Ma

Ma

3b

a

ARR2A

KFFCSC

1

nhoiiac[vcnbaScath

m

0d

Materials Science and Engineering A 527 (2010) 1314–1338

Contents lists available at ScienceDirect

Materials Science and Engineering A

journa l homepage: www.e lsev ier .com/ locate /msea

echanical properties of silicon nitride-based ceramics and its use in structuralpplications at high temperatures

.H. Bocanegra-Bernala,∗, B. Matovicb

Centro de Investigación en Materiales Avanzados, CIMAV S.C. Departamento de Física de Materiales, Miguel de Cervantes # 120 Complejo Industrial Chihuahua,1109 Chihuahua, Chihuahua, MexicoVinca Institute of Nuclear Sciences, Materials Science Laboratory, Belgrado, Serbia

r t i c l e i n f o

rticle history:eceived 14 August 2009eceived in revised form9 September 2009ccepted 30 September 2009

a b s t r a c t

Silicon nitride (Si3N4) based ceramics are gaining more and more attention due to their promising high-temperature thermal and mechanical properties. They have been expected to be the main candidatesfor applications such as turbocharger rotors and gas turbine engine components which can withstandsevere conditions of temperature and heavy loads. Although a big number of studies on silicon nitrideare published, a continuous progress in monolithic Si3N4 as well as Si3N4/Si3N4 composites (seeded

eywords:racture toughnesslexural strengthreep

materials) leads to new scientific and technological data providing new insight that should be reviewedtaking into account their excellent properties at high temperatures. Silicon nitride possesses a bunchvariety of interesting properties that can be specifically designed to produce a given behavior profile.That is why the room temperature and high-temperature properties are discussed and described in moredetail.

uperplasticityavitation

. Introduction

Considering their unique combination of properties, siliconitride is an excellent material for applications where are requiredigh strength at elevated temperatures, abrasion resistance, goodxidation resistance and relatively high toughness. Silicon nitrides probably the most thoroughly characterized non-oxide ceram-cs with wide applications including heat exchangers, turbine andutomotive engine components subjected to high-temperatureonditions, engine valves and other wear-resistant components1–12]. Although the market potential for these applications isery high; there are still several difficulties that must be over-ome before the full potential of structural ceramics based siliconitride is realized. Si3N4 is a basically covalent compound and itsulk diffusion is too low to be consolidated [12–18]. Therefore, theddition of some oxides or nitride sintering additives to the purei3N4 is required in order to fabricate high-density silicon nitrideeramics by different sintering routes [2–6,12–16]. However, these

dditives remain as grain boundary glassy phase, which deterioratehe high-temperature properties of the ceramics such as creep andigh-temperature strength [19,20].

∗ Corresponding author. Tel.: +52 614 439 4801; fax: +52 614 439 4823.E-mail addresses: miguel.bocanegra@cimav.edu.mx (M.H. Bocanegra-Bernal),

ato@vin.bg.ac.yu (B. Matovic).

921-5093/$ – see front matter © 2009 Elsevier B.V. All rights reserved.oi:10.1016/j.msea.2009.09.064

© 2009 Elsevier B.V. All rights reserved.

For designers and manufacturers of advanced gas turbine enginecomponents, silicon nitride is an attractive material for many rea-sons. These materials typically have potentially high temperature ofusefulness (up to 1400 ◦C approximately) [21]. To take advantage ofthe inherent properties of Si3N4, the high-temperature mechanicalperformance of the material must first be characterized. The appli-cation of this excellent material as structural components requiresa more rigorous reliability theory [22] because the mechanicalresponse of silicon nitride to static, dynamic and cyclic conditionsat elevated temperatures, along with reliable and representativedata, is a critical information that gas turbine designers and manu-facturers are required for the confident insertion of silicon nitridecomponents into gas turbine engines.

Today, silicon nitride ceramics can be regarded as a classof material comparable to steel. The different qualities dependon size and shape of the silicon nitride grains and the amountand chemistry of the grain boundary phase [23,24]. For example,high strength Si3N4 ceramics exhibit a fined-grained elongatedmicrostructure, Si3N4 with high fracture toughness are morecoarse grained [25]. In both cases a weak interface is requiredin order to provide a transgranular fracture mode [26]. There-fore, silicon nitride ceramics are prime candidates for such diverse

high-temperature applications items as rotors and stator vanes foradvanced gas turbines, valves and cam roller followers for gasolineand diesel engines [27,28]. On the other hand, the most importantmilestones achieved include the incorporation of silicon nitrideceramic as (i) hot section components produced by AlliedSignal

Scien

foicbctmpwsm

dfatSepptaarsgotw

2

rtnVrac[ta

2

tFhtTmsstpsaAShr1b

M.H. Bocanegra-Bernal, B. Matovic / Materials

or aircraft and industrial auxiliary turbo-power unit and vari-us components for aircraft turbine engines [29] and (ii) valvesn selected automotive diesel engines by Daimler–Benz. Recenteramic gas turbine programs at Roll Royce Allison and Solar tur-ines [30,31] have done much to increase the experience baseoncerning to behavior of ceramic components in industrial gasurbines. A key lesson learned in both programs is that environ-

ental effects may severely limit the long-term reliability anderformance of silicon-based structural components [30]. Evenith these successes, ceramics that have greater toughness and

trength will be necessary to meet demands of potential futurearkets for advanced ceramics [28].The properties of silicon nitride materials can vary widely,

epending on the starting materials and on the method used fororming and sintering of the green bodies [11,12,24,27,32–35]s well as the heat treatment developed in order to crystallizehe intergranular glass phases [32]. Today, commercially availablei3N4 powders are prepared by means of various routes describedlsewhere [36]. Therefore, the realization of a material with goodroperties can be achieved through the control of factors such ashase composition of the starting powders, type and amount of sin-ering aids, processing parameters and microstructure (grain sizend morphology) [4,37–39]. It is well known that the room temper-ture properties of Si3N4 ceramics are mainly determined by aspectatio and grain size of the �-phase, and that the high-temperaturetrength is controlled particularly by the characteristics of therain boundary phase [39–42]. Some of the more important resultsbtained in different investigations concerning to the fractureoughness, flexural strength, and creep of silicon nitride ceramicsill be briefly examinated.

. Fracture toughness

Fracture toughness values are used to characterize the fractureesistance of ceramics and brittle materials. Sample preparation isime consuming and expensive and methods based on single edgeotched beam (SENB), single edge precracked beam (SEPB), edge-notched beam (SEVNB), chevron-notched beam (CNB) are cur-ently used and require precise notch geometry control. However,method that is frequently used to determine fracture toughness oferamic materials is the so-called indentation fracture (IF) method43]. Fracture toughness of silicon nitride ceramics has been inves-igated by numerous researchers considering different methodsnd a summary of their results is as follows.

.1. Indentation fracture (IF)

In sintered Si3N4 along with oxide additives, the fractureoughness is known to be dependent on its grain morphology [44].or example, a higher fracture toughness is obtained by using aigher �-phase-content Si3N4 powder as a raw material in ordero develop the fibrous microstructure during sintering [23,45,46].he increase in toughness could be explained by the bimodalicrostructure of large elongated �-Si3N4 and fine-grained matrix

intered silicon nitride [12]. Liu et al. [39] reported interestingtudies on the influence of ball-milling methods on the microstruc-ure and mechanical properties of Si3N4 ceramics prepared byressureless sintering with an additive from the MgO–Al2O3–SiO2ystem. These authors found promising mechanical propertiesnd homogeneous microstructures with high-energy ball-milling.

composition of Si3N4 with 3 wt% MgO + 1.5 wt% Al2O3 + 5 wt%

iO2 was blended by conventional ball-milling and planetaryigh-energy ball-milling and subsequently sintered in a graphiteesistance furnace under atmosphere pressure of nitrogen at780 ◦C for 1.5 and 3 h. The values of fracture toughness obtainedy the authors for these Si3N4 ceramics prepared at different

ce and Engineering A 527 (2010) 1314–1338 1315

conditions and measured by the indentation fracture method were5.8 ± 0.6, 6.2 ± 0.5, 6.4 ± 0.5 and 6.3 ± 0.5 MPa m1/2 for samplessintered at 1750 ◦C-1.5 h-P*, 1780 ◦C-1.5 h-P*, 1780 ◦C-3 h-P* and1780 ◦C-3 h-G*, respectively, being P* planetary high-energy ball-milling and G* general ball-milling. Si3N4 ceramics with fracturetoughness of 6.4 ± 0.5 MPa m1/2 were prepared by pressurelesssintering with planetary ball-milling process that is comparableto that achieved by previous hot-press or gas pressure sintering[17,47]. According to the SEM observations obtained by Liu et al.[39], the differences in the fracture toughness of silicon nitrideceramics produced by different processings could be due to thechanges in the homogenization of the grain size and their distri-bution throughout the bodies. It is suggested that the improvedhomogeneity of the sintering additives is responsible for theimprovement of the mechanical properties.

Studies by Ling and Yang [48] in silicon nitride doped with5 wt% MgO, 5 wt% MgO + (1, 2, 3, 4, 5, and 6 wt% Y2O3) and sinteredat temperatures of 1500, 1600, 1700 and 1800 ◦C during 60 minshowed a fracture toughness value (measured by the indentationfracture method) of 7.5 MPa m1/2 which was correspondent to thehigh strength and highest relative density for the composition ofSi3N4 + 5 wt% MgO + 4 wt% Y2O3 (sintered at 1700 ◦C during 60 min)which indicated that the combination of magnesia and yttria sinter-ing aids is very effective [32,48–51]. On the other hand, neither MgOnor Y2O3 nor other crystalline phase were detected by X-ray diffrac-tion patterns suggesting that the experimented additives mightreact with the SiO2 on silicon nitride particles to form silicate liq-uid phase and glassy phase after cooling of ceramic. Elongated Si3N4grains well interconnected were obtained in the microstructure.

Zheng et al. [34] have reported similar studies but self-reinforcing silicon nitride with additions of either yttrium oxideor ytterbium oxide at room temperature after various processingheat treatments. They found a number of toughening mecha-nisms, including crack deflection, bridging, and fiber-like grainpull-out. Pressureless sintered (at 1750 ◦C under nitrogen pressureof 0.1 MPa) and hot-pressed (at 1650 ◦C under a pressure of 25 MPa)silicon nitride with additions of SiO2:Y2O3:Al2O3 (molar ratios85:8:4:3) and SiO2:Yb2O3 (molar ratios 85:10:5) were investigated.Appropriate additions of Y2O3 and Al2O3 densify more readily atlower sintering temperatures than ceramics produced with onlyone of these two oxide additives [52]. The fracture toughness datawere collected by both the indentation fracture (IF) method andindentation strength (IS) method. The authors determined thatsuch techniques can be applied to relatively small ceramic sam-ples and can help to determine the toughness of materials forcracks of the order of grain size in these Si3N4 ceramics. It wasobserved in this study that the fracture toughnesses calculatedfrom indentation fracture method increase as the indentation loadused to initiate cracks increased. Zheng et al. [34] revealed that theset of samples more sensitive to indentation load were the pres-sureless sintered ytterbium oxide-doped silicon nitride ceramics(Yb2O3–Si3N4), in which the measured value of fracture toughnessrose from 6.3 to 9.1 MPa m1/2 with indentation loads of 98 and588 N, respectively. On the other hand, value of fracture tough-ness obtained by indentation strength (IS) method for the sameindentation load used to generate crack lengths for the indentationfracture method is consistently higher. However, it is important topoint out that since indentation techniques are necessarily a sim-ple way of estimating absolute toughness, the agreement betweenthe results of both IF and IS techniques is satisfactory [34,53]. Inthe microstructural observations the elongated grains in all sets

of samples are encouraging, providing evidence for the relevanttoughening processes such as deflection and crack bridging. Withbase of this, it is possible to produce ytterbium doping siliconnitride ceramics which retain high values of fracture toughnesswithout compromising room temperature strength.

1 s Scien

tptfltiitasfii[wfcmlstdttsw

(nlpFatvcAlMtnsp

[mftcctsats

dctbnfcot

316 M.H. Bocanegra-Bernal, B. Matovic / Material

Ytterbium oxide (Yb2O3) has been found to be effective as a sin-ering aid of silicon nitride ceramics in improving the mechanicalroperties. Lee et al. [54] documented results on gas pressure sin-ering of Si3N4 with additions between 2 and 16 wt% Yb2O3 andound typical microstructures for in situ toughened Si3N4 includingarge elongated grains randomly distributed in a fine matrix. Fromheir results, it is noteworthy that when Si3N4 samples contain-ng Yb2O3 as a sintering aid were gas pressure sintered at 1950 ◦Cn 4 MPa of N2 gas, the resulting microstructures were not uniformhroughout the sample with large grains near the surface while rel-tively small grains at the center. These microstructural evolutionsuggested that the amount of additive had a strong influence on thenal microstructure. The occurrence of extensive grain growth only

n the outer region in the Si3N4–Yb2O3 system studied by Lee et al.54] implied that the conditions for abnormal grains growth [55,56]ere satisfied only in that region (�–� phase transformation). The

ast growth of �-nuclei formed by dissolution and posterior repre-ipitation is a direct cause of the formation of in situ tougheningicrostructure [57]. According to this, around 1500 ◦C �-grains

arger than average grains must have been formed in the wholeample. On the other side, the increase in weight loss observed withhe amount of sintering aid Yb2O3 was thought to be due to theepletion of Yb-containing species from the specimen. Moreover,he concentration of Yb in the dense outer region decreased withhe sintering time while that in the inner region remained about theame. Therefore, it is believed that the Yb-containing liquid phaseas depleted from the sample surface by evaporation.

Balazsi et al. [58] reported that multiwall carbon nanotubesMWNTs) may serve as crystallization sites and seeds for siliconitride grain growth. They elucidated in their SEM studies web-

ike connected MWNTs in the structure but individual MWNTs,roperty attached to the �-Si3N4 surfaces can also be observed.ully dense samples with improved mechanical properties werechieved at relatively lower sintering temperatures and shorterime by SPS. 5.2 and 5.3 MPa m1/2 were the fracture toughnessalues obtained by using the indentation fracture method in theompositions Si3N4 + 4 wt% Al2O3 + 6 wt% Y2O3 and Si3N4 + 4 wt%l2O3 + 6 wt% Y2O3 + 1 wt% MWNT, respectively. During complex

iquid phase sintering process, Balaszi et al. [59] demonstrated thatWNTs serve as ideal crystallization sites for �-Si3N4 grains. On

he other side, incorporated MWNTs in the middle of the siliconitride grains [60] evidently serve as seeds for crystal growth pre-erving the carbon nanotubes in the composite structure. Similarhenomenon was reported by Cai et al. [61].

Additions of SiC wisker, ZrO2 and TiN particles as second phase62–65] and self-reinforcement by microstructural control of Si3N4

atrix, have been two typical approaches in order to improve theracture toughness of Si3N4 ceramics. Kim et al. [66] have reportedhat a new approach to obtain the bimodal microstructure of sili-on nitride ceramics is due to the decreasing of oxygen content byarbothermal reduction treatment (CRT). The compositions inves-igated were Si3N4 with additions of 6 wt% Y2O3 + 1 wt% Al2O3 asintering additives. For this case, 0 and 0.5% carbon powders weredded to make the changing of the crystal structure of triple junc-ions in the sintered bodies by CRT. The samples were gas pressureintered at 1850 ◦C for 6 h under 2 MPa nitrogen gas pressure.

From results obtained by Kim et al. [66] it is deduced that theecreasing of oxygen contents at the triple junctions promotes therystallization of junction regions and therefore, the microstruc-ural change in these regions can give an effect on the fractureehavior of gas pressure sintered Si3N4 ceramics. Fracture tough-

ess values of 5.1 and 6.6 MPa m1/2 were obtained by indentation

racture (IF) method for the two tested samples without and witharbothermal reduction, respectively. Considering the observationf microstructures and crack propagation, Lee et al. [65] confirmedhat even though the bimodal microstructure was formed easily

ce and Engineering A 527 (2010) 1314–1338

by carbon addition, the fracture toughness was not so remarkablydue to the existence of Y3AlSi2O7N2 crystalline at junction regionsof Si3N4 grains, while a thin amorphous phase was observed withabout 2 nm in thickness at Si3N4 grain boundaries and interfacesbetween the Si3N4 and Y3AlSi2O7N2 phase. For improving the frac-ture toughness of gas pressure sintered Si3N4 ceramics, all the grainboundaries and junction regions must be kept with an amorphousphase as well as the formation of perfect bimodal microstructure[65].

As outlined above, the addition of TiN particles as second phaseis one of the typical approaches to improve the fracture tough-ness of Si3N4 ceramics. The formation of nanostructures has beeninvestigated for Si3N4–TiN by Hojo et al. [67,68]. Si3N4–TiN com-posite was also produced by in situ process from amorphous Si3N4powder and TiO2-containing sintering aid [69]. The inclusion ofTiN was effective for improving the fracture toughness [70]. Thisauthor studied the composite Si3N4–TiN prepared by a vapor phasereaction from SiCl4–TiCl4–NH3–H2 system. The resulting siliconnitride powders were amorphous whereas TiN was detected byXRD. Besides this, the Si3N4 particles included black dots of about10 nm which were assigned to be TiN by analytical TEM observa-tions. As sintering aids were used Y2O3 (6 wt%) and Al2O3 (2 wt%)and the composites were hot-pressed at 1800 ◦C in N2 atmosphereduring 2 h. Si3N4–TiN composites showed a duplex microstructureconsisting of fine grains and rod-like grains wherein the growth ofSi3N4 was remarkable in Si3N4–TiN system. These results suggestthat the Si3N4 grain growth was enhanced simultaneously withcrystallization of �-phase by TiN inclusion. The fracture toughnessof Si3N4–TiN composite was measured by the Vickers indenta-tion method and Hojo [70] evaluated the effect of TiN inclusionon the fracture toughness revealing that the fracture toughnessincreased by the addition of TiN (6.5 MPa m1/2 for 7.5 vol.% TiN con-tent) but was saturated at TiN content of approximately 18 vol.%.This increase in fracture toughness is attributed to inclusion of TiNparticles as the second phase and rod-like grain growth of Si3N4stimulated by TiN. The saturation of fracture toughness may be dueto the retardation of rod-like grain growth at the large TiN content.This is also supported by the result in the in situ Si3N4–TiN system[71,72].

The most previous work on Si3N4-based ceramics employedY2O3 and Al2O3 as the sintering aids [73,74]. Alternative sinteringadditives have been found to result in very different properties suchas fracture toughness. Consequently, several alternative additiveswere studied in order to see if they provided any advantage overthe Y2O3–Al2O3 system in terms of final properties [75]. This inves-tigation used as the starting materials appropriate amounts of Si,�-Si3N4, Al2O3, MgO, MgAl2O4 and Y2O3 to give final compositionsafter nitriding. Tiegs et al. [75] selected different silicon powdersso that they represented a large variation in impurity content. TheY2O3–Al2O3 additive system was used as a base material becauseof the large database already established and it was also used totest the different silicon powders. Nonetheless, the MgO–Y2O3 andMgAl2O4–Y2O3 additives were also studied taking into account thatthey were reported to produce materials with very high fracturetoughness [76]. The nitridation was carried out with a N2–4% H2gas flow and controlled heating to 1450 ◦C and subsequently thesamples were sintered in a graphite-element furnace with Si3N4packing powder under one atmosphere N2 to 1780–1800 ◦C dur-ing 2 h. Fracture toughness was measured by means of indentationfracture (IF) method.

The fracture toughness values [75] were 6.6, 6.1, 5.7 and

6.1 MPa m1/2 for SRBSN from Elkem Metals Co., KemaNord Grade4C, KemaNord Grade 5C and Albemarle Corp. silicon powders,respectively, (additive composition was Si3N4 + 9 wt% Y2O3 + 3 wt%Al2O3). It is clear that the fracture toughness did not show any sig-nificant differences for the samples with different Si types. The

Scien

vi9(crfow

cInp(vto[atsowpttfc5gcaspiTf

malm[dniooopaA1(1ftftet2t

t

M.H. Bocanegra-Bernal, B. Matovic / Materials

alues obtained of fracture toughness for the alternate sinter-ng additives were 6.6, 4.8, 7.5 and 6.6 MPa m1/2 for SRBSN with% Y2O3 + 3% Al2O3 (10 wt% �-Si3N4 content), 9% Y2O3 + 3% Al2O30 wt% �-Si3N4 content), 6.4% Y2O3 + 3.2% MgO (10 wt% �-Si3N4ontent) and 5% Y2O3 + 5% MgAl2O4 (10 wt% �-Si3N4 content),espectively. As can be noted, high fracture toughness was observedor the Y2O3–MgO containing materials. Similarly, Tiegs et al. [75]bserved that large grains were obtained with this system whichould also contribute to the high toughness.

Tiegs et al. [77] reported investigations with sintered sili-on nitride powder compacts manufactured by St. Gobain/Nortonndustrial Ceramics (SG/NIC), sintered reaction bonded siliconitride (SRBSN) powder compacts fabricated by SG/NIC and SRBSNowder compacts elaborated by Oak Ridge National LaboratoryORNL). All samples were processed by microwave and con-entional sintering in order to identify differences in fractureoughness. The details of the powder processing of these typesf materials for the investigation have been reported previously78–81]. The nitridation was performed with N2–4% H2–5% Het ∼0.1 MPa (approximately 16 psi) with additional N2 added ashe reaction proceeded. Subsequently, the nitrided samples wereintered at 1800–1825 ◦C with entire heating cycle of 27 h. Tobtain fracture toughness values, indentation fracture (IF) methodas used with load up to 10 kg. The results of fracture toughnessresented several differences between microwave and conven-ional sintering process and Tiegs et al. [77] clearly concludedhat microwave processed materials exhibit significantly improvedracture toughness (ranging from 5.3 ± 0.2 to 8.8 ± 0.4 MPa m1/2)ompared to the conventionally sintered samples (values between.0 ± 0.3 and 7.7 ± 0.4 MPa m1/2) as a result of the increased elon-ated grain growth observed in the samples. Moreover, the samplesontaining high additive levels showed higher fracture toughnessesnd larger differences between the microwave and conventionalamples. This behavior is due to the larger amounts of liquid phaseresent in the samples with high additive contents, which facil-

tates the elongated grain growth associated with silicon nitride.herefore, the microwave sintering appears to be more appropriateor silicon nitride compositions at high additive levels.

In works developed years ago [82,83] have been developedethods to use microwave sintering for the fabrication of SSN

nd SRBSN, which have demonstrated that microwave heating canead to accelerated nitridation of silicon performs and to improve

echanical properties such as fracture toughness. Kiggans et al.27] investigated silicon nitride compacts fabricated from sprayried powder mixtures using standard die press and isopress tech-iques and silicon-based performs manufactured using gelcast and

sopress techniques with experimental details reported in previ-us publication [84]. The tested samples were SRB-1 (compositionf appropriate amounts of silicon (Elkem metallurgical grade) tobtain a final composition of Si3N4 + 9% Y2O3 + 3% Al2O3) and gasressure sintered composition named as AY6 of �-Si3N4, Al2O3nd Y2O3 to obtain a final composition of Si3N4 + 6% Y2O3 + 2%l2O3. Once the compositions were obtained, they were sintered at950 ◦C for 3 h in N2 gas at a pressure of 1.3 MPa. The dense GPSSNAY6) and SRBSN (SRB-1) samples were then annealed at 1150 or600 ◦C for either 5 or 20 h in a graphite-element or a microwaveurnace. More details of the experimental procedures for this inves-igation have been reported in other work [85]. For these samples,racture toughness was determined by means of indentation frac-ure (IF) method using 20 kg load and applying the Chantikulquation [86]. The conventional and the microwave annealing of

he SBR-1 sample at 1150 ◦C for 5 and 20 h caused approximately a0% reduction in fracture toughness (4.5 MPa m1/2) compared withhe obtained value in the untreated sample (5.6 MPa m1/2).

The annealing treatments at 1600 ◦C had little effect on the roomemperature fracture toughness. The reductions in fracture tough-

ce and Engineering A 527 (2010) 1314–1338 1317

ness when compared to the untreated controls can be caused by thedevelopment of thermal expansion coefficient (CTE) mismatchesbetween the host silicon nitride materials and the crystallinephases that form during the annealing treatments. On the otherside, no significant differences in the room temperature fracturetoughness values between the control sample and the SRB-1 mate-rials annealed at 1600 ◦C by conventional or microwave heating(approximately 8 MPa m1/2 for both cases) suggested that thesematerials were annealed above the glass melting point (∼1280 ◦C),where the crystallization is not favored. Similar behavior was foundin AY6 ceramics [27].

Some results have confirmed a relationship between siliconnitride grain size and fracture toughness using hot-pressed sili-con nitride as a model system. For example, Neil et al. [87] havereported investigations utilizing two base compositions of siliconnitride designed AY6 (Si3N4 + 6% Y2O3 + 1.5% Al2O3) which has beenused for initial AGT-5 rotor fabrication. The PY6 composition isidentical but without the Al2O3 additive. Results from this inves-tigation demonstrated that the fracture toughness of AY6 siliconnitride could be increased by increasing the average grain size ofthe material by extending the time allowed for grain growth tooccur during densification (longer hot-pressing times). Taking intoaccount that the fracture toughness can be defined as the ability ofa material to resist crack growth which, for silicon nitride, happenspredominantly through the grain boundary phase, by increasingthe grain size of silicon nitride, resistance to crack growth increasesbecause the crack is deflected more to propagate around the grains.However, exist an upper limit to increasing the fracture tough-ness of monolithic Si3N4 by grain growth and therefore, the aimof the present activities was to identify that upper limit of frac-ture toughness for the monolithic PY6 silicon nitride and then bringthe densified rotors to this upper level. Nonetheless, if the grainsbecome too large, less energy will be required to propagate a crackthrough a grain than around the grain which will limit the degreeof crack deflection and thus limit the fracture toughness [87].

Although the enhancement of fracture toughness was initiallyexamined by hot-pressing, this method is not viable for produc-ing rotors. However, Helms [88], Buljan et al. [89] and Hefter etal. [90] gathered results from the hot-pressing studies to provideinformation necessary to design experimental HIP cycles to evalu-ate the effects of different HIP conditions on the fracture toughnessof PY6 silicon nitride ceramic. Nowadays the knowledge gainedfrom the hot-pressing studies in understanding the microstructuraldevelopment during densification must be translated to complexshapes via hot isostatic pressing (HIP) cycle providing a greaterdegree of flexibility for producing a large-grained, high aspect ratiomicrostructure to increase fracture toughness [87]. The incompleteconversion to �-Si3N4 during densification with the silicon nitridestarting material, was a difficulty found with this experimentation.Based on these results, to produce complete conversion, the HIPcycle was modified by increasing both the densification time andtemperature. The fracture toughness measured by indentation frac-ture (IF) technique with the modified HIP cycle increased from 3.3to 4.6 MPa m1/2 with uniform and acicular microstructure with highaspect ratio �-Si3N4 grains instead of primarily equiaxed �-Si3N4grains.

Intention to increase further the reliability of Si3N4 has shiftedthe emphasis toward development of specific types of Si3N4/Si3N4composite ceramics. By incorporating a controlled amount of elon-gated �-Si3N4 single-crystal particles into the matrix (seeding) inan attempt to grow further the elongated uniformly distributed

grains in a matrix of equiaxed or slightly elongated grains [91–93],toughening mechanisms such as crack deflection and/or bridgingvia interfacial debonding, are activated [94] similar to those inwhisker-reinforced ceramics. It is now well accepted that seeding isa useful method of providing an effective way to improve the frac-

1318 M.H. Bocanegra-Bernal, B. Matovic / Materials Science and Engineering A 527 (2010) 1314–1338

ing t

tcp

pSaoapgawdFwbemsntwa

Fig. 1. SEM micrographs show

ure resistance while retaining high strength, provided that the size,ontent and distribution of the elongated �-Si3N4 single-crystalarticles are carefully controlled [95,96].

Silicon nitride seeds in beta form are usually obtained by gasressure sintering procedure at 1850 ◦C [97]. Preparation of �-i3N4 seeds under flowing nitrogen using powder bed method islso demonstrated [98]. Since, the additives are precondition forbtaining �-Si3N4 seeds, subsequent acid treatment to remove thedditives is necessary. There is no difference in �-Si3N4 seeds pre-ared from these two methods. Example of large elongated rod-likerains synthesized by pressureless sintering is shown in Fig. 1(a)nd (b). They were identified as �-Si3N4 single-crystal particlesith mean diameter of 2.22 �m and mean length of 5.43 �m. Theistribution of seeds in terms of their width and length is shown inig. 2. The lengths of seeds vary from 2.91 to 11.04 �m, while theiridth varies from 0.97 to 3.62 �m. The amount of the seeds has to

e moderate to avoid impingement of grains. As an example, theffect of seed particles concentration on relative density of speci-ens sintered with ceria as sintering additive for various times is

hown in Fig. 3. It should be pointed out that the green densities do

ot differ to a great extent for sample with different seed concentra-ion in the range studied. As can be inferred from Fig. 3, full densityas achieved in samples denoted as 0 (non-seeded), 3 (1 wt% seeds)

nd 4 (3 wt% seeds) after sintering at 1800 ◦C for 4 h [99]. This is

Fig. 2. Distribution of width and length of synthesized rod-like seeds [99].

he as-synthesized seeds [99].

significantly shorter sintering time than used in other works [98].The results also show that the particle rearrangement process pro-ceeds smoothly with non-seeded samples, in which particles areequiaxed with the size of less than 0.5 �m which was the startingpowder particle size as provided by the supplier. With increas-ing seeds concentration (0–3 wt%) sintered density decreases forshorter sintering times (up to 4 h) while for longer sintering times(4–6 h) sintered density kept on being constant independent ofseeds content within the studied range. Clearly, elongated seeds donot rearrange smoothly during consolidation process due to theirgeometry and are an obstacle to particles rearrangement process.That is why samples with increasing seeds concentration (1, 3 and5 wt% seeds) show density decrease. Densities are lower than forunseeded samples. It can be seen from Fig. 3, that seeding with5 wt% seed particles, caused considerable density decrease, for allsoaking times.

Seeded samples, as expected, exhibit a bimodal microstructurecomposed of small matrix grains and large rod-like grains approx-imately 2 �m in diameter and over 10 �m long. The large rod-likegrains which are fairly uniformly distributed throughout the sam-ple morphologically correspond to the grains developed from theseed particles during sintering. In comparison with non-seeded

sample, matrix grains in seeded samples seem to be smaller. Theaddition of seeds to samples sintered leads to an increase in fracturetoughness with an increase in seeds concentration of up to someamount and thereafter it decreases (Fig. 4).

Fig. 3. Effect of sintering time (at 1800 ◦C) and the amount of seed particles onrelative density of silicon nitride. Theoretical density was calculated from the ruleof mixtures [99].

M.H. Bocanegra-Bernal, B. Matovic / Materials Scien

Fo

stmodcc

agtoitigosp

isisiaA[aiatsnwftf8t

et al. [104].

ig. 4. Effect of sintering time and amount of seed particles on fracture toughnessf silicon nitride sintered at 1800 ◦C with CeO2 [99].

However, an increase in fracture toughness with an increase inintering time was observed for all samples. It is quite clear fromhe densification results in Fig. 3 and fracture toughness measure-

ent in Fig. 4 that the addition of seeds had the dominant effectn fracture toughness and causes its increase despite the fact thatensity remained unchanged. The addition of seeds provided theonditions for activating grain bridging and pulls out mechanismsapable of enhancing fracture toughness.

The decrease in fracture toughness in samples containing seedsbove approximately 3 wt% suggests that the growth of elongatedrains may be limited by their concentration. It seems that underhe present experimental conditions, there is an optimum amountf seeds beyond which the growth of elongated grains is lim-ted. This reasoning is in line with the kinetics of nucleation ofhe elongated grains. As the seed concentration (or their number)ncreases the number of nucleation sites also increases leading torain growth inhibition. The variation in fracture toughness valuesbserved in this investigation depends on the morphology of theilicon nitride grains as well as the character of the intergranularhase [95].

Good mechanical properties of silicon nitride ceramics can bemproved by controlling the amount of sintering additives andintering conditions. With these expectations, Si3N4 has beennvestigated to obtain ceramics for high-temperature applicationsuch as turbocharger rotors, turbine wheels and furthermore, sil-con nitride ceramics have also been successfully demonstrateds hot section components in auxiliary power units produced bylliedSignal for aircraft as well as components for turbine engines

100]. Takatori et al. [101] reported fracture toughness valuesround 7 MPa m1/2 measured by indentation fracture method in sil-con nitride ceramics with 5 wt% Y2O3 and 5 wt% MgAl2O4 (spinel)dditives for fabrication of radial wheels; Ozawa et al. [102] inves-igated the fabrication of ceramic radial turbine rotors using theilicon nitride identified as SN-84H and obtained a fracture tough-ess of 6.4 MN m−3/2 at room temperature; Kobayashi et al. [103]orked also with ceramic radial turbine rotors made of Si3N4 of dif-

erent compositions and labeled as SN-50, SN-81 and SN-84 (for the

hree compositions, the forming technique was injection moldingor the blades and isopressed for the hub). With SN-50 and SN-1 they obtained ceramics by pressureless sintering with fractureoughness values of 5.2 and 5.8 MN m−3/2 at room temperature,

ce and Engineering A 527 (2010) 1314–1338 1319

respectively, meanwhile with the material SN-84 was manufac-tured post-HIPed ceramics obtaining fracture toughness around6.4 MN m−3/2 at room temperature. From these results, it is clearthat post-HIPed Si3N4 offers about 22% higher fracture toughnesscompared with pressureless sintered SN-50 silicon nitride. Similarstudies have been developed by Kobayashi et al. [104] with compo-sitions named SN-84H and SN-88 with fracture toughness of 6 and7 MN m−3/2, respectively, at room temperature. These values wereobtained with post-HIPed in SN-84H and pressureless sintered inSN-88. Natansohn and Pasto [105] reported 4.2 MPa m1/2 as frac-ture toughness value in the composition Si3N4 with additions of6 w/o of Y2O3 (designed as PY6) which was injection molded anddensified by hot isostatic pressing (HIP) within the Advanced Tur-bine Technology Applications Project (ATTAP) program selected byGTE. It is important to stress that this composition and fabricationwere selected because they have been shown to yield silicon nitrideceramics with potential for meeting the challenging goals of thisprogram.

Hecht et al. [106] worked on two types of Si3N4 for heat engineapplications. Their results of fracture toughness were obtained bymeans of two different techniques such as the controlled flowmethod and the microindent method. For the first case, fracturetoughness (KIC) was determined from the relationship

KIC = 2.06�f

√c

�(1)

where �f is the flexure strength measured by three-point bend test,and c the radial crack extension of the microindent. On the otherside, the second case follows the following equation:

KIC =(

c

a

)−3/2( H

E�

)2/5 (H

√A

�

)(2)

where � = 3 (constraint factor), H is the measured hardness, a isone-half the indent diagonal, c is the radial crack extension, and Eis the elastic modulus. With the application of Eqs. (1) and (2), thefracture toughness values reported were 3.0 and 6.1 MPa m1/2 forthe silicon nitride GTE-PY6 with controlled flaw and microindentmethod, respectively. Likewise, for the Norton XL144 silicon nitrideceramic, the reported values for fracture toughness were 3.1 and3.8 MPa m1/2 applying controlled flaw and microindent techniques.From these results, the authors concluded that the fracture tough-ness values measured by the microindent technique tended to behigher than the toughness values measured by the controlled flawmethod. Nonetheless, almost all the values measured were lowerthan published or reported by the manufacturer.

In the use of silicon nitride ceramics for heat engine applications,it is very important to estimate the life time of gas turbine engines orceramic turbocharger rotors in operation considering some factorssuch as stress-temperature distributions in a component, Weibullstatistical theory of fracture as well as the properties of startingSi3N4. Matsui et al. [107] reported fracture toughness values of 6.3and 6.1 MN m−3/2 at room temperature and 5.7 and 5.4 MN m−3/2

at 1200 ◦C in two kinds of sintered silicon nitride (SSN) denomi-nated as SSN-A and SSN-B, respectively. Similarly, Miyauchi andKobayashi [29] studied the fabrication by pressureless sintering ofsilicon nitride radial turbine wheels 90 mm in diameter to be oper-ated at temperatures around 1250 ◦C. The element obtained wassubjected to hot spin test with a turbine inlet gas temperature ofapproximately 1000 ◦C. The fracture toughness reported at roomtemperature was about 5.2 MN m−3/2 which is low compared tovalues reported by Kawase et al. [108], Hecht et al. [106], Kobayashi

Nowadays, the development of joining technologies of ceramicwheels with metal shafts is very important for putting Si3N4ceramic rotors in practical use. Ito et al. [109] presented thebrazing technology of the ceramic rotors wherein reliable test-

1 s Scien

iiitausptsKKTitSgfwbsYg

2

YtaotsaYRitbu725Stsnimlbtten

2

bSbiflnm

320 M.H. Bocanegra-Bernal, B. Matovic / Material

ng of the bondings, using good ceramic materials, is of majormportance in the development of bonding technologies becauset is essential in evaluating any bonded body to identify whetherhe cause of failure was the material itself or its bonding. Theuthors concluded from their investigations that the materialsed for the turbine wheel in the bonding test was gas pres-ure sintered silicon nitride (GPSSN) which was sintered in overressure of nitrogen gas at high temperature obtaining a fractureoughness value of 6.5 MN m−3/2 at room temperature, which isuperior to the values reported by Miyauchi and Kobayashi [29],obayashi et al. [104], Kawase et al. [108], Matsui et al. [107],obayashi et al. [103], and Ozawa et al. [102] to mention some.iegs et al. [85] reported that by means of gas pressure sinter-ng (GPS) the fracture toughness was increased and improvedhe reliability in the compositions Si3N4 + 6% Y2O3 + 2% Al2O3 andi3N4 + Sr2La4Yb4(SiO4)6O2 (at 8 equiv.% oxygen), due to the elon-ated grain growth which leads to the materials to develop a highracture toughness [110]. The achieved fracture toughness valuesere 7.5 ± 0.3 to 8.4 ± 0.3 MPa m1/2 (final sintering temperatures

etween 1900 and 2000 ◦C) and 4.3 ± 0.2 to 7.5 ± 0.7 MPa m1/2 (finalintering temperatures between 1850 and 1900 ◦C) for Si3N4 + 6%2O3 + 2% Al2O3 and Si3N4 + Sr2La4Yb4(SiO4)6O2 (at 8 equiv.% oxy-en), respectively.

.2. Indentation strength in bending (ISB)

An interesting investigation concerning to the effect of2O3/SiO2 and R2O3(ss)/SiO2 as sintering aids on the fractureoughness of silicon nitride ceramics has been reported by Ribeirond Strecker [111]. R2O3(ss) is an oxide mixture in the formf a solid solution and can be used as a sintering additive inhe liquid phase sintering of silicon nitride ceramics as possibleubstitute of Y2O3 [112,113]. The compositions evaluated by theuthors were SNY14 (84.99 wt% Si3N4 + 5.21 wt% SiO2 + 9.79 wt%2O3) and SNTR14 (82.54 wt% Si3N4 + 5.04 wt% SiO2 + 12.42 wt%2O3(ss) which were sintered at 1800 ◦C during 30, 60 and 240 min

n a nitrogen atmosphere under pressure of 1.8 MPa. The frac-ure toughness was calculated by means of indentation strength inending (ISB) method described by Ribeiro [113]. The reached val-es of fracture toughness in the composition SNY14 were 7.3 ± 0.7,.1 ± 0.5 and 5.7 ± 0.1 MPa m1/2 for sintering during 30, 60 and40 min, respectively. On the other hand, 7.5 ± 0.6, 6.5 ± 0.3 and.8 ± 0.6 MPa m1/2 were the values measured for the compositionNTR14 for sintering time of 30, 60 and 240 min. It is clear thathe fracture toughness in the tested samples were similar, albeithowing decreasing values as the sintering time increased, a phe-omenon that can be explained by the grain morphology resulting

n the microstructure. As the sintering time increased, the grainorphology became more equiaxial, with a smaller aspect ratio,

eading to less pronounced effects of crack deflection and cracking-ridging mechanisms. Ribeiro and Strecker [111] concluded thathe fracture toughness was identical or better with R2O3(ss)/SiO2han with Y2O3/SiO2 additive. Table 1 summarizes a number ofxperimental results on the indentation fracture (IF) of siliconitride.

.3. Single edge notched beam (SENB)

As was previously pointed out, many forming methods haveeen developed in order to facilitate densification and improvei3N4 properties and mainly the fracture toughness [37,42]. With

ase in this, slip casting and pressureless sintering process is an

deal forming method because of this low cost, its simplicity, itsexibility and its suitability to form complex shapes. However,owadays this process is not very developed because it leads tooderate thermomechanical properties [114]. Penas et al. [37]

ce and Engineering A 527 (2010) 1314–1338

have studied slip casting and the pressureless sintering route forthe manufacture of Si3N4 ceramics with good mechanical proper-ties similar to those obtained by the common processes includingthe effect of attrition on the final microstructure. The authors havecompared their results with those obtained by hot-pressing a sil-icon nitride powder as reported by Kalantar et al. [115]. Si3N4containing 95% �-Si3N4 and 5% �-Si3N4 with additions of 8 wt%Y2O3 and 1.5 wt% Al2O3 as sintering aids were the compositionsinvestigated. The fracture toughness was evaluated by the singleedge notched beam (SENB) at room temperature with an initialnotch of 2 mm depth and 0.3 mm wide. With the purpose to inves-tigate the effects of attrition on the liquid phase sintering twoways have been used: (i) powder blends were premixed by attri-tor milling in ethanol for 4.5 h and once evaporated the solvent,the dried blend was again dispersed in water with 64.5 wt% ofdried blend and 0.07 wt% of deflocculant (for example Darvan C)in a horizontal rolling mill for 24 h (hereafter named A) and (ii)powder blends were directly dispersed in water with 64.5 wt% ofdried blend and 0.07 wt% of deflocculant in a horizontal rolling millduring 24 h (hereafter named B). Once both samples A and B werecast on a plaster of Paris mould and dried, they were pressurelesssintered in nitrogen atmosphere (1 atm) at 1800 ◦C for 1, 1.5, and2 h.

The fracture toughness values were 8.4 ± 0.3, 8.7 ± 0.7 and9.5 ± 0.1 MPa m1/2 for the samples named A1, A1.5 and A2 (sinter-ing time of 1, 1.5 and 2 h, respectively) and 6.8 ± 0.5, 7.9 ± 0.6 and9.3 ± 0.3 MPa m1/2 for the samples named B1, B1.5 and B2 (sinter-ing time of 1, 1.5 and 2 h, respectively). It is clear that the fracturetoughness increased with the sintering time for the compositionsfrom the batches A and B, respectively suggesting that the longersintering time, the more the �-grains grow, and so increase the rein-forcement of the material. The fracture toughness obtained with asample hot-pressed at 1800 ◦C at a constant applied pressure of45 MPa under N2 atmosphere for 1 h and reported by Kalantar etal. [115] was of 7 MPa m1/2 which is comparable with the valuesreported by Penas et al. [37] in the slip casting and pressurelesssintering process where was demonstrated that the influence ofattrition is particularly important for the formation and the growthof �-grains. On the other hand, the sintering time is an importantparameter in the evolution of the self-reinforced microstructureby the development of acicular �-grains contributing to combinehigh fracture toughness and high flexural strength of Si3N4 ceram-ics manufactured by slip casting and pressureless sintering. Table 2summarizes some experimental results on fracture toughness bythe single edge notched beam (SENB) method in silicon nitrideceramics.

2.4. Single edge precracked beam (SEPB)

Salem et al. [22] conducted studies of fracture toughness in asilicon nitride Kyocera SN251 used to manufacture a gas turbinecombustor. The test specimens for fracture toughness measure-ments were machined from the combustor liner and subsequentlyheat treated in air at temperature of 1150 ◦C during 1.5 h inorder to eliminate machining damage. The single edge precrackedbeam (SEPB) [116] and chevron-notch [116,117] methods wereused to obtain fracture toughness values of 7.9, 7.1, 6.9, 6.0 and10.4 MPa m1/2 for temperatures of 25, 800, 1000, 1200 (with SEPB)and 1371 ◦C with chevron-notch method. On the other side, valuesof 7.4 and 6.1 MPa m1/2 at temperatures of 25 and 1371 ◦C wereobtained applying the SEPB method. In this investigation was con-

firmed that the samples tested by the chevron-notch techniqueexhibited nonlinear load-displacement diagrams suggesting stablecrack extension, meanwhile fracture toughness values calculatedfrom maximum load and minimum stress intensity coefficient,decreased with temperature up to 1200 ◦C. However, at 1371 ◦C

M.H. Bocanegra-Bernal, B. Matovic / Materials Science and Engineering A 527 (2010) 1314–1338 1321

Table 1Summary of experimental results on fracture toughness by the indentation fracture (IF) method in silicon nitride-based ceramics.

Composition Sintering conditions Fracture toughness Method Reference

Si3N4 with 3 wt% MgO + 1.5 wt%Al2O3 + 5 wt% SiO2

1780 ◦C for 1.5 and 3 h undernitrogen pressure

5.8 ± 0.6 MPa m1/2,1750 ◦C–1.5 h–P*

Indentation fracture [39]

6.2 ± 0.5 MPa m1/2,1780 ◦C–1.5 h–P*,6.4 ± 0.5 MPa m1/2,1780 ◦C–3 h–P*6.3 ± 0.5 MPa m1/2,1780 ◦C–3 h–GP* planetary high-energyball-millingG* general ball-milling

Si3N4 + 5 wt% MgO + 4 wt% Y2O3 Pressureless sintered at 1700 ◦Cduring 60 min

7.5 MPa m1/2 Indentation fracture [48]

Si3N4 + SiO2 + Yb2O3 (molar ratios 85:10:5) Pressureless sintered at1750 ◦C under nitrogenpressure of 0.1 MPa)

6.3 MPa m1/2 with load of 98 N Indentation fracture [34]9.1 MPa m1/2 with load 588 N

Si3N4 + 4 wt% Al2O3 + 6 wt% Y2O3 SPS 1500 ◦C, 3 min, 50 MPapressure

5.2 MPa m1/2 Indentation fracture [58]

Si3N4 + 4 wt% Al2O3 + 6 wt% Y2O3 + 1 wt%MWNT

SPS 1500 ◦C, 5 min, 50 MPapressure

5.3 MPa m1/2

6 wt% Y2O3 + 1 wt% Al2O3 + 0 and 0.5%carbon powders

Gas pressure sintered at1850 ◦C for 6 h under 2 MPanitrogen gas pressure

5.1 MPa m1/2 for 0% carbonpowders

Indentation fracture [66]

6.6 MPa m1/2 for 0.5% carbonpowders

Si3N4–TiN + Y2O3 (6 wt%) and Al2O3 (2 wt%) Hot-pressed at 1800 ◦C in N2

atmosphere during 2 h6.5 MPa m1/2 for 7.5 vol.% TiNcontent

Indentation fracture [70]

Si3N4–9 wt% Y2O3–3 wt% Al2O3 Sintered in graphite-elementfurnace with Si3N4 packingpowder under one atmosphereN2 to 1780–1800 ◦C during 2 h

6.6 MPa m1/2 for Elkem MetalsCo.

Indentation fracture [75]

6.1 MPa m1/2 for KemaNordGrade 4C5.7 MPa m1/2 for KemaNordGrade 5C6.1 MPa m1/2 for AlbemarleCorp.

9% Y2O3 + 3% Al2O3 (10 wt% �-Si3N4

content)6.6 MPa m1/2

9% Y2O3 + 3% Al2O3 (0 wt% �-Si3N4 content 4.5 MPa m1/2

6.4% Y2O3 + 3.2% MgO (10 wt% �-Si3N4

content7.5 MPa m1/2

Si3N4–9 wt% Y2O3–3 wt% Al2O3 Microwave processing 2.45 GHz,1800–1825 ◦C, 27 h

5.3 ± 0.2 to 8.8 ± 0.4 MPa m1/2 Indentation fracture [77]

Conventional sintering at1800-1825 ◦C 27 h

5.0 ± 0.3 and 7.7 ± 0.4 MPa m1/2

Si3N4 + 9% Y2O3 + 3% Al2O3 (sinteredreaction bonded, Elkem metallurgicalgrade)

Sintered at 1950 ◦C for 3 h in N2

gas under pressure of 1.3 MPaand annealed at 1150 or1600 ◦C for either 5 or 20 h in agraphite-element or amicrowave furnace

4.5 MPa m1/2 for conventionaland microwave annealing ofthe SBR-1 sample at 1150 ◦C for5 and 20 h

Indentation fracture [82,83]

Si3N4 + 6% Y2O3 + 2% Al2O3 (gas pressuresintered (AY6)

∼8 MPa m1/2 for conventionalor microwave heating of SBR-1and AY6 samples

Si3N4 + 6%Y2O3 + 1.5% Al2O3 (AY6) Injection molded and sintered at1750 ◦C in overpressure ofnitrogen or glass-encapsulatedHIPing process

4.6 MPa m1/2 Indentation fracture [87]

Si3N4 + 10 wt% CeO2 + �-Si3N4 seeds (0, 1,3, 5 wt%)

Sintered at 1800 ◦C 1, 2, 4 and6 h

8.4 MPa m1/2 for sintering at 4 h(3 wt% seeds)

Indentation fracture [99]

7.5 MPa m1/2 for sintering at 6 h(3 wt% seeds)6.1 MPa m1/2 for sintering at 2 h(3 wt% seeds)5.0 MPa m1/2 for sintering at 4 h(3 wt% seeds)

Si3N4 + 6 w/o Y2O3 (PY6) Injection molded and densifiedby hot isostatic pressing (HIP) at1750 ◦C

4.2 MPa m1/2 Indentation fracture [156]

1322 M.H. Bocanegra-Bernal, B. Matovic / Materials Science and Engineering A 527 (2010) 1314–1338

Table 1 (Continued )

Composition Sintering conditions Fracture toughness Method Reference

Si3N4–6% Y2O3–2% Al2O3 Gas pressure sintering (GPS)between 1850 and 2000 ◦C

7.5 ± 0.3 to 8.4 ± 0.3 MPa m1/2

(final sintering temperaturesbetween 1900 and 2000 ◦C,composition Si3N4–6%Y2O3–2% Al2O3)

Indentation fracture [85]

Si3N4–Sr2La4Yb4(SiO4)6O2 4.3 ± 0.2 to 7.5 ± 0.7 MPa m1/2

(final sintering temperaturesbetween 1850 and 1900 ◦C,compositionSi3N4–Sr2La4Yb4(SiO4)6O2

SNY14 (84.99 wt% Si3N4 + 5.21 wt%SiO2 + 9.79 wt% Y2O3)

Sintered at 1800 ◦C in nitrogenunder 1.8 MPa pressure

7.3 ± 0.7 MPa m1/2 (30 min) Indentation strength inbending

[111]7.1 ± 0.5 MPa m1/2 (60 min)5.7 ± 0.1 MPa m1/2 (240 min)

7.56.55.8

ttfrrw

gawlstwaowmg[fptaueo

TS

SNTR14 (82.54 wt% Si3N4 + 5.04 wt%SiO2 + 12.42 wt% R2O3(ss)

Sintered at 1800 ◦C in nitrogenunder 1.8 MPa pressure

he load-displacement diagrams became severely nonlinear andhe fracture toughness appeared to increase substantially. There-ore, the combination of slow stroke rate and high temperatureesulted in creep behavior instead of stable fracture. In regards tooom temperature, the fracture toughness measured with the SEPBas comparable to that of the chevron-notch.

Rundgren [118] reported studies of densification and grainrowth in low-doped Y2O3–Al2O3-based silicon nitride ceramicss well as to make a low-doped/low-cost silicon nitride materialith full density. The composition investigated was specifically a

ow-doped 3 wt% Y2O3 + 1 wt% Al2O3, resulting an in situ reinforcedilicon nitride ceramic with extremely high fracture toughness upo 14.8 MPa m1/2 by means of SEPB method [116]. This materialas made using gas pressure sintering at 1 MPa nitrogen pressure

t 1900 ◦C for 10 h. The slow densification rate (the pinning effectf porosity on the grain growth of �-Si3N4 during densification)as determined to be the most important factor for the develop-ent of this tough ceramic. This behavior gives some of the �-Si3N4

rains time to grow before full density is reached. Vuckovic et al.119] in an interesting investigation reported the effect of two dif-erent sintering aids and amounts of �-Si3N4 on the mechanicalroperties of self-reinforced Si3N4-based composites obtained byhe pressureless sintering technique. �-Si3N4 with 10 wt% CeO2

nd �-Si3N4 with 7 wt% Y2O3 + 3 wt% Al2O3 were the compositionssed in their study. Due to its importance as structural material, theffect of composition and microstructure on mechanical propertiesf Si3N4 ceramic has been widely studied by many authors. Becher

able 2ummary of experimental results on fracture toughness of silicon nitride-based ceramics

Composition Sintering conditions

Si3N4 containing 95% �-Si3N4 and 5% �-Si3N4 withadditions of 8 wt% Y2O3 and 1.5 wt% Al2O3

Powder blends premixed bmilling and then horizontamilled for 24 h (sample A)

Powder blends were directhorizontal rolling milled fo(sample B)

Pressureless sintered in nitatmosphere (1 atm) at 18001.5, and 2 h

Si3N4 + 8% Y2O3 + 1.5% Al2O3 Hot-pressed at 1800 ◦C undof pressure in N2 atmosphe

± 0.6 MPa m1/2 (30 min) Indentation strength inbending

[111]± 0.3 MPa m1/2 (60 min)± 0.6 MPa m1/2 (240 min)

et al. [120] have shown that large elongated grains within a fine-grained matrix yield silicon nitride ceramics with both high fracturestrength (1–1.4 GPa) and high fracture toughness (approximately11 MPa m1/2).

An increase in the aspect ratio (length/diameter) of siliconnitride grains allows increasing the fracture toughness of Si3N4 toup 10 MPa m1/2 in conventional sintering according to investiga-tions carried out by Lee et al. [54], Becher et al. [120], Sun et al.[121]. Many years ago Lange [45] showed that the developmentof elongated grains is related to the �–� phase transformationin silicon nitride. However, like all ceramics, they are brittle andflaw intolerant and taking into account these limitations, existan interest in toughening silicon nitride ceramics. Therefore, insitu toughened silicon nitride ceramics or self-reinforced ceram-ics, have received much attention in recent years as they are easierto process than whisker or fiber reinforced silicon nitride [34,122].With this technique, the elongated grains grow in a fine, uniformmatrix during sintering, so that the resultant microstructures aresimilar to those of whisker-reinforced ceramics, in which whiskersare added separately to the silicon nitride powders [34]. Dependingon the composition and amount of the grain boundary phase, goodmechanical properties at high temperatures as high as 1350 ◦C isattained. However, prolonged heat treatment at high temperatures

is required in order to promote �-Si3N4 grain growth [122].

Considering that these heat treatments are expensive and alsolead to decomposition of Si3N4 to its constitutive elements, Siand N resulting in lower density and poor mechanical properties,

by the single edge notched beam (SENB) method.

Fracture toughness Method Reference

y attritorl rolling

8.4 ± 0.3 MPa m1/2 1 hsample A

Single edge notchedbeam

[37]

8.7 ± 0.7 MPa m1/2

1.5 h sample A9.5 ± 0.1 MPa m1/2 2 hsample A

lyr 24 h

6.8 ± 0.5 MPa m1/2 1 hsample B7.9 ± 0.6 MPa m1/2

1.5 h sample B9.3 ± 0.3 MPa m1/2

1.5 h sample Brogen

◦C for 1,

er 45 MPare for 1 h

7 MPa m1/2 Single edge notchedbeam

[115]

Scien

stss[tsi8tTeasat(ootw

asarhwfttT[t[twfmapfTwsco

TS

M.H. Bocanegra-Bernal, B. Matovic / Materials

ome investigations have been developed to process inexpensiveechniques which facilitate the processing of in situ toughenedilicon nitrides by conventional sintering, hot-pressing or hot iso-tatic pressing (HIP) with small amounts of grain boundary phase34,54,120,122]. Tikare et al. [122] reported results on fractureoughness in silicon nitride ceramics with two different compo-itions and with and without additions of �-Si3N4 whiskers. Theydentified the samples as 6Y (Si3N4 + 6 wt% Y2O3 + 2 wt% SiO2) andSc (Si3N4 + 8.6 wt% Sc2O3) which were sintered in a tungsten resis-ance furnace for 4 h at 2140 ◦C and 2.5 MPa nitrogen overpressure.he fracture toughness values were obtained applying the singledge precrack beam (SEPB) technique. To achieve higher densitynd higher fracture toughness with these compositions was neces-ary a good dispersion of the whiskers in the starting powders tovoid the clumping of them. 8.7 ± 0.6 and 4.5 ± 0.5 MPa m1/2 werehe fracture toughness values achieved for 6Y-Si3N4 with whiskers6Y-w) and 6Y-Si3N4 without whiskers (6Y), respectively. On thether hand, the fracture toughness of the 8Sc–Si3N4 with additionsf whiskers was higher than that of 8Sc without whiskers althoughhe increase was not as dramatic as between the 6Y and 6Y withhiskers.

An explanation for the difference in fracture toughness of the 6Ynd 8Sc (both without whiskers additions) ceramics in spite of theirimilar grain size could be due to differences in their grain bound-ry phases. Analyses of X-ray diffraction in the different samplesevealed that the grain boundary phase in 8Sc (without whiskers)ad devitrified to Sc2Si2O7 meanwhile no crystalline silicate phaseas detectable in 6Y (without whiskers). However, fracture sur-

aces observations revealed that the lengths of the large grains wereypically 30–60 �m and the diameters 10–12 �m suggesting thathe whiskers grew faster in diameter than they did in length [122].hese results are in agreement with the finding of Lai and Tien123]. The length of the reinforcement must be a length to diame-er ratio >4:1 based on the mechanics of the silicon nitride system28]. When large elongated grains are generated but with little con-rol of their number and size, there is a modest toughening effecthich rises rapidly for even small amounts of crack growth and the

racture strength increases to 850 MPa. Similarly, these reinforce-ents only work when interfacial debonding process occurs and

llows the reinforcement to survive as a growing crack reaches andasses it [28,91]. Moreover, frictional motion of the debonded rein-orcement against matrix is a significant toughening contribution.his is enhanced with increase in the reinforcement’s diameter as

ell as the length of the debonded interface over which frictional

liding occurs. The frictional contribution to the toughening pro-ess depends not only on the size but also on the number fractionf reinforcements [28,98].

able 3ummary of experimental results on fracture toughness of silicon nitride-based ceramics

Composition Sintering conditions

Silicon nitride Kyocera SN251(Si3N4 + sintering aids + polymeric binder)

Green machined to near final shdewaxed and sintered in nitrogeatmosphere

Si3N4 + 8% Y2O3 + 1.5% Al2O3 Hot-pressed at 1800 ◦C under 45pressure in N2 atmosphere for 1

Si3N4 + 3 wt% Y2O3 + 1 wt% Al2O3 Gas pressure sintering at 1 MPapressure at 1900 ◦C for 10 h

6Y (Si3N4 + 6 wt% Y2O3 + 2 wt% SiO2) Sintered in a tungsten resistancefor 4 h at 2140 ◦C and 2.5 MPa nioverpressure6Y-w (Si3N4 + 6 wt% Y2O3 + 2 wt% SiO2 + �-Si3N4)

AS800 Si3N4 (fabricated by Honeywell Engines) Gelcast and gas pressure sintere

SN282 Si3N4 (fabricated by Kyocera)

ce and Engineering A 527 (2010) 1314–1338 1323

An interesting work was carried out by Kawase et al. [108] in thedevelopment of ceramic turbocharger rotors for high temperatureuse with silicon nitride material fabricated by NGK and designedas SN-60 in its first generation, adopting a silicon nitride materiallabeled as SN-84EC in the second generation taking into account thehigh mechanical strength up to 1200 ◦C. In order to ensure the relia-bility of the ceramic turbocharger rotor (CTR), it is essential to main-tain high resistance against foreign object damage (FOD). The resis-tance against FOD depends on the mechanical strength and fracturetoughness of the material [124]. As stated previously, the mechan-ical strength of SN-84EC is superior to that of SN-60. On the otherhand, the fracture toughness values resulting with sintered sili-con nitride SN-84EC were 6 MN m−3/2 for room temperature, 700,900, and 1000 ◦C and 7 MN m−3/2 at temperatures of 1200 ◦C. Mean-while, values of fracture toughness of 6 MN m−3/2 were obtainedwith SN-60 at room temperature and 700 ◦C and 5 MN m−3/2 attemperatures around 900 ◦C. From these results it is noteworthythat the KIC value of SN-84EC remains high at even 1200 ◦C.

With regard to the aforementioned subject, gas turbine enginefield tests have shown that silicon nitride failures were most likelydue to impact damage according to studies carried out by Choiet al. [125] where they reported that after separate impact testslow toughness values were a major factor. Their investigationswere based on the effect of fracture toughness on the foreignobject damage (FOD) behavior of two commercially available gaspressure sintered gas turbine grade silicon nitride AS800 (fabri-cated by Honeywell Engines) and SN282 (fabricated by Kyocera)at ambient temperature. Additionally, a conventional, hot-pressed,equiaxed, fine-grained silicon nitride NC132 (manufactured byNorton Advanced Ceramics) [126] was used for comparison. TheNC132 Si3N4 exhibited the lowest fracture toughness of the threematerials tested, providing further evidence that KIC is a key mate-rial parameter affecting FOD resistance. The above mentionedsilicon nitride ceramics are currently considered strong candidatematerials for gas turbine applications in view of their substan-tially improved elevated-temperature properties [127–129]. In allexperiments, the authors followed the single edge precracked beam(SEPB) method in accordance with ASTM C 1421 [130] to obtainthe fracture toughness values at ambient temperature with AS800and SN282 silicon nitride starting materials. The resulting valueswere 8.1 ± 0.3 and 5.5 ± 0.2 MPa m1/2 for AS800 and SN282 siliconnitrides, respectively, meanwhile the value previously evaluatedfor the conventional fine-grained NC132 silicon nitride [126] exhib-

ited the lowest value of fracture toughness around 4.6 MPa m1/2

suggesting low FOD resistance compared with the other two sili-con nitride tested materials which are known as in situ toughenedmaterials tailored to achieve elongated grain structures result-

by the single edge precracked beam (SEPB) method.

Fracture toughness Method Reference

ape,n

7.4 MPa m1/2 attemperature of 25 ◦C

Single edge precrackedbeam

[22]

6.1 MPa m1/2 attemperature of 1371 ◦C

MPa ofh

7 MPa m1/2 Single edge precrackedbeam

[115]

nitrogen 14.8 MPa m1/2 Single edge precrackedbeam

[118]

furnacetrogen

4.5 ± 0.5 MPa m1/2

sample 6YSingle edge precrackedbeam

[122]

8.7 ± 0.6 MPa m1/2

sample 6Y-wd 8.1 ± 0.3 MPa m1/2 for

AS800Single edge precrackedbeam

[125]

5.5 ± 0.2 MPa m1/2 forSN282

1 s Scien

ie

iteinrstn(1acm

3

ioFappnaopsvtcst

3

fiv(rdwMdtYapogfs

auodis

324 M.H. Bocanegra-Bernal, B. Matovic / Material

ng in increased fracture toughness, as compared to conventionalquiaxed, fine-grained silicon nitrides.

It is noteworthy that the values of fracture toughness for ceram-cs quoted in the literature tend to be sensitive to the measurementechnique used, independent of whether or not the materialsxhibit R-curve behavior. As example, four different proceduresn order to estimate the fracture toughness of sintered siliconitride samples were examined by Mukhopadhyay et al. [131] andeported that the value of KIC obtained is sensitive to the mea-urement technique and to the experimental parameters of a givenechnique. The greater value of KIC was attained with the chevron-otched beam method compared to the single edge notched beamSENB) technique. The value of KIC from the SENB method was about.7 times higher than that deduced from IF [131]. Table 3 showssummary of experimental results on fracture toughness of sili-

on nitride ceramics by the single edge precracked beam (SENP)ethod.

. Flexural strength

The utilization of engineering ceramics such as silicon nitriden structural applications is inhibited partly because of a lackf knowledge of mechanical behavior under realistic conditions.lexural strength testing of ceramics is the simplest form to testvailable but is prone to errors of several types being the mainroblem is that the stressed volume of the specimen is small com-ared to the total volume (in tensile or compressive tests this isot so) [132]. In silicon nitride ceramics, the flexural strength usu-lly decreases at temperatures >1000 ◦C depending on the naturef the grain boundary phases. Likewise, depending also on powderroperties, processing and, in particular, the type and amount ofintering aids a wide variety of high-temperature flexural strengthalues can be obtained [19,133]. Taking into account that dueo the lower softening temperature of glassy Mg–Si–O–N phasesompared with Y–Si–O–N phases, the drop in high-temperaturetrength with MgO-doped materials occurs at lower temperatureshan those which are Y2O3-doped.

.1. Three-point flexural strength

Flexural strength has been extensively used with silicon nitrideor different applications [134]. For example, Shimizu et al. [135]n interesting investigation, reported results concerning to thearious durabilities and resistances for foreign objects damageFOD) of Si3N4 ceramic rotor and they demonstrated that theseotors showed enough strength and reliability under various con-itions in engine operation. Compositions based silicon nitrideith additions of 2 wt%/2 wt%, 3 wt%/3 wt% and 4 wt% Y2O3/4 wt%gAl2O4 of oxide additives were selected to obtain both enough

ensification and sufficient strength at high temperatures. Forhe authors the sintered Si3N4 ceramics with additions of 3 wt%2O3/3 wt% MgAl2O4 were favorable both at room temperaturend high-temperature flexural strength (measured by using three-oint bending with load span 30 mm), mainly around 900 ◦C. TEMbservations conducted to conclude that relatively small amount oflassy phases existing in triple points are presumably responsibleor improvement of high-temperature strength as for the compo-ition of 3 wt% Y2O3/3 wt% MgAl2O4.

With the purpose to control the strength, oxidation resistancend dimensions, injection molding and derived techniques were

sed [136–139]. However, the development effort for fabricationf radial turbine rotors is continuing. It is important to note that toate it is already possible to the fabrication of silicon nitride ceram-

cs parts free of cracks in the hub portion [49,140]. The flexuraltrength with the aforementioned compositions showed that the

ce and Engineering A 527 (2010) 1314–1338

strength is constant at about 700 MPa from room temperature upto 1000 ◦C. It is also observed that beyond 1000 ◦C the grain bound-ary phase softens, yielding occurs and strength decreases. Strengthmeasurements are generally evaluated at room temperature in tur-bine rotating components because the highest part stresses arelocated in the hub or dovetail region where temperature is lessthan 1000 ◦C.

3.2. Four-point flexural strength

The application of silicon nitride-based ceramics for hot partshas become an alternative with promise. Research and develop-ment of combustors, stator vanes and rotating blades using siliconnitride ceramics has been progressing for hot gas turbines of 20 MWpower output and 1300 ◦C turbine inlet gas temperature [141]because the improvement in heat resistance metallic materials hasreached its limits, and cooling air has increased proportional to thegas turbine inlet temperature elevation, so significant improve-ment in thermal efficiency cannot be expected. Tsuji et al. [142]and Lin et al. [143] reported the use of silicon nitride SN-88 (manu-factured by NGK), SN-252 (manufactured by Kyocera) for 1st stagestator vane and SN-84 (manufactured by NGK) for the 2nd with anaverage turbine inlet gas temperature of 1300 ◦C. The highest gastemperatures at the stator vanes, considering the radial tempera-ture distribution, are 1388 ◦C in the first stage and 1082 ◦C in thesecond stage. The Si3N4 vanes were made by pressureless sinteringin a nitrogen gas atmosphere at temperature range of 1600–1800 ◦Cand the flexural strength values were 790, 615 and 938 MPa for SN-88, SN-252 and SN-84 at room temperature. At high temperature,the values were 760 MPa (1400 ◦C), 494 MPa (1371 ◦C) and 846 MPa(1200 ◦C) in SN-88, SN-252 and SN-84, respectively.

The strength of the silicon nitride turbine rotors at high temper-ature has been gradually improved along several years [103,104].For example, the four-point flexural strength values reported forcommercial silicon nitride SN-50 (pressureless sintered) were 520,330 MPa at room temperature and 1000 ◦C, respectively [103]; 690,690, 590, 270 MPa for SN-81 silicon nitride (pressureless sintered)at room temperature, 1000, 1200 and 1300 ◦C, respectively; 930,900, 900, 600 for SN-84H silicon nitride (post-HIPed) at roomtemperature, 1000, 1200, and 1300 ◦C, respectively; 790, 770,770, 760 for SN-88 silicon nitride at room temperature, 1000,1200, and 1400 ◦C, respectively. With these values, the Weibullcoefficient which is a measure of the strength variations, has beenimproved from an initial 10 to about 20, which is better than castiron (approximately 15) and now closer to steel (around 20–45).The durability of the turbine rotors at elevated temperatureshas been also steadily improved [30,48,102–104] obtaining inturbine rotors tip speed of more than 700 m s−1 with the turbineinlet gas temperature of 1300 ◦C. It is very interesting to note thebehavior of the SN-84H silicon nitride in turbine rotors. From thisinvestigation, it is clearly observed that the strength of the turbinewheel test piece is nearly the same as that of a piece test taken froma block manufactured by the same method as the turbine wheel(isopressed and post-HIPed). The temperature at which turbinewheel material strength started to decrease sharply (knee pointtemperature) was around 1250 ◦C. Miyauchi and Kobayashi [29] insimilar studies reported a knee point temperature of 800 ◦C with afour-point flexural strength of around 550 MPa following the sameprocedures carried out by Kobayashi et al. [103]. Kawase et al. [108]reported also that the above mentioned SN-84H silicon nitridewithstands a new ceramic–metal joint at high temperatures for a

ceramic turbocharger rotor (CTR) where an Incoloy of low thermalexpansion and nickel–chromium–molybdenum steel (SNCM steel)where friction welded to form the metal shaft, which is thensubjected to heat treatment and machining, followed by pressfitting with the ceramic shaft. Machining, balancing and inspection

Science and Engineering A 527 (2010) 1314–1338 1325

atcctitditsaapuaaaTf4b

v[cOwicoawa8waooFaRdahiet

atatflamAodmadt

ta

Fig. 5. Comparison creep displacement as function of time determined at five nom-◦

M.H. Bocanegra-Bernal, B. Matovic / Materials

re carried out to complete the product after joining. In additiono the above mentioned, development of joining technologies oferamic wheels with metal shafts is very important for puttingeramic rotors into massive practical use. Ito et al. [109] reportedhe brazing technology which was successfully developed and firstntroduced into the market in 1985. It is important to point outhat for bonding ceramics to metals; even small thermal expansionifference between them tends to cause significant residual stress

n the bonded body [144,145]. Therefore, considering that thehermal expansion of silicon nitride is only one-fourth that of theteel which is the shaft material of the rotor, the intervention ofbuffer layer is unavoidable for the bond between silicon nitride

nd steel. The selected material for this investigation was a gasressure sintered silicon nitride (GPSSN-NTK EC-141) with a flex-ral strength of 900, 850 and 740 MPa, at room temperature, 800nd 1000 ◦C. In fact was NTK (NGK SPARK PLUG) from its positions a leading ceramic manufacturer and with practical experienceccumulated from oxide ceramics challenged this theme [109].hey found that the results of the bending strength were satis-actory taking into consideration the maximum temperature of50 ◦C under a loading rate of 2.0 mm min−1 during the heat soakack when mounted on the engine [146].

The use of Si3N4 in automotive intake and exhaust ceramicalves subjected to cyclic loading has been reported by Sonsino147]. The design and evaluation methodology applied to theseeramic valves have been carried out in collaboration with Adampel AG and DaimlerChrysler AG [148,149]. Advantages such aseight reduction by about 60%, improvement of valve dynam-

cs and valve train efficiency by about 20% and saving of fuelonsumption by 3–4% and reduction of CO emission by 20% arebtained with the use of ceramic valves instead of steel valvess follows. The valves manufactured were gas pressure sinteredith tensile bending strength of 1000–1300 MPa at room temper-

ture and 850–950 MPa at temperature of 1000 ◦C compared to80–1030 MPa at room temperature and 70 MPa at 800 ◦C reachedith a valve steel X45CrSi93 (AISI HNV3) [150]. For these same

pplications, Ribeiro and Strecker [111] have reported the usef Si3N4 with Y2O3/SiO2 and R2O3(ss)/SiO2 where the rare-earthxide R2O3(ss) was obtained from the mineral Xenotime by Demar-aenquil [112,113,151]. It was found that the flexural strengtht room temperature and at 1200 ◦C was identical or better with2O3(ss)/SiO2 additive than with Y2O3/SiO2 additive and due to theebonding that should occur at the interface between the grainsnd the grain boundary phase inasmuch as the debonding energyas been reported to be directly influenced by chemical bond-

ng between the grain boundary phase and the grains. The joiningnergy in the case of the R2O3(ss)/SiO2-enriched system is higherhan in the Y2O3/SiO2 system [152].

Takatori et al. [101] worked on the composition of Si3N4 withdditions of 5 wt% Y2O3 and 5 wt% MgAl2O4 (spinel). It is importanto point out that Mg-spinel is one of the most effective sinteringids for the densification of silicon nitride at relatively low sin-ering temperature. The authors in their investigations, used forexural strength rectangular bars sliced from the sintered platesnd with dimensions of 3 mm × 4 mm × 40 mm obtaining values byeans of four-point flexure at a crosshead speed of 0.5 mm min−1.t room temperature, the samples had about the same strengthf 700 MPa. However, the strength of the as-sintered specimenropped to 450 MPa at 1000 ◦C, which was improved to approxi-ately 50% with the devitrification of intergranular phases at 1250

nd 1350 ◦C. At 1200 ◦C an increase in strength was not so drastic

ue to the softening of residual glassy phase even in the heat-reated ceramics.

Si3N4 with additions of Y2O3 and Al2O3 in different propor-ions have been investigated by numerous authors due to thedvantages obtained at high temperatures. Balazsi et al. [58] inves-

inal compressive stresses for NC132 silicon nitride at 1300 C. From Ref. [176],reproduced with permission of NASA from “Silicon Nitride Creep Under variousSpecimen-Loading Configurations”, NASA/TM-2000-210026, November 2000, bySung R. Choi and Frederic A. Holland.

tigated a composition of Si3N4 with additions of 4 wt% Al2O3 + 6 wt%Y2O3 + 1 wt% MWNTs (multiwall carbon nanotubes) which wasHIPed at 1700 ◦C in high purity nitrogen using BN embedding pow-der. The bending strength of this composition was measured bymeans of four-point and three-point bending test with spans of40 and 20 mm. The authors reported that the three-point bendingstrength showed an increase by approximately 45% compared toreference with different compositions but with the same level ofporosity. The bending strength values obtained in HIPed samplewere ∼200 and 280 MPa for four-point and three-point bend-ing strength, respectively. The outlined composition has been ofprimary interest to GTE which investigated two specific mate-rials as follows: AY6 (Si3N4 + 6 w/o Y2O3 + 2 w/o Al2O3) and PY6(Si3N4 + 6 w/o Y2O3) which in turn have been selected based onan extensive optimization effort in relation to composition, densi-fication as well as strength [153,154].

Typical room temperature and high-temperature strength val-ues for injection molded and sintered AY6 and PY6 MOR barsrevealed that the bending strength is approximately constant up to900 ◦C (100–105 ksi) to decrease drastically until 1400 ◦C (∼45 ksi).However, the PY6 sample exhibits improved strength to temper-atures at or above 1200 ◦C [155]. Similar behavior was found bythe same authors in HIPed rotor for GTE isopressed sintered AY6and PY6 MOR bars. More significant improvement observed inthe Weibull modulus indicating the impact on the reliability ofmaterial properties in AY6 and PY6 compositions was reached byBandyopadhyay and Neil [156]. To appreciate this affirmation, seefor example Figs. 6 and 7 in the same reference. Nevertheless,Rundgren [118] obtained an extremely high fracture toughnessceramic with composition Si3N4 + 3 wt% Y2O3 + 1 wt% Al2O3, a flex-ural strength value of 564 MPa due to the presence of large �-Si3N4grains acting as defects. Hecht et al. [106] reported flexural strengthvalues of 750, 415, 200 MPa for hot-pressed, sintered and reactionbonded Si3N4 at 20 ◦C. At 1400 ◦C, the flexural strength values were300, 70, and 250 MPa, respectively. It is clear from these results that,several materials failed to those temperatures.

3.3. Biaxial flexure strength

Some works have been undertaken in United TechnologyResearch Center (UTRC) along with Pratt & Whitney Canada(P&WC) in order to develop and produce a 400 kW microturbine

1326 M.H. Bocanegra-Bernal, B. Matovic / Materials Scien

Fig. 6. Lifetime results of Norton nitride materials NT164 silicon nitride turbine andd ◦

fPo

smsIDabdA6rtsTcs

sift

FtwLa

evelopmental billets at 1370 C in air, From Ref. [243] “Verification of Creep Per-ormance of a Ceramic Gas Turbine Blade” by H.T. Lin, P.F. Becher, M.K. Ferber and V.arthasarathy, Key Eng. Mater. vols. 161–163 (1999) pp. 671–674. With permissionf Trans Tech Publications, Copyright 1999.

ystem based on P&WC’s PW207 helicopter engine modified toeet the needs of industrial power generation [157] using as

tarting materials, microturbine rotors manufactured by Kyocerandustrial Ceramic Corp., US (SN-237) and Kyocera Automotiveivision, Kagoshima, Japan (SN-281). Lin and Ferber [157] gener-ted strength values by means of biaxial flexure strength using aall-on-ring arrangement following the ASTM C1161-2002 stan-ard [158]. The bars were tested at 20 ◦C and 30 MPa s−1 perSTM C1368 [159]. Flexural strength values about 939–1019 and87–725 MPa were reported by the authors for SN-237 and SN-281,espectively. By means of SEM, it revealed significant change in theype of strength limiting flaw between as-processed and machinedample, resulting in different flexural performance and reliability.herefore, the mechanical properties of complex-shaped ceramicomponents are often quite different from those determined fromtandardized simple-shaped test specimens.

The machining techniques in ceramics are very important to

tudy because there are two reasons to take into account: the firsts that the machining process for the turbocharger rotor accountsor the greatest position of the total rotor production cost andhe second reason is that the residual damage introduced dur-

ig. 7. Tensile creep displacement as function of time determined at five nominalensile stresses for NC132 silicon nitride at 1300 ◦C. From Ref. [176], reproducedith permission of NASA from “Silicon Nitride Creep Under various Specimen-

oading Configurations”, NASA/TM-2000-210026, November 2000, by Sung R. Choind Frederic A. Holland.

ce and Engineering A 527 (2010) 1314–1338

ing machining degrades the mechanical strength of silicon nitrideceramics. Katano et al. [160] conducted experiments with Si3N4using Al2O3 and Y2O3 as additives to examine the relationshipbetween machining and material strength. Silicon nitride sampleswere ground along their longitudinal and lateral axes using a dia-mond grinding wheels of different grain sizes. Subsequently bothsurface roughness and strength were measured. From these experi-ments was evident an increase in strength degradation with a largergrain size for grinding in the lateral direction. The fracture originin the samples was thought to be the grinding tracks and/or crackgenerated just below the grinding track. The track length and/orcrack length became larger.

Moreover, the authors reported some residual strength asa function of surface roughness as consequence of machin-ing process. The residual strength is an important factor instress simulation. Machined samples of different thicknesses fromcommercially available gas pressure sintered silicon nitride man-ufactured by NGK Spark Plug Co., LTD., were polished with twogrades of diamond paste (6 and 3 �m) to eliminate surface machin-ing damage and studied by Katano et al. [160] to estimate the effectof the post impact bending strength as a function of the impactvelocity for turbine blade thicknesses of 1, 2, and 3 �m. Examplesof this behavior include the failure of aircraft turbine blades causedby ingestion of solid particles [161] and Hertzian cone cracks in heli-copter windshields resulting from flotation of solid particles [162].The results showed that the impact damage initiated Hertzian conecracks more easily in the 1 mm thick specimen than in the 2 or 3 mmthick samples. The mechanical properties of silicon nitride ceram-ics are sensitive to defects and therefore, it is essential to eliminatethem from the fabrication process as much as possible.

Lin et al. [30] studied the mechanical reliability of silicon nitrideAS800 (Honeywell Ceramic Components, Torrance, CA) and SN282(Kyocera Industrial Ceramics Corp., Vancouver, WA) vanes afterfield tests in an industrial gas turbine. Both materials were den-sified by gas pressure sintering and with rare-earth as additivesfor sintering. By means of biaxial flexure strength [150,163], theyobtained the values for selected vanes using a ball-on-ring arrange-ment where the specimens were machined from both the airfoiland platform surfaces by first diamond core drilling small cylindershaving nominal diameters of 6 mm. The strength Sb, was calculatedfrom the following equation:

Sb = 3P1 + n

4nt2

[1 + 2 ln

(a

b

)+

(1 − n

1 + n

)(1 − b2

2a2

)(a2

R2

)](3)

where P is the ultimate sustained load, a is the radius of the supportring, b is the effective radius of contact of the loading ball on thespecimen, R is the specimen radius, t is the specimen thickness, andn is the Poisson’s ratio. As the first approximation, b was taken ast/3 [30]. The strength measured for biaxial disks machined from theSN282 airfoil surfaces did not change significantly with time. How-ever, the strength of samples machined from AS800 airfoil surfacesdecreased with an increase in exposure time due to the formation ofa very rough surface and/or subsurface damage zone. Indeed, fromthis investigation was reported an interesting conclusion relatedto that the lifetime of the Si3N4 vanes is controlled by dimensionaland not mechanical considerations.

Silicon nitride-based ceramics exposed in engine tests have notexhibited the required reliability and stability [164,165] due to thatthe properties measured in the laboratory may not truly reflectthose of the components manufactured in the production line. Lin

and Ferber [164] have cited that silicon nitride ceramics particularlyairfoils, can contain as-processed surfaces which may exhibit dra-matically different properties compared with the bulk material dueto the differences in microstructure as well as chemical composi-tion. During the early 1990s Pratt and Whitney began developing an

M.H. Bocanegra-Bernal, B. Matovic / Materials Science and Engineering A 527 (2010) 1314–1338 1327

Table 4Summary of experimental results on flexural strength of silicon nitride-based ceramics.

Composition Sintering conditions Flexural strength Method Reference

Si3N4 + 3 wt% Y2O3/3 wt% MgAl2O4 Sintered at temperatures between1700 and 1760 ◦C

820 MPa from room temperatureup to 1000 ◦C (Si3N4 + 3 wt%Y2O3/3 wt% MgAl2O4)

Three-point flexuralstrength

[135]

Si3N4 + 4 wt% Y2O3/4 wt% MgAl2O4 700MPa from room temperatureup to 1000 ◦C (Si3N4 + 4 wt%Y2O3/4 wt% MgAl2O4)

SN-88 (NGK) Pressureless sintering in a nitrogen gasatmosphere at 1600–1800 ◦C

790 MPa (RT); 760 MPa (1400 ◦C) Four-point flexural strength [142]SN-252 (Kyocera) 615 MPa (RT); 491 MPa (1371 ◦C)SN-84 938 MPa (RT); 846 MPa (1200 ◦C)

SN-50 Pressureless sintered 520 MPa (RT) Four-point flexural strength [103,104]330 MPa (1000 ◦C)

SN-81 Pressureless sintered 690 MPa (RT) Four-point flexural strength [103,104]690 MPa (1000 ◦C)590 MPa (1200 ◦C)270 MPa (1300 ◦C)

SN-84H Pressureless sintered andpost-HIPed

930 MPa (RT) Four-point flexural strength [103,104]900 MPa (1000 ◦C)900 MPa (1200 ◦C)600 MPa (1300 ◦C)

SN-88 Pressureless sintered 790 MPa (RT) Four-point flexural strength [103,104]770 MPa (1000 ◦C)770 MPa (1200 ◦C)760 MPa (1400 ◦C)

GPSSN-NTK EC-141 Gas pressure sintered 900 MPa (RT) Four-point flexural strength [109]850 MPa (800 ◦C)740 MPa (1000 ◦C)

Si3N4 + Y2O3/SiO2 Sintered at 1800 ◦C 30, 60, and 240min in nitrogen under 1.8 MPapressure

402 MPa (RT); 241 MPa (1200 ◦C)30 min

Four-point flexural strength [111]

437 MPa (RT); 231 MPa (1200 ◦C)60 min469 MPa (RT); 219 MPa (1200 ◦C)240 min

Si3N4 + R2O3(ss)/SiO2 Sintered at 1800 ◦C 30, 60, and240 min in nitrogen under1.8 MPa pressure

394 MPa (RT); 289 MPa (1200 ◦C)30 min

Four-point flexuralstrength

[111]

491 MPa (RT); 315 MPa (1200 ◦C)60 min500 MPa (RT); 303 MPa (1200 ◦C)240 min

Si3N4 + 4 wt% Al2O3 + 6 wt%Y2O3 + 1 wt% MWNTs

Sintered and HIPed at 1700 ◦C in highpurity nitrogen

200 MPa (RT) Four-point flexural strength [58]

Si3N4 + 3 wt% Y2O3 + 1 wt% Al2O3 Gas pressure sintered at 1 MPa nitrogenpressure at 1900 ◦C for 10 h

564 MPa (RT) Four-point flexural strength [118]

Si3N4 (SN-237) Gas pressure sintered + hotisostatic pressing

939–1019 MPa at 20 ◦C Biaxial flexure strength [160]Si3N4 + Lu2O3 (SN-281) 687–725 MPa at 20 ◦C

Si3N4 + Lu2O3 (SN-282) Gas pressure sintered + hotisostatic pressing

225 MPa suction side-inside Biaxial flexure strength [164]470 MPa suction side-outside720 MPa suction side-bulk650 MPa pressure side-inside557676

aDamvstos[iiflii

ir-cooled silicon nitride turbine vane under funding provided byARPA, US Department of Energy, Office of Industrial Technologies,nd Office of Naval Research [144]. With the purpose to study theechanical consistency of an air-cooled SN282 [30] silicon nitride

ane, the miniature biaxial disk sample was used to measure thetrength for specimens cut from the pressure (concave) and suc-ion (convex) surfaces. Three groups of samples were made for eachbtained surface, as follows: (i) inside surface machined, (ii) outerurface machined, and (iii) both sides machined. Lin and Ferber164] clearly observed in their studies that the strength of SN282 sil-

con nitride vane strongly depended upon the sample location. Thencrease of roughness, for example, as well as population of surfaceaws on the inner surfaces arising from the green-state process-

ng and machining, are factors that influence in the strength of thenside surfaces mainly on the suction side. A summary of exper-

0 MPa pressure side-outside0 MPa pressure side-bulk0 MPa production rod

imental results on flexural strength of silicon nitride ceramics isshown in Table 4.

4. Creep

The revolutionary developments during the past 70 years inboth design and efficiency of power plants have emphasized theneed for materials of great strength and durability at very hightemperatures. This is a natural consequence of the fact that, in any

device for the conversion of heat energy into work, the efficiency ofthat device may be increased by increasing the temperature differ-ential between the beginning and end of the conversion [166]. Theadvantages and limitations of ceramics for service at high tempera-tures have been discussed from several years ago [167,168]. Briefly,

1 s Scien

tawaaowcmp

cfmbSmttohr

ndrmi(pggdiflnMs

me[ac[tmirtdmtfarai

spiaape

328 M.H. Bocanegra-Bernal, B. Matovic / Material

he advantages are their high resistance to deformation, fusionnd chemical changes and their relatively low specific gravity,hereas the disadvantages of ceramics when compared to metals

nd metallic alloys, are their relatively low thermal conductivitynd their brittleness. As a generalization, the research in the fieldf ceramics should be aimed toward the development of materialsith excellent thermomechanical properties such as high thermal

onductivity, low thermal dilation, and high ratio of strength toodulus of elasticity as well as high mechanical strength at tem-

eratures >1000 ◦C.As was highlighted above, the manufacture of silicon nitride

eramics requires the use of sintering aids in order to achieveull density, which results in the formation of secondary, com-

only vitreous phases [11,12,19,23,133,134] at high temperaturesy reaction with SiO2, which is always present in the startingi3N4 powder [169], and the degradation of high-temperatureechanical properties of these ceramics is generally caused by

his secondary or grain boundary phase [9,170,171] which limithe creep resistance [11,40,172]. However, it is important to pointut that those silicon nitride ceramic materials exhibiting excellentigh-temperature strength use in most of the cases Y2O3 or otherare-earth oxides as additives for sintering [173,174].

Assuming that a single secondary vitreous phase accompa-ies silicon nitride ceramics [19] from the processing route, theeformation at high-temperature creep behavior in these mate-ials can occur by three ways as follows: (i) the second phaseay act as a ‘lubricant’ that allows a possible grain boundary slid-

ng process due to the viscous movement of this glassy phase,ii) improving the diffusion process due to the higher diffusivityathways of the grain boundaries in the presence of secondarylassy phases, and (iii) the preferential location for nucleation androwth of cavities that provides the secondary phase, during theeformation. However, considering the predominant features dur-

ng the deformation, the mechanism to consider will be viscousow, solution–precipitation, or cavitation-creep. These mecha-isms have been broadly discussed in the elegant review ofelendez-Martinez and Dominguez-Rodriguez [19] on creep of

ilicon nitride.Silicon-based ceramics such as Si3N4, have been the pri-

ary candidates for both advanced automotive and gas turbinengine applications at temperatures between 1300 and 1500 ◦C9,28,32,132,175–177] reaching in the last 10 years the status ofn engineering material [178] by virtue of its high-temperaturereep resistance where a high strength-to-weight ratio is required179]. Being Si3N4 a non-oxide structural ceramic, is susceptibleo creep-assisted microstructural changes (i.e. damage) when is

echanically loaded at high temperatures in ambient air [180]nfluenced by interaction of the material with the oxidative envi-onment. Successful utilization of silicon nitride ceramics in gasurbine engine designs requires extensive characterization of theireformation behavior at such high temperatures where highechanical and thermal stresses are involved and the creep resis-

ance is of particular concern [174]. In addition, it is advantageousor designers to know how the component material will react tostresses at elevated temperatures and whether the material will

elax fast enough at temperature to avoid the initiation of dam-ge associated with the stresses and its associated higher stressntensity [181].

In a complex polycrystalline system such as silicon nitride, mayimultaneously occur phase transformation, sliding of secondaryhase by viscous flow, grain rearrangement, nucleation of cav-

ties as well as partial or total devitrification, notwithstanding,ny single process is unable to contribute for creep characteristicsnd therefore, can be considered that concurrent mechanisms takelace, but one or some of them will predominate under certainxperimental conditions exhibiting greatly different microstruc-

ce and Engineering A 527 (2010) 1314–1338

tural characteristics depending on composition and processingroute [19]. Extensive investigations have been undertaken aboutcreep behavior of Si3N4 under both compression and tensile indi-cating a strong dependence of creep on the sign of the applied stress[129,182–185]. As was reported by Choi and Holland [176], themost of silicon nitride ceramics are subjected to tensile and com-pressive stresses and therefore, in order to predict or estimate creepdeformation and rupture of multiaxially stressed components, itis very necessary to know the creep (rupture) parameters of thematerial individually both in tension and in compression.

4.1. Creep under compression

It is well known that most of studies carried out on creep of sil-icon nitride have been conducted in either tension or flexure [186]and of the few compression studies developed, a majority reportedtesting at low temperatures (<1400 ◦C) where deformation isinvariably accompanied by damage accumulation. According to theabove mentioned and taking in to account that the accumulation ofdamage ultimately leads to the failure of ceramic component, it isimportant to understand the nature of the damage process in orderto determine the strategies to improve high-temperature life time[182], since typical high temperature, high-stress applications forSi3N4, (turbine rotors for example) will involve tensile stresses, andtherefore can be expected a different creep response in tension aswell as in compression [171]. Nevertheless, from the point of viewof design and when silicon nitride will be used in the field of form-ing die tools, research effort has to be directed toward separatetension and compression testing [187]. During compressive creepcavity nucleation and growth also has been observed as being animportant accommodation mechanism of grain boundary sliding[188]. On the other hand, for the same applied stress level in ten-sion and compression, the maximum local tensile stress in tensiletest has been shown to be approximately twice that in compressivetests [189,190]. At the same temperature and applied stress, Si3N4can creep 100 times faster in tension than in compression beingin compression the creep rate linearly proportional to the appliedstress [182,184,185], meanwhile in tension is distinctly nonlinear[182,184,185,191].

The compression creep behavior is conceptually simpler toexplain and offers a baseline to measure the tension creep results.Several investigations have determined that a stress exponentthat increases with stress (2 < n < 7) characterizes the tensile creepresponse, while the compressive creep response exhibits a stressdependence of unity [19,172,176,182,185–187,192–195]. Otherauthors have reported a stress exponent in the range between n = 1and n = 2 [172,186,196,197] and some researchers also find stressexponent n > 2 [198–200]. An example of compressive creep for var-ious Si3N4 ceramics exhibiting a stress exponent n = 1 in the plot oftemperature-compensated steady strain rate as a function of theapplied compressive stress, is clearly illustrated in Fig. 10 in Ref.[19].

The Norton equation [11,19,185,200] is probably the most com-mon of the equations used to represent creep rate of Si3N4:

ε̇N = ε̇No exp −QN

RT�n (4)

where ε̇No is an empirical constant, QN the apparent activationenergy, � the applied stress, R the constant gas, T the absolute tem-perature and n the stress exponent. The apparent activation energyranges from ∼275 to 1000 kJ/mol [19,185,187,199,201,202]. Vari-

ations in activation energy and stress exponent suggest that thehigher values of n are the result of cavitation-creep according toappreciations from Lange et al. [203] and Crampon et al. [187].The deformation mechanisms in creep require a description of themicrostructural dynamical evolution [19]. For mentioning some,

Scien

tprs�brt

s[n[a3raprEoiaCp1�scsaiootoaacpcts[tc

rRo[stapsmpcnw≥dTat

M.H. Bocanegra-Bernal, B. Matovic / Materials

he partial devitrification at high temperatures of the secondaryhase produces an increase in its viscosity which in turn may giveise to strain hardening. This effect has been studied by Wilkin-on [204]. In a similar way, Rouxel et al. [205] have suggested that–� silicon nitride phase transformation affects creep due proba-ly to microstructural reinforcement that it causes and that the �/�atio in Si3N4 seems not to have influence since, in most cases, theransformation is complete after densification.

Several interesting investigations have been done on compres-ive creep of silicon nitride ceramics. For example, Yoon et al.185] selected for their study a commercial silicon nitride denomi-ated SN–Si3N4 due to its high reproducibility in creep behavior206,207]. The compression tests experiments were carried outt 1300, 1350 and 1400 ◦C in air and at stresses between 20 and00 MPa with samples of 2.5 mm × 2.5 mm × 8 mm. The authorseported that the compression data fit the Norton equation withstress exponent of ∼1 indicating that cavity formation is not therimary deformation mode for silicon nitride compression. Theseesults are consistent with the reported by Luecke et al. [182] andvans and Rana [208] for creep controlled by solution–precipitationf Si3N4. A material designated as RAY 38 SM prepared by mechan-cally mixing an �-Si3N4 powder with additions of 8 wt% Y2O3nd 3 wt% Al2O3 as-sintered aids, was the selected material byrampton et al. [187] for their investigations. The mentioned com-osition was hot-pressed under a pressure of 30 MPa at 1710 ◦C forh. The as-sintered ceramic revealed �-Si3N4 with ∼10% residual-Si3N4 as main phases. Creep studies on samples of dimen-ions 3 mm × 3 mm × 9 mm showed that the strain–time curves areomposed of a limited transient regime followed by an apparentteady-state creep whose extension depends on the applied stressnd finally a tertiary stage preceding the rupture. These results aren agreement with previous investigations on compressive creepf Si3N4/MgO ceramics [203]. Applying Eq. (1) in Ref. [209], wasbtained an activation energy value, Q = 650 kJ/mol which is consis-ent with a viscosity-controlled solution–precipitation mechanismf creep. A stress exponent for creep n = 1.5, is indicative of a notice-ble contribution de cavitational creep to the deformation. In thenalysis developed by Evans and Rana [208], the failure under creepavitation state in a material containing an amorphous secondhase involves the nucleation of holes and growth of grain-sizedavities within a damage zone. Crampton et al. [187] concluded thatheir experimental material obtained a time-to-failure in compres-ive creep test at 1300 ◦C, following the Monkman-Grant model210], of m = 1.8 (creep rate exponent), confirming that the straino failure is not constant in this tested material indicating also thatavity growth was controlled by the creep response of the material.

The effect of microstructure on compressive creep of self-einforced hot-pressed silicon nitride was investigated by Boling-isser et al. [186]. Many creep investigations have been developedn self-reinforced Si3N4 containing an oxide grain boundary phase182,211–214]. The starting material used was a hot-pressedilicon nitride with additions of Y2O3, MgO and TiO2 manufac-ured by Dow Chemical, USA. In order to promote grain growth,dditional heat treatments were done at 1900 ◦C under nitrogenressure of 5.17 MPa for 2.67 and 24 h. The authors designed theiramples as as-received (AR) ceramic, short-heat-treatment (SH)aterial and long-heat-treatment (LH) ceramic. Likewise, sam-

les ∼2.3 mm × 2.3 mm × 6 mm were machined to compressivereep tests at temperatures of 1450–1625 ◦C. The stress expo-ent obtained for the AR material was temperature dependentith an n∼1.5–2 at 1450–1525 ◦C, being decreased at temperature

1575 ◦C at n ∼1. A stress exponent of 1 is very consistent with aiffusional grain boundary sliding (GBS) creep mechanism [215].he increase in stress exponent observed with decrease in temper-ture indicated that another creep mechanism can be active or aemperature-dependent threshold stress existed [216]. It is note-

ce and Engineering A 527 (2010) 1314–1338 1329

worthy that the SH and LH ε̇ (s−1) vs � data were nearly identical(see Fig. 3 in Ref. [186]). The AS material deformed by the mecha-nism of grain boundary sliding (GBS) at 1575–1625 ◦C with a stressexponent of n = 1 and the activation energy was 610 ± 110 kJ mol−1.

Considering the obtained microstructures, it is important tostress the evidence of dislocations in the deformed samples, forexample AR (as-received) material. TEM studies on AR sample creptat 1450 ◦C revealed grain translations as well as little evidence ofcavitations. However, in the same material (AR) but deformed at1600 ◦C, few dislocations were observed. Nevertheless, evidence ofgrain boundary sliding in this material, was also elucidated. In theAR materials at 1450–1525 ◦C and the SH and LH materials, GBSprobably could not accommodate all the imposed stress. Stresslevels were high, as indicated by the generation of dislocationsand cavitation occurred. Boling-Risser et al. [186] concluded thatin these tests, steady state was not achieved and the apparentplasticity was due in part to the damage accumulation. The grainsize exponent (Eq. (1) in Ref. 186) was determined to be zero at1525 ◦C for all materials indicating that the elevated-temperaturefracture stress of this tested Si3N4 was a very weak function ofgrain size [211].

A study on the measurements of grain boundary film widthsbefore and after creep in high purity silicon nitride with and with-out additions of Ba, was undertaken by Jin et al. [172]. Recently,grain boundary film thickness and intergranular film chemistryhave been measured in several silicon nitride ceramics [217–219]using high-resolution and analytical electron microscopy [172].High purity Si3N4 powder was doped with 800 wt ppm Ba and thesamples were hot-pressed at temperature of 1925 ◦C during 1 hfor the undoped material and 0.5 h for doped one. The as-sinteredmaterials showed a microstructure composed of equiaxed hexag-onal �-grains with dimensions ranging from 0.2 to 0.3 �m. Someelongated grains also were observed. Compressive creep testingin this study was performed on samples whose dimensions were2 mm × 5 mm × 7 mm in size. The compressive stresses used were50, 100, and 200 MPa at 1400 ◦C in air. Both materials exhibited twocreep stages which differ in strain rate by approximately 1 order ofmagnitude. Jin et al. [172] reported a redistribution of grain bound-ary glass phase after creep suggesting a viscous flow controlledcreep mechanism. Likewise, was observed an increase of 40% in theintergranular film when Si3N4 is doped with Ba, providing furthersupport for the hypothesis that the intergranular film thickness insilicon nitride ceramics depends on the film chemistry being con-sistent this appreciation with the observed by Tanaka et al. [219].The calculated stress exponent n for doped and undoped Si3N4 was3.1 ± 0.6 and 1.7 ± 0.4 indicating that the creep behavior in its firststage is affected by the composition of the starting materials.

Compressive creep of a member of the group of silicon nitridescommonly called in situ reinforced composite materials, such asSi3N4 Norton/TRW NT154 with additions of 4 wt% Y2O3 as sinter-ing aid, hot isostatic pressed and characterized by a microstructurecontaining acicular �-Si3N4 grains (10 �m long by 1 �m diameter),was investigated extensively by Luecke et al. [182]. In this siliconnitride, the second phase glass in interstitial regions is nearly com-pletely crystallized to form �-Y2Si2O7 [220] and Y5(SiO4)3N [182].

Specimens with no more than 5.5 mm tall and 5 mm2 cross-sectional areas, came from sections of the flanges of already-testedtension samples, were selected for compressive creep tests. It isvery difficult to assess the true shape of the compressive creepcurves, due to the slow creep rates as well as the limited resolu-tion of the used telescope focused near the specimen in order tomeasure creep. This measuring technique has been described in

detail in Ref. [221]. Clearly the creep rate in compression is muchlowers that in tension according to the studies also undertakenon tension samples by Luecke et al. [182]. The compression datashow a stress exponent of about unity in samples tested at 1430 ◦C,

1 s Scien

scttocoYt

sntCrFstmsc[

vidsthacer[s(sltampiieatCf1Ttwc

4

atnluipp

330 M.H. Bocanegra-Bernal, B. Matovic / Material

uggesting that there is negligible cavitation during compressivereep and therefore, this stress exponent is a powerful evidencehat creep occurs by a diffusional mechanism, at least for theseested Si3N4 ceramics. On the other hand, the compressive creepf virgin specimens only occurred at temperatures ≥1430 ◦C andreep at temperatures lower than this was not observed. However,ther commercial Si3N4 Kyocera SN220 powder doped with 4 wt%2O3 and 4 wt% Al2O3 as sintering aids crept in compression atemperatures of 1100–1400 ◦C [222].

An extensive creep testing at 1300 ◦C under different appliedtresses to determine the creep behavior of hot-pressed siliconitride (NC132, Norton Advanced Ceramics, Northboro, MA) con-aining MgO as the primary sintering aid was undertaken byhoi and Holland [176]. Their compressive creep displacementesults were plotted as a function of time as can be seen inig. 5. At compressive stress of 630 MPa, the sample failed in aimilar manner to fast fracture producing many fragments. Onhe other hand, from this figure it is clearly observed that the

ost of the curves (except the one at 500 MPa) exhibited theteady-state creep region which in turn is in somewhat goodontrast with the tensile creep curves reported by other authors223–227].

In compressive creep of silicon nitride ceramics, the obser-ation of contrast fringes, also denominated as “strain whorls”s commonly present in the intergranular region in sampleseformed and cooled under load [19,186,198,202,223]. Thesetrain whorls have been seen under compressive creep condi-ions in “pure” and 0.5 wt% Al2O3 + 0.5 wt% Y2O3 Si3N4 materialsot-pressed at temperatures of 1500 or 1600 ◦C [189] as wells in samples of Si3N4/MgO alloys. It is noteworthy that thereep of silicon nitride under compressive conditions can bexplained in terms of grain boundary sliding governed by differentate-controlling mechanisms such as: (i) solution–precipitation185,186,192,193,195,196,202,224], (ii) grain boundary diffu-ion [196], (iii) bulk diffusion [225], (iv) viscous flow [172,197],v) solution–precipitation + shear thickening [223,226], (vi)olution–precipitation + cavitation [187,227], and (vii) accumu-ated microstructural damage [198,199]. Indeed, it is very difficulto conclude that, for some set of experimental conditions suchs temperature, applied stress or strain rate, environment andicrostructure [19], a single deformation mechanism can com-

letely account for the plastic deformation of materials, takingn to account the complexity of the polycrystalline system. Its also important to stress that few works have reported stressxponents n less than one under certain experimental conditionss mentioned above, and such values for the stress exponent seemso be limited to superplastic silicon nitrides as was first reported byhen and Hwang [228] in compressive creep of SiAlONs processedrom � powders with additions between 7.77 and 14.69%, 1.01 and6.56% and 3.70 and 8.91% of AlN, Al2O3 and Y2O3, respectively.heir experiments were carried out under nitrogen atmosphere atemperatures between 1500 and 1600 ◦C. The data collected over aide range of deformation conditions for several conditions [228]

an be seen in Fig. 11(a) and (b) in Ref. [19].

.2. Creep under tension

The use of Si3N4 instead of superalloys as the components thatre exposed at high temperature and stress allows the turbine inletemperature to be increased to 1350 ◦C and cooling can be elimi-ated [229]. The tensile creep of this special ceramic because the

ong-term durability of the rotating parts at high temperatures isltimately controlled by creep, and therefore, a detailed character-

zation of the creep behavior is an essential part for the undertakenrojects. A principal goal of these projects is to develop a life-rediction method up to 10,000 h.

ce and Engineering A 527 (2010) 1314–1338

Different investigations have reported a substantial body of ten-sile creep data on commercial grades of Si3N4 [40,182,230–232]. Itis noteworthy that the most of them summarized their creep databy the classical Norton equation [200] written as follows:

ε̇s = ε̇o(�

�o)n

exp(−�H

RT) (5)

where ε̇s is a secondary or minimum creep rate, � is the appliedtensile stress, T is the temperature expressed in Kelvin, and ε̇o,n, and H empirical constants of the respective fit. In this equa-tion, the linear processes n = 1, are usually interpreted as resultingfrom diffusional [233–235] or solution–precipitation [236] mecha-nisms. Silicon nitride ceramics in tensile creep presents noticeablefeatures in relation to compressive creep [19,168,185,187]. Siliconnitride is significantly less creep resistant under tension than undercompression by between 1 and 2 order of magnitude approximatelyand under similar conditions of applied stress and temperature[19] and according to the reported by Kossowsky [237], Lueckeet al. [182], Gasdaska [191], Yoon et al. [185], Hockey et al. [209],and Whalen et al. [238]. In other words, the mentioned above sug-gests that different deformation mechanisms occur in the two casessuch as illustrated in Fig. 12 in Ref. [19] where is plotted the creeprate vs stress for tensile and compressive creep of Si3N4 under dif-ferent experimental conditions. It is interesting to note from thatfigure the approach of the tensile and compressive strain rate atlow stresses indicating this behavior that the operative deformationmechanism becomes common under those conditions.

The stress exponent reported by some researchers for tensilecreep (for example Refs. [191,205,209]) ranging between 2 and 13,while the values for the activation energies have been reportedbetween 600 and 1645 kJ mol−1 [231,239]. A description in termsof the classical creep equation is less successful at representing ten-sile creep of silicon nitride [185] where the data often exhibit morecurvature than can be captured by a single power-law function[182,184,191]. In effect, the value of n increases with increasingstress [19,185] (Fig. 13 in Ref. [19]). Because n varies, some authorshave suggested that the mechanism of creep deformation differs athigh and low stresses [230,232]. On the other hand, other authors,for example Gasdaska [191] have suggested that the Norton equa-tion (Eq. (5)) does not fit in the creep data well and assumed that theviscosity of the secondary, rate-controlling phase could decreasewith increasing stress. The Gasdaska’s model, predicts a large stressdependence of the creep rate. He reported that the creep occurs bygrain boundary sliding accommodated by solution–precipitationof silicon nitride controlled by the viscosity of glass at the sili-con nitride two-grain interfaces. At high sliding rates, the viscositydepends on the hyperbolic sine of the applied stress. The result wasderived by Luecke and Wiederhorn [230] as follows:

ε = ε3(ı

l) exp(− Q3

RT) sinh(

�˝

RT) (6)

where ı is the thickness of the glass at the two-grain interfaces, lis the distance between sliding interfaces, ˝ is the activation vol-ume for deformation of glass at the interface. Instead of a slidingprocess, it is argued that the cavity formation in and subsequentredistribution of silicate phase limits the creep rate.

Recent studies on commercial Si3N4 (NT154, Saint Gob-ain/Norton Industrial Ceramics Corp., Northboro, MA [182,220],SN-88, NGK Insulators, Ltd., Nagoya, Japan [240] and AS800,AlliedSignal, Torrence, CA [241] have shown curvature in log(strainrate) − log(stress) plots. Values of n = 5 and n = 6 have been obtained

for some of these materials wherein at low stresses, extrapolationsof the two curves converge supporting the outlined above that theoperative deformation mechanism becomes common for compres-sive and tensile creep [230]. It is established that, under tension,the major contribution to the macroscopic strain is attributable to

Scien

cbtwcaftiin[catdic

atasoirssist

apmtdtvsbauNcawoAtircheittsbinsijC

M.H. Bocanegra-Bernal, B. Matovic / Materials

avity formation and nucleation [19]. Likewise, the cavitation iselieved to be part of the net creep process [174]. Cavity forma-ion produces the bulk of the tensile creep strain in silicon nitridehere the cavity volume fraction increases linearly with the tensile

reep strain, with a slope ranging from ∼0.8 to 1. A clear examplebout this can be seen in Fig. 3 in Ref. [230] or Fig. 15 in Ref. [19]or different silicon nitrides ceramics crept under different condi-ions. These cavities normally nucleate and grow primarily in thenterstitial pockets located at multigrain junctions as was reportedn a commercial available hot isostatically pressed (HIPed) siliconitride designated as NT154 and tested at 1316 ◦C by Menon et al.174] who also reported an increase in size of the triple-junctionavity with increase in temperature, strain, or creep rate. Luecke etl. [242] suggested that if cavitation is responsible for the most ofhe creep strain, the temperature dependence of the creep rate isue to this matter transport and therefore, the grain boundary slid-

ng mechanism is not expected to be rate-controlling in the presentase.

Considering that the grain boundary sliding and viscous flowre considered the principal mechanisms for intergranular cavita-ion and are controlled by the amount and viscosity of the residualmorphous phase, the improvement of the creep resistance inilicon nitride ceramics can be expected from the engineeringf the residual glass [229]. Several researchers have participatedn studies in order to establish the repeatability of tensile creepupture in silicon nitride developing a wide variety of specimenhapes and sizes, gripping systems, extensometry techniques, ando forth [178]. An interesting investigation concerning creep of sil-con nitride was carried out by Lin et al. [243] on a Norton NT164ilicon nitride ceramic turbine with 4 wt% Y2O3 as a sintering aid atemperature of 1370 ◦C in air and under selected stress levels.

Various ceramic components such as turbine blades, nozzles,nd turbo compressors have been mass produced for actual com-onent application or field tests. A fundamental factor is theanufacturing of ceramic components in the ability to document

he mechanical properties of complex-shaped components andetermine if these are comparable to those obtained from simpleest coupons utilized in the development of ceramic materials. Iniew of this, Lin et al. [243] measured the creep properties of testpecimens extracted from a complex-shaped ceramic gas turbinelade with the purpose to verify the response of actual componentsnd compare with those obtained from developmental billets. Theysed as starting material a silicon nitride designated as NortonT164 and the turbine blade was manufactured by means of slip-ast plus hot isostatic pressing (HIP) processes with a 4 wt% Y2O3s sintering additive. A smaller density or large elongated grains asell as a finer matrix grain size compared with the microstructure

f the developmental materials, was exhibited by the turbine blade.small dog-bone type tensile creep specimen was extracted from

he as-received turbine blade following a similar model accord-ng to French and Wiederhorn [206]. Although previously reportedesults suggest that the use of small dog-bone specimens yieldsreep results comparable with those obtained from large button-ead specimens using a super grip system, the investigations of Lint al. [243] showed that specimens from the airfoil section exhib-ted the highest creep rates and shortest lifetimes as compared withhose obtained from specimens machined from either (i) the dove-ail section of a turbine blade or (ii) the developmental billets. It isuggested that a separate database of ceramic components needs toe developed for end users and advanced structural design capabil-

ties. These important results are shown in Fig. 6. Stress exponent

∼10 was similar for specimens from both airfoil and dovetail

ections suggesting for this silicon nitride with liquid phase, anncrease in creep cavitation localized [244,245] in multiple grainunctions within fine-grained regions and along grain boundaries.reep deformation in HIP silicon nitride ceramics can be due to the

ce and Engineering A 527 (2010) 1314–1338 1331

diffusional mechanism of solution/precipitation, viscous flow andgrain boundary sliding with creep cavities formed when the creepstrain cannot be fully accommodated [182].

Luecke et al. [182] also studied during tensile creep (compres-sive creep was also investigated by them as outlined above) a hotisostatically pressed (HIPed) silicon nitride with additions of Y2O3and found that the volume fraction of cavities increases linearlywith strain producing nearly all of the measured strain. The ten-sile creep specimens that they used were similar to the developedby Carroll et al. [246] but with some variations as a second reduc-tion in width in the gauge length, giving them a 2 mm × 2.5 mmcross-section. Depending of temperature, the authors reported twodifferent types of cavities. For example, a TEM study revealed for1370 ◦C the presence of small lens-shaped cavities from range100 to 200 nm located commonly on two-grain boundaries withnormals approximately parallel to the tensile axis. On the otherhand, the same TEM analysis showed cavities larger 0.2–0.7 �mlocated at multigrain junctions, typically in the pockets of equiaxedsub-micrometer-sized grains of Si3N4 and crystalline silicate. Incontrast, at 1400 ◦C and higher very few of the lens-shaped cavitieswere revealed. Fig. 7 of [182] shows a collection of TEM micro-graphs where the most of the features of the cavity size and shapedistributions can be clearly observed, supporting that the additionof cavities is central to the production of tensile strain.

Similar studies were undertaken by Wereszczak et al. [180] oncommercially available hot isostatic pressed (HIPed) silicon nitride(NCX-5102) under ambient air and argon environments. It is wellknown that silicon nitride is susceptible to both oxidation andcreep assisted microstructural changes when it is mechanicallyloaded at elevated temperatures in ambient air. Taking into accountthat silicon nitride subjected to high-temperature mechanical testsin an inert atmosphere permits the observation of creep assistedmicrostructural changes without the effects from oxidation beingactive. In reason of this, the authors carried out creep and oxida-tion tests via dynamic fatigue with argon atmosphere at 1370 ◦C[247,248]. The composition experimented was Si3N4 + 4 wt% Y2O3and the specimens were obtained via slip-cast, glass encapsulatedand then HIPed at a pressure of 210 MPa (30 ksi) at temperaturesbetween 1700 and 1950 ◦C. An additional post-HIPing treatmentwas developed in order to reach the crystallization and promotethe devitrification of the second phase. The tensile creep testingwas carried out on button-head specimens at 110, 125 and 140 MPaat 25 ◦C and 40–60% RH or alternately in an argon (99.999% purity)environment at 1370 ◦C. The authors concluded that NCX-5102 sili-con nitride creeps faster and exhibits shorter lifetimes in argon thanin ambient air at the testing temperature being the stress depen-dence of the minimum creep rate of specimens tested in air wasabout 60% of that for specimens tested in argon (stress exponent nwas of 3.6 and 5.9).

Similarly, Kossoswky [237] reported that the creep rate of ahot-pressed silicon nitride was greater and the lifetime shorter inhelium than in an air environment. A relatively high value for thecreep exponent suggests that creep can be due to diffussion (n = 1)accompanied by cavitation as well as grain separation by viscousflow and probably grain boundary sliding. Macroscopically largestress-corrosion cracking (SCC) damage zone caused failure in allspecimens creep-ruptured in air and argon. At a microscopic scale,the SCC damage zone was characterized by a high concentrationof multigrain junction cavities and pores suggesting a coalescenceof pores and cavities to form a macroscopically large SCC damagezone leading finally at the fracture of the tested specimen. Other

results concerning to tensile creep of hot isostatic pressed siliconnitride ceramics have been summarized in Table 7 of the interest-ing review by Melendez and Dominguez [19]. Choi and Holland[176] and Kleebe et al. [40] also developed tensile creep studieson hot-pressed silicon nitride (NC132, Norton Advanced Ceramics,

1 s Science and Engineering A 527 (2010) 1314–1338

NuTrtsitstp

poieioasa[cbutwfObabttbst

Fig. 8. Nominal creep strain rate as function of time determined for NC132 sili-con ntride at 1300 ◦C. (a) Pure tension and (b) pure compression. From Ref. [176],reproduced with permission of NASA from “Silicon Nitride Creep Under variousSpecimen-Loading Configurations”, NASA/TM-2000-210026, November 2000, by

Fm“

332 M.H. Bocanegra-Bernal, B. Matovic / Material

orthboro, MA) with additions of MgO as sintering aid at 1300 ◦Cnder different stresses in order to determine the creep behavior.heir results are shown in Fig. 7 where the end point of each curveepresents to the end of the tests since no specimen failed before theest interruption. On the other side, they reported after SEM studiesome degree of oxidation, depending on the test time period. It isnteresting to extract from this figure that the curves correspondingo the samples tested at 57 and 81 MPa exhibited no well-definedteady-state region, but rather a change in displacement rates withhe time suggesting as result, no distinct region corresponding torimary creep.

A summary of the strain rates determined in tensile and com-ressive creep by Choi and Holland [176] are illustrated in Fig. 8. It isbserved that the strain rates range from 10−10 to 10−8 s−1, whichs typical for most advanced ceramics. It is noteworthy in this figure,xcept for compression tests, the strain rate decreases with increas-ng time which in turn can be due to the applied stress and typef loading. The nominal creep strain rate as a function of nominalpplied stress is summarized in Fig. 9 for tension and compressionpecimen-loading configuration for the tested NC132 silicon nitridet 1300 ◦C and fitted to Power-law (Norton) [200], hyperbolic sine191], step [248], and redistribution [230,240] creep models. Foromparison effects, results corresponding to uniaxial flexure andall-on-ring biaxial flexure are also included. Analyzing the fig-re, there was no significant difference in curve fitting betweenhe experimented models and the stress exponent (n) calculatedere 1.72 ± 0.34 in tension, 2.51 ± 0.36 for compression, 2.01 ± 0.45

or uniaxial flexure, and 2.49 ± 0.28 for ball-on-ring biaxial flexure.n the other hand, the negligible difference in stress exponent netween the four loading configurations suggests that the mech-nisms associated with creep of this specific material would note significantly different with any specimen-loading configura-

ion. Choi and Holland [176] concluded from their investigationshat although the stress exponent was not significantly differentetween tension and compression, the difference in strain (andtrain rate) between the two for a given applied stress was impor-ant being in tension 1.5–2 orders of magnitude larger than in

Sung R. Choi and Frederic A. Holland.

ig. 9. Nominal creep strain rate as function of nominal applied stress for four specimen-loading configurations for NC132 silicon nitride at 1300 ◦C and fitted to four creepodels. (a) Power-law (Norton) model, (b) Hyperbolic sine model, (c) Step model, and (d) redistribution model. From Ref. [176], reproduced with permission of NASA from

Silicon Nitride Creep Under various Specimen-Loading Configurations”, NASA/TM-2000-210026, November 2000, by Sung R. Choi and Frederic A. Holland.

Scien

cs

riassaaaTwTqofpmatw[[ttutr1

lrTfsamrctsthlVcfc

sticsotWatfactsY

M.H. Bocanegra-Bernal, B. Matovic / Materials

ompression if the compression data are extrapolated toward lowertresses.

As outlined above, there is an important interest in the creepesistance and creep failure of several silicon nitride-based ceram-cs in advanced gas turbine engines operating at high temperaturesnd stress levels [249,250]. Likewise, among several silicon nitridestudied for this aim, a green-state injection molded and gas pres-ure sintered silicon nitride using ytterbium silicate as a sinteringid, is a promising candidate for the high-temperature gas turbinepplications [240,251]. Krause and Wiederhorn [250] investigatedn in situ reinforced composite of silicon nitride and silicon carbide.hey compared their results on gas pressure sintered silicon nitrideith the obtained in a hot isostatically pressed silicon nitride [182].

he material was processed by aqueous slipcasting and subse-uently gas pressure sintering with a final sintering temperaturef >1900 ◦C. Posterior heat treatment (1500 ◦C) was carried outor crystallization and formation of the M5Si3O12N and M4Si2N2O7hases being M the rare-earth elements. All the creep tests with thisaterial (denominated by authors as EXPSN) exhibited primary

nd secondary creep, but only two of the tests also involved ter-iary creep. The creep behavior of the EXPSN ceramic was comparedith two structural ceramics such as NT154 (labeled as HIPSN)

182] and SN88 (labeled as GPSSN) [240]. Krause and Wiederhorn250] reported that the EXPSN ceramic is more creep resistant thanhe HIPSN and GPSSN, and the HIPSN is more creep resistant thanhe GPSSN. Some experimental points were extrapoled to 150 MPasing least-squares fitted parameters of the logarithmic version ofhe expression reported by Luecke and Wiederhorn [230] to all theate data of that material (see in more detail Figs. 7 and 8 and Tablesand 2 in Ref. [250]).

It is shown in this investigation that the EXPSN has a creepifetime that lasts longer than that of the other two ceramic mate-ials over a range of stress at temperatures lower than 1425 ◦C.he fracture mechanism map can provides another perspectiveor comparing the creep rupture behavior of different materialsuch as reported Wiederhorn et al. [252]. Following this appreci-tion, Krause and Wiederhorn [250] using the fitted parametersentioned in Table 1 of [250] predicted stresses that would cor-

espond to a creep lifetime of 1000 h for each of the silicon nitrideeramic materials at temperatures between 1300 and 1400 ◦C werehe stresses are well below those required for fast fracture andtatic fatigue. According to these studies, for selected creep life-imes at elevated temperatures, the EXPSN ceramic could sustainigher applied stresses, almost twice as high at 1300 ◦C for a creep

ifetime of 1000 h, compared with the other two ceramic materials.ery interesting to note that at temperatures >1400 ◦C the stressurve for the other two materials tended to converge with the oneor EXPSN suggesting that at high temperatures the differences inreep behavior became insignificant.

As was aforementioned, silicon nitride containing ytterbiumilicate as a sintering aid is a promising candidate for the high-emperature gas turbine applications [240,251,253]. However, it ismportant to stress that the effect of the grain size on the tensilereep of silicon nitride has not been extensively studied. However,ome investigations about this theme have been undertaken, bothf them in flexure [254,255] and compression [238,256]. Unfor-unately, these results do not apply to tensile creep. Therefore,

iederhorn et al. [253] reported experimental results on materi-ls prepared by conventional powder processing techniques addingo silicon nitride powder (E10, UBE Industries, Japan) 13.14 massraction Yb2O3 (supplied by Ventron, Germany) as sintering aid

chieving complete densification. Additional experiments werearried out in a second set of samples which contained an addi-ional 0.5% mass fraction Al2O3 (AKP 53 Sumitomo, Japan). Thepecimens were denoted as SN5Yb for samples containing onlyb2O3 and SiO2 as additives, whereas the Al2O3-containing sam-

ce and Engineering A 527 (2010) 1314–1338 1333

ples are labeled SN5YbAl. The tensile creep tests used a dog-bone(SR51) specimen according to specified by French and Wiederhorn[206].

In order to avoid any microstructural change not related tocreep, the specimens were heat treated in air after machiningto crystallize the grain boundary phase at 1350 ◦C during 250 hfor the samples SN5Yb and 325 h and 1250 ◦C for SN5YbAl speci-mens. The creep tests to failure for SN5Yb material were conductedat 1300 ◦C < T < 1370 ◦C in a stress range 75 MPa < � < 125 MPa.For the SN5YbAl material, the creep tests were carried out at1100 ◦C < T < 1250 ◦C and a stress range of 75 MPa < � < 100 MPa. Ingeneral, the creep curves exhibited only primary and secondarycreep with a tensile stress exponents much larger than 1. Like-wise, Wiederhorn et al. [253] concluded that a factor of 2 differencein grain size had no effect on the tensile creep behavior of siliconnitride sintered with additions of Yb2O3 suggesting an agreementwith the new theories of tensile creep of silicon nitride, whichpredict grain size exponents p = 0 or −1. Taking into account thenegligible effect of grain size on creep rate, the addition of 0.5%mass fraction of aluminum oxide has a substantial effect on thetensile creep of silicon nitride (increasing the creep rate by overfour orders of magnitude), supporting that the chemical composi-tion of the second phase is more relevant than grain morphologyin controlling tensile creep behavior.

Very similar to the creep under compression and outlined above,the deformation mechanisms currently invoked to explain thecreep [19] of silicon nitride under tension can be explained in termsof: cavitation mechanism governed by (i) cavitation [22,257–259],(ii) grain boundary sliding [185,205,209,240,259–261], (iii) grainboundary sliding + viscous flow [262,263], (iv) diffusion [232]; grainboundary sliding mechanism governed by (i) cavitation + viscousflow + diffusion [191,264], (ii) viscous flow + diffusion [266], (iii)viscous flow + solution–precipitation [266]; grain boundary dif-fusion mechanism governed by cavitation + viscous flow + grainboundary sliding [252] and solution–precipitation mechanism gov-erned by the same [257].

In order to study in more depth the themes here summarized,the reader is referred to consult some other important experimen-tal results on creep of silicon nitride ceramics shown in Tables6 and 7 from Ref. [19], where information about the processingroute, microstructure and other parameters of importance for theevaluation of creep are indicated.

4.3. Superplasticity of silicon nitride

Since superplasticity in Si3N4 was discovered by Wakai et al.[267] in 1990 in a Si3N4/20% SiC composite, some interesting litera-ture concerning superplastic flow in the field of advanced structuralceramics is nowadays available [19,268,269]. It is well known thatsilicon nitride-based ceramics are promising structural materialsfor mechanical applications. However, the excellent mechanicalproperties that make these materials desirable also make themvery difficult and costly to machine into complex shapes [179].These problems can be solved by the in situ formation of reinforce-ment, instead of adding the secondary phase to the starting powdermixture, applying superplastic deformation as a novel method forthe net-shape fabrication of components in ceramics [270]. Manycovalently bonded ceramics based on fine-grained Si3N4 alloys canbe superplastically deformed at temperatures between 1500 and1600 ◦C [265,267,269,271]. Fine-grain superplasticity was broadlyreported for yttria-stabilized tetragonal zirconia ceramics [272],

alumina [273], and mullite [274] between others.

Other clear examples of superplastic Si3N4-based ceramics areSiC–Si3N4 and �-Si3N4 [275], SiAlONs (�, �, and �/� phases) [226],�-Si3N4 [205]. Two types of materials previously cited are veryinteresting inasmuch as they possess high strain rate (ε̇ >10−4 s−1)

1 s Scien

aiswp[h[te

�tp1floafiaumwr>pot(lsAflibs(sabc

ctetw∼pasboftfii4tIiac

w

334 M.H. Bocanegra-Bernal, B. Matovic / Material

nd relatively low flow stress (�): (i) SiAlONs ceramics [271,272]n which the superplasticity have been enhanced greatly by tran-ient liquid phases and (ii) fine-grained �-Si3N4 ceramics [275,276]hich showed excellent formability at the relatively low tem-erature of 1500 ◦C according to the reported by Mitomo et al.275], Nishimura et al. [277] and Wang et al. [276], exhibiting aigher grain size stability after superplastic forming at 1500 ◦C179] and with the possibility to modify the equiaxed microstruc-ure by means of a postforming annealing obtaining a ceramic withxcellent mechanical properties [278].

Xie et al. [270] conducted studies with a mixture of ultrafine-Si3N4 and a SiO2-containing additive and produced a superplas-

ic Si3N4-based composite using the concept of a transient liquidhase, which is formed at temperatures in the range of 1500 and700 ◦C, just within the temperature regime for the superplasticow of Si3N4-based ceramics. The starting materials comprisedf 63 wt% ultrafine �-Si3N4 powder with a particle size of 0.2 �mnd additions of 7 wt% cordierite (2MgO·2Al2O3·5SiO2). The densi-cation of samples was performed via hot-pressing in a nitrogentmosphere of 0.1 MPa in range of 1650–1850 ◦C during 5 minnder 25 MPa of applied pressure. All materials prepared for defor-ation were sintered at temperature of 1750 ◦C, since accordingith Xie et al. [270] there is a narrow processing window in the

ange of 1725–1775 ◦C where can be reached sintered densities97% theoretical density and small amounts of Si2N2O. The liquidsroduced at the deformation temperatures may act as lubricants inrder to allow the grains to slide over each other more easily. Forhe sintered �-Si3N4 samples, relatively flat stress–strain curvesdeformed at 1600 ◦C at various constant ε̇) are maintained over aarge range of strain for all strain rates suggesting extensive steady-tate deformation meanwhile, the strain hardening was negligible.

value of 1 for strain-rate sensitivity (m) indicates Newtonianow behavior. The activation energy Q calculated for this compos-

te was 450 kJ mol−1 determined by means of an Arrhenius plot,eing this calculated value somewhat less than that reported foruperplastic SiAlON (650–980 kJ mol−1) [226,271,272] and Si3N4750–950 kJ mol−1) [205]. It is important to note that the equiaxedhape of the � grains along with the transient liquid that formst the intermediate stage of the �–� phase transformation haveeen used to enhance superplastic deformation in silicon nitrideeramics from the � powders [19].

Rouxel et al. [205] found in their investigations that for �ontents lower than 30 vol.%, the � phase controls the deforma-ion whereas � contents higher than 40 vol.%, transient liquid andquiaxed shape of the predominant � phase control de superplas-ic flow. The authors also observed that even at high strain ratesere achieved in compression large strain without cracking (at1600 ◦C) suggesting therefore, that the tested material is super-lastic. During deformation at high temperature, Si2N2O nucleatednd grew from a SiO2-rich liquid and developed into elongatedhapes, which in turn led to an in situ Si2N2O reinforced Si3N4-ased composite with improved mechanical properties as a resultf superplastic deformation. For example, the strength increasedrom 559 ± 96 MPa for the as-sintered material to 740 ± 65 MPa forhe deformed material. Similarly, the fracture toughness increasedrom 3 ± 0.2 to 4.4. ± 0.2 MPa m1/2. On the other hand, via anneal-ng treatment, the finished composite showed a further increasen bending strength and fracture toughness of 957 ± 51 MPa and.8 ± 0.3 MPa m1/2. These results were consistent with the observa-ions of Nishimura et al. [278] and Kondo and co-workers [279,280].t is believe with these results that the deformation mechanism that

s operative is solution–precipitation-accommodated grain bound-ry sliding, which is common among fine-grained materials thatontain glassy phase [19,270].

Similar investigation was undertaken by Zhan et al. [179]ith a hot-pressed fine-grained �-Si3N4 powder with a parti-

ce and Engineering A 527 (2010) 1314–1338

cle size of approximately 0.2 �m and additions of 5 wt% Y2O3(99.9% pure) and 2 wt% MgO as sintering aids. The hot-pressedsamples were tested at temperatures between 1450 and 1650 ◦Cunder compression achieving a high deformation strain rateof 1 × 10−2 s−1 and under punch stretching, the strain rateachieved was 1.2 × 10−3 s−1 [226] suggesting that SiAlONs mate-rials exhibit excellent formabilities via transient phase. Zhan etal. [179] compared their experimental results with some resultsreported by other researchers under different experimental con-ditions [205,226,228,265,275,277,281,282] where observed thatflow stress under tension is consistently lower than flow stressunder compression, reflecting therefore a tension-compressionflow asymmetry [205]. Likewise, the authors concluded that thegrain boundary sliding and grain rotation, accommodated by vis-cous flow, might be the mechanism of superplasticty for the testedmaterial.

The transient liquid, which is formed by the eutectic reactionbetween oxide additives and nitrides during the manufacturingof SiAlONs materials (�- and �-silicon nitride materials that mayform solid solutions with some aluminum-based compounds ormixtures) [19] remain in the as-fabricated material in order to facil-itate subsequent superplastic deformation. Guided by this concept,Hwang and Chen [272] in an important investigation, fabricatedseveral SiAlON materials with large amounts of unreacted �-Si3N4after fabrication. Superplastic SiAlON’s of the nominal compositionYm/3–Si12−(m+n)Alm+nOnN16−n (in other words: the generic com-position for �′SiAlON) which lie on the �′-plane of the Janeckeprism, were investigated by these authors. The compositions wereprepared from �-Si3N4 (UBE, E10), AlN (Tokuyama Soda Co., Ltd.,Type F), Al2O3 (Sumitomo Chemical America Inc., AKP50) and Y2O3(Aldrich Chemical Co.). The compression tests were developed atstrain rates ranging from 1 × 10−4 to 1 × 10−3 at temperature of1550 ◦C using a single specimen for each strain rate. At higher strainrates, a constant flow stress is maintained over a large range ofstrain indicating therefore, steady-state deformation for the mate-rial labeled as 0610 with a phase assemblages on �′-plane of 64%�, 27% �′, 9% �; 32% �, 41% �′, 27% �′and 5% �, 48% �′, 47% �′

for hot-pressed at 1550 ◦C, deformed at 1550 ◦C, and deformed at1600 ◦C, respectively. On the other hand, in the plot of stress vsstrain rate for the cited material can be observed two regimenswhere at lower strain rate range the slope is close to 1 suggestingNewtonian flow, whereas at higher strain rates, the slope increasesto around 2, indicating shear thickening (also known as Newtonian-shear-thickening transition [228].

Regarding to punch-stretching tests also carried out by Hwangand Chen [272], these were conducted at a constant crossheadspeed of 0.2 mm min−1 until failure. From the different tested spec-imens, the material identified as 1010 (44% �, 56% �′; 30% �, 70%�′ and 5% �, 90% �′, 5% �′ for hot-pressed at 1550 ◦C, deformedat 1550 ◦C, and deformed at 1600 ◦C, respectively) has relativelylow formability with an essentially �/�′phase assemblage, whileall other materials contain a considerable amount of �′-SiAlON(more details in Table 1 of [272]). From this investigation was con-cluded that the role of �′-SiAlON in formability lies in its tendencyto form elongated grains and to align on the biaxial stretching planeproducing a fiber-strengthening effect which in turn retard the frac-ture process [281,283]. Likewise, all the SiAlON’s ceramics on the�′-plane in the “Janecke prism” have demonstrated their super-plasticity by punch-stretching tests with a microstructure formedby very-fine grains.

Rosenflanz and Chen [226] also studied the “classical” super-

plasticity (generally assumed to proceed by fully accommodation-grain boundary sliding) [19] of SiAlON ceramics with the nominalcomposition (YxLi)0.6/(3x+y)Si8.9Al3.1O2.5N13.5. The starting materi-als used for this investigation were the same used by Hwang andChen [270] besides of Li2O (Aldrich Chemical Co.). The composition

Scien

ooocfpidditigisNmpd1

aio[dAamaetiawtaItvsaabSatit

R

M.H. Bocanegra-Bernal, B. Matovic / Materials

f powder mixtures as well as hot-pressing parameters and amountf �′-SiAlON formation are described elsewhere [226]. In order tobtain full densification it is necessary a small fraction of �-phaseonverted to �′-SiAlON phase and the resulting microstructure isormed by fine, equiaxed grains during the entire history of super-lastic forming, exploiting the window for classical superplasticity

n Si3N4 materials (the left-hand side of the transformation curverawn in Fig. 1 of [226]. In this investigation, the authors haveemonstrated that during superplastic deformation, these ceram-

cs undergo only a small amount of � → � phase transformation andhe resulting �′-SiAlON remains mostly equiaxed. Considering themportance of the liquid phase in silicon nitride ceramics, it is sug-ested that the superplastic deformation in silicon nitride ceramicss controlled by the liquid surrounding the grains. The value oftress exponent n was approximately unity corresponding to theewtonian regime. An important observation was that high defor-ation rates (10−3 s−1) in compression and 1.2 × 10−3 s−1 in biaxial

unch stretching were achieved in the tested SiAlONs at interme-iate lithium/yttrium compositions in the temperature range of450–1625 ◦C.

To allow superplasticity in Si3N4 it is important to reduce themount of the glass phase [269]. To achieve this, hot isostatic press-ng (HIP) is a very effective way to consolidate Si3N4 powder withr without a small amount of additives as reported Honma et al.284]. Burger et al. [223] investigated the superplasticity in Si3N4ensified by means of HIP and with additions of Y2O3 (0.5 wt%) andl2O3 (0.5 wt%) as sintering aids. This study revealed a glassy phaset two-grain and triple-point junctions being obtained a startingaterial with a quasi-ideal distribution of the amorphous phase

s a continuous film at two-grain junctions as well as very smallxcess at the triple-point junctions. It is noteworthy that a shear-hickening behavior (for these systems, it corresponds to a decreasen viscosity with increasing applied stress) [19] has been observedt 1643 ◦C during compression tests under stresses of 10–100 MPaith a stress exponent decreasing from unity to ∼0.5 at a transi-

ion from mild to a strong strain hardening at 20 MPa which wasttributed to the occurrence of rigid contacts between the grains.t is assumed that, when the normal stress on a volume element athe grain boundary exceeds a certain critical value �c, the associateolume element becomes rigid. Therefore, the material can be con-idered to be a “composite” formed by the two types (deformablend rigid) phases [19]. The authors also concluded that grain bound-ry sliding accommodated by solution–precipitation creep coulde the mechanism of superplasticity in the present fine-grainedi3N4 with the presence of a thin intergranular liquid phase and thebsence of dislocation activity. The observed strain worls were usedo evidence the increase of rigid contacts between the grains withncreases in compressive stress as consequence of the expulsion ofhe wetting liquid film.

eferences

[1] A. Pfeiffer, A. Schulz, S. Witting, ASME Paper 91-GT-63, 1991.[2] S.C. Danforth, W. Symons, K.J. Nilsen, R.E. Riman, in: J.G.P. Binner (Ed.),

Advanced Ceramic Processing and Technology, Willian Andrew Publishing,Noyes, 1990, pp. 39–71.

[3] C. Santos, S. Ribeiro, K. Strecker, L.E.G. Da, Silva Acta Microsc. 2 (2003) 83–86.[4] L.J. Bowen, R.J. Weston, T.G. Carruthers, R.J. Brook, J. Mater. Sci. 13 (1978)

341–350.[5] D.R. Messier, F.L. Riley, R.J. Brook, J. Mater. Sci. 13 (1978) 1199–1205.[6] P. Sajgalik, K. Rajan, P. Warbichler, F. Hoffer, J. Dusza, Key Eng. Mater. 159–160

(1999) 405–410.[7] H. Emoto, H. Hirotsuru, M. Mitomo, Key Eng. Mater. 159–160 (1999) 215–220.

[8] K.H. Jack, J. Mater. Sci. 11 (1973) 1135–1158.[9] N. Hirosaki, A. Okada, M. Mitomo, J. Mater. Sci. 25 (1990) 1872–1876.

[10] S. Hampshire, M. Leigh, V.J. Morrissey, M.J. Pomeroy, B. Saruhan, 3rd Interna-tional Symposium Ceramic Materials & Components for Engines, November27–30, Las Vegas, USA, 1998.

[11] K.M. Fox, J.R. Hellmann, Int. J. Appl. Ceram. Technol. 5 (2008) 138–154.

ce and Engineering A 527 (2010) 1314–1338 1335

[12] P. Gui-hua, L. Xiang-guo, L. Min, L. Zhen-hua, L. Quian, L. Wen-lan, Scr. Mater.61 (2009) 347–350.

[13] O. Yeheskel, Y. Gefen, M. Talianker, J. Mater. Sci. 19 (1984) 745–752.[14] T. Fujikawa, Y. Inoue, K. Homma, T. Tatsuno, H. Okada, N. Kondo, Kobelco

Technol. Rev. 3 (1988) 1–8.[15] M.J. Hoffmann, G. Petzow, Pure Appl. Chem. 66 (1994) 1807–1814.[16] K. Negita, J. Mater. Sci. Lett. 4 (1985) 755–758.[17] N. Hirosaki, A. Okada, K. Matoba, J. Am. Ceram. Soc. 71 (1988), C-144-C-147.[18] S. Hampshire, M.J. Pomeroy, B. Saruham, in: P. Vincenzini (Ed.), High Tech

Ceramics, Elsevier Science Publishers B.V., Amsterdam, The Netherlands,1987, pp. 941–951.

[19] J.J. Meléndez-Martínez, A. Domínguez-Rodríguez, Progr. Mater. Sci. 49 (2004)19–107.

[20] D.W. Richerson, Am. Ceram. Soc. Bull. 52 (1973) 560–569.[21] A.A. Wereszczak, M.K. Ferber, M.G. Jenkins, C.K.J. Lin, Report ORNL/TM-12943,

1996.[22] J.A. Salem, J.M. Manderscheid, M.R. Freedman, J.P. Gyekenyesi, ASME 91-GT-

155, 1991.[23] S. Chockalingam, D.A. Earl, J. Eur. Ceram. Soc. 29 (2009) 2037–2043.[24] M.J. Hoffmann, Pure Appl. Chem. 67 (1995) 939–946.[25] Y. Tajima, in: I.W. Chen (Ed.), Silicon Nitride-Scientific and Technological

Advances, vol. 287, MRS Symposium Proceedings, MRS Pittsburgh, USA, 1993,p. 189.

[26] K. Urashima, Y. Tajima, M. Watanabe, in: R.C. Bradt, et al. (Eds.), FractureMechanics of Ceramics, vol. 9, Plenum Press, New York, 1992, p. 235.

[27] J.O. Kiggans Jr, F.C. Montgomery, T.N. Tiegs, M.K. Ferber, H-T. Lin, CONF-970568-9, 1997.

[28] P.F. Becher, E.Y. Sun, C.H. Hsueh, G.S. Painter, K.L. Morel, E00007852, ORNL/CP-102662, 1999.

[29] J. Miyauchi, Y. Kobayashi, SAE Technical Papers 850313, 1985.[30] H.-T. Lin, M.K. Ferber, W. Westphal, F. Macri, Proceedings of ASME Turbo Expo

2002, June 3–6, Amsterdam, The Netherlands, 2002, pp. 1–8, GT-2002-30629.[31] V. Parthasarathy, M. van Roode, J. Price, S. Gates, S. Waslo, P. Hoffman,

Proceedings of the 6th International Sympsium on Ceramic Materials andComponents for Engines, October 19–23, Arita, Japan, 1997, pp. 259–264.

[32] G. Ziegler, J. Heinrich, G. Wötting, J. Mater. Sci. 22 (1987) 3041–3086.[33] Silicon Nitride, Technical Brochure, Ceradyne, Inc., Advanced Technical

Ceramics, Costa Mesa, CA, 2005.[34] Y.S. Zheng, K.M. Knowles, J.M. Vieira, A.B. Lopes, F.J. Oliveira, J. Microsc. 201

(2001) 238–249.[35] M.H. Lewis, G. Leng-Ward, C. Jasper, in: G.L. Messing, E.R. Fuller Jr., H. Hausner

(Eds.), Ceramic Transactions, Ceramic Powder Science, II, B, American CeramicSociety, Westerville, OH, 1998, pp. 1019–1033.

[36] M.H. Bocanegra-Bernal, B. Matovic, Mater. Sci. Eng. A 500 (2009) 130–149.[37] O. Penas, R. Zenati, J. Dubois, G. Fanzotti, Ceram. Int. 27 (2001) 591–596.[38] H.J. Kleebe, G. Pezzotti, G. Ziegler, J. Am. Ceram. Soc. 82 (1999) 1857–1867.[39] X.-J. Liu, Z.-Y. Huang, Q.-M. Ge, X.-W. Sun, L.-P. Huang, J. Eur. Ceram. Soc. 25

(2005) 3353–3359.[40] H.J. Kleebe, W. Braue, G.D. Quinn, J. Am. Ceram. Soc. 90 (2007) 1286–1290.[41] J. Wan, R.-G. Duan, M.J. Gasch, A.K. Mukherjee, J. Am. Ceram. Soc. 89 (2006)

274–280.[42] W. Wotting, G. Ziegler, Ceram. Int. 10 (1984) 18–22.[43] F. Sergejev, M. Antonov, Proc. Estonian Acad. Sci. Eng. 12 (2006) 388–398.[44] A. Okada, N. Hirosaki, J. Mater. Sci. 25 (1990) 1656–1661.[45] F.F. Lange, J. Am. Ceram. Soc. 62 (1979) 428–430.[46] R.W. Rice, K.R. McKinney, C.C. Wu, S.W. Freiman, W.J.M. Donough, J. Mater.

Sci. 20 (1985) 1392–1406.[47] Z. Lencces, S. Guicciardi, C. Melandri, A. Bellosi, Br. Ceram. Trans. 94 (1995)

138–145.[48] G. Ling, H. Yang, Mater. Chem. Phys. 90 (2005) 31–34.[49] M. Liu, S. Nemat-Nasser, J. Mater. Res. 14 (1999) 4621–4629.[50] L.K.L. Falk, J. Mater. Sci. 39 (2004) 6655–6673.[51] I. IturrizaI, J. Echeberría, I. Gutierrez, F. Castro, J. Mater. Sci. 25 (1990)

2539–2548.[52] E.M. Knutson-Wedel, L.K.L. Falk, H. Björklund, T. Ekström, J. Mater. Sci. 26

(1991) 5575–5584.[53] B. Lawn, Fracture of Brittle Solids, 2nd ed., Cambridge University Press, Cam-

bridge, 1993.[54] W.-H. Lee, H.-E. Kim, S.-J. Cho, J. Am. Ceram. Soc. 80 (1997) 2737–2740.[55] C. Kooy, in: G.W. Stewart (Ed.), Science of Ceramics, vol. 1, Academic Press,

New York, 1962, pp. 21–34.[56] D. Kolar, Ceramic transactions, in: C.A. Handwerker, J.E. Blendell, W.A. Kaysser

(Eds.), in: Sintering of Advanced Ceramics, vol. 7, American Ceramic Society,Westerville, OH, 1990, pp. 529–545.

[57] M. Mitomo, S. Uenosono, J. Am. Ceram. Soc. 75 (1992) 103–108.[58] Cs. Balazsi, F. Weber, Zs. Kover, Z. Shen, Z. Konya, Zs. Kasztovszky, et al., Curr.

Appl. Phys. 6 (2006) 124–130.[59] Cs. Balaszi, Z. Konya, Zs. Kasztovszky, F. Weber, Z. Vertesy, L.P. Viro, et al., Fifth

International Conference on High Temperature Ceramic Matriz Composites(HTCMC-5) New Frontiers and Horizons, Seattle, Washington, USA, Ceramic

Transactions, The American Ceramic Society, 2004, pp. 107–112.

[60] N. Shibata, S.J. Pennycook, T.R. Gosnell, G.S. Painter, W.A. Shelton, P.F. Becher,Nature 428 (2004) 730–732.

[61] Y. Cai, S.R. Shah, A. Zimmermann, M. Weinmann, R. Raj, F. Aldinger, Phys.Status Solidi 193 (2002) R13–R15.

[62] L.K.L. Falk, K. Rundgren, J. Am. Ceram. Soc. 75 (1992) 28–35.

1 s Scien

[

[

[

[[[[

[

[[

336 M.H. Bocanegra-Bernal, B. Matovic / Material

[63] B.T. Lee, T. Koyama, A. Nishiyama, K. Hiraga, Scr. Metall. Mater. 32 (1995)1073–1077.

[64] B.T. Lee, S. Hayashi, T. Hirai, K. Hiraga, Mater. Trans. JIM 34 (1993) 573–579.[65] B.T. Lee, B.D. Han, H.D. Kim, Mater. Lett. 58 (2003) 74–79.[66] H.D. Kim, B.D. Han, D.S. Park, B.T. Lee, P.F. Becher, J. Am. Ceram. Soc. 85 (2002)

245–252.[67] J. Hojo, M. Uehara, H. Maeda, Adv. Tech. Mater. Mater. Process. J. 2 (2000)

17–23.[68] J. Hojo, K. Eto, M. Umezu, K. Kishi, S. Umebayashi, J. Jpn. Soc. Powder Powder

Metall. 45 (1998) 25–30.[69] H. Kugimoto, M. Uehara, N. Enomoto, J. Hojo, J. Ceram. Process. Res. 4 (2003)

6–9.[70] J. Hojo. J. Ceram. Soc. Jpn. Supp. 112-1, PacRim5 Special Issue 112 (2004)

S1002–S1008.[71] K. Ueno, T. Inoue, S. Sodeoka, M. Suzuki, M. Ishikawa, K. Uchiyama, et al., J.

Ceram. Soc. Jpn. 105 (1997) 304–307.[72] T. Yano, J. Tatami, K. Komeya, T. Meguro, J. Ceram. Soc. Jpn. 109 (2001)

396–400.[73] J.A. Mangels, G.J. Tennenhouse, Am. Ceram. Soc. Bull. 59 (1980) 1216–1218.[74] L.K.L. Falk, R. Pompe, G.L. Dunlop, J. Mater. Sci. 20 (1985) 3545–3556.[75] T.N. Tiegs, J.O. Kiggans, F.C. Montgomery, H-T. Lin, D.L. Barker, J.D. Snodgrass,

et al., CONF-960106-5, 1996.[76] A.J. Pyzic, D.R. Beamen, J. Am. Ceram. Soc. 76 (1993) 2737–2744.[77] T.N. Tiegs, J.O. Kiggans Jr, H.-T. Lin, DOE/OR/21400-7490, 1990.[78] J.O. Kiggans, C.R. Hubbard, R.R. Steele, H.D. Kimrey, C.E. Holcombe, T.N. Tiegs,

Ceramic transactions, in: Microwaves, Theory and Applications in MaterialsProcessing, American Ceramic Society, Westerville, OH, 1988, pp. 403–410.

[79] T.N. Tiegs, J.O. Kiggans, H.D. Kimrey, Ceram. Eng. Sci. Proc. 12 (1991)1981–1992.

[80] T.N. Tiegs, M.K. Ferber, J.O. Kiggans, K.L. More, C.M. Hubbard, D.W. Coffey,Ceramic transactions, in: D.E. Clark, F.D. Gac, W.H. Sutton (Eds.), Microwaves:Theory and Application in Materials Processing, American Ceramic Society,Westerville, OH, 1988, pp. 411–420.

[81] J.O. Kiggans, T.N. Tiegs, H.D. Kimrey, C.E. Holcombe, Processing of complex sin-tered reaction bonded silicon nitride parts by microwave heating, in: CeramicTransactions, Microwaves: Theory and Application in Materials Processing II,American Ceramic Society, Westerville, OH, 1993, pp. 269–276.

[82] J.O. Kiggans Jr., N.T. Tiegs, Proceedings Spring Meeting Mater. Res. Soc., SanFrancisco, CA, 1992.

[83] T.N. Tiegs, J.O. Kiggans Jr., H.-T. Lin, C.A. Wilkens, Comparison of the prop-erties of sintered and sintered reaction-bonded silicon nitride fabricated bymicrowave and conventional heating, in: Materials Research Society Proc-cedings, San Francisco, CA, 1994.

[84] W.J. Walker Jr., J.S. Reed, Ceram. Trans. 62 (1996) 141–148.[85] T.N. Tiegs, S.D. Nunn, T.M. Breavers, P.A. Menchhofer, D.L. Barker, D.W. Coffey,

Ceram. Eng. Sci. Proc. 16 (1995) 467–473.[86] P. Chantikul, G.R. Anstis, B.R. Lawn, D.B. Marshal, J. Am. Ceram. Soc. 64 (1981)

539–543.[87] J. Neil, G. Bandyopadhyay, D. Sordelet, M. Mahoney, Proceedings of 27th

Automotive Technology Development Contractors’ Coordination Meeting,Dearborn, MI, 1990, pp. 303–309.

[88] H.E. Helms, AGT 100 Project Summary, ASME Paper No. 88-GT-223, 1988.[89] S.T. Buljan, J.G. Baldoni, M.L. Huckabee, Am. Ceram. Soc. Bull. 66 (1987)

347–352.[90] J. Hefter, M. Mahoney, J. Harris, First International Ceramic Science and Tech-

nology Congress, November 3, Anaheim, CA, USA, 1989.[91] K. Hirao, T. Nagaoka, M.E. Brito, S. Kanzaki, J. Am. Ceram. Soc. 77 (1994)

1857–1862.[92] G. Petzow, M. Herrmann, Silicon nitride ceramics. High performance non-

oxide ceramics II, in: M. Jansen (Ed.), Structure and Bonding, vol. 102,Springer-Verlag, Berlin, 2002, pp. 47–167.

[93] M. Sawitz, Am. Ceram. Soc. Bull. 78 (1999) 53–56.[94] H. Mandel, M.J. Hoffmann, in: S. Hampshire, M.J. Pomeroy (Eds.), Nitrides and

Oxynitrides, Trans. Tech. Publications Ltd., Zelene, 2000, pp. 219–224.[95] P. Sajgalik, J. Dusza, M. Hoffman, J. Am. Ceram. Soc. 78 (1995) 2619–2624.[96] H. Emito, M. Mimoto, J. Eur. Ceram. Soc. 17 (1997) 797–804.[97] S. Kanzaki, M. Brito, M. Valecillos, K. Hirao, M. Toriyama, J. Eur. Ceram. Soc.

17 (1997) 1841–1847.[98] H. Imamura, K. Hirao, M.E. Brito, M. Toriyama, S. Kanzaki, J. Am. Ceram. Soc.

83 (2000) 495–500.[99] A. Vuckovic, Master Thesis, University Belgrade, 2005.100] Z.S. Rak, Polish Ceramics 2000 Conference, ECN-RX-00-003, 29-31 of May,

Spala, 2000.101] T. Takatori, T. Honma, N. Kamiya, H. Masaki, S. Sasaki, S. Wada, ASME Paper

91-GT-142, 1991.102] T. Ozawa, T. Matsuhisa, Y. Kobayashi, E. Matsuo, T. Inagaki. SAE Technical

Papers 890427, 1989.103] Y. Kobayashi, E. Matsuo, K. Kato, SAE Paper 880727, 1988.104] Y. Kobayashi, E. Matsuo, T. Inagaki, T. Ozawa, SAE Paper 910401, 1991.105] S. Natansohn, A.E. Pasto, ASME 91-GT-316, 1991.

106] N.L. Hecht, D.E. McCullum, G.A. Graves, Ceram. Eng. Sci. Proc. 9 (1988)

1313–1332.107] M. Matsui, T. Soma, Y. Ishida, I. Oda. SAE Technical Paper Series 860443, 1986,

141–149.108] H. Kawase, T. Matsuhisa, K. Kato, T. Mizumo. AME Paper 91-GT-270, 1991.109] M. Ito, N. Ishida, N. Kato. SAE Technical Paper Series 880704, 1988.

ce and Engineering A 527 (2010) 1314–1338

[110] T.N. Tiegs, S.D. Nunn, P.A. Menchhofer, D.L. Barker, E.D. Larson, T.L. Bravers, etal, in: Ceramic Technology Project, Semiannual Progress Report for April 1994Through September 1994, DE96000294, ORNL/TM 12924, 1995, pp. 65–79.

[111] S. Ribeiro, K. Strecker, Mater. Res. 7 (2004) 377–383.[112] K. Strecker, R. Gonzaga, S. Ribeiro, M.J. Hoffmann, Mater. Lett. 45 (2000) 39–42.[113] S. Ribeiro, Ph.D. Thesis, Faculdade de Engenharia Quimica de Lorena-

Departamento de Engenharia de Materiais, 1997, p. 194.[114] E.M. Rabinovich, Sh. Leitner, A. Goldenberg, J. Mater. Sci. 17 (1982) 323–328.[115] M. Kalantar, J. Dubois, C. Olagnon, G. Fantozzi, Key Eng. Mater. 132–136 (1997)

615–618.[116] T. Nose, T. Fujii, J. Am. Ceram. Soc. 71 (1988) 328–333.[117] D. Munz, T. Budsey, J.L. Shannon Jr., J. Am. Ceram. Soc. 63 (1980) 300–305.[118] K. Rundgren, Ceramic Development Programme Research 1993–1996, Sven-

ska Keraminstitulet, Swedish Ceramic Institute, 1996, p. 12.[119] A. Vuckovic, S. Boskovic, L. Zivkovic, J. Serb. Chem. Soc. 69 (2004) 59–67.[120] P.F. Becher, E.Y. Sun, K.P. Plucknett, K.B. Alexander, C.-H. Hsueh, H.-T. Lin, et

al., J. Am. Ceram. Soc. 81 (1998) 2811–2830.[121] E.Y. Sun, P.F. Becher, K.P. Plucknett, C.-H. Hsueh, K.B. Alexander, S.B. Waters,

et al., J. Am. Ceram. Soc. 81 (1998) 2811–2830.[122] V. Tikare, W.A. Sanders, S.R. Choi, NASA Contractor Report 191190, 1993.[123] K.R. Lai, T.Y. Tien, J. Am. Ceram. Soc. 76 (1993) 91–96.[124] H. Tsuruta, M. Masuda, T. Soma, M. Matsui, J. Am. Ceram. Soc. 73 (1990)

1714–1718.[125] S.R. Choi, J.M. Pereira, L.A. Janosik, R.T. Bhatt, NASA/TM-2002-211821, August

2002.[126] S.R. Choi, J.P. Gyekenyesi, ASME J. Eng. Gas Turb. Powers 121 (1999) 18–24.[127] T. Ohji, Ceram. Eng. Sci. Proc. 22 (2001) 159–166.[128] F. Lofaj, S.M. Wiederhorn, P.R. Jemian, Ceram. Eng. Sci. Proc. 22 (2001)

167–174.[129] H.-T. Lin, S.B. Waters, K.L. More, J. Wimmer, C.W. Li, Ceram. Eng. Sci. Proc. 22

(2001) 175–182.[130] ASTM C 1327, Test method for Vickers indentation hardness of advanced

ceramics, in: Annual Book of ASTM Standards, vol. 15.01, ASTM, West Con-shohocken, PA, 2002.

[131] A.K. Mukhopadhyay, S.K. Datta, D. Chakraborty, Ceram. Int. 25 (1999)447–454.

[132] D.A. Bell. in: C.C. Sorrel, B. Ben-Nissan (Ed.), Ceramic Developments, MaterialsScience Forum Volumens 34–36, 1988, pp. 1087–1091.

[133] N. Karakus, A. Osman Kurt, H. Ozkan Toplan, Ceram. Int. 35 (2009) 2381–2385.[134] G.-P. Jiang, J.-F. Yang, J.-Q. Gao, K. Niihara, Mater. Charact. 60 (2009) 456–460.[135] T. Shimizu, K. Takama, H. Enokishima, K. Mikame, S. Tsuji, SAE Technical

Papers 900656, 1990.[136] G. He, D.A. Hirschfeld, J. Cesarano III, SAND2000-0284C, February 2000.[137] M.K.M. Hruschka, W. Si, S. Tosatti, T.J. Graute, L.J. Gauckler, J. Am. Ceram. Soc.

82 (1999) 2039–2043.[138] M.A. Janney, O.O. Omatete. Method for molding ceramic powders using water-

based gelcasing, US Patent No. 5,028,362, July 2, 1991.[139] S.D. Nunn, J.O. Kiggans Jr, R.E. Simpson II, J.P. Maria, CONF-950401-32, 1995.[140] D.W. Richerson, D.W. Freitag. DOE/ORO 2076, 2–1 (1998) 2–25.[141] S. Kang, J.P. Johnston, T. Arima, M. Matsunaga, H. Tsuru, F.B. Prinz, ASME Paper

GT2003-38933, 2003.[142] I. Tsuji, S. Aoki, S. Miyazaki, Y. Hara, Y. Furuse, S. Hamada, ASME Paper 91-GT-

311, 1991.[143] H.-T. Lin, M.K. Ferber, T.P. Kirkland, J. Am. Ceram. Soc. 86 (2003) 1176–1181.[144] Y. Kobayashi, J. Mayauchi, E. Matsuo, T. Inagaki, 87-TOKYO-IGTC-122 (1987)

I-177.[145] P. Rottenkolber, M. Langer, R.S. Storm, F. Frechette, SAE Paper 810523, 1981.[146] K. Katayama, T. Watanabe, K. Matoba, N. Katoh, SAE Paper 861128, 1986.[147] C.M. Sonsino, Int. J. Fatigue 25 (2003) 107–116.[148] C.M. Sonsino, U. Brandt, K. Storzel, Symposium 2 of Materials Week’98, Octo-

ber 12–15, Munich, Germany, 1998.[149] K. Morgenthaler, Proceedings of the 6th International Symposium on Ceramic

Materials & Components for Engines, October 19–24, Arita, Japan, 1997.[150] D.K. Shetty, A.R. Rosenfield, P. McGuire, B.K. Bansal, W.H. Duckworth, Am.

Ceram. Soc. Bull. 59 (1980) 1193–1197.[151] C.A. Kelly, P.A. Suzuki, S. Ribeiro, S. Kycia, Ceramica 51 (2005) 168–172.[152] H.H. Lu, J.L. Huang, Ceram. Int. 27 (2001) 621–628.[153] G. Bandyopadhyay, K.W. French, J. Eng. Gas Turb. Power 108 (1986) 536–539.[154] C.L. Quackenbush, J.T. Neil, J.T. Smith, ASME Paper 81-GT-220, 1981.[155] J.T. Smith, C.L. Quackenbuch, Am. Ceram. Soc. Bull. 59 (1980) 529–532.[156] G. Bandyopadhyay, J.T. Neil, in: D.J. Viechnicki (Ed.), 37th Sagamore Army

Materials Research Conference, Publ., Dept. of Army, 1991, pp. 232–245.[157] H.-T. Lin, M.K. Ferber, Key Eng. Mater. 287 (2005) 393–403.[158] “Standard Test Method for Flexural Strength of Advanced Ceramics at Ambi-

ent Temperatures”, ASTM C1161, Annual Book of ASTM Standards, Vol 15.01,West Conshohocken, PA: American Society for Testing and Materials; 2002.

[159] Standard Test Method for Determination of Slow Crack Growth Parameters ofAdvanced Ceramics by Constant Stress-Rate Flexural Testing at Room Tem-peratures, vol. 15.01, ASTM C1368, Annual Book of ASTM Standards, AmericanSociety for Testing and Materials, West Conshohocken, PA, 1999.

[160] Y. Katano, M. Ando, T. Itoh, M. Sasaki, ASME paper 91-GT-264, 1991.[161] K.C. Dao, D.A. Shockey, L. Seaman, D.R. Curran, D.J. Rowcliffe, In Annual

Reports, Part III, Office of Naval Research, Contract No. 0001476-057, May1979.

[162] M. Abou el leil, F. Camaratta, R.R. Digenova, J. Am. Ceram. Soc. 68 (1985)C18–C21.

Scien

M.H. Bocanegra-Bernal, B. Matovic / Materials

[163] R. Thiruvengadaswamy, R.O. Scattergood, Scr. Metall. Mater. 25 (1991)2529–2532.

[164] H.-T. Lin, M.K. Ferber, J. Eur. Ceram. Soc. 22 (2002) 2789–2797.[165] J.E. Holowzak, T.G. Bornemisza, W.H. Day, Proccedings of the ASME Tur-

boExpo, Paper Number 2000-GT-662, Munich, Germany, May 8–11, 2000.[166] N.M.D. Burdick, R.E. Moreland, R.F. Geller, NACA Technical Note No. 1561,

1949, 1–53.[167] R.F. Geller, P.J. Yavorsky, B.L. Steierman, A.S. Creamer, J. Res. Nat. Bur. Stand.

36 (1946) 277–312.[168] R.F. Geller, M.D. Burdick, NACA ARR No. 6D24, 1946.[169] F.F. Lange, Am. Ceram. Soc. Bull. 62 (1983) 1369–1374.[170] C.T. Bodur, J. Mater. Sci. 30 (1995) 1511–1515.[171] A.A. Wereszczak, M.K. Ferber, T.P. Kirkland, Ceram. Eng. Sci. Proc. 15 (1994)

49–56.[172] Q. Jin, X.G. Ning, D.S. Wilkinson, G.C. Weatherly, J. Am. Ceram. Soc. 3 (1997)

685–691.[173] G.E. Gazza, J. Am. Ceram. Soc. 56 (1973) 662–1662.[174] M.N. Menon, H.T. Fang, D.C. Wu, M.G. Yenkins, M.K. Ferber, K.L. More, et al., J.

Am. Ceram. Soc. 77 (1994) 1217–1227.[175] H. Klemm, M. Herrmann, C. Schubert, J. Eng. Gas Turb. Power 122 (2000)

13–18.[176] S.R. Choi, F.A. Holand, NASA/TM-2000-210026, 2000.[177] D.A. Woodford, A.A. Wereszczak, W.T. Bakker, J. Eng. Gas Turb. Power 122

(2000) 206–211.[178] W.E. Luecke, J. Am. Ceram. Soc. 85 (2002) 408–414.[179] G.D. Zhan, M. Mitomo, T. Nishimura, R.J. Xie, T. Sakuma, Y. Ikuhara, J. Am.

Ceram. Soc. 83 (2000) 841–847.[180] A.A. Wereszczak, T.P. Kirkland, M.K. Ferber, DE95009172 CONF-950129-3,

1995.[181] A.A. Wereszczak, M.K. Ferber, T.P. Kirkland, E. Lara-Curzio, V. Parthasarathy,

T.T. Gribb, CONF-950129-4, DE95009177, 1995.[182] W.E. Luecke, S.M. Wiederhorn, B.J. Hockey, R.F. Krause Jr., G.G. Long, J. Am.

Ceram. Soc. 78 (1995) 2085–2096.[183] A.A. Wereszczak, T.P. Kirkland, H.-T. Lin, M.K. Ferber, C.-W. Li, J.A. Goldacker,

Ceram. Eng. Sci. Proc. 18 (1997) 45–55.[184] D. Chen, M.E. Sixta, X.F. Zhang, L.C. de Jonghe, R.O. Ritchie, Acta Mater. 48

(2000) 4599–4608.[185] K.J. Yoon, S.M. Wiederhorn, W.E. Luecke, J. Am. Ceram. Soc. 83 (2000)

2017–2022.[186] M.A. Boling-Risser, K.C. Goretta, J.L. Routbort, K.T. Faber, J. Am. Ceram. Soc. 83

(2000) 3065–3069.[187] J. Crampon, R. Duclos, F. Peni, S. Guicciardi, G. De Portu, J. Am. Ceram. Soc. 80

(1997) 85–91.[188] J. Crampon, R. Duclos, N. Rakotoharisoa, J. Mater. Sci. 28 (1993) 909–916.[189] W. Beere, Philos. Trans. R. Soc. Lond. A 288 (1978) 177–196.[190] R. Baddi, R. Duclos, J. Crampon. Mater. Sci. Eng. A165 (1993) L1–L3.[191] C.J. Gasdaska, J. Am. Ceram. Soc. 77 (1994) 2408–2418.[192] S.Y. Yoon, T. Akatsu, E. Yasuda, J. Mater. Res. 11 (1996) 120–126.[193] R.J. Xie, M. Mitomo, G.D. Zhan, Acta Mater. 48 (2000) 2049–2058.[194] G.D. Zhan, M. Mitomo, Y. Ikuhara, T. Sakuma, J. Am. Ceram. Soc. 83 (2000)

3179–3184.[195] G.D. Zhan, M. Mitomo, R.J. Xie, K. Kurashima, Acta Mater. 48 (2000)

2373–2382.[196] M. Backhaus-Ricoult, P. Eveno, J. Castaing, H.J. Kleebe, in: R.C. Bradt, C.A.

Brookes, J.L. Routbord (Eds.), Plastic Deformation of Ceramics, Plenum Press,New York, 1995, pp. 555–565.

[197] G.D. Zhan, M. Mitomo, T. Nishimura, R.J. Xie, J. Am. Ceram. Soc. 83 (2000)841–847.

[198] M. Backhaus-Ricoult, J. Castaing, J.L. Routbort, Rev. Phys. Appl. 23 (1998)239–249.

[199] J.M. Birch, B. Wilshire, J. Mater. Sci. 13 (1978) 2627–2636.[200] F.H. Norton, The Creep of Steel at High Temperatures, McGraw Hill, New York,

1929 (Chapter 8).[201] F.F. Lange, B.I. Davis, J. Mater. Sci. 17 (1982) 3637–3640.[202] J. Crampon, R. Duclos, N. Rakotoharisoa, J. Mater. Sci. 25 (1990) 1203–1208.[203] F.F. Lange, B.I. Davis, D.R. Clarke, J. Mater. Sci. 15 (1980) 601–610.[204] D.S. Wilkinson, J. Am. Ceram. Soc. 71 (1988) 562–565.[205] T. Rouxel, F. Rossignol, J.L. Besson, P. Goursat, J. Mater. Res. 12 (1997) 480–492.[206] J.D. French, S.M. Wiederhorn, J. Am. Ceram. Soc. 79 (1996) 550–552.[207] W.E. Luecke, S.M. Wiederhorn, J. Am. Ceram. Soc. 80 (1997) 831–838.[208] A.G. Evans, A. Rana, Acta Mater. 28 (1980) 129–141.[209] B.J. Hockey, S.M. Wiederhorn, W. Liu, J.G. Baldoni, S.-T. Buljan, J. Mater. Sci. 26

(1991) 3931–3939.[210] F.C. Monkman, N.J. Grant. Proc. Am. Soc. Test. Mater. 56 (1956) 593–620.[211] M.A. Boling-Risser, The effect of microstructure on the high-temperature

mechanical behavior of a self-reinforced, hot-pressed silicon nitride, Ph.D.Thesis, Northwestern University, Evanston, IL, 1998.

[212] A.A. Wereszczak, M.K. Ferber, T.P. Kirkland, K.L. Moore, in: R.C. Bradt, C.A.Brookes, J.L. Routbort (Eds.), Plastic Deformation of Ceramics, Plenum Press,New York, 1995, pp. 457–466.

[213] A.A. Wereszczak, M.K. Ferber, T.P. Kirkland, M.R. Foley, R.L. Yeckley, J. Am.Ceram. Soc. 78 (1995) 2129–2140.

[214] J.E. Marion, A.G. Evans, M.D. Drory, D.R. Clarke, Acta Metall. 31 (1983)1445–1457.

[215] G.D. Quin, W.R. Braue, J. Mater. Sci. 25 (1990) 4377–4392.[216] B. Burton, G.L. Reynolds, Acta Metall. 21 (1973) 1073–1078.

ce and Engineering A 527 (2010) 1314–1338 1337

[217] H.-J. Kleebe, M.K. Cinibulk, J. Mater. Sci. Lett. 12 (1993) 70–72.[218] H.-J. Kleebe, M.K. Cinibulk, R.M. Cannon, M. Ruhle, J. Am. Ceram. Soc. 76 (1993)

1969–1977.[219] I. Tanaka, H.-J. Kleebe, M.K. Cinibulk, J. Bruley, D.R. Clarke, M. Ruhle, J. Am.

Ceram. Soc. 77 (1994) 911–914.[220] S. Haig, W.R. Cannon, P.J. Whalen, Ceram. Eng. Sci. Proc. 13 (1992) 1008–1023.[221] S.M. Wiederhorn, D.E. Roberts, T.-J. Chuang, L. Chuck, J. Am. Ceram. Soc. 71

(1988) 602–608.[222] M.M. Chadwick, R.S. Jupp, D.S. Wilkinson, J. Am. Ceram. Soc. 76 (1994)

385–396.[223] P. Burger, R. Duclos, J. Crampon. Mater. Sci. Eng. A222 (1997) 175–181.[224] S.Y. Yoon, T. Akatsu, E. Yasuda, J. Mater. Sci. 32 (1997) 3813–3819.[225] G. Bernard-Granger, J. Crampon, R. Duclos, B. Cales, J. Eur. Ceram. Soc. 17

(1997) 1647–1654.[226] A.W. Rosenflanz, I.W. Chen, J. Am. Ceram. Soc. 80 (1997) 1341–1352.[227] A.A. Wereszczak, M.K. Ferber, T.P. Kirkland, A.S. Barnes, E.L. Frome, M.N.

Menon, J. Eur. Ceram. Soc. 19 (1999) 227–237.[228] I.W. Chen, S.-L. Hwang, J. Am. Ceram. Soc. 75 (1992) 1073–1079.[229] F. Lofaj, A. Okada, H. Usami, H. Kawamoto, J. Am. Ceram. Soc. 82 (1999)

1009–1019.[230] W.E. Luecke, S.M. Wiederhorn, J. Am. Ceram. Soc. 82 (1999) 2769–2778.[231] J.-L. Ding, K.C. Liu, K.L. More, C.R. Brinkman, J. Am. Ceram. Soc. 77 (1994)

867–874.[232] M.N. Menon, H.T. Fang, D.C. Wu, M.G. Yenkins, M.K. Ferber, J. Am. Ceram. Soc.

77 (1994) 1228–1234.[233] F.R.N. Nabarro, Deformation of crystals by the motion of single ions, Report of

a Conference on Strength of Solids (H.H. Wills Physical Laboratory, Universityof Bristol, July 7–9, 1947), London, UK: Physical Society, 1948, pp. 75–90.

[234] C. Herring, J. Appl. Phys. 21 (1950) 437–445.[235] R.L. Coble, J. Appl. Phys. 34 (1963) 1679–1682.[236] R. Raj, C.K. Chyung, Acta Metall. 29 (1981) 159–166.[237] R. Kossowsky, J. Mater. Sci. 8 (1973) 1603–1615.[238] P.J. Whalen, C.J. Gasdaska, R.D. Silvers, Ceram. Eng. Sci. Proc. 11 (1990)

633–640.[239] N. Kondo, Y. Suzuki, M.E. Brito, T. Ohji, J. Mater. Res. 16 (2001)

2182–2185.[240] R.F. Krause Jr., W.E. Luecke, J.D. French, B.J. Hockey, S.M. Wiederhorn, J. Am.

Ceram. Soc. 82 (1999) 1233–1241.[241] H.A. Maupas, Personal communication on creep of AlliedSignal AS800 Silicon

Nitride, 1998.[242] W.E. Luecke, S.M. Wiederhorn, B.J. Kockey, G.G. Long, Mater. Res. Soc. Symp.

Proc. 287 (1993) 467–472.[243] H.T. Lin, P.F. Becher, M.K. Ferber, V. Parthasarathy, Key Eng. Mater. 161–163

(1999) 671–674.[244] S.L. Hwang, H.T. Lin, P.F. Becher, J. Mater. Sci. 30 (1995) 6023–6027.[245] H.T. Lin, J.O. Kiggans Jr., T.N. Tiegs, J. Mater. Sci. 31 (1996) 6477–6483.[246] D.F. Carroll, S.M. Wiederhorn, D.E. Roberts, J. Am. Ceram. Soc. 72 (1989)

1610–1614.[247] A.A. Wereszczak, K. Breder, M.K. Ferber, J. Am. Ceram. Soc. 76 (1993)

2919–2922.[248] A.A. Wereszczak, T.P. Kirkland, K. Breder, M.K. Ferber, P. Khandelwal, Mater.

Sci. Eng. A191 (1995) 257–266.[249] W.E. Lee, C.H. Drumond III, G.E. Hilmas, S. Kumar, J. Am. Ceram. Soc. 73 (1990)

3575–3579.[250] R.F. Krause Jr., S.M. Wiederhorn, J. Am. Ceram. Soc. 84 (2001) 2394–2400.[251] A.A. Wereszczak, M.K. Ferber, A.S. Barnes, T.P. Kirkland, Ceram. Eng. Sci. Proc.

20 (1999) 535–544.[252] S.M. Wiederhorn, G.D. Quinn, R.F. Krause Jr., in: C.R. Brinkman, S.F. Duffy (Eds.),

Life Prediction Methodologies and Data for Ceramic Materials ASTM STP 1201,American Ceramic for Testing and Materials, West Conshohocken, PA, 1994,pp. 36–61.

[253] S.M. Wiederhorn, A.R. de Arellano Lopez, W.E. Luecke, M.J. Hoffmann, B.J.Hockey, J.D. French, et al., J. Am. Ceram. Soc. 87 (2004) 421–430.

[254] W. Acchar, W. Rendtel, H. Hubner, in: P. Duran, FernandezF J.J. (Eds.), ThirdEuro-Ceramics, vol. 3, Faenza Editrice Iberica, SL, Castellon de la Plana, Spain,1993, pp. 405–410.

[255] T. Hirano, T. Ohji, K. Niihara, Mater. Lett. 27 (1996) 53–58.[256] S.-Y. Yoon, T. Akatsu, E. Yasuda, Y. Tanabe, J. Ceram. Soc. Jpn. 104 (1996)

239–242.[257] T. Ohji, Y. Yamauchi, J. Am. Ceram. Soc. 76 (1993) 3105–3112.[258] S. Testu, T. Rouxel, J.L. Besson, Acta Mater. 49 (2001) 3589–3596.[259] A. Rendtel, H. Hubner, J. Eur. Ceram. Soc. 22 (2002) 2517–2525.[260] Q. Wei, J. Sankar, J. Narayan. Mat. Sci. Eng. A299 (2001) 141–151.[261] J.W. Cao, A. Okada, N. Hirosaki, J. Eur. Ceram. Soc. 22 (2002) 237–245.[262] F. Lofaj, Mat. Sci. Eng. A279 (2000) 61–72.[263] F. Lofaj, J.W. Cao, A. Okada, H. Kawamoto, in: K. Niihara, K. Kanzaki, K. Komeya,

S. Hirano, K. Morinaga (Eds.), Proceedings of the 6th International Sympo-sium on Ceramic Materials and Components for Engines. Technoplaza, Tokyo,Japan, 1998, pp. 713–718.

[264] M. Mayne, T. Rouxel, D. Bahloul-Hourlier, J.L. Besson, J. Eur. Ceram. Soc. 18

(1998) 1985–1994.

[265] T. Rouxel, F. Wakai, K. Izaki, J. Am. Ceram. Soc. 75 (1992) 2363–2372.[266] N. Kondo, Y. Suzuki, T. Ohji, E. Sato, F. Wakai, Mater. Sci. Eng. A268 (1999)

141–146.[267] F. Wakai, Y. Kodama, S. Sakaguchi, N. Murayama, K. Izaki, K. Niihara, Nature

(London) 344 (1990) 421–423.

1 s Scien

[[[

[[[[[

[

338 M.H. Bocanegra-Bernal, B. Matovic / Material

268] Q. Jin, D.S. Wilkinson, G.C. Weatherly, J. Eur. Ceram. Soc. 18 (1998) 2281–2286.269] I.W. Chen, L.A. Xue, J. Am. Ceram. Soc. 73 (1990) 2585–2609.270] R.J. Xie, M. Mitomo, G.D. Zhan, H. Emoto, J. Am. Ceram. Soc. 83 (2000)

2529–2535.271] X. Xu, I.W. Chen, J. Am. Ceram. Soc. 75 (1992) 2733–2741.272] S.L. Hwang, I.W. Chen, J. Am. Ceram. Soc. 10 (1994) 2575–2585.

273] L.A. Xue, X. Wu, I.W. Chen, J. Am. Ceram. Soc. 74 (1991) 842–845.274] L.A. Xue, I.W. Chen, J. Am. Ceram. Soc. 75 (1992) 1185–1191.275] M. Mitomo, H. Hirotsuru, H. Suetmatsu, T. Nishimura, J. Am. Ceram. Soc. 78

(1995) 211–214.276] C.M. Wang, M. Mitomo, T. Nishimura, Y. Bando, J. Am. Ceram. Soc. 80 (1997)

1213–1221.

ce and Engineering A 527 (2010) 1314–1338

[277] T. Nishimura, Y. Bando, M. Mitomo, H. Hirotsuru, in: A. Bellosi (Ed.), FourthEuropean Ceramics Society Conference vol. 4, Gruppo Editoriale FaenzaEditrice, Riccione, Italy, 1995, pp. 265–270.

[278] T. Nishimura, G.D. Zhan, M. Mitomo, H. Sato, Mater. Sci. Forum 304–306(1999) 477–482.

[279] N. Kondo, T. Ohji, F. Wakai, J. Am. Ceram. Soc. 81 (1998) 713–716.

[280] N. Kondo, Y. Suzuki, T. Ohji, J. Am. Ceram. Soc. 82 (1999) 1067–1069.[281] N. Hirosaki, Y. Kimono, M. Mitomo, J. Ceram. Soc. Jpn. 101 (1993) 1239–1243.[282] J.A. Schneider, A.K. Mukherjee, Ceram. Eng. Sci. Proc. 17 (1996) 341–353.[283] C.K. Yoon, I.W. Chen, J. Am. Ceram. Soc. 73 (1990) 1555–1565.[284] K. Honma, H. Okada, T. Fujikama, T. Tatuno, J. Ceram. Soc. Jpn. Int. Ed. 95 (1987)

195–199.