to the fiber axis.




Fig. 7. Matrix structures of SiC/Si3N4 composites af

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wofcracks were healed and a homogeneous matrix with both matrix

ig. 8. X-ray diffraction patterns of SiC/Si3N4 composites after sintering atifferent temperatures.

valuated in directions perpendicular and parallel to the fiberxis. The coefficient of thermal expansion was determinedccording to DIN 51045 and DIN EN 821-1 with a NetzschiL 402C. The thermal conductivity of the CMC material wasetermined using the laser flash method according to DIN EN


21-2 with a Netzsch LFA 427.The microstructures of the composite material and the joints

ere analyzed by means of ceramographically prepared sections

e(c

Fig. 9. Circular ring test (a) and four-point bendin

ter sintering at 1700 ◦C (a) and at 1800 ◦C (b).

n the SEM (Leica S 260). Elemental distribution analysis waserformed using energy-dispersive X-ray spectroscopy (ISIS SiLi) detector).

. Results and discussion

.1. Structures and properties of the CMCs

The open porosity of the pyrolyzed material decreased withncreasing number of infiltration cycles. After five infiltrationycles an open porosity of 6% was reached (versus 22% afterwo pyrolysis steps). These data were comparable with thosebtained by PIP of 0/90◦ UD cross-ply layers of SiC/SiCNomposites.25 The last sequential weight gain after five pyrolysisteps was about 2%. A comparable manufacturing technologysing BlackglasTM as a precursor26–28 in the PIP process neededeven cycles to reach a weight gain of 2% with similar porosity.

After pyrolysis the fibers were embedded in a glassy matrixith dispersed Si3N4 particles. A lattice of small cracks devel-ped as a result of precursor shrinkage during conversionrom the liquid to the amorphous phase. During sintering the


lements, i.e., precursor and powder, dissolved was formedFig. 5). During the conversion from amorphous to crystallineeramic the mass loss of the precursor was about 16 wt%.

g test (b) results for SiC/Si3N4 composites.




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sallel and perpendicular to the fibers. The resulting CTE wasmainly determined by the matrix material and was in the rangeof that of monolithic Si3N4.3

Table 1Thermal properties of SiCf/Si3N4.

α [10−6 K−1] λ [W m−1 K−1]

SiCf/Si3N4, perpendicular 50–1400 ◦C 2.5–5.2 3.1–2.2SiCf/Si3N4, parallel 50–1400 ◦C 2.5–5.2 18.3–10.3

Fig. 10. Fracture surfaces after circular ring test – path o

he overall ceramic yield of the precursor was about 48%. Aetailed description of the precursor conversion is given in theiterature.29

A sintering aid was added to the Si3N4 powder in the slurry tonable liquid-phase sintering. The sintering aid reacted with theiO2 on the powder surface to form a silicate phase, which wasolten at the sintering temperature.30 The liquid phase promoted

ensification, accelerating mass transport through solution andeprecipitation processes.31 Due to the stable fiber framework,hich did not change over the course of sintering, only theatrix material (powder and precursor) densified, thereby cre-

ting additional voids. The extent of matrix densification andhe accompanying pore formation were found to depend on theintering temperature. Fig. 6 shows the increase in open porosityith increasing sintering temperature.The matrix structure and crystalline phases were affected by

intering conditions such as atmosphere, pressure, and temper-ture. Experiments were performed at a nitrogen pressure of.1–1.0 MPa and temperatures between 1500 ◦C and 1800 ◦C.t temperatures up to 1650 ◦C, the conversion of the precursoras incomplete. At 1700 ◦C, a leaf-like matrix structure with-Si3N4 as the main crystalline phase was obtained. At higher

emperatures, a needle-like microstructure was observed; thisas the result of reactions that occurred in the gas phase30:

iO2(l) + 2SiC(s) + 2N2(g) ↔ Si3N4(s) + 2CO(g) (1)

SiO(g) + 2N2(g) ↔ Si3N4(s) + 3SiO2(l) (2)

The deposited Si3N4 exhibited a needle-like structureFig. 7); �-Si3N4 was found to be the main crystalline phase.he conversion from �-Si3N4 to �-Si3N4 crystals is mapped inig. 8. SiC was detected through the SiC fibers at all tempera-

ures. �-Si3N4 occurred as a second crystalline phase at 1500 ◦C.he phase transition from �- to �-Si3N4 started at 1600 ◦C, andoth phases existed in parallel until 1700 ◦C, when only �-Si3N4ccurred and the transition was finished. Above this temperature,o crystal conversion occurred, as indicated by the similarities

◦ ◦

Please cite this article in press as: Herrmann M, et al. Laser-supported jprecursor infiltration. J Eur Ceram Soc (2014), http://dx.doi.org/10.1016/j

etween the diffraction patterns at 1700 C and at 1800 C.A circular ring test was performed according to DIN EN 295-

for the characterization of component strength.32 A notablenfluence of the matrix structure on the strength of the material

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k after four-point bending test (a) and fracture path (b).

as observed. Due to the needle-like structure and interlockingf the �-Si3N4 grains, the strength was much higher (Fig. 9a).

In addition to the circular ring test, a four-point bendingest was performed. The samples for this test were prepared byarm pressing of cut wound laminates. As reported previously

n the literature,33 the composites exhibited damage toleranceFig. 9b). The crack was deflected along the fiber axis at eachber bundle (Fig. 10a). The fracture surfaces clearly showed

hat the crack path was deflected by the fibers and fiber bundles.ecause of the weak fiber–matrix interface, some fibers were

eparated from the matrix in the wake of the crack and fiberullout was observed (Fig. 10b).

The results of thermophysical characterization are listed inable 1. The properties of pressureless-sintered SiC (SSiC) are

ncluded for the sake of comparison.The thermal conductivity was found to be strongly depend-

nt on the measurement direction. Parallel to the fiber axis,he free path of phonons was not interrupted. Perpendicularo the fiber axis, the phonon oscillation was interrupted forvery fiber–matrix crossover. The nominal thermal conductiv-ty of the Tyranno fibers was about 65 W/mK at 20 ◦C.35 The

uch lower thermal conductivity of the composite in compar-son with that of the monolith ceramic reflected the high poreontent in the matrix of the composite. The thermal conductivityas lower than that of SiC/SiC composites35,37 due to the lower

onductivity of the matrix material.The thermal expansion (CTE) of SiCf/Si3N4 materials

howed similar, only slightly temperature-dependent values par-

oining of SiC-fiber/SiCN ceramic matrix composites fabricated by.jeurceramsoc.2014.03.016

i3N4, reaction-bonded34 10–1000 ◦C 2.1–3.0 4–15iC fiber Tyranno SA35 20 ◦C 65SiC36 20–400 ◦C 3.5 120



ropean Ceramic Society xxx (2014) xxx–xxx 7

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Fig. 11. IR thermogram of the CMC specimen at different stages of the lasert

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.2. Thermal response of CMCs to heating with a diodeaser beam

Sintered silicon carbide (SSiC) absorbs the laser wavelengths08 nm and 940 nm on the surface.38 The further heating behav-or is dependent on the thermal conductivity of the material.ence, the material properties dictate the duration of the heatinghase and the homogeneity of the temperature distribution overhe joining region. The investigated CMC material consistedf SiC fibers and SiC in the matrix material. Hence, similarbsorption behavior was expected. However, because the ther-al conductivity of the CMC material was lower than that ofSiC, homogeneous heating was expected to be more time-onsuming. Furthermore, the effect of the anisotropy of thehermal conductivity on the thermal response to laser radiationas unknown.Fig. 11 shows the CMC components to be joined at differ-

nt times during the process; in the initial (untreated) state, athe beginning of the process, after 2, 3, and 4 s of laser beamreatment, and at the end of the heating process (after 60 s).t is notable that the laser beam was absorbed by the singleiC fibers first, i.e., before homogeneous heating of the mate-ial commenced. The anisotropic thermal conductivity of theMC affected the heating process; parallel to the fiber axis,

he temperatures increased very rapidly, whereas perpendic-lar to the fiber axis, the heating process was clearly moreluggish. Heating of the bulk resulted from the thermal con-uctivity in both directions. Approximately 60 s were needed toeat the components to 1400 ◦C. This short heating period washe result of the relatively small wall thickness of about 1 mmnd should increase with more compact samples with thickeralls. The experiments showed that the thermal response to laser

reatment was determined by the absorption of the laser beamy the SiC fibers as well as by the thermal conductivity deter-ined by the fiber orientation.

.3. Process parameters

Defect-free heating of the joining components to the filleroftening point and generation of homogeneous thermal fieldsn the joining zone were found to be the crucial steps in theaser-supported process. Both were adjusted by variation of thearameters “laser power” and “duration of laser beam treat-ent.” Fig. 12 shows the correlation between adjacent laser

eam power and temperature on the sample surface. When erro-eous measurements were excluded, the temperature appearedo increase linearly with increasing laser beam power. Hence, itan be concluded that the heating of the samples up to 1420 ◦Cequired a laser energy of approx. 140 W cm−2.

One advantage of the laser-supported technology is the shortrocess time, which, however, can induce high thermal gradi-nts and transients. These gradients and transients should beinimized for a high joint quality. Homogeneous temperatures


n the surfaces to be joined were obtained by rotation of theomponents during the laser process.

Fig. 13a shows the spot on the CMC component at whichhe temperature was measured and the temperature profile as a

gao1

reatment process.

unction of time for the duration of the laser treatment process.bsorption of the laser beam on the surface led to rapid heat-

ng. After 5 s the temperature reached 1000 ◦C, and after 20 she temperature reached 1420 ◦C, sufficient for softening of theller. The dwell time was accordingly adjusted to ensure homo-eneous temperature distribution over the joining region as wells homogeneous softening and distribution of the filler. A time


f 40 s was chosen for this process step to accommodate the-mm wall thickness.



8 M. Herrmann et al. / Journal of the Europea

F

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jwmatrix–filler interfaces. The fast cooling process after the laser

ig. 12. Correlation between adjacent laser energy and surface temperature.

Fig. 13b presents the line scan used for temperature mea-urement along the z-axis of the laser-heated component afterhermal treatment for 60 s. The regions in which the tempera-ure reached 1000 ◦C on both sides of the joint were not widerhan 5 mm. This reflected the dimensions of the laser beam onhe sample surface. The temperature decreased linearly from the

iddle to the area outside of the laser beam and finally to theample holder as a result of the low thermal conductivity of theaterial.An important criterion for the quality of the joints was the

rocess atmosphere. A temperature of at least 1420 ◦C wasequired for softening of the glass–ceramic filler. At this temper-ture, oxidation of the CMC occurred. The effect of the processtmosphere should be investigated in the future with the aim ofinimizing possible oxidation processes.Bubbles formed during laser heating in ambient air

Fig. 14a). SiC reacted with the oxygen in the air to formaseous products according to reaction (3). This reaction wasnhanced with increasing temperatures and led to the formation


f significant quantities of gas at temperatures above 1400 ◦C.t these temperatures, the viscosity of the glassy filler was low

nough to allow the bubbles to escape from the material surface.

tat

Fig. 13. IR temperature measurement during the laser process: surface tempe


he viscosity of the glassy filler increased as the temperatureecreased after the laser was switched off. This resulted inntrapment of gas bubbles in the solidified glass–ceramic,hich had a detrimental effect on the final joint quality.

iC(s) + O2(g) → SiO(g) + CO(g) (3)

The formation of gaseous species according to reaction (3)an be suppressed through use of an inert atmosphere for therocess. Fig. 14b shows the microstructure of a joint producedn a nitrogen atmosphere. Bubbles did not form under these con-itions. The CMC components were heated up to temperaturesequired for joining of the components using the glass–ceramicller composition described here without significant defect for-ation.

.4. Microstructures of the joints

The good wetting behavior of the glass–ceramic filler on theSiC surfaces and the formation of gastight and mechanicalroof joints were described by Herrmann et al.23 However, theood wettability was connected with the existence of a SiO2ayer on the material surface. Because a continuous SiO2 layerould not exist in the porous CMC matrix, simple transfer of theSiC wetting behavior to that of CMC was not feasible.

The positioning of the filler on the CMC material differedecause of the differences in porosity of the CMC samplesFig. 15). The small wall thickness and the single butt joint geom-try led to a relatively low contact surface between the two partso be joined, necessitating accurate positioning during the join-ng process. Additionally, because the filler had a low viscosity athe process temperatures, it flowed into the CMC matrix voidsnd significantly complicated the process of forming uniformoints.

A cross section of a typical microstructure of a CMC sampleoined with the glass–ceramic filler is given in Fig. 16. The filleret the CMC matrix very well, resulting in mostly well-formed


reatment led to the solidification of the filler in a glassy state. At higher magnification (see Fig. 17a), several light crystals lesshan 1 �m in size could be discerned. EDX analysis of these

ratures over time (a) and temperature distribution along the z-axis (b).




Fig. 14. Cross sections through the surfaces of CMC test samples coated with yttrium aluminosilicate filler after laser treatment at approx. 1420 ◦C in air (a) and innitrogen (b).

MC p

cs

odg

bft

Fig. 15. Joining surfaces on the C

rystals yielded the elements Y and Si, suggestive of yttriumilicates.

A joint fabricated using the glass–ceramic filler with addition


f SiC particles is shown in Fig. 17b. The added SiC particlesid not influence the wettability of the CMC by the filler. Theap between the two CMC parts was filled, and the interface

Fig. 16. Cross section through a joined CMC sample.

SSfichbo

cfivowsiglshtc

ipe sections (with filler in place).

etween the CMC and the filler was also found to be nearlyree of defects. The SiC grains were homogeneously distributedhroughout the seam (Fig. 17b). In contrast to the filler withoutiC particles, this filler showed no crystallization. Addition ofiC particles increased the viscosity of the filler, leading to lessller flow into the pores of the CMC material. The higher thermalonductivity of the SiC particles in the filler could lead to moreomogeneous temperatures in the filler, but this thesis could note proven through the experiments conducted within the scopef this study.

The aim of the fiber reinforcement in ceramic materials is toounteract the inherent brittleness of ceramics. Glass–ceramicllers are also brittle, but their chemical compositions can bearied to tailor their properties. Baron et al. studied the effectf addition of SiC grains to Y–Al–Si–O–N glasses.39 The effectas twofold: both the fracture toughness and the mechanical

trength increased. The SiC grains can affect the laser join-ng process by increasing the absorption of the filler if theirrain size is larger than the laser wavelengths used.40 In theaser-supported joining process, the homogeneous heating of theubstrate material is important. This leads to filler softening and


omogeneous adhesion to the substrate surface. Higher absorp-ion of the filler is undesirable, and hence a SiC powder mostlyonsisting of grains with sizes smaller than the laser wavelength




t with

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Fig. 17. Joint with yttrium aluminosilicate filler (a) and join

i.e., <1 �m) was chosen for the experiments described in theresent paper.

Because of the high joint quality achieved with them,2O3–Al2O3–SiO2 glass–ceramic compositions with SiC pow-er additions should be investigated further in the future.

.5. Mechanical strengths of the joints

The potential of the filler composition in the2O3–Al2O3–SiO2 system for laser-supported joining ofonolithic SiC ceramic (SSiC) materials was demonstrated in

ecent investigations in which a four-point bending strengthf 122 MPa with a Weibull modulus of 5.6 was obtained.21

he importance of the parallelism and absolute flatness ofhe two surfaces to be joined and a small seam thicknessor joint strength was emphasized. Because CMC materialsxhibit rough surfaces, parallelism and flatness of the joiningurfaces and thin joints are difficult to achieve. A maximumn the mechanical strength of joined SiCf/SiC componentsas achieved with a maximum overlap in the joining area.11

his effect was utilized in the samples for mechanical testing


Fig. 18a). Unfortunately, the edges in the joining area couldot be made to be fully rectangular, and hence the two partsid not fit together perfectly. As a result, the joining zone onlyovered an area of about 5 mm × 5 mm.

cmdo

Fig. 18. Joined sample for the bending test (a)

yttrium aluminosilicate filler containing SiC particles (b).

In contrast to the components rotated during laser processing,he components for mechanical testing were fixed under the laseream. The structure in the bulk of the material was comparableith the structure perpendicular to the fiber axis. Hence, the

hermal conductivities were expected to be similar, and longereating cycles were expected. Double-sided heating was alsomplemented. Despite these two adjustments, softening of theller at a depth of 2.5 mm from the component surface was onlychieved with higher laser powers, which led to higher thermalradients in the fixed samples than in the rotating ones.

On joined samples, four-point bending strength was mea-ured at intervals of less than 10 MPa; however, the samplesever fractured inside the joining zones. Typically, fractureccurred in the CMC material parallel to the joint (Fig. 18b). Theeasured strength was insufficient and could not be explained

y the joining procedure. For comparison purposes, samples ofeometry 5 mm × 5 mm × 50 mm and with notches similar tohe gaps in the joined samples were prepared. Four-point bendingtrength was also determined for these samples. For all speci-ens, the measured strength had the same order of magnitude

s that of the joined test bars (<10 MPa). Under the given test


onditions, the mechanical strength was determined by the CMCaterial and not by the joint. Apparently, the applied stress led to

elamination of the CMC material caused by the lower strengthf the matrix material between the SiC fibers.

and cross section of the joint seam (b).



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Although the joint in Fig. 18b was not responsible forailure under mechanical loading conditions, it was inhomo-eneous and included a large number of bubbles. This effectould be attributed to the higher process temperatures neces-ary for softening the filler in the bulk. Apparently, reactionsther than reaction (3) could also facilitate bubble formation.ickel and Quirmbach41 and Jacobson42 reported gas formation

s a result of the reaction between SiC and SiO2, indepen-ently of the atmosphere (reaction (4)). The start of this reactionas specified as lying in the range between 1470 ◦C41 and527 ◦C,42 making it possible that it occurred in the present laserrocess.

iC(s) + SiO2(s) → 3SiO(g) + CO(g) (4)

Stationary samples were less advantageous for homogeneouseat input into the bulk of the CMC components than rotatingamples were and led to the development of thermal gradients.

The results of microstructural characterization presented inection 3.4 showed completely filled joints and good wettingehavior of the filler. After optimization a defect-free inter-ace between the filler and the matrix was achieved; hence, theechanical strength of CMC joints was expected to be com-

arable with that of monolithic SSiC joints. The measured lowour-point bending strength of the joined CMC was caused byhe low mechanical strength of the CMC.

. Conclusions

SiC fiber-reinforced CMCs were prepared using the PIProcess. This technology enables control of process param-ters such as porosity, microstructure, and strength of theMC material. After parameter optimization the CMC mate-

ial exhibited a porosity of approx. 30%. The thermal treatmentetermined the structure of the resultant material. Conversion ofhe SiCN precursor to the amorphous state started in the pyrol-sis stage. In a second step, conversion from the amorphous tohe crystalline state occurred, first with formation of �-Si3N4nd then with formation of �-Si3N4 at higher temperatures.ith optimized temperatures, needle-like �-Si3N4 formed.MCs with these structures yielded strengths in the range of50 MPa in circular ring tests performed according to DIN EN95-3.

The presented results demonstrated the possibility ofoining SiC fiber-reinforced CMCs using a laser-supportedechnology with a glass–ceramic filler composition in the

2O3–Al2O3–SiO2 system. However, processing in ambient aired to bubble formation, making it necessary to perform the pro-ess in an inert atmosphere. This, in turn, limited the possibleizes of the components to be joined because of the necessity ofoining in a closed chamber. Changing the process atmosphereas also time-consuming.The applied SiC fiber-reinforced materials exhibited mul-


iphase structures and anisotropic properties. Laser-supportedeating was influenced by the differences in absorption of theaser wavelength and the anisotropic thermal conductivity. The

aterial had the highest thermal conductivity in the direction of

1

n Ceramic Society xxx (2014) xxx–xxx 11

he fiber axis. Thus, the fiber direction in the CMC componentss an important factor to be considered in the optimization of theeating process.

Use of the laser-supported technology for joining CMC mate-ials with high-melting-point glass–ceramic fillers has certainimitations, but the properties of the joints produced are similaro those of conventionally produced joints. Hence, the presentedesults are important for the further development of the laser-upported technology.

cknowledgments

This work was performed within the scope of the projectCeramic components based on nonoxide fiber compositeaterials for advanced systems in high-temperature energy tech-

ology” and as part of the work of the Dresdner Innovationenter Energy Efficiency. The project was funded by the Euro-ean Regional Development Fund (ERDF) and the Free State ofaxony (Sächsische Aufbaubank), Grant no. SAB 14258/2423.

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